B-Site Metal Cation Exchange in Halide Perovskites - ACS Publications

May 2, 2017 - The activation energies for defect-mediated diffusion of the A- and X-site components have been calculated to be lower than those for B-...
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B‑Site Metal Cation Exchange in Halide Perovskites Giles E. Eperon and David S. Ginger* Department of Chemistry, University of Washington, Seattle, Washington 98105, United States S Supporting Information *

ABSTRACT: We demonstrate exchange of the B-site metal cation in hybrid organic−inorganic halide perovskite thin films. We exchange tin in formamidinium tin triiodide (CH(NH2)2SnI3, or FASnI3) with lead at controllable levels, forming CH(NH2)2SnxPb1−xI3 alloys with partial substitution and fully converting the film to CH(NH2)2PbI3 with a large excess of Pb2+. We observe no evidence for phase segregation or bilayered films, indicating that conversion is uniform throughout the film. This facile technique provides a new way to control composition independently from the crystallization processes, allowing formation of the black phase of CH(NH2)2PbI3 at much lower temperatures than those previously reported while also opening the door to new morphology−composition combinations. The surprising observation that the B-site metal cations are mobile may also provide insight into the nature of transient processes in these materials, suggesting that they may be involved in ionic conduction, and will be a critical consideration for long-term stability.

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containing the desired cation or anion, either driven by diffusive mass action or by a favorable chemical reaction. Ion exchange is often used to retain the morphology of a material while changing the composition, thereby forming a product that is not otherwise attainable via straightforward synthesis. Ion exchange is also used to create alloys with interesting properties and to introduce dopant ions with careful control over the dopant density.28−30 Halide perovskites have an ABX3 crystal structure where A is a 1+ cation (usually formamidinium, methylammonium, or Cs), B is a 2+ metal cation (usually Pb or Sn), and X is a 1− halide anion (iodide, bromide, chloride). The activation energies for defect-mediated diffusion of the A- and X-site components have been calculated to be lower than those for B-sites, indicating that the A- and X-sites should be most easily exchanged.9,14,31,32 Indeed, rapid X-site halide exchange has been demonstrated and is now widely used in both thin films and nanocrystals of halide perovskites.15−18 A-site cation exchange has a higher activation energy, occurring on a somewhat slower time scale than X-site anion exchange,19 and several groups have recently reported the use of A-site exchange for compositional engineering.19−22 However, we are unaware of any reports of complete B-site metal cation exchange on thin films to date, perhaps because the B-site binding energy is larger. Exchange of the metal cation could potentially be the most important and

alide perovskites, solution-processable semiconductors with an ABX3 crystal structure, have sprung to the forefront of optoelectronic materials research in the past few years, with the demonstration of highly efficient solar cells, light-emitting diodes (LEDs), and other devices.1,2 However, fundamental understanding of these materials lags behind the fabrication of efficient devices, and only recently are many of the properties underpinning the impressive performance of these materials coming to light. One such property is the fundamental “softness” of the perovskite lattice; recent work has shown that the crystal lattice shows significant motion at room temperature.3−6 This softness is in stark contrast to more “classical” high-quality optoelectronic materials such as Si and GaAs, which have relatively static lattices. The soft crystal lattice has several ramifications for device applications. First, this property means that it is relatively easy for constituent ions or defects in the lattice to move around.7−9 This motion has been used to explain the presence of electronic hysteresis in the current−voltage response of solar cells as ions move to compensate an applied field, temporarily changing the resultant energy landscape in the device.10−13 A second implication is that the activation energy for removal or replacement of lattice ions is low.9,14 As well as having implications for long-term stability, this lattice softness has been leveraged to carry out ion exchange reactions of both the A-cations and X-anions with compatible species.15−22 Ion exchange reactions are a well-studied concept in the field of III−V and II−VI semiconductor nanocrystals.23−27 Ion exchange can occur upon exposure of nanocrystals to a solution © XXXX American Chemical Society

Received: April 3, 2017 Accepted: April 24, 2017

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ACS Energy Letters useful ion exchange process because the metal imparts the majority of the properties of the material and is the most interesting to replace with dopant ions or to alloy with. In this Letter, we report complete metal site cation exchange in halide perovskite thin films. We design a reaction protocol allowing us to exchange tin for lead in a controlled manner in CH(NH2)2SnI3 films to form CH(NH2)2(Pb,Sn)I3 and CH(NH2)2PbI3. This technique provides a new way to control material composition that is detached from the often complex processes involved in crystallization and should enable facile and controllable doping of perovskite materials with other metal ions. It also suggests that the metal ions may not be exempt from participation in ionic motion under electronic bias and may contribute to slow transient effects, having implications for long-term stability under operating conditions. We find that the alloyed Pb−Sn materials show a previously reported anomalous bandgap bowing in thin films.33,34 We studied cation exchange in thin films, most relevant for photovoltaic and LED device applications, and focused on the exchange between Sn2+ and Pb2+. Both Sn2+ and Pb2+ are known to both form perovskite structures and are similar enough in size to form an alloy.35,36 In order to carry out B-site cation exchange, we sought a solvent that would solubilize the appropriate metal salt yet for which the resulting solution would not dissolve the perovskite material. Following the use of phosphines to passivate perovskite surfaces,37 we found that trioctylphosphine (TOP) was able to dissolve PbI2 and SnI2 in reasonable concentrations and that these solutions did not dissolve the perovskite films upon exposure. Neat TOP, on the other hand, appeared to dissolve Sn-based materials within minutes but not Pb-based ones, as we show in the Supporting Information (SI) (Figure S1); we will discuss implications of this difference later. To carry out the cation exchange, we immersed CH(NH2)2MI3 (M = Pb,Sn) perovskite thin films, fabricated via a previously reported precursor-phase antisolvent immersion method (PAI),34 in solutions of MI2 dissolved in TOP and diluted in toluene at varying concentrations, as illustrated in Figure 1a. We observed a gradual color change in the exchange of FASnI3 with PbI2, progressing slowly over the course of days when heated at 70 °C. After ∼72 h, no further color change was observed by eye and no more changes in the absorbance spectra of the films were observed. Figure 1b shows the absorbance of CH(NH2)2SnI3 thin films when exposed to Pb2+ of varying concentrations, denoted as the equivalent number of moles of Pb in solution to the Sn in the film, for a sufficient time for the absorbance spectrum to have reached a stable state, and Figure 1c shows the corresponding photoluminescence (PL) spectra for some of the compositions. The initial FASnI3 films showed a broad absorption onset and PL peak at ∼900 nm. We found that when adding only a small amount of Pb to the solution the film converted mainly to a yellow−orange phase with low absorption, as we show in the SI (Figure S2). However, when we added higher concentrations of Pb (1−10 mol equiv of Pb, determining the moles of Sn from the number of Sn atoms in the volume of the film and using concentrations such that 1×Pb = 0.1 mM Pb2+) to the solution, the exchange process induced a clear red shift in both the absorption onset and PL spectrum. Interestingly, at higher concentrations of Pb (corresponding to a roughly 50×Pb/Sn ratio, 5 mM Pb2+), the converted films showed a significantly blue-shifted absorption onset and PL peak compared to the original film. Tin−lead alloys have been reported to show an anomalous bandgap bowing, with intermediate compositions

Figure 1. B-site cation exchange in thin films. (a) Illustration of the cation exchange process. (b) Absorbance spectra of CH(NH2)2SnI3 thin films reacted with the equivalent molar concentration of Pb shown in the legend. Labels refer to the relative molar amount of Pb in solution compared to the Sn in the film, that is, “1×Pb” refers to an addition of an equivalent number of moles of Pb as Sn in the film. Films were heated at 70 °C for 96 h in the exchange solution. CH(NH2)2PbI3 and CH(NH2)2Sn0.5Pb0.5I3 spectra were measured on non-ion exchanged films of those materials, formed directly from the appropriate solution, as references. (c) Photoluminescence (PL) spectra for three of the films shown in (a). (d) Plot of the estimated bandgap as a function of the relative Pb addition concentration, also showing bandgaps of the reference nonexchanged materials as gray regions.

having a lower bandgap than either the pure Pb or Sn end points.36,38 The red-shifted absorption spectrum that we observe at intermediate Pb concentration appears similar to the absorbance of a reference film of CH(NH2)2Pb0.5Sn0.5I3, fabricated via PAI, though we note that the bandgap of CH(NH2)2PbxSn1−xI3 is fairly insensitive to the x value over the range x = 0.1−0.7,34 (we determine exact Pb/Sn ratios independently below). The blue-shifted spectrum obtained at high (50×) Pb concentration is consistent with CH(NH2)2PbI3. We conclude that partial Sn−Pb substitution is achieved with 1−10× molar equivalents and full Sn-to-Pb conversion is attained with the larger 50× excess. Figure 1d shows the estimated bandgap of the films, extracted from the absorbance via Tauc plots, as compared to the measured bandgaps of CH(NH2)2PbI3 and CH(NH2)2Sn0.5Pb0.5I3 reference films, showing that the 1−10× Pb addition results in a material with a bandgap similar to that of CH(NH2)2Sn0.5Pb0.5I3 and the 50× excess has a bandgap similar to that of CH(NH2)2PbI3. In order to confirm the composition of the cation-exchanged materials, we carried out X-ray diffraction (XRD) and elemental analysis via Glow Discharge Optical Emission Spectroscopy (GDOES). Figure 2a shows XRD spectra of films equivalent to those in Figure 1, with magnification of the peak centered near ∼14° in Figure 2b and that near 28° in Figure 2c. These peaks 1191

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Figure 2. XRD spectra of the CH(NH2)2SnI3 films exchanged with varying concentrations of PbI2. (a) Complete spectra showing the whole region of interest. (b,c) Magnifications of the peak centered at around 14° (b) and 28° (c). Peaks assigned a # correspond to peaks from PbI2, and peaks assigned a * correspond to the ITO substrate.

experimental details). We find that the 1× and 10×Pb treatments result in films with similar lead incorporation, at 63 ± 2 and 65 ± 2%, respectively, and that the 50×Pb treatment results in 94 ± 4% Pb. The compositions determined via atomic emission agree well with the XRD peak positions. The data also confirm that with the 50×Pb treatment we replace almost all of the Sn, except trace residue, in the film to form CH(NH2)2PbI3 and that the compositions of the intermediates are CH(NH2)2PbxSn1−xI3, where x = 0.63 and 0.65 for the 1× and 10×Pb excess, respectively. The determined bandgaps agree well with literature values for these compositions, as we show in the SI (Figure S3). We take the fact that the XRD spectra agree well with compositions inferred from the elemental analysis as evidence that most of the Sn or Pb detected comprises the crystalline films, rather than noncrystalline residuals at the grain boundaries or surfaces. The complete conversion may indicate that the TOP plays a significant role in the cation exchange reaction because van der Stam et al. reported that they were unable to attain complete Bsite cation exchange in nanocrystals of the purely inorganic perovskite CsPbBr3 while using solutions of B-site cations in oleylamine and toluene.40 To begin to elucidate a mechanism for the conversion reaction, we studied the morphology of the converted films. The conversion could take place either via a morphologypreserving ion exchange reaction or via dissolution and reprecipitation of the new material. Scanning electron microscope (SEM) images of the top surface are shown in Figure 3. While there is a small amount of damage observed in the exchanged films, the general structure of the films, including the grain size, remains intact. Some precipitation of PbI2 (higher contrast) is observed on the surface, consistent with XRD data. Measurements of film thicknesses and cross-sectional SEMs

correspond to the (111) and (222) peaks of trigonal (P3m1) CH(NH 2 ) 2 PbI 3 and the (100) and (200) peaks of orthorhombic (Amm2) CH(NH2)2SnI3. The XRD spectra of our reference films agree well with literature reports.35,39 As shown in Figure 2a, we observe only a single perovskite phase for all films, with no evidence for two separate CH(NH2)2PbI3 and CH(NH2)2SnI3 phases, which would appear as split peaks. In the films with higher Pb concentration, the peak at 12.6°, associated with PbI2, indicates that there is likely some precipitation of PbI2 out of solution. Magnification of the two most prominent perovskite peaks, shown in Figure 2b,c, reveals that as the Pb concentration increases above 0.2×, the peaks shift from the CH(NH2)2SnI3 position toward CH(NH2)2PbI3, with 50×Pb excess overlaying the CH(NH2)2PbI3 spectrum well and 1× and 10×Pb lying in between the two end compositions. Previous literature suggests that Sn−Pb alloys have lattice parameters intermediate to the end point compositions, as we observe, reinforcing our hypothesis that we form crystalline alloys here.33,34 To determine the exact composition of the films, we determined the lead−tin percentage via GDOES, as shown in Table 1 (see the SI for Table 1. Tin−Lead Composition Determined from GDOES for the Neat and Exchanged Filmsa film/treatment FASnI3 1×Pb 10×Pb 50×Pb

Pb/(Pb + Sn) % 0 63 65 94

± ± ± ±

2 2 2 4

a

Errors are determined from repeat measurements and instrument noise. 1192

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Figure 3. SEM images of the surface of pristine and exchanged films. The white features on the surface are assumed to be precipitated lead iodide. The scale bar is 1 μm in all cases.

reduced bandgap with only a few lattice unit cells, but it may be larger in practice, with a large enough volume needed to ensure that there are sufficient regions with a low bandgap to actually affect the bulk properties. While we do not conduct an extensive investigation into this effect here, we speculate that the apparent lack of band bowing during metal cation exchange with the nanocrystals could provide an experimental route by which to probe the size of the crystalline domains required to achieve the anomalous low bandgaps. In this case, our results would suggest that the relevant size must be at least several nm because the nanocrystals studied herein contain on the order of tens to hundreds of unit cells, being between ∼2 and 6 nm diameter. Another possibility is that the increased surface area/ volume ratio present may be relevant in “locking in” a particular Sn−Pb arrangement from the surface, which may not exhibit the lower bandgap. We suggest that future work on better sizecontrolled samples of nanocrystals could provide more definitive answers to these questions. Having observed significantly faster Sn−Pb exchange reactions in nanocrystals than thin films, we sought to reverse the reaction in nanocrystals. Upon adding a large molar excess of SnI2/TOP to the finished exchange reaction, we found that after 12 h of stirring at 70 °C the absorbance and PL had redshifted once more, as we show in Figure S9. These data suggest that we have been successful in reversing the cation exchange and restoring the original composition. Elemental analysis (Table ST2) is in agreement. As mentioned earlier, addition of pure TOP appears to dissolve tin-based perovskites rapidly but not lead-based materials. The general structure of tin-based thin films is preserved upon addition of the PbI2−TOP solution. This observation allows us to speculate about the mechanism for ion exchange and explain the difference in rates of exchange between Sn and Pb. If the TOP favors removal of the Sn from the lattice, as suggested by the solubility of Sn-based perovskites, this solubility would make the substitution of Sn ions easier in the presence of TOP. If there are Pb ions present in solution, these could then replace the removed Sn ions, preventing extensive morphological damage of the film. In other words, we suggest that TOP acts to facilitate the replacement of Sn with Pb, much as differences in cation solubility in tributylphosphine have been proposed to facilitate replacement exchange of Cd2+ for Ag+ when interconverting CdSe and Ag2Se nanocrystals.43 Once some Pb replacement has occurred, its reduced solubility in TOP should protect the

(see SI Table ST1 and Figure S4) indicate that there is effectively no change in film thickness, which remains at ∼200 nm, within error margins. This conservation of the film structure and thickness suggests that the exchange occurs via a morphology-conserving ion exchange, rather than dissolution and reprecipitation, which would be likely to incur significant changes to the grain structure, thickness, and roughness. While we were able to replace Sn with Pb in thin films, we were unable to achieve exchange in the opposite direction, from Pb to Sn, on reasonable time scales. Even with very large excesses of Sn in solution, we observed no obvious change after days of reaction (SI Figure S5), suggesting that Sn replacement with Pb is more favorable. We thus sought a way to further understand the replacement of Sn for Pb in thin films, and we turned to explore perovskite nanocrystals as ion exchange reactions in II−VI nanocrystals have been reported to occur in seconds for reactions that do not proceed appreciably, if at all, in microscale or bulk samples.27We fabricated quantumconfined CH3NH3SnI3 (MASnI3) nanocrystals by adapting the methods of Jellicoe et al.41 and Vybornyi et al.,42 as detailed in the SI. We then carried out B-site cation exchange on these nanocrystals in the same way as for thin films, stirring dilute solutions at 70 °C with PbI2/TOP solutions added. We observed that the color of the nanocrystal solutions changed much faster than did the thin films, with the reaction appearing to be fully complete within ∼30 min. As we show in the SI (Figure S6), we observed a monotonic blue shift in the absorbance and PL when adding PbI2/TOP to a MASnI3 nanocrystal solution. XRD and elemental analysis (Figure S6 and Table ST2) confirm that exchange from CH3NH3SnI3 nanocrystals to a composition much closer to that of CH3NH3PbI3 has taken place. During the exchange, we observe an increase in the average size of the nanoparticles, which we attribute to Ostwald ripening under heating (SI Figure S7). We also note that, in contrast to the thin-film spectra, the absorbance of the nanoparticles during cation exchange from Sn to Pb appears to blue shift monotonically (despite the increase in size), with no evidence of the initial red shifting expected from the Pb−Sn alloy’s anomalous band bowing. The exact mechanism responsible for the origin of this compositional band bowing in thin films is still under debate,33 but one theory is that a particular repeated arrangement of tin−lead unit cells in the lattice can result in the red-shifted bandgap.34 However, the necessary size of the motif has not been experimentally determined; theoretical calculations find a 1193

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layer over time, the material will change over time, representing a degradation pathway that has not been previously considered.48 It may be the case that many contact layers containing metals are fundamentally unsuitable for contact with the perovskite, although instability would also depend on the relative binding energy of the metals in the contact materials. To conclude, we have demonstrated exchange of the B-site metal cation in both halide perovskite thin films and nanocrystals, converting between Sn- and Pb-based perovskites. We observed formation of crystalline alloys rather than bilayered or phase-segregated regions. This process points to new opportunities for controlling composition as a postsynthetic process, detached from the often complex processes of crystallization, and could enable more precise control over substitutional doping in these interesting semiconductors. This work also provides strong evidence for long-time-scale mobility of the B-site metal cations in perovskites, indicating that they may play a role in ionic conduction and suggesting that care must be taken to ensure compositional stability over an operational time scale. We hope that the observation of B-site cation exchange in thin films will also motivate future work to explore cation exchange with isovalent and aliovalent cations such as Mn2+, Co2+, Hg2+, Bi3+, and Al3+, which have been previously reported to affect perovskite material properties via substitution of Pb2+ in the lattice.49−53 However, careful control over the dopant quantities has thus far been impossible because they have been added in the initial crystallization. Ion exchange could provide a facile method to control the dopant density, having been used in other material systems to dope even as precisely as one dopant per nanocrystal by simply varying the concentration of the exchange solution.30

structure from further damage. This mechanism is also consistent with the faster exchange of Sn for Pb than the reverse; TOP effectively lowers the lattice binding energy of the Sn but not the Pb. To further elucidate the mechanistic details of the exchange, we carried out elemental depth profiling on FASnI3 thin films during an exchange reaction with Pb, as we show in Figure S14. We observe that even at the shortest times, some Pb has penetrated completely throughout the depth of the film, showing a fairly flat line profile. As time goes on, the Pb profile increases in magnitude and Sn decreases, without significantly changing distribution line shapes. The fact that we do not see a higher concentration of Pb at the surface suggests that the reaction is a surface-exchange limited one; Pb ions appear to be able to diffuse fully throughout the film more quickly than entering the film in the first place. The observation of metal cation exchange in halide perovskites has several implications for this material family. First, it shows that B-site cation exchange is possible in both nanocrystals and thin films, despite large activation energies for B-site motion that have been predicted by theory.9 Second, the method provides a new postsynthetic low-temperature process to control composition independent of crystallization, which provides new control over composition and could allow access to novel structure−composition combinations that are not easily attainable in a straightforward synthesis. For instance, we highlight that it has previously been impossible to fabricate black-phase FAPbI3 at temperatures below 150 °C, whereas here we have realized such films using processing at temperatures at or below 70 °C.44,45 We do however note that extending this approach to other halide perovskites may not be straightforward. Previously, the dynamic response of perovskites to electrical stimuli has been ascribed to motion of the halides13 and, to some extent, the A-site cation.46 The fact that the B-site metal cation can be exchanged via a diffusion-based process also suggests that, contrary to what has been assumed, it may not be exempt from movement under external stimuli. Cation motion may thus also be partly responsible for the observation of transient processes in these materials. Comparing the rate of exchange of the B-site cation with the A-site and X-site ions, it is a much slower process; the A-site cation is exchanged in a matter of hours and the X-site in a matter of minutes.16,19 This difference is in agreement with the expectation that the activation energy for substitution on the B-site is higher,14 though still accessible. Our observation of exchange at moderate temperatures also suggests that the activation energy for motion of the B-site ion may be toward the lower end of the range theoretically predicted by various groups (0.8−2.3 eV).14 It is possible that motion of the metal cations could be responsible for previously unexplained slow-time-scale effects such as a reversible “burn-in” during long-term aging.47 Finally, we note that the facile exchange of the metal cation has implications for device stability. In operation, devices are held at maximum power point bias, which continually changes over the course of a day. It is thus likely that the metal cations may move around in the material under operation as the field across the device changes. This motion may not be problematic because the other A- and X-site ions will also move, but it could contribute to a long-term instability of the material. Moreover, it imposes new criteria upon suitable contact materials; tin oxide is often used as an n-type contact layer, for example, but if the tin from that layer can exchange with the Pb in the active

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ASSOCIATED CONTENT

* Supporting Information S

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsenergylett.7b00290. Materials and methods, supplementary data including control film experiments, photoluminescence, absorbance, transmission electron microscopy, and elemental analysis data showing reversible exchange in nanocrystals, and depth profiling of films (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Giles E. Eperon: 0000-0001-9600-4847 David S. Ginger: 0000-0002-9759-5447 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We thank Ravishankar Madhu for assistance with nanocrystal synthesis, Erin Jedlicka for assistance with GDOES, and Prof. Brandi Cossairt for fruitful discussions. This Letter is based primarily on work supported by the DOE BES DE-SC0013957. We acknowledge additional support for instrumentation, infrastructure, and travel funding from the University of Washington Clean Energy Institute, the Alvin L. and Verla R. Kwiram Endowment from the Department of Chemistry at the 1194

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University of Washington, and the Washington Research Foundation. Part of this work was conducted at the Molecular Analysis Facility, a National Nanotechnology Coordinated Infrastructure site at the University of Washington, which is supported in part by the National Science Foundation (Grant ECC-1542101), the University of Washington, the Molecular Engineering & Sciences Institute, the Clean Energy Institute, and the National Institutes of Health.

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