Perovskite Solar Cells
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Long Electron–Hole Diffusion Length in High-Quality Lead-Free Double Perovskite Films
Weihua Ning, Feng Wang, Bo Wu, Jun Lu, Zhibo Yan, Xianjie Liu, Youtian Tao, Jun-Ming Liu, Wei Huang, Mats Fahlman, Lars Hultman, Tze Chien Sum,* and Feng Gao*
Solution-processed lead halide perovskites have shown supe-rior optoelectronic properties, including strong and tunable light absorption/emission, long carrier diffusion lengths, and high carrier mobilities.[1,2] As a result, the power conversion
Dr. W. Ning, Dr. F. Wang, Dr. J. Lu, Dr. Z. Yan, Dr.
X. Liu, Prof. M. Fahlman, Prof. L. Hultman, Prof. F.
Gao Department of Physics, Chemistry, and Biology (IFM) Linköping University
Linköping SE-581 83, Sweden E-mail: [email protected]
Dr. W. Ning, Prof. Y. Tao, Prof. W. Huang
Key Lab for Flexible Electronics & Institute of Advanced Materials Jiangsu National Synergistic Innovation Center for Advanced Materials (SICAM)
Nanjing Tech University
30 South Puzhu Road, Nanjing 211816, P. R. China Dr. B. Wu, Prof. T. C. Sum
Division of Physics and Applied Physics School of Physical and Mathematical Sciences Nanyang Technological University (NTU)
21 Nanyang Link, Singapore 637371, Singapore E-mail: [email protected] Dr. Z. Yan, Prof. J.-M. Liu
Laboratory of Solid State Microstructures and Innovation Center of Advanced
Microstructures Nanjing University
Nanjing 210093, P. R. China
The ORCID identification number(s) for the author(s) of this article can be found under https://doi.org/10.1002/adma.201706246.
DOI: 10.1002/adma.201706246
efficiencies of perovskite solar cells have increased from 3.8% to 22.1% within only a few years, making perovskites the fastest- advancing technology in the photovoltaic history.[3,4] To ensure the sustainability of the perovskite photovol-taic technology, the number of studies to address the lead (Pb) toxicity and device stability issues has increased.[5–7] The most obvious option for lead-free perovskites is the substitution of Pb2 with another diva-lent cation (e.g., germanium (Ge2) or tin (Sn2)).[8,9]
Unfortunately, the resulting perovskites based on Sn2 or Ge2 are easily oxidized by O2, limiting their prac-tical applications.[10] Bismuth (Bi)-based organic–inorganic metal halides have also been studied as an alternative for solar cell applications.[11] Different from the 3D lead-based perovskites, the 0D to 2D structures of Bi-based organic–inorganic halides lead to strongly bound excitons with low mobilities.[5]
A new generation of perovskites, lead-free halide double per- ovskites with a general formula of A2MM3X6, where both A and M are monovalent cations, M3 is a trivalent cation, and X is a halide, provide rich substitutional chemistry and promising optoelectronic properties.[12] Several groups have successfully syn- thesized double perovskite powders and single crystals, and car-ried out crystal characterizations and fundamental studies.[13–15] Double perovskites show tunable band gaps spanning the visible to near- infrared spectra and possess relatively low carrier effec-tive masses that are favorable for efficient charge transport and extraction, similar to 3D lead-based perovskites.[14,16] Moreover, these materials provide rich substitutional chemistry, which can dramatically change their photophysical properties.[17,18] For example, Tl-doped Cs2(Ag1−aBi1
−b)TlxBr6 (x 0.075) results in a decrease in the bandgap of ≈0.5 eV.[18] Recent first-principle calculations also indicate that halide double perovskites are promising candidates for photovoltaic applications.[16,19,20] Fur-thermore, these double perovskites are much more stable than Ge or Sn perovskites in repelling the attacks by O2 and H2O.[15,21] However, since the precursors of double perovskites cannot dis-solve in common solvents (for example, dimethylformamide) which are frequently for lead-based perovskites, it is still a chal-lenge to fabricate double perovskite solar cells.[22]
And also, most of the fundamental questions concerning the photophysics of double perovskite films remain unexplored and unknown due in part to the lack of uniform and high-quality films.
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Developing environmentally friendly perovskites has become important in solving the toxicity issue of lead-based perovskite solar cells. Here, the first double perovskite (Cs2AgBiBr6) solar cells using the planar structure are demonstrated. The prepared Cs2AgBiBr6 films are composed of high-crystal- quality grains with diameters equal to the film thickness, thus minimizing the grain boundary length and the carrier recombination. These high-quality double perovskite films show long electron–hole diffusion lengths greater than 100 nm, enabling the fabrication of planar structure double perovskite solar cells. The resulting solar cells based on planar TiO2 exhibit an average power conversion efficiency over 1%. This work represents an important step forward toward the realization of environmentally friendly solar cells and also has important implications for the applications of double perovskites in other optoelectronic devices.
Figure 1. a) Low-magnification and b) high-magnification SEM images, c) the TEM image and SAED pattern, and d) the XRD pattern of the prepared Cs2AgBiBr6 films annealed at 250 C for 5 min.
In this work, we demonstrate the first double perovskite solar cells using the planar structure. We prepare high-quality films with single-layer Cs2AgBiBr6 crystals. Through photo-physical investigations, we find the coexistence of excitons and free carriers in the material. These Cs2AgBiBr6 films show a long photoexcited carrier diffusion length of ≈110 nm. The resulting solar cells based on planar TiO2 exhibit an average power conversion efficiency (PCE) over 1%.
The high quality Cs2AgBiBr6 films are prepared through a one- step spin-coating process from single-crystal Cs2AgBiBr6 solutions. Figure 1a,b shows typical scanning electron micro- scopy (SEM) images of the perovskite films from a 0.5 m solu- tion. The surface roughness (Rq) is only ≈24 nm (Figure S1, Supporting Information). The smooth film is essential for the following photoluminescence (PL)-quenching measure-ments and photovoltaic performance. The films are composed of closely packed polycrystalline grains with diameters of 100–500 nm. To examine each individuate grains, transmis-sion electron microscopy (TEM) and selected electron diffrac-tion (SAED) are performed. Impressively, both TEM and SAED reveal that each grain is a single crystal (Figure 1c), although the films are polycrystalline. This feature is beneficial for the photovoltaic performance since there is no grain boundary in between from top to bottom of the film. The limited grain boundary in the vertical direction would be important for efficient carrier transfer in devices. In addition, the X-ray dif-fraction (XRD) pattern confirms the pure phase in the Cs2Ag-BiBr6 films, matching well with the results of the simulation (Figure 1d).
Figure 2 shows the UV–vis absorption of Cs2AgBiBr6 thin film. There are three parts in the absorption spectrum: below 400 nm, with a flat absorption feature; an excitonic absorp-tion band in the region from 400 to 500 nm; and a very weak indirect absorption band between 500 and 538 nm, similar to that in single crystals.[14] The absorption coefficients at 439 nm reach up to 1 105 cm−1. By using the Elliott formula, the direct bandgap (Egd) is
≈3.26 eV (See Figure S2 and the Supporting Information). The indirect band gap cannot be determined by the thin films due to its weak absorption, but was extracted pre-viously to be ≈1.95 eV from single crystals.[14]
We obtain PL spectra and time-resolved photolumines-cence (TRPL) studies to understand the photoexcited species. As shown in Figure 2a, the broad PL peak centered at ≈2.0 eV (620 nm) can be attributed to the indirect bandgap emission, as it corresponds well with the indirect absorption and emission found previously in single crystals.[14] The excitation intensity dependence of PL intensity just after photoexcitation is gene- rally a good indicator of the nature of the radiative recombi-nation processes.[23] Briefly, the initial PL intensity exhibits a quadratic dependence on the photoexcitation density for emission by free-carrier band edge recombination: PLt0n2, where n is the photoexcitation density (See the Supporting Information for detailed explanation). For emission by radia-tive recombination of excitons or free carriers with doped car-riers, PLt0 n.[24–26] It is noted that the PL signal is too weak to be detected if the carrier density below 1016 cm−3. When the carrier density is between 1016 and 1017 cm−3, the PL intensity shows power-dependence on the carrier density with a scaling
Figure 2. a) UV–vis absorption and PL spectra of a 100 nm Cs2AgBiBr6 film. b) PL intensity as a function of carrier density, the PL intensities were the values just after photoexcitation (PLt0), rather than the integrated PL intensity. c) PL decay dynamics of the Cs2AgBiBr6, Cs2AgBiBr6/PC61BM, and Cs2AgBiBr6/Spiro-OMeTAD films. d) The photocurrent of the Cs2AgBiBr6 devices as a function of solution concentration.
factor (λ) of 1.34 (Figure 2b). This value suggests either the coexistence of radiative recombination of the free electrons– holes (recombination order 2) and the free carriers with doped carriers (recombination order 1), or the coexistence of exci-tons (recombination order 1) and free electrons–holes in the measured excitation density range. However the doped carrier density measured by Hall Effect is on the order of 1013 cm−3, which is not comparable to the photoexcitation density. Hence, it would be more reasonable to conclude that excitons and free carriers coexist in the double perovskite films, and the exciton possesses a higher ratio in the measured excitation range. Fur-ther increasing the carrier density to above 1017 cm−3 causes the linear scale factor (λ) to decrease to 0.85. Meanwhile, the effective PL lifetime (τeff, the time at which the PL intensity drops to 1/e of its maximum value) also decreases continu-ously with increasing carrier densities (Figure S3, Supporting Information), implying strong high-order recombination such as exciton–carrier and exciton–exciton Auger recombination at high carrier densities.
This unusual photophysical behavior contrasts from that of lead- based perovskites and further inves-tigations are warranted.
To quantify the carrier diffusion length of the Cs2AgBiBr6 film, we carry out TRPL experiments. The thicknesses of perovskite films are ≈100 nm, and the quenching samples are prepared by spin-coating layers of either a hole-transporting acceptor (Spiro- MeOTAD) or an electron-accepting fullerene (PC61BM) on top of the perovskite film. The results show that the PL decay dynamics of Cs2AgBiBr6 film can be fitted by a biexponential decay function at low excitation density, with time constants being τ1 2.5 0.4 ns (52%), τ2 35 1 ns (48%), which arises from the crystal size inhomogeneity. As a result, the effective excitation lifetime τeff (the time for the PL decaying
t t
to 1/e of its initial intensity: ∑ Ai exp− ∑Ai exp− ,
i τ eff i
τ
i
where τi is the ith fitted lifetime component of the decay curve and Ai is its weighted amplitude) in pristine film is 13.7 0.4 ns (Figure 2c). After the film is coated with PC61BM, the PL decays much faster, with τeff being 2.4 0.4 ns. This indi-cates highly efficient electron transfer from the perovskite to PC61BM. A similar value of 2.6 0.2 ns was obtained for Spiro-MeOTAD coated perovskite films. Based on the PL quenching model,[25,26]
we estimate an average photoexcitation mobility of 0.37 0.15 cm2 V−1 s−1, and the photoexcitation diffusion length for electrons and holes ≈110 20 nm. The similarity of diffusion length is consistent with the dominance of excitons as the primary photoexcitation species in the excitation density range. With photoexcitation diffusion lengths above 100 nm, the double perovskite polycrystalline film already shows excel-lent carrier diffusion properties comparable to those of lead-based perovskite films possessing typical carrier diffusion lengths of 100 nm–1
m.[25–27] It is noted that the diffusion length measurements were performed at a carrier density of ≈5 1016 cm−3, which is expected to be higher than that of devices under solar illumination (1015–1016 cm−3).[27] It is dif-ficult to make sure whether there is a trap-filling effect during the TRPL measurements, as we did not observe an increase of carrier lifetime with increasing the carrier density (Figure S4, Supporting Information), possibly due to low trap density in the devices. Therefore, trap-filling as well as the tendency of the photoexcitation species changing from excitons to free carriers were not considered in the measurements.
The long diffusion lengths and high mobility of the Cs2Ag- BiBr6 films implies that carriers can travel through the film; thus solar cells are expected to operate well on planar-structured
devices. The detailed information about the fabrication of an indium tin oxide (ITO)/compact TiO2/Cs2AgBiBr6/spiro- MeOTAD/Au device is described in the Supporting Information.
The thickness-dependent photocurrent is compared by varying the concentration of the solution as follows: 0.4, 0.45, 0.5, and 0.55 m. The corresponding thicknesses are 145 12, 170 15, 205 10, and 223 10 nm, respectively. J–V results show that the photocurrent initially increases with increasing thickness, with an optimised Jsc ≈ 1.7 mA cm−2 at a ≈205 nm thickness (Figure 2d; Figures S5 and S6, Supporting Information). Fur-ther increasing the thickness to ≈223 nm decreases the photo-current to ≈1.1 mA cm−2. The optimum thickness of ≈205 nm confirms the long diffusion length of Cs2AgBiBr6 films.
The annealing temperature also has a significant effect on the device performance. The Voc values are all ≈1.0 V for dif-ferent annealing temperatures (Figure S7, Supporting Infor-mation).
However, Jsc and fill factor (FF) values first gradually increase with increasing annealing temperatures, and then drop down when the temperature is above 250 C. This result can be rationalised by considering the characteristics of crys-tallinity, grain size, and pinholes of films together (Figures S8 and S9, Supporting Information). The grain size and crystal-linity gradually increase with increasing annealing tempera-ture. Unfortunately, a large number of pinholes appear when the annealing temperature increases above 300 C. Further increase in the annealing temperature to 400 C results in the degradation of the films (Figure S8, Supporting Information).
The optimized devices exhibit an average PCE of 1.05%
(averaged from 40 devices from five different batches) and outstanding PCE of 1.22% with a Voc of 1.06 V (Figure 3c).
Figure 3d shows the corresponding stabilized power output with a bias of 0.82 V. The device exhibits a rapid response after
illumination, resulting in a stable PCE of 1.17% over 600 s of illumination. Notably, there is almost no hysteresis behavior in the J–V curves, implying less trapping/detrapping or ion migration in Cs2AgBiBr6 compared with lead-based hybrid per-ovskites. Very recently, Scanlon and co-workers estimated the spectroscopic limited maximum efficiency (SLME) of Cs2Ag-BiBr6 (200 nm) to be 7.92%.[28] The SLME takes into account the strength of optical absorption and the nature of the band gap in the overall theoretical efficiency of an absorber material.[19] As an indirect band gap semiconductor, the SLME of Cs2Ag-BiBr6 is significantly dependent on the thickness of the films, and hence future approaches are required either to enhance the thickness of Cs2AgBiBr6 films or to make Cs2AgBiBr6 into direct bandgap semiconductor, e.g., through doping.[18]
We note that the efficiency we obtained (up to 1.22%) is much lower than the SLME of Cs2AgBiBr6 at 200 nm, in spite of high crystal quality. One of the reasons might be due to the fact that the charge extraction efficiency of TiO2 for the Cs2Ag-BiBr6 films is not as efficient as those of Spiro-MeOTAD and PC61BM. PL intensity and decay of TiO2/Cs2AgBiBr6 show no obvious difference compared with those of pure Cs2Ag-BiBr6 (Figure S10, Supporting Information). Since the band energy of Cs2AgBiBr6 matches with those of TiO2 and Spiro-OMeTAD (Figure S11, Supporting Information), the reason is possibly due to the presence of an interfacial barrier from sur-face defects, similar to the MAPbI3:TiO2 heterojunction.[29] We also estimate the theoretical JSC value for Cs2AgBiBr6 devices based on the diffusion length values and the absorption coef-ficient.[30] The calculated theoretical Jsc value for Cs2AgBiBr6 devices can be around 5.2 mA cm−2, which is higher than 1.7 mA cm−2 in our results. This also implies poor charge extraction of TiO2 in the devices. We thus further fabricate
Figure 3. a) Planar Cs2AgBiBr6 solar cell structure, ITO/compact TiO2/Cs2AgBiBr6/Spiro-MeOTAD/Au. b) Cross-sectional SEM image of Cs2AgBiBr6 solar cells. c) J–V curve of the optimized device. d) The stabilized power output of the Cs2AgBiBr6 solar cells with a bias of 0.82 V.
an ITO/perovskite/Spiro-MeOTAD/Au device without TiO2 to understand the charge extraction process. The resulting Jsc is
≈1.7 mA cm−2 (Figure S12, Supporting Information), which is almost the same as that of devices with TiO2. This result con- firms less effective electron extraction by TiO2. However, the Voc of the devices without TiO 2 is only ≈0.5 V, much lower than the
≈1.0 V shown with TiO2. The high Voc suggests that TiO2 can prevent the carrier recombination at the ITO side.
In summary, we fabricate a uniform Cs2AgBiBr6 thin solid film of high crystal quality through a one-step spin-coating process from single-crystal Cs2AgBiBr6 solution. Upon excita-tion, excitons and free carriers coexist in double perovskite films, with a long diffusion length of ≈110 nm. We achieve an average PCE over 1% based on a planar device structure with a maximum value of 1.22%. The photovoltaic performance is expected to be further boosted by replacing the TiO2 (used presently in this study) with more suitable ETL materials and by increasing the film thickness while maintaining the film quality. In addition, it will be favorable to develop direct bandgap double perovskites, e.g., through doping. The long carrier diffusion length of high-quality double perovskite films opens a route toward developing environmentally friendly per-ovskite-based solar cells.
Experimental Section
Cs2AgBiBr6 Single Crystals Synthesis: CsBr (213 mg, 1.00 mmol), BiBr3
(225 mg, 0.5 mmol), and AgBr (94 mg, 0.5 mmol) were dissolved in 3 mL of 47% HBr. The solution was transferred to a Teflon-lined reactor. After reacting at 120 C for 24 h, and cooling to room temperature slowly, red Cs2AgBiBr6 octahedral single crystals with the size of 2–5 mm can be collected (Figure S13, Supporting Information). The yield is ≈85% calculated from Ag.
Cs2AgBiBr6 Solar Cell Fabrication: The TiO2 compact film precursor solution in ethanol consists of 0.3 m titanium isopropoxide (Sigma- Aldrich, 99.999%) and 0.01 m HCl. A ≈35 nm dense TiO2 film was coated onto an ITO substrate by spinning a titanium precursor at 5000 rpm, followed by annealing at 200 C for 2 h. The synthesized Cs2AgBiBr6 single crystals were dissolved in dimethyl sulfoxide (DMSO) with a temperature of 100—130 C. After the crystals were completely dissolved, the solution was cooled to room temperature, and then deposited onto the TiO2/ITO substrate by spin-coating at 3000 rpm for 60 s. The films were annealed at 250 C for 5 min in order to obtain better crystallization. The thickness of the Cs2AgBiBr6 films was controlled by varying the concentration of the precursor solution from 0.4–0.55 m. The highest PCE value of Cs2AgBiBr6 solar cells was achieved from the 0.5 m solution. The spiro-MeOTAD based hole- transfer layer was prepared by dissolving 60 mg spiro-MeOTAD, 17.5
L lithium-bis(trifluoromethanesulfonyl)imide (Li-TFSI) solution (520 mg Li-TFSI in 1 mL acetonitrile), and 28.8 L 4-tert-butylpyridine in 1 mL chlorobenzene. The devices were put into a dry cabinet for 15 h for the oxidization of Spiro-MeOTAD. The hole-transfer layer was deposited by spin-coating at 5000 rpm for 30 s. Finally, a 100 nm gold layer was deposited by thermal evaporation at a pressure of 1 10−4 mbar. All device fabrication steps were carried out in an N2-purged glovebox.
Measurement and Characterization: The XRD patterns of the products were recorded with a X’Pert PRO X-ray diffractometer using Cu Kα1 irradiation (λ 1.5406 Å). The Ultraviolet–visible absorption spectra were measured on a Shimadzu spectrophotometer (UV-2450).
The general morphologies of the films were characterized by SEM (LEO 1550). The Atomic force microscope measurement was carried out using a Dimension 3100/NanoScope IV system equipped with a C-AFM module (Veeco, Bruker). TEM was performed in the FEI Tecnai
G2 TF20 UT with a field emission gun operated at 200 kV and a point resolution of 0.19 nm. Sample thicknesses were measured using an Alpha step 500 Surface profilometer. The current density–voltage (J–V) curves were measured (Keithley Instruments, 2400 Series SourceMeter) under simulated AM 1.5 Solar Simulator. The effective area of the cell was defined as 0.075 cm2. The external quantum efficiency (EQE) data were obtained using a solar cell spectral response measurement system (QE-R3011, Enli Technology Co. Ltd), and the light intensity at each wavelength was calibrated with a standard single-crystal Si photovoltaic cell. PL and TRPL measurements were performed using 400 nm femtosecond excitation pulses (50 fs). The laser pulses were generated by passing the strong 800 nm femtosecond laser beam (Coherent Libra, 50 fs) through a beta barium borate (BBO) crystal (frequency doubler).
The emitted light was collected at a backscattering angle by a spectrometer (Acton, Spectra Pro 2500i) and CCD (Princeton Instruments, Pixis 400B) in PL measurements and by an Optronis Optoscope streak camera system which has an ultimate temporal resolution of 6 ps in TRPL measurements.
Supporting Information
Supporting Information is available from the Wiley Online Library or from the author.
Acknowledgements
W.N. and F.W. contributed equally to this work. The authors thank Lijun Zhang (Jilin University, China) for insightful discussions. The work was financially supported by the Joint NTU-LiU PhD programme on Materials- and Nanoscience, the Swedish Government Strategic Research Area in Materials Science on Functional Materials at Linköping University (Faculty Grant SFO Mat LiU No. 200900971), the European Commission Marie Skłodowska-Curie actions (No. 691210); the European Commission SOLAR-ERA.NET, the Swedish Energy Agency (Energimyndigheten), and the Swedish Research Council (FORMAS). F.G. is a Wallenberg Academy Fellow. Both F.W. and Z.Y. are VINNMER Marie Skłodowska-Curie Fellows. W.N. is supported by the China Scholarship Council. T.C.S. acknowledges the financial support from the Singapore Ministry of Education Academic Research Fund Tier 1 grants RG101/15 and RG173/16, and Tier 2 grants MOE2014-T2-1-044, MOE2015-T2-2-015, and MOE2016-T2-1-034; and from the Singapore National Research Foundation through the Competitive Research Program NRF-CRP14-2014. L.H. acknowledges support from the Knut and Alice Wallenberg (KAW) Foundation for a Scholar Grant 2016.0358, and for support to the Linköping Ultra Electron Microscopy Laboratory.
Conflict of Interest
The authors declare no conflict of interest.
Keywords
double perovskite, lead free, long diffusion length, planar solar cell
Received: October 27, 2017 Revised: February 1, 2018 Published online: March 30, 2018
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