Tuning the bandgap of Cs<sub>2</sub>AgBiBr<sub>6</sub> through dilute tin alloying
The promise of lead halide hybrid perovskites for optoelectronic applications makes finding less-toxic alternatives a priority. The double perovskite Cs 2 AgBiBr 6 (1) represents one such alternative, offering long carrier lifetimes and greater stability under ambient conditions. However, the large and indirect 1.95 eV bandgap hinders its potential as a solar absorber. Here we report that alloying crystals of 1 with up to 1 atom% Sn results in a bandgap reduction of up to ca. 0.5 eV while maintaining low toxicity. Crystals can be alloyed with up to 1 atom% Sn and the predominant substitution pathway appears to be a $2 : 1 substitution of Sn 2+ and Sn 4+ for Ag + and Bi 3+ , respectively, with Ag + vacancies providing charge compensation. Spincoated films of 1 accommodate a higher Sn loading, up to 4 atom% Sn, where we see mostly Sn 2+ substitution for both Ag + and Bi 3+ . Density functional theory (DFT) calculations ascribe the bandgap redshift to the introduction of Sn impurity bands below the conduction band minimum of the host lattice. Using optical absorption spectroscopy, photothermal deflection spectroscopy, X-ray absorption spectroscopy, 119 Sn NMR, redox titration, single-crystal and powder X-ray diffraction, multiple elemental analysis and imaging techniques, and DFT calculations, we provide a detailed analysis of the Sn content and oxidation state, dominant substitution sites, and charge-compensating defects in Snalloyed Cs 2 AgBiBr 6 (1:Sn) crystals and films. An understanding of heterovalent alloying in halide double perovskites opens the door to a wider breadth of potential alloying agents for manipulating their band structures in a predictable manner.
Introduction<!>Synthesis and stability<!>Optical effects of Sn alloying<!>XANES analysis<!>Redox titration<!>119 Sn NMR<!>Sn 2+ oxidation<!>Elemental composition<!>Lattice vacancies<!>Band structure<!>Sn content in lms<!>Film imaging<!>Sn speciation in lms<!>Conclusions<!>Conflicts of interest
<p>Lead halide perovskites, with the general formula APbX 3 (A ¼ monovalent cation, X ¼ Br or I), have exhibited remarkable properties for use as solar absorbers; 1,2 however, concerns regarding their long-term stability and the toxicity of watersoluble Pb 2+ salts still need to be addressed. 3,4 In order to identify lead-free materials that show similar optoelectronic properties to APbX 3 , we, 5 and others, 6,7 introduced halide double perovskites as potential absorbers. In particular, Cs 2 -AgBiBr 6 (1) displayed a long carrier lifetime, which is benecial for charge extraction in a solar cell, and higher stability to heat and moisture compared to (CH 3 NH 3 )PbI 3 . 5 However, the large and indirect bandgap of 1.95 eV in 1 affords weak sunlight absorption. Recent work by our group showed that the bandgap of 1 could be reconstructed through dilute Tl alloying. 8 Here, incorporation of Tl + resulted in a modest reduction in bandgap, although the transition was calculated to be direct. Incorporation of less than 1 atom% of Tl 3+ , on the other hand, resulted in a ca. 0.5 eV bandgap reduction while retaining the indirect bandgap of the host perovskite. Indeed Tl 3+ alloying makes the photophysical properties of 1 competitive with those of the APbX 3 absorbers. Although the use of toxic Tl, even at small concentrations, is undesirable for large-scale applications, this study provided the orbital basis for reconstructing the bandgap of 1 through impurity alloying.</p><p>Building on our understanding of the effects of Tl alloying, we sought a less toxic element that could provide similar bandgap reconstruction in 1. Our prior computational studies revealed that bandgap reconstruction can be affected by both the lled 6s shell of Tl + and the empty 6s shell of Tl 3+ . 8 With this knowledge guiding our search, we sought to incorporate Sn into 1 because of the similar electronic congurations of Tl + and Tl 3+ with Sn 2+ and Sn 4+ , respectively. Additionally, both Sn 2+ and Sn 4+ are known to form halide perovskites. 9,10 Because both Sn 2+ and Sn 4+ are heterovalent with respect to Ag + and Bi 3+ , this study presented the opportunity of understanding both possible substitution sites and charge-compensating defects in 1, and how they depended on sample morphology and synthetic conditions. Heterovalent alloying has been studied in many other materials and is known to have a variety of effects on the host material, including modulating carrier concentrations, 11 altering the kinetics of phase transformation, 12,13 lowering the ferroelectric transition temperature, 13 inducing disorder and vacancies, 14,15 and affecting ionic conductivity. 12,14 Although there have been numerous attempts at incorporating small amounts of monovalent and trivalent metals into the lead perovskites, [16][17][18][19][20][21][22][23][24][25][26][27] detailed experimental characterization of the resulting materials that provide a thorough understanding of the structural and electronic changes in the alloys are still needed. A recent study with a computational focus explored Pb 2+ alloying of 1, which resulted in a ca. 0.1 eV redshi of the absorption onset. 28 In order to probe the complexities of heterovalent alloying in 1, here we use a combination of single-crystal and powder X-ray diffraction, optical absorption and photothermal deection spectroscopy, elemental analysis and redox titration, nuclear magnetic resonance spectroscopy, X-ray absorption spectroscopy, scanning electron and atomic force microscopy, and band structure calculations to elucidate the structural and electronic effects of Sn alloying in 1. We show here that Sn alloying substantially reduces the bandgap of 1, affording a low-bandgap halide perovskite free of highly toxic elements.</p><!><p>Addition of SnBr 2 to the precursor solution 5 of 1 under inert atmosphere resulted in truncated octahedral crystals of Snalloyed Cs 2 AgBiBr 6 (1:Sn) with inductively coupled plasma mass spectroscopy and optical emission spectroscopy (ICP-MS/ ICP-OES) showing the Sn content ranging from 0.023(1)-1.0(2) atom% (Table S1 and Fig. S1 †). Thin lms of 1:Sn were prepared by spincoating a solution of the pre-synthesized solids of CsSn II Br 3 and 1 dissolved in dimethylsulfoxide (DMSO) in various ratios, maintaining a xed concentration of Cs and Br in solution; this method effectively enabled direct replacement of Ag and Bi with Sn (see ESI † for detailed synthetic procedures and material characterization). The thin lms accommodated a higher Sn content, ranging from 1-4 atom%, as estimated from X-ray photoelectron spectroscopy.</p><p>The stability of the parent structure of 1 in 1:Sn was tested by exposing pulverized crystals of 1:Sn (1 atom% Sn) to 90 C heat in air for 30 days, 55% relative humidity in air for 30 days, and 0.5 sun under N 2 at ca. 45 C for 30 days. The powder X-ray diffraction (PXRD) patterns of these samples did not change signicantly, indicating that the parent structure of 1 in 1:Sn is structurally sound (Fig. S2 †).</p><!><p>Crystals of 1:Sn range from red to black in color (Fig. 1B). To quantify this color change, the Kubelka-Munk theorem was used to transform UV-visible-NIR (UV-vis) diffuse reectance spectra measured on crystals of 1:Sn to absorbance values (Fig. 1C). 29 These spectra clearly show the absorption onsets, although the signal likely saturates at higher energies. Bandgaps were determined by plotting a r as a function of photon energy (E), where a is the pseudo-absorption coefficient and r ¼ 0.5 for an indirect bandgap and r ¼ 2 for a direct bandgap (Fig. S3 and S4 †). Although the results show the previously reported indirect bandgap for 1, 5 they support either a direct bandgap of 1.71 eV or an indirect bandgap of 1.48 eV for 1:Sn (at 1 atom% Sn). The absorption onset energy initially drops quickly with increasing Sn content at low alloying concentrations, then shows a slower reduction at high Sn concentrations, reaching a maximum redshi of ca. 0.5 eV (Fig. S4 †). To test the effect of particle size on apparent absorption onset, 30 UV-vis diffuse reectance spectra were measured on ball-milled powders of 1 and 1:Sn with sub-mm particle sizes, both of which show absorption onsets at the same energy as the crystals (Fig. S5 and S6 †).</p><p>Photothermal deection spectroscopy (PDS) was used to address the possibility that Sn alloying may, instead of reconstructing the bandgap, cause lattice disorder that creates weakly absorbing sub-bandgap trap states. PDS offers a wider dynamic range of absorption than UV-vis spectroscopy but requires samples with smooth surfaces with areas of several mm 2 . Thin lms of 1 and 1:Sn were therefore used for these measurements. The lms varied in color from yellow to brown with increasing Sn content (lm thickness ¼ 190 AE 16 nm). UV-vis transmission and reectance measurements track this color change, showing increased low-energy absorption as a function of Sn content (Fig. S7 †). PDS measurements on thin lms revealed similarly large Urbach energies (70-100 meV) for both 1 and 1:Sn (4 atom%; Fig. 1D). Because 1 has an indirect bandgap, the absorption coefficient at the bandgap onset is small; subbandgap trap states that may form upon alloying may have similarly small absorption coefficients. Thus a redshi of the absorption onset, i.e., at low values of absorption coefficient (<10 4 cm À1 ), does not necessarily indicate a bandgap reduction, in contrast to lead perovskites with direct bandgaps. 16,17,20,22,24,26 We therefore looked deeper into the bandgap, at energies where the absorption coefficient reaches ca. 10 4 cm À1 . The PDS spectrum of 1:Sn reaches an absorption coefficient of 3 Â 10 4 cm À1 at an energy ca. 0.7 eV lower than for 1, indicating that the observed absorption redshi arises from a bandgap reduction, 31 as opposed to sub-bandgap trap states induced by the Sn impurities, which are not expected to affect these higher-energy transitions.</p><!><p>Because Sn 2+ is known to be oxidatively unstable in CsSn II Br 3 , 32 we sought to determine the oxidation state of Sn in 1:Sn. Comparing X-ray absorption near-edge structure (XANES) spectra collected at the Sn L 3 -and K-edges in 1:Sn (1 atom% Sn) to those of CsSn II Br 3 and Cs 2 Sn IV Br 6 perovskite standards indicated that Sn in 1:Sn was present in a mixture of 2+ and 4+ oxidation states in pulverized crystal samples that had been prepared and measured under inert atmosphere. Linear combination tting of the XANES spectrum of 1:Sn to the XANES spectra of CsSn II Br 3 and Cs 2 Sn IV Br 6 standards indicated a mixture of Sn 2+ and Sn 4+ in ratios ranging from 7 : 1 to 3 : 2 Sn 2+ to Sn 4+ in several samples, with an average ratio of 2 : 1 Sn 2+ to Sn 4+ (Fig. 2A and S8 †). The low signal-to-noise ratio at the Sn L 3 -edge due to the small concentration of Sn and lifetime broadening combined with the low energy resolution at the Sn K-edge precluded a more precise assignment of the ratio of Sn 2+ to Sn 4+ . Although the host lattice of 1 in 1:Sn remains structurally unchanged with exposure to ambient atmosphere, the Sn 2+ in 1:Sn is unstable to post-synthetic oxidation: exposure of a sample of pulverized crystals of 1:Sn to ambient atmosphere for 1 week resulted in a reduction of the ratio of Sn 2+ to Sn 4+ from 3 : 2 to 2 : 9 (Fig. S9 †).</p><!><p>The Sn 2+ content of 1:Sn (1 atom% Sn) was determined to be 0.59(4) atom% by iodimetric titration on crystals digested in puried hydrobromic acid. Assuming a total Sn content of 1 atom%, as given by ICP analysis (Table S1 and Fig. S1 †), gives a Sn 2+ to Sn 4+ ratio of ca. 3 : 2, in agreement with the XANES results range. Due to the variance of the total Sn content given by ICP analysis for the highest alloying level (1.0(2) atom% Sn, averaging over ve samples), this ratio is taken to be an approximation.</p><!><p>As an additional determination of the oxidation state ratio of Sn in 1:Sn we turned to 119 Sn NMR. Magic-angle spinning solidstate 119 Sn nuclear magnetic resonance (NMR) of CsSn II Br 3 and Cs 2 Sn IV Br 6 gave clear, relatively broad (FWHM $ 30 ppm) resonances at À370 and À1965 ppm, respectively. However, despite long (25-33 h) acquisition times, no signals were observed from crystals of 1:Sn (1 atom% Sn), likely due to the low concentration of Sn (Fig. S10 †). We, therefore, collected solution-state 119 Sn NMR spectra on crystals of 1:Sn (1 atom% Sn) dissolved in DMSO, which showed one broad peak at À520 ppm (FWHM $ 10 ppm) and two narrow peaks at À1010 and À1270 ppm (FWHM $ 1.5 ppm; Fig. 2B). The chemical shis of the two upeld peaks match those of the two most intense peaks in the NMR spectrum of Cs 2 Sn IV Br 6 in DMSO. Though its chemical shi does not match that of the broad peak in the NMR spectrum of CsSn II Br 3 in DMSO, the broad peak is shied downeld of the Sn 4+ peaks, suggesting its origin to be Sn 2+ . Indeed, intentionally adding SnBr 2 and SnBr 4 to a solution  of 1 in DMSO gives a spectrum that is a close match to that of 1:Sn (Fig. 2B). The change in chemical shi of the Sn 2+ signal from the CsSn II Br 3 standard to the sample of 1:Sn spiked with SnBr 2 and SnBr 4 suggests that Ag and/or Bi interact with the Sn 2+ in DMSO solution, altering the observed chemical shi. Peak integration of the spectrum of 1:Sn in DMSO gives an approximate ratio of 2 : 1 Sn 2+ to Sn 4+ , in agreement with the results from XANES and redox titration. Given the agreement amongst these techniques, we assign an average ratio of 2 : 1 Sn 2+ to Sn 4+ in crystals of 1:Sn (1 atom% Sn). This ratio appears to vary from 7 : 1 to 3 : 2, likely based on subtle synthetic changes such as rates of crystallization.</p><!><p>Although our syntheses were performed with Sn 2+ precursors in inert atmosphere, we consistently see evidence for Sn 4+ in crystals of 1:Sn. Attempts to prevent oxidation of Sn 2+ by reagent purication and addition of hypophosphorous acid as a reducing agent were unsuccessful, indicating that the partial oxidation of Sn 2+ is thermodynamically favorable in this system. Although the oxidant for Sn 2+ in 1:Sn is yet unknown, the oxidized product Cs 2 Sn IV X 6 is reported to form in syntheses of CsSn II X 3 in hydrohalic acid, regardless of reagent purity. 33 Indeed, the Sn 6s 2 orbitals lie at the valence band maximum in 1:Sn and is most easily oxidized, similar to the case for ASnI 3 . 34,35 At the other extreme, addition of SnBr 4 to the crystallization solution for 1 resulted in the formation of 1 and Cs 2 Sn IV Br 6 and did not yield 1:Sn in the absence of SnBr 2 .</p><!><p>Given that heterovalent substitution of Sn in 1 was expected to produce charge-compensating vacancies, ICP-MS/ICP-OES were used to quantify the elemental composition of 1:Sn (Table S1 and Fig. S1 †). The results showed a tunable quantity of Sn inclusion, up to $1 atom% in crystals, proportional to the concentration of SnBr 2 in the precursor solution. As indicated by PXRD, the formation of a biphasic mixture of 1:Sn and the perovskite CsSn II Br 3 gives an upper bound to the Sn concentration achievable in phase-pure 1:Sn crystals (Fig. S11 †). A slight shi of the PXRD pattern to higher angles is apparent in 1:Sn crystals, corresponding to a lattice parameter increase of up to $0.14%. The observed expansion is consistent with Sn alloying into 1, as the doubled lattice parameter of CsSn II Br 3 is 3% larger than the lattice parameter of 1. The ICP results also showed a modest decrease in Bi content and a substantial decrease in Ag content, more than the predicted value for a 1 : 1 substitution of Sn for Ag, suggesting that Ag vacancies are being formed as a charge-compensating mechanism with Sn alloying.</p><!><p>Single-crystal X-ray diffraction (SC-XRD) also corroborates the formation of Ag vacancies upon Sn alloying. SC-XRD of 1:Sn (1 atom% Sn) showed a similar structure to that of 1 but with signicant missing electron density at the Ag site when modelled with full Ag occupancy (Fig. 3A and Table S2 †). Because Sn has a larger scattering factor than Ag, the missing electron density could not be attributed to Sn substitution at the Ag site. The missing electron density was instead assigned to Ag vacancies, consistent with the ICP results showing greater Ag loss than expected for a 1 : 1 substitution of Sn for Ag. Though the ICP results additionally show Bi loss with Sn alloying in 1:Sn, the low magnitude of this loss precludes detection by SC-XRD. Rening the Ag site in this model as having an equal concentration of Sn impurity atoms and vacant sites gave occupancy values of 86% Ag, 7.0% Sn, and 7.0% vacancies, similar to the experimentally determined concentrations of Ag and Sn in 1:Sn (1 atom% Sn). In order to probe for evidence of Bi vacancies, we attempted to obtain a single crystal with higher Sn content. Crystallizations with high concentrations of SnBr 2 lead to mixtures of 1:Sn and CsSn II Br 3 crystals. Manually separating a single crystal of 1:Sn from this mixture allowed us to obtain the SC-XRD data of a crystal with a higher content of Sn than we could isolate from a phase-pure mixture. Renement of these data gave a structure solution with signicant missing electron density at the Ag site and, to a lesser extent, at the Bi site (Fig. 3B and Table S2 †). Modelling mixed occupancy at the Ag site gave a Ag occupancy of 65% and Sn and vacancy occupancies each of 18%. Modelling the Bi site as an atomically mixed site with Sn gives occupancies of 84% for Bi and 16% for Sn. Alternatively, the missing electron density at the Bi site could be assigned to vacancy formation at that site or a combination of Sn substitution and vacancy formation.</p><p>The XANES, redox titration, NMR, ICP, and SC-XRD results, taken together, indicate that Sn 2+ substitutes at the Ag + site while Sn 4+ substitutes at the Bi 3+ site with an approximate ratio of 2 : 1 Sn 2+ to Sn 4+ , collectively generating Ag + vacancies as the primary charge-compensating defect (Fig. 1A). Similar chargecompensating metal cation vacancies have been observed in oxide perovskites upon heterovalent substitution. 36 Though the participation of competing substitution mechanisms cannot be rigorously excluded, this model provides the best t with the experimental data (Table S3 †). We assume substitution of Sn 2+ at the Ag + site and Sn 4+ at the Bi 3+ site due to the large mismatch in valence between Sn 4+ and Ag + . Additionally, substitution of Sn 2+ at the Bi 3+ site would require a positively charged compensating defect, such as a Br À vacancy, which is not supported by the ICP nor the SC-XRD results. The modest decrease of Bi content with Sn alloying suggests that there is no signicant density of Bi vacancies in 1:Sn. We therefore propose the formula Cs 2 (Ag 1À(2a+b) Sn II a )(Bi 1Àb Sn IV b )Br 6 for 1:Sn crystals, where 0.0023(1) < a + b < 0.10(2).</p><p>Similar to the host lattice 1, which exhibits only very weak photoluminescence, the alloyed material 1:Sn does not photoluminescence, consistent with its computed indirect bandgap. It is also possible that the lattice vacancies required for heterovalent alloying may act as nonradiative carrier recombination sites. Therefore, the carrier dynamics of 1:Sn should be studied to assess the feasibility of charge extraction from lattices containing metal cation vacancies.</p><!><p>We calculated the band structures for 1:Sn to understand the electronic consequences of Sn 2+ and Sn 4+ alloying in 1. Our calculations were performed using density functional theory (DFT) within the generalized gradient approximation of Perdew, Burke, and Ernzerhof (PBE) as implemented in the VASP code. 37,38 Spin-orbit coupling (SOC) effects were treated selfconsistently. The calculated bandgap of an 80-atom supercell of 1 using this method is 1.11 eV (Fig. 4A), underestimating the experimental value of 1.95 eV, as is expected for DFT-PBE-SOC and consistent with past calculations from similar computations for (CH 3 NH 3 )PbI 3 . 39 We note that in the band structure of the primitive unit cell, the valence band maximum is unfolded from G to X. Accurate prediction of bandgap energies requires a more rigorous treatment of exchange and correlation effects, including electron-hole interactions, that is currently prohibitive for the large unit cells considered here; however, our DFT-PBE-SOC calculations are expected to capture trends, sufficient for the present study (see ESI † for details).</p><p>For our DFT calculations of the bandgap of 1:Sn, we constructed supercells in which one or more of the B-site cations were substituted with Sn, where B-site refers to the octahedrally coordinated cation in the perovskite (Ag and Bi in 1). The insertion of heterovalent Sn-which may be considered to assume nominal charges of Sn 2+ or Sn 4+ -into the structure of 1 necessitates a charge-compensating mechanism to maintain charge neutrality, such as the formation of cation vacancies. Because of the numerous permutations of possible substitution scenarios and the computational expense of using such large supercells, we used our experimental results to guide our choice of calculations: we considered four separate cases that could arise in nominal Sn 2+ and Sn 4+ substitutions, assuming the Sn atoms and vacancies substitute only at the Ag and Bi sites, consistent with our experimental evidence.</p><p>Case 1: two Ag atoms were removed and replaced with one Sn and one vacancy, respectively, to model nominal Sn 2+ substitution at the Ag site with Ag vacancies as the chargecompensating defect. The supercell consisted of 80 atoms, corresponding to 1.25 atom% Sn substitution. In this supercell, there were 8 different Ag sites for substitution, thus 7 different relative arrangements of the Sn and vacancy on these sites were considered. For all arrangements, the lattice parameters and internal coordinates were relaxed (without spin-orbit coupling) until the forces were converged to 10 meV ÅÀ1 (Tables S4-S6 †). The arrangement with Sn and the vacancy occupying adjacent B-sites was computed to have the lowest energy (Table S7 and Fig. S12; see ESI † for details). The calculated band structure of this arrangement, including spin-orbit coupling, revealed a direct bandgap with a modest reduction of <0.1 eV from the DFT-PBE-SOC bandgap of 1 (Fig. 4B and S13 †).</p><p>Case 2: one Bi atom was replaced with one Sn atom and one Ag vacancy was introduced to model nominal Sn 4+ substitution at the Bi site with Ag vacancies as the charge-compensating (1/2, 1/2, 1/2) corresponds to L and R in the double perovskites and in CsPbBr 3 , respectively. The range of k-points plotted is from (0, 0, 0) to (0.1, 0.1, 0.1), corresponding to the direction from G towards L or R, and from (0, 0, 0) to (0.1, 0.0, 0.1), corresponding to the direction from G towards X.</p><p>defect. The supercell consisted of 80 atoms for a 1.25 atom% concentration of Sn. Aer considering several arrangements of the Sn and vacancy (see ESI † for details), the structure placing Sn and the vacancy at adjacent B-sites was computed to be the lowest-energy structure, similar to Case 1. In contrast to Case 1, Case 2 resulted in an indirect but substantially reduced DFT-PBE-SOC bandgap (by 0.44 eV), in reasonable agreement with the experimentally determined bandgap reduction of 0.5 eV in crystals of 1:Sn (1 atom% Sn; Fig. 4C and S14 †). Note that whereas these calculations largely underestimate the absolute bandgap values, as expected, relative changes in bandgap can be well captured, and that is indeed what we nd here.</p><p>Case 3: four Bi atoms were replaced with three Sn atoms and one vacancy to model nominal Sn 4+ substitution at the Bi site with Bi vacancies as the charge-compensating defect. The supercell consisted of 160 atoms for a 1.875 atom% concentration of Sn. Aer comparing the DFT-PBE energies of 8 relaxed structures with substitutions on different Bi sites in the supercell, the lowest-energy conguration was computed to be one in which the Sn atoms and vacancy occupy adjacent sites, similar to the previous cases. The calculated DFT-PBE-SOC band structure in this scenario had a slightly indirect bandgap, while the conduction band had a very narrow bandwidth compared to the other cases. The bandgap was reduced by 0.61 eV compared to 1, in fair agreement with the experimentally determined reduction (Fig. S15 †).</p><p>Case 4: one Ag and one Bi atom were each replaced with Sn atoms to model nominal Sn 2+ substitution in equivalent concentrations at each B-site. The supercell consisted of 80 atoms, giving an overall concentration of 2.5 atom% Sn. The lowest-energy arrangement contained Sn atoms on adjacent Ag and Bi sites (see ESI † for details). A largely computational study of Pb 2+ alloying in 1 proposes a similar mechanism. 28 The DFT-PBE-SOC bandgap in this case was calculated to be direct, and reduced by 0.47 eV relative to the bandgap of 1 (Fig. S16 †).</p><p>In Case 1, our DFT calculations indicate the direct bandgap results from the introduction of Sn 5p 0 character and Sn 5s 2 character at G. In contrast, in Case 2, the Sn substitution at the Bi site results in the introduction of a band below the CBM with Sn 5s 0 character at L, resulting in a more reduced but slightly indirect bandgap (Fig. 4 and Table S8 †). To conrm that the direct bandgap afforded by Sn 2+ alloying contributes to optical absorption, we calculated the magnitudes of the transition dipole matrix elements in the independent particle approximation for 1, Sn 2+ -alloyed 1 (Cases 1 & 4), and the directbandgap semiconductor CsPbBr 3 for comparison. Although the magnitudes of the transition dipole matrix elements for 1:Sn are smaller than those of CsPbBr 3 , they are signicantly greater than the negligible magnitude for 1, supporting the presence of an optically active direct bandgap in Sn 2+ -alloyed 1 (Fig. 4D). We additionally calculated the imaginary part of the dielectric function as a function of energy in the independent particle approximation for 1 and Sn 2+ -alloyed 1 (Cases 1 & 4) to approximate their absorption spectra. These calculations show a clear redshi in the onset energy of the dielectric functions, offering further support for the bandgap reconstruction in Sn 2+alloyed 1 (Fig. S17 †).</p><p>Because the 5s 0 orbital from nominal Sn 4+ substitutions is expected to be lower in energy than the 5p 0 orbital from Sn 2+ substitutions, a sample comprising both Sn 2+ and Sn 4+ in comparable amounts, as is the case for 1:Sn crystals, would therefore be expected to have a substantially reduced indirect bandgap. This was conrmed by a computation performed on a 320-atom supercell modelled to match this experimentally observed substitution pattern (Table S9 and Fig. S18 †). At very low alloying concentrations, Sn impurities should act as isolated sub-bandgap trap states. With increasing concentration, impurity bands quickly form from the hybridization of Sn atoms with the host lattice. At higher concentrations of Sn, the defect bands become more dispersive, thus serving to fundamentally alter the electronic structure of 1, as corroborated through PDS measurements of 1:Sn lms.</p><!><p>The nature of the alloy may depend on material morphology and synthetic conditions. Therefore, in order to understand the effects of Sn alloying in 1 in the morphology relevant for devices, we studied thin lms of 1:Sn. X-ray photoelectron spectroscopy (XPS) measurements indicated that a higher concentration of Sn could be alloyed into thin lms compared to crystals, with PXRD showing that phase-purity was maintained up to ca. 4 atom% Sn-a 4-fold increase relative to crystals (Table S10, Fig. S19 and S20 †). The PXRD data also revealed a slight lattice expansion of up to $0.52% with increasing Sn content, consistent with a $4-fold lattice expansion relative to crystals of 1:Sn (1 atom% Sn). The higher Sn content achievable in thin lms of 1:Sn relative to crystals may arise from increased kinetic control offered by the faster lm deposition methods, increased allowance for strain in these thin lms, and the effect of changing the solvent from aqueous hydrobromic acid to DMSO.</p><!><p>Thin lms of 1 and 1:Sn were further characterized using scanning electron microscopy (SEM) and atomic force microscopy (AFM), revealing a change in morphology with inclusion of Sn (Fig. 5A, B and S21 †). Whereas lms of 1 have roughly spherical grains with an average size of 90 AE 30 nm and a rootmean-square (RMS) surface roughness of 15 nm, lms of 1:Sn (4 atom% Sn) have irregularly shaped polydisperse grains with an average size of 100 AE 70 nm and a lower RMS surface roughness of 4 nm. Such changes in lm morphology with alloying have been previously observed. 40,41 Film thickness was measured by prolometry to be 190 AE 16 nm, independent of Sn content. To address the possibility of phase segregation of 1 and CsSn II Br 3 at the nanoscale, energy-dispersive X-ray spectroscopy (EDX) elemental mapping (Fig. 5C, S22, and S23 †) and concentric backscatter imaging (CBS; Fig. S24 †) were employed. EDX revealed a homogeneous distribution of Sn within a resolution of $500 nm while CBS and AFM showed homogeneous phases within resolutions of 25 nm and 8 nm, respectively.</p><!><p>As a further comparison between crystals and thin lms of 1:Sn, XANES spectra were measured on thin lms of 1:Sn of various alloying concentrations (Fig. S25 †). Sn L 3 -edge XANES results indicated that the ratio of Sn 2+ to Sn 4+ decreased from nearly 1 : 0 at the highest achievable Sn concentration of 4 atom% to 4 : 1 at a Sn concentration of 2 atom% to 2 : 1 at a Sn concentration of 1 atom%, the same total Sn concentration as the maximum Sn concentration achievable in crystals. The agreement of the Sn 2+ to Sn 4+ ratios between thin lm and crystal samples with 1 atom% alloyed Sn indicates that there is a thermodynamic driving force for partial oxidation of Sn 2+ in 1:Sn, regardless of synthetic conditions and material morphology. Though the possibility of partial oxidation of Sn in 1:Sn from adventitious oxygen cannot be eliminated, rigorous precautions were taken to prevent such an occurrence. Intentional oxidation of a thin lm of 1:Sn (4 atom% Sn) by exposure to ambient atmosphere overnight results in complete oxidation of the Sn 2+ as determined by XANES (Fig. S26 †). This oxidation occurs at an elevated rate relative to that for pulverized crystals likely due to an increased surface-area-to-volume ratio in lms. SEM imaging of intentionally air-exposed thin lms of 1:Sn revealed the presence of pillar-like surface features, identied by EDX elemental mapping to contain a highly elevated concentration of Ag and a moderately elevated concentration of Cs and O relative to the surrounding sample (Fig. S27 and S28 †). Exposure of samples to dry air causes the formation of the pillar-like surface features, while exposure to wet nitrogen gas instead causes the formation of pinholes and the separation of grains, indicating that the formation of the surface features is connected to the oxidation of 1:Sn (Fig. S29 †). Exposure of thin lms of 1:Sn (4 atom% Sn) to ambient atmosphere over a period of hours results in a reversion of its brown color to the yellow color of thin lms of 1.</p><!><p>We provide a detailed analysis of how the photophysical properties of crystals and thin lms of the halide double perovskite Cs 2 AgBiBr 6 (1) can be tuned through Sn alloying. XANES, redox titration, 119 Sn NMR, ICP, and SC-XRD indicate that a $2 : 1 ratio of Sn 2+ and Sn 4+ is present in crystals of the alloyed perovskite (1:Sn; 1 atom% Sn) along with Ag + vacancies. Our results are consistent with the following primary substitution pathway for crystals: Sn 2+ substitutes at the Ag + site while Sn 4+ substitutes at the Bi 3+ site, collectively generating anionic Ag + vacancies as the dominant charge-compensating defect (see note on the Kröger-Vink notation below):</p><p>DFT calculations of the band structure show that the introduction of new bands with Sn 4+ 5s 0 impurity character below the conduction band minimum of the host lattice effects the large bandgap reduction seen in 1:Sn (1 atom% Sn), successfully mimicking the bandgap reduction resulting from Tl 3+ incorporation in 1 8 in a nontoxic composition (Fig. 4).</p><p>Thin lms of 1 can accommodate a higher Sn content (4 atom% Sn). In contrast to crystals, XANES measurements reveal the dominance of Sn 2+ in thin lms of 1:Sn (4 atom% Sn). This suggests a distinct alloying mechanism in thin lms at high concentrations of Sn wherein only Sn 2+ substitutes at both Ag + and Bi 3+ sites in approximately equal concentrations, thus maintaining overall charge neutrality (see Section 5, Case 4). Our computational results for this substitution pattern predominant in thin lms show that, with a Sn concentration of 2.5 atom%, substitution of only Sn 2+ at both Ag + and Bi 3+ sites in equal concentrations affords a direct bandgap (although an indirect transition is nearby, only 40 meV higher in energy) reduced by 0.5 eV from that of 1 (Fig. S16 †), mimicking the indirect-to-direct bandgap change produced by Tl + substitution 8 in a nontoxic composition. Thus, heterovalent alloying further expands the considerable compositional diversity of double perovskites for nding functional analogs to the lead perovskites.</p><!><p>There are no conicts to declare.</p>