Binary-Metal Perovskites Toward High-Performance Planar-Heterojunction Hybrid Solar Cells

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ON Binary-Metal Perovskites Toward High-Performance

Planar-Heterojunction Hybrid Solar Cells

Fan Zuo , Spencer T. Williams , Po-Wei Liang , Chu-Chen Chueh , Chien-Yi Liao , and Alex K.-Y. Jen*

Dr. F. Zuo, S. T. Williams, P.-W. Liang, Dr. C.-C. Chueh, C.-Y. Liao, and Prof. A. K.-Y. Jen Department of Materials Science and Engineering University of Washington Seattle , WA 98195 , USA E-mail: [email protected] Prof. A. K.-Y. Jen Department of Chemistry University of Washington Seattle , WA 98195 , USA C.-Y. Liao Institute of Polymer Science and Engineering National Taiwan University Taipei 106 , Taiwan

DOI: 10.1002/adma.201401641

stability is especially critical because Pb toxicity hampers the practicality of perovskite-based photovoltaics.

Replacing Pb with Sn is promising since both Sn and Pb belong to the IVA group and have similar ionic radii due to relativistic effects (Sn 2+ 1.35 Å and Pb 2+ 1.49 Å), which may enable substitution without signifi cant lattice perturbation. The Sn-based perovskites have been reported to possess a nar-rower band-gap (as small as ∼1 eV) and much higher charge mobility (10 2 – 10 3 cm 2 /V·s) compared to pure Pb perovksites (10–10 2 cm 2 /V·s). [ 21 ] Although this is encouraging, the PCE of solution processed planar hetrojunction solar cells fabricated from pure tin or mixed metal perovskites are typically lower than those achievable through vapor deposition, primarily owing to poor fi lm morphology and coverage. This low cov-erage can cause severe shorting, poor charge transport, and fast electron/hole recombination; therefore rational control of per-ovskite nucleation and fi lm growth is critical for improving per-formance. Directly addressing the diffi culties associated with controlling the crystallization of solution cast perovskite fi lms is a necessary component of tackling these issues. Relevant studies on this topic are very rare, and little has been done to characterize the effect of Sn alloying on the phase transforma-tion of organolead halide perovskites.

Very recently, Sn-based perovskite solar cells have been reported to demonstrate PCEs around 5–6%. [ 30–33 ] However, all these devices are constructed on a mesoporous TiO 2 scaf-fold, which requires a high temperature (500 °C) sintering process. In order to develop a simple and low temperature solution processed Sn-based perovskite solar cell, we are inter-ested in exploring the planar heterojunction architecture. Herein, we report the synthesis of binary Pb-Sn perovskites (CH 3 NH 3 Pb 1–a Sn a I 3−x Cl x ) and conduct a detailed investigation of the effect of Sn 2+ substitution on perovskite fi lm growth and optoelectronic properties. By optimizing synthetic conditions, we demonstrate high quality perovskite fi lms that can be grown with coverage as high as 97% on PEDOT:PSS coated ITO sub-strate. Sn substitution is also found to red-shift the absorption onset from 800 nm for pure Pb perovskite to 900 nm without signifi cantly affecting absorptivity. This is currently one of the broadest absorption ranges reported for high-performing perovskite-based devices. [ 11,23,30–33 ] We’ve achieved an exciting PCE of 10.1% in a planar heterojunction device based on this system compared to 7.3% in our purely lead based control (CH 3 NH 3 PbI 3−x Cl x ) as a result of the broadened absorption and improved fi lm quality induced by the inclusion of Sn. The impact of enhanced fi lm quality is evident in the greatly improved fi ll factor (0.67 vs. 0.55) and short-circuit current density (19.5 mA/cm 2 vs. 14.5 mA/cm 2 ). To the best of our

Great effort has been made to fi nd promising alternatives for silicon solar cells to drive the widespread implementation of photovoltaic energy production forward. [ 1–6 ] Among the innova-tions in this emerging fi eld, hybrid organolead halide perovskite solar cells have attracted rapidly increasing attention with power conversion effi ciencies (PCE) rocketing to 15.7% in less than one year. [ 7–17 ] Hybrid organolead perovskites are a class of semicon-ductors with the chemical formula AMX 3 (A = organic molecule, M = Pb, X = Cl, Br, and I). [ 18–20 ] Halide anions form a network of corner sharing MX 6 octahedra with Pb 2+ cations located at the octahedrons’ center. Organic cations occupy the cuboctahedral voids between adjacent octahedra ( Figure 1 ). The great interest in using these hybrid perovskites stems from their unique combination of properties that are critical for high photovoltaic performance: 1) direct band gap with gap size ranging from sem-iconducting to metallic, tunable through choice of metals, [ 21,22 ] halogens, [ 11 ] and organic cations; [ 23 ] 2) large dielectric coeffi cient compared to organic polymers leading to small exciton binding energy (∼20 meV), [ 24 ] long diffusion lengths (100–1000 nm), and long lifetimes (∼100 ns); [ 25,26 ] 3) low temperature (<100 °C) solution processability due to a self-assembly driven phase trans-formation; [ 19,27,28 ] and 4) prominent and consistent absorption throughout the visible spectrum. [ 11 ]

There are several key challenges that need to be addressed before this technology becomes feasible for practical applica-tions. The accurate determination of perovskite’s electron/hole mobility, chemical structure and its infl uence on charge transport behavior, precise function of interfaces in device architectures, factors infl uencing stability, and strategies to replace lead as the divalent metal cation without sacrifi cing device performance are the pressing concerns facing the fi eld. [ 29 ] Finding ways to create lead-free perovskites with good

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knowledge, this is the fi rst reported PCE above 10% for solar cells based on Pb/Sn alloy perovskites.

Unlike recently reported high-performance planar hetero-junction perovskite solar cells requiring metal oxide layers and/or vapor deposition, [ 13,34,35 ] our method is a simple solution pro-cess. Regardless of the improved performance, development of this system provides a better understanding of the effect of Pb substitution in organometal perovskites and it may eventually lead to completely Pb-free systems to achieve the cost-effective production of high effi ciency and environmentally benign photovoltaics.

Understanding the effect of Sn on perovskite crystalliza-tion is critical for controlling fi lm morphology and thus device performance. Therefore, we have systematically studied the structure and morphology of organolead perovskites with dif-fering concentrations of Sn. X-ray diffraction patterns of pure and mixed Pb/Sn perovskites, CH 3 NH 3 Pb 1−a Sn a I 3−x Cl x , depos-ited on PEDOTS:PSS/ITO substrates are shown in Figure S1, where a = 0, 0.15, 0.5, 0.75, and 1. Calculated space groups and lattice parameters are summarized in Table 1 .

Within the full range of composition, all samples can be indexed according to the perovskite structure. As a increases from 0 to 1, corresponding to CH 3 NH 3 PbI 3−x Cl x (MAPbX) and CH 3 NH 3 SnI 3−x Cl x (MASnX), a phase transition is clearly observed. It has been reported that pure MAPbX and MASnX have tetragonal I4cm and P4mm symmetry respectively at room temperature. [ 21 ] Figure S2 shows an enlarged view of 22°–25° 2θ. There are two peaks within this range when a ≤ 0.5, which could be indexed to (211) and (202) planes in the tetragonal I4cm space group, agreeing with the reported MAPbI 3 structure at room temperature. [ 36 ] The two peaks gradually disappear and a single peak rises up corresponding to the (113) plane in the P4mm space group when a becomes greater than 0.5 due to

the increased symmetry of P4mm . This phenomenon has also recently been reported by Kanatzidis et al. [ 21 ]

There is a gradual and systematic peak shift to higher 2θ when a changes from 0 to 1, demonstrating that as Sn 2+ concen-tration increases the unit cell size consistently decreases. From the calculated lattice parameters presented in Table 1 , it is clear that higher Sn 2+ content results in a smaller unit cell, which is anticipated since Pb 2+ (1.49 Å) is larger than Sn 2+ (1.35 Å). It is worth noting that unit cell volume only changes slightly in this series, which is consistent with previous reports asserting that perovskite unit cell dimensions are primarily determined by the size of halide anion/s, Cl − , Br − , and/or I − . [ 37 ] Over the entire composition range from one end member to the other, Sn and Pb form a substitutional solid solution without inducing pre-cipitation of a secondary non-perovskite phase. This solubility is instrumental in the success of our binary system.

The XRD patterns of MAPbX, our control, and MAPb 0.85 Sn 0.15 X, the optimum composition for photovoltaic performance, both annealed for 120 min at 90 °C, are compared in Figure 2 a and b. Upon introducing 15 mole% of Sn 2+ , peak intensities signifi cantly increase to ∼30 times those of MAPbX suggesting a dramatic increase in crystallinity. More importantly, unlike the MAPbX pattern which noticeably displays many crys-tallographic planes including (110), (200), (211), (202), (220), (310), (312), (224), (314), (404), and (440), the MAPb 0.85 Sn 0.15 X diffraction pattern is strikingly dominated by the (110), (220), (330), and (440) peaks. This implies that the MAPb 0.85 Sn 0.15 X has signifi cantly greater texture and preferentially grows along <110>. MAPb 0.85 Sn 0.15 X has a tetragonal structure with lattice parameters a = b = 8.871 Å and c = 12.434 Å, close to those reported for MAPbI 3−x Cl x . [ 38 ] It is worth noting that there is only a weak peak at 2θ = 12.66 corresponding to the (001) plane of the PbI 2 , suggesting high phase purity.

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Figure 1. From left to right, SEM image for cross section, device confi guration of planar heterojunction solar cell, and the crystalline structure for MAPb 1−a Sn a X.

Table 1. Elemental analysis and Structural Analysis for MAPb 1−a Sn a X .

Sample ID Sn at% Pb at% Space Group a (Å) b (Å) c (Å)

MAPb 1 Sn 0 X 0 1 I4 cm 8.920 8.920 12.568

MAPb 0.85 Sn 0.15 X 0.12 0.88 I4 cm/P4mm 8.871 8.871 12.510

MAPb 0.5 Sn 0.5 X 0.44 0.56 I4 cm/P4mm 8.838 8.838 12.434

MAPb 0.25 Sn 0.75 X 0.75 0.25 P4 mm 6.258 6.258 6.681

MAPb 0 Sn 1 X 1 0 P4 mm 6.220 6.220 6.708

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Characterization of MAPbX and MAPb 0.85 Sn 0.15 X fi lms with scanning electron microscopy (SEM) supports conclusions drawn from the XRD data. The two top images in Figure 2 c show the morphology of MAPbX and MAPb 0.85 Sn 0.15 X fi lms, respectively. While the MAPbX fi lm shows decent fi lm quality with good continuity and high coverage (∼87%), the MAPb 0.85 Sn 0.15 X fi lm displays far superior quality and coverage (∼97%). This markedly increased coverage ensures excellent connectivity between grains, which is crucial for carrier trans-port. The virtually pin-hole free fi lm mitigates short-circuiting, charge leaking, and large series resistance.

In the bottom images of Figure 2 c, morphological differences between MAPbX and MAPb 0.85 Sn 0.15 X fi lms are clearly evident. The MAPbX fi lm is composed of interconnected nanoscale domains with sizes ranging from 200 nm to 500 nm with no obvious faceting. In the MAPb 0.85 Sn 0.15 X fi lm, however, we observe strong faceting with comparatively smooth surfaces sug-gesting enhanced, and potentially more defect free, crystal devel-opment. Apparent domain size is signifi cantly larger (500 nm – 1 µm) than that expressed in the MAPbX fi lm suggesting an alteration of transformation kinetics. Regular faceting and orientation of the MAPb 0.85 Sn 0.15 X crystallites reveal texture suggested by the XRD data, demonstrating that {110} planes are preferentially developed and exposed. Regular faceting sur-rounding interconnected features suggests that the fi lm is com-posed of an oriented network of 500 nm to 1 µm single crystals.

To elucidate the growth kinetics of two fi lms, we have studied their morphologies and phase compositions at different annealing times with SEM ( Figure 3 ) and XRD (Figure S3). After spin-coating the precursor solutions onto the substrate,

ions in the precursor solutions immediately begin to self-assemble into the perovskite structure as is evidenced in the XRD patterns for both MAPb 0.85 Sn 0.15 X and MAPbX measured before annealing (Figure S3). The alteration in transforma-tion kinetics induced by tin inclusion is evident in the massive increase in the (110) refl ection 14.2° 2θ after being annealed for 30 min in MAPb 0.85 Sn 0.15 X relative to MAPbX, which indicates more rapid transformation. Peaks corresponding to both iodide and chloride based organometal perovskites (at 2θ = 14.2° and ∼15.9°, respectively) [ 13,38 ] are evident in MAPb 0.85 Sn 0.15 X and MAPbX at initial stage and after 30 min of annealing while only the refl ection corresponding to the iodide based organometal perovskite remains after annealing for 120 min.

This indicates that in addition to being exist in the fi nal iodide based perovskite as substitutional impurities, chloride ions also alter the transformation pathway during annealing by introducing the competing chloride based perovskite phase. This alteration in growth pathway likely infl uences overall transformation kinetics in addition to modulating the resulting morphology as compared to pure lead systems, but as of now the exact mechanism by which chloride inclusion enhances ultimate material properties is inadequately understood and an area of active research in the fi eld. [ 25,38,39 ] It is understood from current literature that chloride ions preferentially occupy axial positions in the mixed I/Cl lead octahedra within the perovskite lattice while iodide ions may occupy both equatorial and axial positions, [ 13,38 ] however it is diffi cult for us to make any conclu-sion to this end in this work. It seems apparent that tin inclu-sion suppresses the formation of the chloride based perovskite lattice by virtue of the decreased peak height at ∼15.9° 2θ at

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Figure 2. (a) and (b) XRD spectra for MAPbX and MAPb 0.85 Sn 0.15 X fi lms respectively. (c) SEM images for MAPbX and MAPb 0.85 Sn 0.15 X fi lms. (d) EDS elemental maps of Pb and Sn in MAPb 0.85 Sn 0.15 X.


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both 0 and 30 minutes in MAPb 0.85 Sn 0.15 X relative to MAPbX (Figure S3).

Even before annealing, MAPb 0.85 Sn 0.15 X and MAPbX show great morphological differences (Figure 3 ). Holes in both fi lms may result from the evaporation of dimethylformamide (DMF) during spin-coating. Compared to the MAPbX sample, contrast features apparent in the MAPb 0.85 Sn 0.15 X fi lm suggest the for-mation of larger ordered domains, again implying a more rapid transformation via self-assembly. Even in this as cast state, the MAPb 0.85 Sn 0.15 X fi lm demonstrates greater coverage and conti-nuity, possibly due to more favorable interfacial energy in addi-tion to the faster assembly. After annealing at 90 °C for 30 min, the differences between the systems become more pronounced. MAPb 0.85 Sn 0.15 X crystals begin to show well-defi ned facets with smooth surfaces, sharp edges, and preferential orientations not apparent in the MAPbX sample. After 120 min of annealing, the two fi lms fully develop. The MAPb 0.85 Sn 0.15 X perovs-kite crystals fuse and grow together to share certain facets (Figure 2 c), signifi cantly depressing crystallite boundaries. On the contrary, the MAPbX crystals do interconnect but retain more voids and a greater degree of discontinuity. The rapid per-ovskite conversion rate for MAPb 0.85 Sn 0.15 X relative to MAPbX is also evidenced by the time-resolved XRD study as discussed above and shown in Figure S3.

The effect of composition on morphology is shown in Figure S4. As the mole fraction of Sn increases, domains become larger. Surfaces of all Sn-containing fi lms are relatively smooth, but sharp edges and corners disappear along with the widening of voids at concentrations greater than 15%, possibly because of increasingly rapid transformation and unfavorable interfacial energy.

Figure 2 d shows an elemental map of MAPb 0.85 Sn 0.15 X obtained via SEM energy-dispersive X-ray spectroscopy (EDS). To avoid any interference from Sn in the ITO, these fi lms are spin-coated on bare glass coated with PEDOT:PSS. Sn and Pb are homogeneously distributed throughout the fi lm with no apparent phase separation, which further supports the pro-posed solid solution. EDS analysis shows the real Pb:Sn ratio to be 88:12 in the MAPb 0.85 Sn 0.15 X fi lm, which is close to the

theoretical 85:15 ratio given by stoichiometry of the precursor solutions. The EDS results for all perovskite compositions are summarized in Table 1 . All perovskites exhibit reasonable con-sistency between real and theoretical Pb:Sn ratios.

Chemical differences between Sn 2+ and Pb 2+ explain Sn’s signifi cant effect on nucleation and growth behavior. Their dif-ferences in size results in differing affi nities toward halogen anions, which is refl ected by bond lengths and energies listed in Table S1. Sn-Cl has the largest bond energy (350 kJ/mol), which suggests that during nucleation Sn-Cl bonds form fi rst fol-lowed by Pb-Cl, Sn-I, and Pb-I bonds. Although this is a kinetic argument derived from thermodynamic data, the hypothesis is supported by our EDS analysis of MAPb 0.85 Sn 0.15 X at differing annealing times. Before annealing, the Pb:Sn and I:Cl ratios are 2.3:1 and 1:1, respectively, which are much lower than the same ratios produced after 2 h of annealing: 7.3:1 and 5.8:1. After annealing for 30 min, the Pb:Sn and I:Cl ratios reach intermediate values of 5.6:1 and 3.2:1 respectively, demon-strating a consistent trend.

It should be noted that it may appear from this data as if tin is being steadily evolved as a volatile species during annealing, but the initial Pb:Sn ratio noted above is artifi cially high. From this EDS data, at the beginning of the annealing pro-cess the composition appears to be ∼30 at.% Sn with respect to Pb, which is much higher than the stoichiometry of the pre-cursor solution used (15 at.% Sn). The fi nal composition after annealing for 2h is ∼12 at.% Sn which is within the reasonable range compared to the initial solution composition, considering errors that may be introduced by surface topology and chem-ical environment in EDS characterization. Additionally, a sig-nifi cant loss of chloride is generally observed by both us and others because of its relatively low concentration allowed in the iodide based organolead perovskite lattice (∼4%). [ 38 ] Since we have observed similar degree of chloride loss in both MAPbX and MAPb 1−a Sn a X systems, we think it should not be caused by the presence of tin. With regard to the loss mechanism, it has been found in the pure organolead iodide perovskite system that iodide is lost through the decomposition of meth-ylammonium iodide into volatile methylamine and hydroiodic

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Figure 3. Top images show the morphologies of MAPb 0.85 Sn 0.15 X fi lms for annealing times of 0 min, 30 min, and 120 min. Bottom images show the morphologies of MAPbX in the same manner. The scale bars indicate 5 µm.


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ON acid. [ 40 ] We surmise that an analogous loss mechanism is rel-

evant for excess chloride in both MAPbX and MAPb 1−a Sn a X and we present this in Scheme S1. Thus, we conclude the trend of morphological changes should not be relevant to the loss of chloride since it is a common phenomenon observed in all systems.

To appreciate the implications of these compositional trends, it is necessary to understand the physical mechanism behind SEM EDS, which is shown in Figure S5. The generated char-acteristic X-ray emission enables elemental analysis. The pear-shaped character of the detection volume in EDS means that this technique acquires more signals from a sample’s bulk than its surface. Note that the electron accelerating voltage in our EDS measurement is 15 keV, which means a penetration depth of few micrometers, much greater than the thickness of our perovskite fi lms (350–450 nm). Therefore, it is reason-able to propose that the elemental ratios reported above and in Table 1 are most representative of the composition at the base of the growing perovskite fi lm, an event that is preferentially initiated at the substrate/perovskite interface. Summarizing the above data with this in mind, we observe an abnormally high concentration of Sn and Cl at the substrate interface in the as cast specimens which we thought it is due to initiation of the nucleation process, suggesting a Sn-Cl bond generation event before other bonds formation. This explains why the bulk of the sample in Figure 3 a has a much higher concentration of Sn 2+ and Cl − than the theoretical value.

All the above demonstrates the unique infl uence of Sn 2+ on nucleation and growth behavior of these hybrid perovskites.

It regulates perovskite nucleation rate and ensuing crystal development. At the optimal Sn 2+ concentration a high quality and almost completely continuous fi lm is achieved.

Earlier, Kanatzidis et al. reported that the introduction of Sn into organolead halide perovskites narrows the band gap, red-shifting the absorption onset. [ 21 ] Band-gaps of pure iodide end members MAPbI and MASnI are 1.6 and 1.2 eV, respectively. Precisely determining the energy levels of MAPb 1–a Sn a I 3 is still challenging and the reported values so far are quite scat-tering, [ 30–33 ] potentially due to the ease of oxidation of the Sn-containing samples. Regardless of this, the band gap of our resulting perovskites match well with their proposed electronic structure. Figure 4 a compares UV-visible absorption spectra for the mixed halide MAPbX and MAPb 0.85 Sn 0.15 X. Upon partial substitution with Sn, the absorption edge red-shifts to ∼900 nm. More importantly, the absorption from 600 nm to 800 nm is greatly enhanced, demonstrating the binary-metal system's potential for expanding light harvesting capacity of current perovskites. The absorption spectra for all compositions are plotted in Figure S6 and a diminishing return is apparent upon increasing Sn 2+ substitution signifi cantly beyond 15%, resulting in dramatically decreased absorption below ∼700 nm due to the low absorption coeffi cient of Sn. [ 21 ]

Photovoltaic performance of devices fabricated with perovs-kites of all compositions studied is presented in Table 2 and external quantum effi ciency (EQE) spectra for devices fabricated with MAPbX and MAPb 0.85 Sn 0.15 X are shown in Figure 4 b. From 330 nm to 800 nm, MAPbX exhibits an average EQE of ∼60%. The EQE range of MAPb 0.85 Sn 0.15 X extends up to 900 nm,

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Figure 4. (a) UV-vis absorption spectra, (b) EQE results and (c) J – V curves for MAPbX and MAPb 0.85 Sn 0.15 X measured under simulated AM1.5 illumina-tion and in the dark, (d) The relationship between 1/ R s and light intensity for MAPbX and MAPb 0.85 Sn 0.15 X devices.


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which is one of the broadest EQE ranges achieved in high-performance hybrid perovsktie solar cell. [ 11,23,30–33 ] Peak EQE exceeds 80% near 450 nm, and across the entire range tested (330 nm to 900 nm) the MAPb 0.85 Sn 0.15 X device shows marked enhancement compared to the MAPbX device.

Figure 4 c shows current density-voltage ( J − V ) curves meas-ured under simulated AM1.5 illumination (100 mW/cm 2 ) and in the dark, respectively. Both perovskites show good diode behavior with relatively low leakage current in the dark J−V curves. The MAPb 0.85 Sn 0.15 X perovskite solar cell exhibits a short-circuit current dentsity ( J sc ) of 19.5 mA/cm 2 , an open-circuit voltage ( V oc ) of 0.77 V, and a fi ll factor (FF) of 0.67, yielding an overall power conversion effi ciency (PCE) of 10.1%. Relevant results for the MAPbX perovskite solar cell are J sc = 14.5 mA/cm 2 , V oc = 0.92 V, and FF = 0.55. The MAPb 0.85 Sn 0.15 X device shows a remarkably improved J sc , which agrees well with its broadened absorption and superior EQE. In addition to the substantial increase in J sc , Sn substitution enables a signifi cant FF improvement from 0.55 to 0.67. As both devices are fabri-cated with the same architecture, electrode and base contact resistances are consistent. Thus, the increase in FF points to a decrease in internal and interfacial resistances of the perovskite absorber in MAPb 0.85 Sn 0.15 X due to the greatly improved per-ovskite fi lm quality and coverage as compared to the pure lead based control.

As evidenced in the results from XRD and SEM, the MAPb 0.85 Sn 0.15 X fi lm possesses much enhanced crystallinity and surface coverage compared to those from the pure Pb sample, suggesting a more ordered arrangement of atoms and less defects. Shown in Figure 4 d is the plot 1/ R s vs. light inten-sity, where the series resistance ( R s ) is estimated from the slope of the J−V curve near V oc . As can be seen, MAPb 0.85 Sn 0.15 X has a steeper slope than MAPbX, suggesting that carrier transport is signifi cantly improved via partially replacement of Pb 2+ with Sn 2+ . It can be speculated that the improved fi lm coverage and homogeneity enable more effective exciton dissociation and facilitate the charge transport across the interfacial layers to the electrodes while the poor fi lm formation and coverage will induce severe charge recombination. When a ≥ 0.5, V oc dra-matically decreases. Takahashi et al. has attributed this to Sn 2+ oxidation and the associated p-type doping. [ 22 ] Moreover, poor fi ll factors with a ≥ 0.5 can be easily explained by consistently inferior fi lm quality as is evident in Figure S4.

In conclusion, we have demonstrated the success of par-tial Sn substitution in organolead perovskites to simultane-ously reduce Pb content and increase performance through broadened absorption and improved solution-cast fi lm

morphology. Time-resolved SEM study of perovskite growth during annealing indicates the dynamic morphological modu-lation offered by Sn 2+ incorporation. A high coverage (∼97%) fi lm with excellent continuity can be grown because of Sn’s effect on nucleation and growth. The markedly improved morphology with well-ordered and relatively large crystalline domains suppresses charge recombination and improves trans-port, resulting in an increased fi ll factor. UV-visible absorption measurements demonstrate the 100 nm red-shift of the binary-metal perovskite’s absorbance onset. The EQE study further verifi es both the red-shifted absorption and effi ciency enhance-ment that spans the entire characterized range from 330 nm to 900 nm. The enhancement in solar light absorption contributes to the large J sc as do the dramatic improvements in fi lm quality uniquely enabled in this binary system. By optimizing Sn con-centration in MAPb 1−a Sn a X, a very promising PCE of 10.1% is achieved. To compliment these fi ndings, the rigorous correla-tion between domain size and perovskite optoelectronic proper-ties will be elucidated in the near future.

Our results are quite encouraging considering that partial Sn substitution not only reduces the content of Pb in hybrid per-ovskite photoactive layers, mitigating their ecological impact, but also enables higher PCE compared to the purely lead based perovskite devices, enhancing their economic viability. This pre-liminary study establishes the foundation for further exploring the benefi ts of alloy perovskite systems to replace lead while simultaneously increasing performance to make these hybrid photovoltaics competitive with Si-based technologies.

Experimental Section Materials and Sample Preparation : Methylammonium iodide (MAI)

was synthesized by reacting 24 mL, 0.20 mole methylamine (33 wt% in absolute ethanol, Aldrich), 10 mL, 0.04 mole hydroiodic acid (57 wt% in water with 1.5% hypophosphorous acid, Alfa Aesar), and 100 mL ethanol in a 250 mL round bottom fl ask under nitrogen protection at 0 °C for 2 h with magnetic stirring. After reaction, the white precipitate of MAI was recovered by rotary evaporating solvents at 40 °C then dissolved in ethanol and sedimentated in diethyl ether by stirring the solution for 30 min. This step was repeated three times then the MAI powder was collected and dried at 50 °C in a vacuum oven for 24 h. To prepare perovskite precursor solution, MAI and lead chloride (PbCl 2 , 98%, Aldrich) powder were mixed in anhydrous dimethylforamide (DMF, Aldrich) with a molar ratio of 3:1. For the Sn-based perovskite precursor solution, MAI, PbCl 2 , and Tin Chloride (SnCl 2 , 98%, Aldrich) powder were mixed with desired molar ratio. The solutions (40 wt%) were stirred overnight at 80 °C and fi ltered through 0.45 µm PVDF fi lters before device fabrication.

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Table 2. Performance of the studied solar cells under AM 1.5G Illumination (100 mW/cm 2 ).

V oc (V) FF J sc (mA/cm 2 ) PCE (%)

Pure Pb perovskite * 0.91 ± 0.02 0.52 ± 0.025 13.3 ± 1.0 7.01 ± 0.22 (7.33)

15% Sn perovskite * 0.76 ± 0.01 0.66 ± 0.008 19.1 ± 0.2 9.77 ± 0.25 (10.10)

50% Sn perovskite 0.24 ± 0.03 0.53 ± 0.010 0.7 ± 0.1 0.11 ± 0.02

75% Sn perovskite 0.06 ± 0.02 0.25 ± 0.035 3.5 ± 0.8 0.06 ± 0.03

100% Sn perovskite 0.03 ± 0.02 0.25 ± 0.016 5.3 ± 2.0 0.04 ± 0.01

*The number in the parenthesis is the highest PCE.


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ON Fabrication of the Thin Film Perovsktie Devices : The thin fi lm perovskite

devices were fabricated with the confi guration: indium tin oxide (ITO)/poly(3,4-ethylenedioxythiophene):poly(p-styrene sulfonate) (PEDOT:PSS, CLEVIOSTM P VP Al 4083)/CH 3 NH 3 Pb 1−a Sn a I 3−x Cl x /[6, 6]-phenyl-C61 -butyric acid methyl ester (PCBM, American Dye Source Inc.)/fullerene surfactant (C60-bis)/Ag. ITO glass substrates (15 ohm/sq) were cleaned sequentially by sonication in detergent and deionized water, acetone, and isopropanol for 10 min. After drying under N 2 stream, substrates were further cleaned by exposing to plasma for 10 s. A PEDOT:PSS (Baytron P VP Al 4083 fi ltered through a 0.45 µm Nylon fi lter) hole-transporting layer with a thickness of ∼45 nm was spin-coated onto ITO substrates at 5k rpm for 30 s and annealed at 150 °C for 10 min in air. The substrates were then transferred into a N 2 -fi lled glove box, spin-coated with prepared precursor solution at 6k rpm for 45 s, and annealed at 90 °C for 2 h. The thickness of the studied perovskite thin fi lms is around 350–450 nm, determined by surface profi ler (Dektak 3030). The electron-transporting layer (PCBM, 15 mg/mL in chloroform) and C 60 -bis surfactant (2 mg/mL in isopropyl alcohol) was then sequentially deposited by spin coating at 1 k rpm for 60 s and 3k rpm for 60 s, respectively. Silver electrodes with a thickness of 150 nm were deposited under high vacuum (<2 × 10 −6 Torr) through a shadow mask, defi ning a device area of 3.14 mm 2 , by thermal evaporation. All J−V curves in this study were recorded using a Keithley 2400 source meter unit. The device photocurrent was measured under AM1.5 illumination condition at an intensity of 100 mW/cm 2 . The illumination intensity of the light source was accurately calibrated with a standard Si photodiode detector equipped with a KG-5 fi lter, which can be traced back to the standard cell of the National Renewable Energy Laboratory (NREL). The EQE spectra performed here were obtained from an IPCE setup consisting of a Xenon lamp (Oriel, 450 W) as the light source, a monochromator, a chopper with a frequency of 100 Hz, a lock-in amplifi er (SR830, Stanford Research Corp), and a Si-based diode (J115711–1-Si detector) for calibration.

Supporting Information Supporting Information is available from the Wiley Online Library or from the authors.

Acknowledgements The authors thank the support from the Air Force of Scientifi c Research (FA9550–09–1–0426), the Offi ce of Naval Research (N00014–11–1–0300), the Asian Offi ce of Aerospace R&D (FA2386–11–1–4072), the Boeing Foundation, and the National Science Foundation Graduate Research Fellowship Program under Grant No. DGE-1256082. Part of this work was conducted at the University of Washington NanoTech User Facility, a member of the NSF National Nanotechnology Infrastructure Network (NNIN).

Received: April 11, 2014 Revised: May 30, 2014

Published online: August 14, 2014

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