RuddlesdenâPopper Hybrid Lead Iodide Perovskite 2D Homologous

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Ruddlesden−Popper Hybrid Lead Iodide Perovskite 2D Homologous Semiconductors Constantinos C. Stoumpos,†,¶ Duyen H. Cao,†,¶ Daniel J. Clark,‡ Joshua Young,§,⊥ James M. Rondinelli,⊥ Joon I. Jang,‡ Joseph T. Hupp,† and Mercouri G. Kanatzidis*,† †

Department of Chemistry, Northwestern University, 2145 Sheridan Road, Evanston, Illinois 60208, United States Department of Physics, Applied Physics and Astronomy, Binghamton University, P.O. Box 6000, Binghamton, New York 13902, United States § Department of Materials Science and Engineering, Drexel University, 3141 Chestnut Street, Philadelphia, Pennsylvania 19102, United States ⊥ Department of Materials Science and Engineering, Northwestern University, 2220 Campus Drive, Evanston, Illinois 60208, United States ‡

S Supporting Information *

ABSTRACT: The hybrid two-dimensional (2D) halide perovskites have recently drawn significant interest because they can serve as excellent photoabsorbers in perovskite solar cells. Here we present the large scale synthesis, crystal structure, and optical characterization of the 2D (CH3(CH2)3NH3)2(CH3NH3)n−1PbnI3n+1 (n = 1, 2, 3, 4, ∞) perovskites, a family of layered compounds with tunable semiconductor characteristics. These materials consist of well-defined inorganic perovskite layers intercalated with bulky butylammonium cations that act as spacers between these fragments, adopting the crystal structure of the Ruddlesden−Popper type. We find that the perovskite thickness (n) can be synthetically controlled by adjusting the ratio between the spacer cation and the small organic cation, thus allowing the isolation of compounds in pure form and large scale. The orthorhombic crystal structures of (CH3(CH2)3NH3)2(CH3NH3)Pb2I7 (n = 2, Cc2m; a = 8.9470(4), b = 39.347(2) Å, c = 8.8589(6)), (CH3(CH2)3NH3)2(CH3NH3)2Pb3I10 (n = 3, C2cb; a = 8.9275(6), b = 51.959(4) Å, c = 8.8777(6)), and (CH3(CH2)3NH3)2(CH3NH3)3Pb4I13 (n = 4, Cc2m; a = 8.9274(4), b = 64.383(4) Å, c = 8.8816(4)) have been solved by single-crystal X-ray diffraction and are reported here for the first time. The compounds are noncentrosymmetric, as supported by measurements of the nonlinear optical properties of the compounds and density functional theory (DFT) calculations. The band gaps of the series change progressively between 2.43 eV for the n = 1 member to 1.50 eV for the n = ∞ adopting intermediate values of 2.17 eV (n = 2), 2.03 eV (n = 3), and 1.91 eV (n = 4) for those between the two compositional extrema. DFT calculations confirm this experimental trend and predict a direct band gap for all the members of the Ruddlesden− Popper series. The estimated effective masses have values of mh = 0.14 m0 and me = 0.08 m0 for holes and electrons, respectively, and are found to be nearly composition independent. The band gaps of higher n members indicate that these compounds can be used as efficient light absorbers in solar cells, which offer better solution processability and good environmental stability. The compounds exhibit intense room-temperature photoluminescence with emission wavelengths consistent with their energy gaps, 2.35 eV (n = 1), 2.12 eV (n = 2), 2.01 eV (n = 3), and 1.90 eV (n = 4) and point to their potential use in light-emitting diodes. In addition, owing to the low dimensionality and the difference in dielectric properties between the organic spacers and the inorganic perovskite layers, these compounds are naturally occurring multiple quantum well structures, which give rise to stable excitons at room temperature.

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INTRODUCTION The class of halide perovskite compounds of the chemical formula AMX3 (A = Cs+, CH3NH3+, or HC(NH2)2+; M = Ge2+, Sn2+, Pb2+; X = Cl−, Br−, I−) has witnessed a spectacular surge in scientific interest in the last five years and has enabled revolutionary achievements in the field of solid-state photovoltaics.1−9 The recent strong scientific activity in the halide perovskites was triggered by the successful demonstration of CH3NH3PbX3 (X = Br, I) as an efficient light absorber1−9 © XXXX American Chemical Society

and CsSnI3 as an efficient hole transporter in dye-sensitized solar cells (DSCs).10,11 These were followed by the utilization of CH3NH3PbI3 as a light absorbing material producing solidstate solar cells with remarkable efficiency in late 2012.12−14 Received: February 28, 2016 Revised: March 29, 2016

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DOI: 10.1021/acs.chemmater.6b00847 Chem. Mater. XXXX, XXX, XXX−XXX

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Chemistry of Materials

single layer.31 The exploration of this exciting phenomenon drew the attention of many researchers who devised novel experiments42−47 and performed extensive theoretical analyses48,49 aiming to understand and control the origins of the unique properties of the 2D perovskites. The most remarkable finding was that the stability of the exciton does not arise from the dimensional confinement alone but that the organic material plays an important role by modulating the dielectric properties of the material.50,51 The stability of the exciton in the n = 1 2D perovskites gives rise to an intense photoluminescence (PL), which persists even at room temperature. As a result, the 2D perovskites are frequently employed in optical and electronic devices such as field-effect transistors (FETs),34 light-emitting diodes (LEDs),31 and hard radiation detectors.52 The higher n members of the homologous series, however, remain less explored.53−55 They are now of intense interest because of their tunable optical properties, which have led to the fabrication of promising solar cells,26,27 and studies of their photophysical properties.56−58 The lack of extensive synthetic and structural information on the higher 2D perovskites further prompted us to investigate the Ruddlesden−Popper homologous series. In this article, we report the scalable synthesis, crystal structure, and optical properties of the homologous 2D series of lead iodide perovskites (CH 3 (CH 2 ) 3 NH 3 ) 2 (CH3 NH3 ) n−1 PbnI3n+1.30,46,53,56 We find that the (CH3(CH2)3NH3)2(CH3NH3)Pb2I7 (n = 2), (CH3(CH2)3NH3)2(CH3NH3)2Pb3I10 (n = 3), and (CH3(CH2)3NH3)2(CH3NH3)3Pb4I13 (n = 4) members combine the structural features of the simple 2D (n = 1) and 3D (n = ∞) perovskite end members. We further find that these intermediate phases exhibit tunable band gaps resulting from quantum confinement due to the dimensional reduction of the perovskite spacer layers. Furthermore, we show that the n = 2−4 compounds display optoelectronic properties that are different from either the n = 1 or n = ∞ end-members. Thus, they define a novel class of naturally forming superlattice semiconductors with thickness-dependent optical properties reminiscent of the complex quantum-wells in AlGaAs/GaAs heterostructures.59,60

Despite the astonishing photovoltaic performance, reaching ∼20% conversion efficiency in 2015,15−17 some drawbacks of CH3NH3PbI3 have become evident, with major concerns focusing on the insufficient long-term stability of the devices incorporating these materials and the toxicity of lead.18,19 Although the latter problem could be potentially resolved by replacing Pb with the more environmentally friendly Sn metal (a technology still in its infancy),20−25 the former has not yet been addressed in a satisfactory manner. One solution to the long-term stability problem has been recently proposed in the form of the layered iodide perovskites, which are two-dimensional (2D) derivatives of the three-dimensional (3D) perovskite formed by “slicing” the 3D frameworks into well-defined 2D slabs.26,27 The 2D derivatives in hybrid lead iodides have not been studied extensively; however, they offer far more tunability and flexibility in terms of being able to control the physical properties. Further motivation for examining the 2D hybrid iodide perovskites in much greater detail comes from recent processingdependent optical responses observed in 3D perovskite films. In particular, we previously observed the formation of red colored films upon spin-coating CH3NH3PbI3 solutions onto mesoporous TiO2 films prior to the formation of the expected black CH3NH3PbI3 film, in addition to slightly varying absorption spectra of CH3NH3PbI3 films with different PbI2 fraction.28 These observations led us to the hypothesis that some intermediate 2D compounds may be formed during the film fabrication process and that these optical responses could be due to quantum confinement effects in either nanoscale variants of perovskite CH3NH3PbI3 or the intrinsic properties of 2D metastable derivative phases. Such confinement effects studied through pair distribution function (PDF) analysis on CH3NH3PbI3 films29 arise from a structurally disordered perovskite dimensionally confined within the nanosized TiO2 pores and are responsible for a significant blue shift (∼50 meV) of the optical absorption edge. Thus, we set out to study in detail the 2D perovskites as a well-defined example of naturally forming semiconductor materials with quantum confinement. The 2D perovskites have the generic chemical formula of (RNH3)2(A)n−1MX3n+1 (n is an integer), where RNH3 is a primary aliphatic or aromatic alkylammonium cation acting as a spacer between the perovskite layers, and the A and M cations and X anions form the perovskite framework. The 2D network consists of inorganic perovskite layers of corner-sharing [MX6]4− octahedra confined between interdigitating bilayers of long chain alkylammonium cations.30 The unit layers are held together by a combination of Coulombic and hydrophobic forces, which maintain the structural integrity. The existence of the 2D homologous M = Pb series of halide perovskites (RNH3)2(A)n−1MX3n+1 has been known for more than 25 years30,31 and was subsequently demonstrated for the M = Sn series.32−34 Although these 2D halide perovskites are analogues to the oxide perovskites described by Ruddlesden and Popper,35−37 the most wellstudied compounds in the class of 2D halide perovskites are the (RNH3)2MI4 (n = 1) family, where M is a group 1438 or lanthanide39 metal, and little is known for the higher members. The initial discovery of the 2D perovskites immediately drew attention because these systems can be regarded as natural multiple-quantum-wells in which the semiconducting inorganic layers act as potential “wells” and the insulating organic layers act as potential “barriers”.40,41 The electronic confinement in the perovskite 2D semiconductors in subnanometer layers induces the generation of stable excitons with unusually high binding energy and a Bohr radius that extends beyond the limits of a

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EXPERIMENTAL SECTION

Starting Materials. All chemicals were purchased from SigmaAldrich and used as received. Methylammonium iodide (MAI) was synthesized by neutralizing equimolar amounts of a 57% w/w aqueous hydriodic acid (HI) and 40% w/w aqueous methylamine (CH3NH2) (pH ≈ 7). The white precipitate was collected by evaporation of the solvent using rotary evaporation at 60 °C under reduced pressure. For convenience, we denote the 3D CH3NH3PbI3 perovskite as MAPbI3 and the 2D (n-CH3(CH2)3NH3)2(CH3NH3)n−1PbnI3n+1 family as (BA)2(MA)n−1PbnI3n+1 (n = 4, 3, 2, 1) throughout. Syntheses. MAPbI3 (n = ∞). PbO powder (2232 mg, 10 mmol) was dissolved in a mixture of 57% w/w aqueous HI solution (10.0 mL, 76 mmol) and 50% aqueous H3PO2 (1.7 mL, 15.5 mmol) by heating to boiling under constant magnetic stirring for about 5 min, which formed a bright yellow solution. Subsequent addition of solid CH3NH3Cl (675 mg, 10 mmol) to the hot yellow solution initially caused the precipitation of a black powder, which rapidly redissolved under stirring to afford a clear bright yellow solution. The stirring was then discontinued, and the solution was left to cool to room temperature and left to stand overnight to afford black polyhedral crystals. The crystals were collected by suction filtration and dried under reduced pressure. Yield 3.8 g (60%). Diffuse reflectance infrared Fourier transformed (DRIFT) spectrum, (KBr, cm−1): 3180br, 2823w, 2711w, 2485w, 2383w, 1820w, 1581s, 1467s, 1248m, 960s, 910s, 490m. (BA)2(MA)3Pb4I13 (n = 4). PbO powder (2232 mg, 10 mmol) was dissolved in a mixture of 57% w/w aqueous HI solution (10.0 mL, 76 mmol) and 50% aqueous H3PO2 (1.7 mL, 15.5 mmol) by heating to B


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Chemistry of Materials boiling under constant magnetic stirring for about 5 min, which formed a bright yellow solution. Subsequent addition of solid CH3NH3Cl (507 mg, 7.5 mmol) to the hot yellow solution initially caused the precipitation of a black powder, which rapidly redissolved under stirring to afford a clear bright yellow solution. In a separate beaker, n-CH3(CH2)3NH2 (248 μL, 2.5 mmol) was neutralized with HI 57% w/w (5 mL, 38 mmol) in an ice bath resulting in a clear pale yellow solution. Addition of the n-CH3(CH2)3NH3I solution to the PbI2 solution initially produced a black precipitate, which subsequently dissolved under heating the combined solution to boiling. The stirring was then discontinued, and the solution was left to cool to room temperature during which time black rectangular-shaped plates started to crystallize. The precipitation was deemed to be complete after ∼2 h. The crystals were isolated by suction filtration and thoroughly dried under reduced pressure. Yield 2.1 g (33% based on total Pb content). DRIFT spectrum, (KBr, cm−1): 3174br, 2962w, 2929w, 2725m, 2485w, 2382w, 1815w, 1577s, 1468s, 1250m, 1149m, 1072m, 1034m, 960m, 912s, 785w, 735w, 476m. (BA)2(MA)2Pb3I10 (n = 3). PbO powder (2232 mg, 10 mmol) was dissolved in a mixture of 57% w/w aqueous HI solution (10.0 mL, 76 mmol) and 50% aqueous H3PO2 (1.7 mL, 15.5 mmol) by heating to boiling under constant magnetic stirring for about 5 min, which formed a bright yellow solution. Subsequent addition of solid CH3NH3Cl (450 mg, 6.67 mmol) to the hot yellow solution initially caused the precipitation of a black powder, which rapidly redissolved under stirring to afford a clear bright yellow solution. In a separate beaker, n-CH3(CH2)3NH2 (327 μL, 3.33 mmol) was neutralized with HI 57% w/w (5 mL, 38 mmol) in an ice bath resulting in a clear pale yellow solution. Addition of the n-CH3(CH2)3NH3I solution to the PbI2 solution initially produced a black precipitate, which was subsequently dissolved under heating the combined solution to boiling. The stirring was then discontinued, and the solution was left to cool to room temperature during which time deep-red/purple rectangular-shaped plates started to crystallize. The precipitation was deemed to be complete after ∼2 h. The crystals were isolated by suction filtration and thoroughly dried under reduced pressure. Yield 2.5 g (36% based on total Pb content). DRIFT spectrum, (KBr, cm−1): 3174br, 2962w, 2929w, 2873w, 2713m, 2463w, 2382w, 1810w, 1573s, 1468s, 1389w, 1335w, 1255w, 1149m, 1072m, 1017m, 1001m, 964m, 914s, 785w, 748w, 735w, 484m. (BA)2(MA)Pb2I7 (n = 2). PbO powder (2232 mg, 10 mmol) was dissolved in a mixture of 57% w/w aqueous HI solution (10.0 mL, 76 mmol) and 50% aqueous H3PO2 (1.7 mL, 15.5 mmol) by heating to boiling under constant magnetic stirring for about 5 min, which formed a bright yellow solution. Subsequent addition of solid CH3NH3Cl (338 mg, 5 mmol) to the hot yellow solution initially caused the precipitation of a black powder, which rapidly redissolved under stirring to afford a clear bright yellow solution. In a separate beaker, n-CH3(CH2)3NH2 (694 μL, 7 mmol) was neutralized with HI 57% w/w (5 mL, 38 mmol) in an ice bath resulting in a clear pale yellow solution. Addition of the n-CH3(CH2)3NH3I solution to the PbI2 solution initially produced a black precipitate, which was subsequently dissolved under heating the combined solution to boiling. The stirring was then discontinued, and the solution was left to cool to room temperature during which time cherry red rectangular-shaped plates started to crystallize. The precipitation was deemed to be complete after ∼2 h. The crystals were isolated by suction filtration and thoroughly dried under reduced pressure. Yield 3.0 g (41% based on total Pb content). DRIFT spectrum, (KBr, cm−1): 3174br, 2962w, 2929w, 2880w, 2721w, 2474w, 2389w, 1808w, 1574s, 1468s, 1385w, 1335w, 1255w, 1149m, 1070m, 1017m, 1002m, 968m, 914s, 785w, 748m, 737m, 476m. (BA)2PbI4 (n = 1). PbO powder (2232 mg, 10 mmol) was dissolved in a mixture of 57% w/w aqueous HI solution (10.0 mL, 76 mmol) and 50% aqueous H3PO2 (1.7 mL, 15.5 mmol) by heating to boiling under constant magnetic stirring for about 5 min, which formed a bright yellow solution. Subsequent addition of liquid. In a separate beaker, n-CH3(CH2)3NH2 (924 μL, 10 mmol) was neutralized with HI 57% w/w (5 mL, 38 mmol) in an ice bath resulting in a clear pale yellow solution. Addition of the n-CH3(CH2)3NH3I solution to the PbI2

solution initially produced a black precipitate, which was subsequently dissolved under heating the combined solution to boiling. The stirring was then discontinued, and the solution was left to cool to room temperature during which time orange rectangular-shaped plates started to crystallize. The precipitation was deemed to be complete after ∼2 h. The crystals were isolated by suction filtration and thoroughly dried under reduced pressure. Yield 3.5 g (49% based on total Pb content). DRIFT spectrum, (KBr, cm−1): 3025br, 2960w, 2929w, 2880w, 2739w, 2472m, 2389w, 1830w, 1572s, 1475s, 1385w, 1390m, 1255w, 1159m, 1081m, 1047m, 1004m, 968m, 920s, 787w, 737m, 474m. Characterization. Single-crystal X-ray diffraction data were collected using an image plate STOE IPDS II diffractometer using Mo Kα radiation (λ = 0.71073 Å), operating at 50 kV and 40 mA. Data reduction and numerical absorption corrections were performed using the X-AREA suite. Single-crystals of the (BA)2(MA)n−1PbnI3n+1 compounds were mounted on the end of the glass tip directly from the mother liquor, and the use of glue or dispersion oil was avoided due to the clear deterioration of the crystals. All structures were solved by direct methods and refined by full-matrix least-squares on F2 using the SHELXTL-2013 program package.61 The PLATON62 functions operating within the WinGX platform63 were used to ensure the validity of the refined crystal structures. The structures were refined based on the following considerations. The perovskite lattice {PbnI3n+1} was refined anisotropically without any constraints on the Pb and I atoms. The highly anisotropic shape of the thermal ellipsoids of coordinated iodide ions appears due to a high thermal motion relative to the Pb−I bond vector. This is typical for the halide perovskites structures at room temperature, and therefore no special treatment was used.64 All organic atoms were refined isotropically. Restraints were applied to the C−C and C−N bond lengths. Each cation was treated as having equivalent thermal parameters. Significant disorder exists in the interlayer cations, particularly for the CH3CH2− tail of butylammonium (the NH3CH2CH2− head is relatively stable), causing the atoms to move and destabilize the refinement. The bond length restraints cannot account for the motion of the interlayer cations. Therefore, the cations were modeled manually on idealized positions based on reasonable bond length and bond angle parameters. Thus, the whole butylammonium cations were constructed on top of the mirror planes imposed by the space group by slightly modifying the positions of the off-plane Q-peaks found in the electron density map. The mirror planes have the (x, y, 1/4)/(x, y, 3/4) and (x, y, 1/2) coordinates for the Ccmm (and Cc2m) and Acam space groups, respectively, whereas in the case of C2cb where no mirror plane is crystallographically imposed, the cations were modeled based on the initial Q-peaks. In the final steps of the refinement, the {(CH3NH3)n−1PbnI3n+1} structure segment was refined without restraints, while the butylammonium cations were fixed. The purpose of this treatment was to generate chemically reasonable models without disorder. The modeled cations increase the R-values slightly compared to the unrestrained or disordered refinement and make the e.s.u.’s unrealistically small. Because of this, a BLOC instruction was used to finalize the refinement, which fixed the positions of the butylammonium cations and refined the rest of the atoms separately, to maintain the proper e.s.u.’s of the perovskite layers. Powder and film X-ray diffraction patterns were collected using a Rigaku MiniFlex 600 X-ray diffractometer (Cu Kα, 1.5406 Å) operating at 40 kV and 15 mA. Scanning electron microscope (SEM) images were acquired at an accelerating voltage of 10−20 kV using either a Hitachi SU8030 or a Hitachi SU3400 instruments equipped with Oxford X-max 80 SDD EDS detectors. Optical diffuse-reflectance spectra were collected at room temperature using a Shimadzu UV-3600 PC double-beam, double-monochromator spectrophotometer on powdered samples using BaSO4 as a 100% reflectance reference. The samples were irradiated with a halogen (NIR/Visible, 50W, 2000H) and a D2 lamp (UV, 2000H), and the spectra were recorded using a PbS photoconductive element (NIR) and a PMT (R928) detectors from 2500−200 nm. Band gaps were determined as described elsewhere.65,66 DRIFT spectra were recorded on a Nicolet 6700 IR spectrometer in the 400−4000 cm−1 spectral region with KBr beam splitter. C


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Chemistry of Materials Table 1. Crystal Data and Structure Refinement for (BA)2(MA)n−1PbnI3n+1 (n = 2−4) at 293(2) Ka empirical formula formula weight temperature wavelength refinement method cryst syst space group color unit cell dimensions a, Å b, Å c, Å α = β = γ, deg volume, Å3 Z density (calc), g/cm3 absorption coeff., mm−1 F(000) cryst size (mm3) θ range completeness to θ index ranges

reflns collected independent reflns data/restraints/parameters GOF final R indices [I > 2σ(I)] R indices [all data] largest diff. peak and hole (e Å−3)

(BA)2(MA)Pb2I7 (n = 2)

(BA)2(MA)2Pb3I10 (n = 3)

(BA)2(MA)3Pb4I13 (n = 4)

1483.04 293(2) K 0.71073 Å full-matrix least-squares on F2 orthorhombic Cc2m red 8.9470(4) 39.347(2) 8.8589(6) 90°

2103.00

2722.95

orthorhombic C2cb dark red 8.9275(6) 51.959(4) 8.8777(6) 90°

orthorhombic Cc2m black 8.9274(4) 64.383(3) 8.8816(4) 90°

3118.7(3) 4 3.159 17.712 2560 0.140 × 0.081 × 0.025 2.521−25.000° 99.7% −10 ≤ h ≤ 10 −46 ≤ k ≤ 46 −10 ≤ l ≤ 10 9452 2732 [Rint = 0.0365] 2732/3/79 1.145 Robs = 0.0529 wRobs = 0.1394 Rall = 0.0638 wRall = 0.1478 0.933 and −1.687



R = Σ∥Fo| − |Fc∥/Σ|Fo|, wR = {Σ[w(|Fo|2 − |Fc|2)2]/Σ[w(|Fo|4)]}1/2 and n = 2, w = 1/[σ2(Fo2) + (0.0768P)2 + 22.2041P]; n = 3, w = 1/[σ2(Fo2) + (0.0916P)2 + 15.8934P]; n = 4, w = 1/[σ2(Fo2) + (0.0736P)2], where P = (Fo2+ 2Fc2)/3. a

Raman spectra were recorded on a DeltaNu Advantage NIR spectrometer equipped with a CW diode laser (785 nm, 60 mW) and a CCD camera detector in a backscattered geometry. The powdered samples were packed in standard melting point capillaries (0.8 mm ID). Photoluminescence spectra were collected on oriented rectangular crystals of (BA)2(MA)n−1PbnI3n+1 (n = 4, 3, 2, 1) and rhombic dodecahedral crystals of MAPbI3 using Horiba LabRam Evolution highresolution confocal Raman microscope spectrometer (600 g/mm diffraction grating) equipped with a diode CW laser (473 nm, 25 mW) and a Synapse CCD camera. The incident laser beam was parallel to the (010) direction of the crystals and focused at ∼1 μm spot size. Unless stated otherwise, the maximum power output of the laser source was filtered to 0.1% of the maximum power output. Second harmonic generation (SHG) and third harmonic generation (THG) spectra were recorded on powdered (BA)2(MA)n−1PbnI3n+1 perovskite samples and a powdered α-SiO2 reference for a particle size distribution from 90 μm−125 μm placed in standard melting point capillaries (0.8 mm ID) using reflection geometry with a fiber-optic bundle as described previously.67,68 The bundle was coupled to a Horiba iHR320 spectrometer (600 g/mm diffraction grating) equipped with a Synapse CCD camera. To prevent absorption effects during the SHG measurement, the fundamental beam from an optical parametric oscillator, pumped by a Nd:YAG laser, was tuned to λ = 1800 nm so that the SHG photon energy (ℏωSHG = 1.38 eV) was below the band gap of all the materials. The incident pulse energy was tuned to ∼30 μJ using a Glan-Thompson polarizer with a beam spot size of ∼0.5 mm. All detector exposure times were scaled to 60 s. The THG photon

energy on the other hand was above the band gap (except (BA)2PbI4) (ℏωTHG = 2.07 eV), and therefore THG was significantly underestimated because of the self-absorption processes in the material. However, this absorption effect was properly taken into account for the comparison between SHG and THG using the absorbance data. Computational Methods. We investigated three of the synthesized compounds using density functional theory: (BA)2PbI4, (BA)2(MA)2Pb3I10, and (BA)2(MA)3Pb4I13. All calculations were performed with the Vienna ab initio Simulation Package (VASP)69,70 using projector augmented-wave (PAW) potentials within the PBEsol exchange-correlation functional.71,72 We initially generated a centrosymmetric and a noncentrosymmetric structure for each of the three compounds through manipulation of the organic cations, which was inferred by the symmetries identified in the experimental structure refinements. Subsequently, we fixed the lattice parameters and cell volume to those determined in our experiments (Table 1) and fully relaxed the internal atomic positions until the forces were less than 1 meV Å−1 using a 500 eV plane-wave cutoff and 5 × 1 × 5 Monkhorst− Pack mesh73 to sample the Brillouin zone. We computed the density of states and band structures in the noncentrosymmetric crystal structures, which were found to be the lowest energy phases in each case using a 5 × 1 × 5 Monkhorst−Pack mesh and increased 550 eV planewave cutoff.

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RESULTS AND DISCUSSION Synthesis. The 2D (CH 3 (CH 2 ) 3 NH 3 ) 2 (CH 3 NH 3 ) n−1 PbnI3n+1 family of perovskite compounds (n = 1−4) was D


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Figure 1. SEM images (top) and photographs of the (BA)2(MA)n−1PbnI3n+1 perovskite crystals (bottom) (scale bars = 200 μm).

The above procedure consists of a scalable and efficient method of preparing these layered compounds in pure form. The reaction scheme described above can be conveniently used to prepare the 2D perovskite members (n = 2, 3, and 4), which contradicts a proposed axiom that Ruddlesden−Popper halide perovskites with n > 3 cannot be isolated in pure form.76 Note that when we attempted to isolate higher members of the 2D series by adjusting the Pb2+/MA+ ratio, the reactions always yielded the n = 4 member. Nevertheless, we believe that n > 4 members of the 2D series may be isolated by varying either the nature of the “spacer” (the long chain amine) or the “perovskitizer” (the small organic cation)a hypothesis whose testing lies outside the scope of the present work. To the best of our knowledge, this is the only method by which the individual compounds with n > 2 can be isolated on a gram scale. The availability of pure bulk materials with n > 2 facilitates in-depth studies of their photophysical properties and the deployment of the title compounds in future solid-state solar cells.27,77 Crystal Structure Description. Basic Structural Characteristics. The crystal structures of the (BA)2(MA)n−1PbnI3n+1 (n = 1−4) compounds are shown in Figure 2, and selected crystallographic information is tabulated in Tables 1 and 2. Detailed crystallographic data are provided in the Supporting Information (Section S2). All compounds crystallize in orthorhombic space groups; (BA)2PbI4 crystallizes in the primitive centrosymmetric Pbca space group, as previously reported,38,78 whereas (BA)2(MA)Pb2I7, (BA)2(MA)2Pb3I10, and (BA)2(MA)3Pb4I13, which have been determined in the present study, crystallize in the polar (C2v), base-centered Cc2m, C2cb, and Cc2m, space groups, respectively. The noncentrosymmetric configurations represent the case where all MA cations are oriented in the same direction or nearly so, which results in an unquenched net dipole moment within the unit cell. The respective centrosymmetric (D2h) space groups Ccmm,

synthesized from a stoichiometric reaction between PbI2, CH3NH3I (MAI), and n-butylamine (BA). A homogeneous solution of concentrated, I2-free hydroiodic acid containing stoichiometric amounts of PbI2 and CH3NH3I (MA), according to the desired composition, was allowed to react with half the stoichiometric amount of BA by addition of the neutralized base into the boiling acid solution under vigorous stirring. This highly exothermic reaction resulted in the formation of a clear, bright yellow solution, which upon cooling to ambient conditions precipitates into the layered perovskite compounds (n = 2−4) in the form of colorful rectangular plates with the spectral range spanning from red to black (Figure 1). We find that the use of BA as the reaction limiting reagent is essential in obtaining the compounds in pure form; at the same time, it is detrimental to the reaction yield, which is limited to ∼50% based on the total Pb content due to the high solubility of CH3NH3PbI3 in the HI/H3PO2 solvent medium. The employment of this concept takes advantage of the different solubility of the 2D perovskite members in the solvent medium, which increases as the number of perovskite layers increases up to a maximum solubility obtained for the CH3NH3PbI3 end-member (∼0.6 M at the boiling point). Interestingly, the solubility trend is inverted in intermediate polarity solvents such as acetone or acetonitrile. In these solvents, the 2D compounds are relatively soluble, where the member with the largest organic fraction becomes the most soluble (i.e., n = 1). Using HI/H3PO2 as a solvent medium, we find that if BA is used stoichiometrically or MA is used in excess, as has been previously recommended,26,54,74,75 then the end product is contaminated either with the (n−1) member in the former case, which acts as a kinetic barrier, or with the (n + 1) member in the latter case, which acts as a thermodynamic sink during synthesis of the compounds. Thus, the driving force to guide the reaction into a single n-member product comes from careful control of the stoichiometry based on the “limiting reagent” principle. E


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Figure 2. Crystal structures of the 2D lead iodide perovskites, (BA)2(MA)n−1PbnI3n+1, extending from n = 1 to n = ∞. The L-value denotes the thickness of the inorganic layer in each compound. The numerical values refer to the distance between the terminal iodide ions of each layer and were determined directly from the refined crystal structures.

BA cations.30 Similar to the parent compound, the layers consist of tilted, corner-sharing [PbI6]4− octahedra that propagate in two directions (the ac plane), whereas in the third dimension (the b-axis), the octahedral sheets are physically disconnected by the intercalated organic bilayers. The separation of the inorganic layers results in an increase of the lattice constants as the incorporation of the BA spacers requires a gap of ∼7.8 Å between the perovskite layers; the simplest case of n = 1 has a thickness of ∼6.4 Å, which roughly corresponds to the sum of the two Pb−I bond lengths. As the 2D perovskite layers grow thicker by introducing MA cations in the crystal structure, the unit cell incrementally expands by addition of a single perovskite layer at a time. These changes in the unit cell can be monitored by X-ray diffraction (XRD), which characteristically reveals an additional low angle reflection for each added perovskite layer (Figure 3a,b). Namely, the n = 2 analogue shows two evenly spaced reflections, the n = 3 analogue three, and the n = 4 analogue four reflections below ∼2θ = 14°. The 2θ cut off for numbering is based on the fact that at this angle the d-spacing matches the distance between the discrete perovskite layers in both the 2D [(111) reflection] and 3D perovskites [(110) reflection]. This simple numerical estimation of the crystal structure could be a useful analytical tool for further study of the (BA)2(MA)n−1PbnI3n+1 or related Ruddlesden−Popper systems. The reported Ruddlesden−Popper perovskites have the same gross symmetry determined by the positions of the heavy atoms (i.e., the [PbnI3n+1]−(n+1) layers), but the space groups differ because the nature of the organic spacer indirectly dictates the details of crystal symmetry of the final layered compound. This occurs because the organic spacer component influences the relative rotation and tilting of the [PbI6]4− octahedra rather than the MA cations in the cages or the “MAn−1PbnI3n+1” layers themselves. A similar point has been previously made for the case of the (CH3(CH2)mNH3)2PbI4 (n = 1, m = 3−17) series,78,84,85 where it was shown that the compounds undergo a series of temperature-induced, structural phase transitions, with the critical temperature of the transition displaying a strong dependence on the length of the alkyl chain (m-value). Interestingly, the set of phase transitions in the 2D (RNH3)2PbI4 perovskites

Acam, and Ccmm, which provide an equally good structural solution, were considered to be less likely on the basis of the wellknown classical crystallographic Hamilton statistical significance criteria,79,80 whereas SHG measurements and DFT calculations, independently, bear out this assignment. Nevertheless, the structural solutions in the centrosymmetric space groups are also provided as a supplement in this work (see Supporting Information) in aid of the quest for addressing the ferroelectric behavior of these compounds by other researchers, as will be discussed in detail in the SHG properties section. We consider the centrosymmetric crystal structures to represent the average structure of the bulk materials, whereas the polar crystal structures may be considered as domains of a polycrystal and may account for the presence of ferroelectric domains where local polarization can be observed. This is supported by our 0 K DFT calculations, presented below, which find the noncentrosymmetric structures to be lower in energy than the centrosymmetric phases. The (BA)2(MA)n−1PbnI3n+1 family joins a handful of other structurally characterized multilayer hybrid halide perovskites (n > 1) with a group 14 metal in the octahedral site. These include the (n = 2) (PEA)2(MA)Pb2I7 (in P1̅)30 and the very recently reported (n = 3) (PEA)2(MA)2Pb3I10 (in P1)26 perovskites (PEA is the phenylethylammonium cation), although the crystal structures have been only tentatively assigned. For the n = 3 Sn compound (BA)2(MA)2Sn3I10 (in Cmca)33 and the n = 2 Pb compounds, (MBA)2(MA)Pb2I7 (MBA is the 4-methylbenzylammonium cation),81 (TMA)2(MA)Pb2I7 (TMA is the 2-thienylmethylammonium cation),82 and (ABA)2(MA)Pb2I7 (ABA is the 4-ammoniobutanoic acid cation),83 the Pbcn, Aba2 (≡ C2cb) and C2cb space groups, respectively, have been rigorously assigned. The unit cell in all compositions contains two inorganic layers and two organic bilayers (4 f.u./cell) with the inorganic layer defining the ac plane (ab plane for n = 1) and the organic bilayers intercalating along the b-axis (c-axis for n = 1). Each 2D inorganic layer relates to the tetragonal parent 3D compound (n = ∞) by slicing the perovskite along the (110) plane such that some of the oriented MA cations are partially (n = 2, 3, 4, ...) or fully (n = 1) substituted by the terminal F


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Table 2. Selected Bond Lengths [Å] and Bond Angles [°] for (BA)2(MA)n−1PbnI3n+1 at 293(2) K with Estimated Standard Deviations in Parentheses (BA)2PbI4 (n = 1)a

(BA)2(MA)Pb2I7 (n = 2)

label

distances

label

distances

label

distances

Pb−I(1) Pb−I(2) Pb−I(2)′

3.205(3) 3.1554(18) 3.2072(18)

Pb(1)−I(7) Pb(1)−I(1) Pb(1)−I(2) Pb(1)−I(5) Pb(2)−I(6) Pb(2)−I(3) Pb(2)−I(4) Pb(2)−I(5)

3.081(7) 3.1561(14) 3.1749(16) 3.249(8) 3.075(6) 3.169(2) 3.1715(19) 3.278(8)

Pb(1)−I(4) Pb(1)−I(1) Pb(1)−I(4)′ Pb(2)−I(5) Pb(2)−I(2) Pb(2)−I(3) Pb(2)−I(2)′ Pb(2)−I(3)′

3.065(8) 3.157(3) 3.249(8) 3.047(4) 3.152(6) 3.160(6) 3.170(6) 3.179(6)

(BA)2PbI4 (n = 1)a label

angles

(BA)2(MA)Pb2I7 (n = 2) label

Pb−I(2)-Pb′ 155.38(8) Pb(1)′-I(1)−Pb(1) Pb(1)−I(2)−Pb(1)′ Pb(2)−I(3)−Pb(2)′ Pb(2)′-I(4)−Pb(2) Pb(1)−I(5)−Pb(2)

a

(BA)2(MA)2Pb3I10 (n = 3)

(BA)2(MA)2Pb3I10 (n = 3)

MAPbI3 (n = ∞)b

(BA)2(MA)3Pb4I13 (n = 4) label

distances

Pb(1)−I(4) 3.12(2) Pb(1)−I(3) 3.14(2) Pb(1)−I(6) 3.154(8) Pb(1)−I(2) 3.164(8) Pb(2)−I(5) 3.122(17) Pb(2)−I(13) 3.156(8) Pb(2)−I(1) 3.170(8) Pb(2)−I(4) 3.26(2) Pb(3)−I(10) 3.06(2) Pb(3)−I(8) 3.161(8) Pb(3)−I(12) 3.172(9) Pb(3)−I(3) 3.27(2) Pb(4)−I(11) 3.01(2) Pb(4)−I(9) 3.164(8) Pb(4)−I(7) 3.175(9) Pb(4)−I(5) 3.330(15) (BA)2(MA)3Pb4I13 (n = 4)

label

distances

Pb−I(1) Pb−I(1)′ Pb−I(2)

3.125(8) 3.196(8) 3.1613(8)

MAPbI3 (n = ∞)b

angles

label

angles

label

angles

173.1(3) 164.3(3) 165.1(4) 167.6(3) 165.62(11)

Pb(1)−I(1)−Pb(2) Pb(2)−I(2)−Pb(2)′ Pb(2)−I(3)−Pb(2)′ Pb(1)−I(4)−Pb(1)′

169.52(12) 172.28(17) 164.57(15) 171.3(3)

Pb(2)′-I(1)−Pb(2) Pb(1)−I(2)−Pb(1)′ Pb(1)−I(3)−Pb(3) Pb(1)−I(4)−Pb(2) Pb(2)−I(5)−Pb(4) Pb(1)−I(6)−Pb(1)′ Pb(4)−I(7)−Pb(4)′ Pb(3)′-I(8)−Pb(3) Pb(4)′-I(9)−Pb(4) Pb(3)−I(12)−Pb(3)′ Pb(2)−I(13)−Pb(2)′

166.1(12) 168.8(14) 167.6(12) 167.2(6) 165.0(9) 172.4(12) 162.6(14) 170.4(14) 172.9(14) 165.5(13) 172.7(14)

label

angles

Pb−I(1)-Pb′ 180 Pb−I(2)-Pb′ 163.55(6)

Centrosymmetric reference.38 bNoncentrosymmetric reference.64

as a guiding principle in predicting the crystal structures of the higher n-members of the (BA)2(MA)n−1PbnI3n+1 family. Comparison between the 2D and 3D Perovskites. To elaborate on the structural differences among the compounds, let us first consider the fundamental monolayer consisting of “[Pb9I36]18‑” units shaded in yellow in Figure 4. In the simplest case of the (BA)2PbI4 (n = 1) compound (Figure 4, left column), a single slab is isolated between the BA ions, which align perpendicular to it as viewed along the [110] direction; the -NH3+ “head” of the ammonium cation points toward the center of the rhombic cavity generated by four adjacent octahedra, while the −CH3 “tail” points toward the interlayer space. The accommodation of the -NH3+ “head” in the near vicinity of the corner-connected octahedra (presumably by electrostatic attraction and aided by NH2−H···I−Pb hydrogen bonding) forming the perovskite slab forces the octahedra to tilt and induces a 155.20° Pb−I−Pb angle in (BA)2PbI4. It is important to point out that the corner-connected octahedra in a single slab always need to rotate in an opposite sense (out-of-phase) to maintain their structural integrity. Interestingly, the other endmember, MAPbI3 (n = ∞), exhibits similar behavior with respect to octahedral tilting although in a more complex fashion as the octahedral tilting also occurs in the third direction. To better illustrate this point, we define the different “[Pb9I36]18‑” units expanding along the c-axis shaded in blue, green, or yellow for clarity purposes (Figure 4, right). The individual “[Pb9I36]18‑”

strongly resembles the properties of the 3D APbI3 counterparts, where small changes in the A+ cation (from Cs to CH3NH3 to HC(NH2)2) result in similar structural changes in the inorganic {PbI3}− frameworks, where α-, β-, and γ-phases evolve as a function of temperature.64 It is very likely that similar phase transitions exist for the higher members of the 2D perovskite series as well, but because of their complexity, we limit our discussion to the room temperature structures of these compounds in the present manuscript. Such considerations allow us to suggest that the role of the A cations in the 3D perovskites is assumed by the spacer cations in the 2D perovskites. Given the complexity in assigning the proper space group in Ruddlesden−Popper perovskites, as discussed above, we generated the simulated precession images of reciprocal space from the raw experimental single-crystal diffraction data for the (BA)2(MA)n−1PbnI3n+1 family (Figure 3c,d). The data confirm the assigned space groups (Cc2m or Ccmm for n = 2, C2cb or Acam for n = 3, and Cc2m or Ccmm for n = 4) based on systematically absent reflection conditions; moreover, they uncover an underlying trend that compounds with an odd number of layers (n = 1, 3) systematically adopt a higher symmetry configuration (in terms of increasing space group number) in comparison with the structures consisting of an even number of layers (n = 2, 4). Such a distinction between crystal structures of the odd and even layers, already established for oxide perovskites of the Ruddlesden−Popper type,86,87 can serve G


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Figure 3. (a, b) X-ray diffraction patterns (Cu Kα) of the (BA)2(MA)n−1PbnI3n+1 perovskites: (a) Bulk powder diffraction and (b) close-up views of the characteristic regions between 2θ = 2−15° (left) and 2θ = 25−30° (right) where some pronounced differences of the 2D versus 3D perovskites occur. Miller hkl indices are included to describe the strong preferred orientation of the 2D materials. (c, d) Precession images of the BA2MAn−1PbnI3n+1 simulated from single-crystal X-ray diffraction (Mo Kα). (c) Images along the (001) direction emphasizing on the characteristic allowed reflections and reflection conditions for the assigned space groups. Note that for n = 1, the corresponding direction is (01̅0). (d) Images along the (101̅) direction emphasizing on the ordering of the 2D perovskites with increasing number of layers and with respect to the parent compounds n = 1 and n = ∞ (top). The green circles correspond to the marked reflections in panel b, whereas the red circles highlight the number of distinct, extra reflections for the layered perovskites with respect to the 3D perovskite. Note that for n = 1 and n = ∞, the corresponding directions is (11̅0), whereas for n = 4, the (101) direction is shown for clarity.

Figure 4. Views of the (BA)2PbI4 single-layer (n = 1), (BA)2(MA)2Pb3I10 three-layer (n = 3) 2D, and MAPbI3 (n = ∞) 3D lead iodide perovskite crystal structures, highlighting their three-dimensional distortions. The defined “[Pb9I36]18‑” slabs represent a 3 × 3 building block of the perovskite in two dimensions and are used as a basic model to visualize the distortion. A single slab is composed of nine corner-connected octahedra (yellow) that distort out-of-phase along the perovskite plane. The tilting scheme in the n = 1 2D perovskite can be solely described by a single slab. As the perovskite thickness increases, the “[Pb9I36]18‑” slabs connect to form “[Pb27I81]27‑” boxes in three dimensions, which are the representative building units for the three-layer structure in the n = 3 2D perovskite and a fragment of the infinite layers in the n = ∞ 3D perovskite. This figure illustrates the difference in the distortion modes of the n = 3 perovskite emphasizing the in-plane and out-of-plane views of the perovskite slabs. Slabs belonging in different layers have been drawn in different colors (yellow, blue, green) to project the connectivity of the slabs in the third dimension. H


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Figure 5. Optical properties of the (BA)2(MA)n−1PbnI3n+1 perovskites (for n = 1, 2, 3 4, ∞). (a) Optical absorption of polycrystalline samples obtained from diffuse reflectance measurements converted using the Kubelka−Munk function (α/S = (1 − R)2/2R).65,66(b) Photoluminescence of oriented crystals with the wide planar facets oriented perpendicular to the laser beam (λexc = 473 nm) with the illumination vector pointing toward the (010) axis for n = 2, 3, 4, (001) for n = 1 and (100) for n = ∞. Plots a and b are scaled to match; the dashed lines in panel a correspond to the photoluminescence maximum in panel b and are used as a guide to the eye to illustrate the absorption features that correspond to the emission maxima. (c) Plot of the band gap and photoluminescence versus layer thickness with the latter expressed in the form of 1/n2. Note that in panels a−c the energy is plotted in a reciprocal scale and the wavelength in a linear scale. (d) Plot of the exciton energy, defined by convention as the difference in energy of the absorption edge to the emission maximum (see Table 3) versus layer thickness expressed in the form of 1/n2.

slabs fuse along the [001] direction to form a “[Pb27I81]27‑” box and 3D octahedral network within which the MA cations are contained. The MA cations align along the crystallographic c-axis with the -NH3+ groups again, as in the case of BA cations, pointing at the center of the I4 rectangle generated by four adjacent octahedra, causing a significant degree of distortion. As a result of this distortion, the Pb−I−Pb angle in MAPbI3 becomes 163.55°. The distinct structural features of MAPbI3, however, become evident when the perovskite begins to self-assemble into “[Pb27I81]27‑ boxes”, as shown in Figure 4. The boxes consist of three stacked “[Pb9I36]18‑” slabs, which themselves display an outof-phase octahedral distortion and connect along the direction defined by the C−N bond while displaying no rotation in the plane perpendicular to the C−N bond direction (the a0a0c− tilt pattern in Glazer notation, Figure 4, right column).88 This pattern is characteristic of the 3D CH3NH3MX3 perovskites in general and is opposite to the CsMX310 and HC(NH2)2MX364 counterparts, which adopt an in-phase tilting (a0a0c+ in Glazer notation) when the “[Pb9I36]18‑” slabs connect along the unique crystallographic direction. On the basis of the above characteristics, it becomes evident that the nature and orientation of the ammonium cation influence the tilting pattern in the hybrid perovskites and strongly impact the degree of distortion of the inorganic component. A unique situation arises in the (BA)2(MA)n−1PbnI3n+1 multilayered compounds where both MA and BA cations are present, owing to the competition between the two cations trying to satisfy their stereochemical demands. Because the situation is similar for n = 2, 3, and 4, we will limit our discussion to the n = 3 compound for simplicity (Figure 4, center column). Continuing with the “[Pb9I36]18‑ slab” concept, see Figure 4, we can clearly

observe that the configuration of the perovskite slab as well as the perovskite box in BA2(MA)2Pb3I10 differs from both of the parent compounds. The most obvious difference arises from the fact that the out-of-phase rotation axis now coincides with the orientation of BA cations’ (the [010] direction), whereas the plane parallel to the MA cation (the [101] direction) remains relatively undistorted. The term “relatively” is used because the decrease in symmetry from tetragonal (n = ∞) to orthorhombic (n = 3), which precludes a 180° Pb−I−Pb angle connectivity between the octahedra, introduces a small relative rotation. An examination of the placement of the -NH3+ “heads” reveals that the heads of the BA cations appear optimally “centered” in (BA)2PbI4 with respect to the perovskite slab, whereas the MA heads lose their ideal alignment by becoming slightly “offcentered” with respect to the perovskite slabs (unlike β-MAPbI3). In that sense, neither BA nor MA is in its preferred configuration, with BA trying to distort the perovskite along the b-axis and MA trying to distort the ac-plane. The net result of this competition is an overall distortion in the intermediate 2D perovskites, which occurs across two different directions (as opposed to (BA)2PbI4 and MAPbI3), but it is less severe compared to the parent compounds due to the competing interactions between MA and BA. For example, in (BA)2(MA)2Pb3I10, the distortion angles are (Pb−I−Pb)010 = 169.52°, (Pb−I−Pb)101 = 164.58°, and (Pb−I−Pb)101 = 172.26°. The mixed cation configuration as well as the competition between the constituting cations renders the multilayered perovskites fundamentally different from both the simple 2D single-layer compounds as well as the 3D perovskites and provides them with some unique properties, which will be discussed in detail below. I


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Chemistry of Materials Table 3. Optical Parameters of the (BA)2(MA)n−1PbnI3n+1 Perovskites compound

band gap Eg (eV)

excitonic absorption (eV)a

photoluminescence PL (eV)

(Eg − PL) (meV)

electron effective mass (me,xz)b

hole effective mass (mh,xz)b

(BA)2PbI4 (BA)2(MA)Pb2I7 (BA)2(MA)2Pb3I10 (BA)2(MA)3Pb4I13 MAPbI3

2.43 2.17 2.03 1.91 1.50

2.35 2.08 1.96 1.85 1.59

2.35 2.12 2.01 1.90 1.60

80 50 20 10 n/a

0.082

0.144

0.097 0.094

0.141 0.153

a

Position of the excitonic peak in the diffuse reflectance spectra. bCalculated at the DFT-PBEsol level in this work. mxz is reported as the average of the effective mass along the x and z directions and is in units of the bare electron mass.

Band Gaps, Electronic Structure, and Photoluminescence. Figure 5 displays the room-temperature absorption and photoluminescence spectra of bulk samples from the (BA)2(MA)n−1PbnI3n+1 series. The optical absorption edges are remarkably sharp and increase in energy with decreasing n value, from 1.50eV (n = ∞) to 2.43eV (n = 1), a property arising from the dimensional reduction of the 3D perovskite lattice into the lower dimensionality homologous structures (Figure 5a).89−91 From the experimental spectra, we assign the higher energy absorption edge to the band gap of the materials (Table 3). In all compounds, the onset of the optical absorption consists of sharp edges, although occasionally an absorption tail can be observed in samples of the n = 3 and n = 4 compounds possibly arising from small impurities of intergrown higher order homologous members. The sharp nature of the absorption edges points toward a direct band gap in all compounds. A similar band gap trend can be also observed from thin-films of these materials deposited by spin coating dimethylformamide (DMF) precursor solutions as we reported previously,27 confirming that thin films of pure 2D compounds can be made by direct deposition onto the substrates. Using DFT, we computed the electronic band structure and density of states for the n = 1, 3, and 4 structures (Figure 6). All compounds are semiconducting, with the valence band almost exclusively consisting of I 5p states with a small amount of Pb 6s character, while the bottom of the conduction band primarily consists of Pb 6p states (Figure 7). The three studied compounds also each exhibit a clear direct band gap. The band gaps are 1.99 eV for (BA)2PbI4, 1.78 eV for (BA)2(MA)2Pb3I10, and 0.96 eV for (BA)2(MA)3Pb4I13. Although consistently underestimated within DFT at the PBEsol level, we find the same general trend of decreasing band gap with increasing number of perovskite layers to be in agreement with our experimental results (Table 3). As with other Pb-based metal−organic halide perovskites,92 the bands near the valence and conduction edges are highly dispersive in reciprocal space along the corresponding real-space axes in which the PbI6 octahedra are corner-connected, indicating small hole and electron effective masses (mh and me, respectively) along the x and z directions. By applying the parabolic band approximation to the computed electronic structure, we find that mh and me are approximately 0.14 m0 and 0.08 m0 along both x and z, where m0 is the rest electron mass, and are nearly independent of the number of perovskite and spacing layers (Table 3). Because in the Ruddlesden− Popper structure the inorganic semiconducting network is disconnected along the y axis (Figure 1), it has a nearly flat bands along this direction, indicating that the effective masses along this direction are extremely large, which would tend to restrict the charge transport to the xz plane. When spin−orbit interactions (SOIs) are included in our calculations, for example,

Figure 6. Electronic band structure of the polar configurations of selected (BA)2(MA)n−1PbnI3n+1 perovskites. (a) (BA)2PbI4 (n = 1), (b) (BA)2(MA)2Pb3I10 (n = 3), and (c) (BA)2(MA)3Pb4I13 (n = 4) along the Γ(0,0,0)-X/S(1/2,0,0)-U/R(1/2,0,1/2)-Z(0,0,1/2)-Γ(0,0,0) path (solid vertical lines) throughout the Brillouin zone. The Fermi level is set to 0 eV and indicated by the horizontal broken red line.

Figure 7. Density of states of the polar configurations of selected (BA)2(MA)n−1PbnI3n+1 perovskites. (a) (BA)2PbI4 (n = 1), (b) (BA)2(MA)2Pb3I10 (n = 3), and (c) (BA)2(MA)3Pb4I13 (n = 4). Iodine 5p (red) and lead 6s (blue) states make up the top of the valence band, while Pb 6p states (green) form the bottom of the conduction band.

in the n = 1 phase, we find the band gap decreases to 1.25 eV, and there are changes to the band degeneracies, but the band curvature remains the same; furthermore, the noncentrosymmetric configuration discussed above remains lower in energy than the centrosymmetric. Unfortunately, the large size of the unit cells prevented us from including SOI in the higher order compounds, but we expect the same trends to persist. J


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Unlike the absorption spectra, which are characterized by an absorption edge and an excitonic peak, the photoluminescence spectra consist of one single emission peak corresponding to the energy value of the excitonic peak in the absorption spectra. This is because the relaxation pathway is dictated by direct radiative recombination of an exciton, which is typical for an excitonic material. As predicted from the above 2D effects, the energy differences between the band gap and the photoluminescence emission values (or the excitonic absorption values) decrease with increased n (80 meV for n = 1 to almost 10 meV for n = 4) as shown in Table 3; here, the emphasis is given to the correct trend in this energy difference as a function of n, but the actual exciton binding energies should be more precisely determined at low temperatures. For the 3D perovskite, it was challenging to extract the correct bandgap energy due to the spectral proximity between the primary absorption edge and the excitonic peak. In fact, the primary absorption onset at ∼1.50 eV in Figure 5 (panel a) should be interpreted as the “low-energy tail” of the excitonic transition, not the fundamental band edge (Eg) that is located slightly higher in energy due to the contribution of the exciton binding energy. This picture is indeed in line with the spectral matching of the photoluminescence maximum (∼1.6 eV) and the absorption saturation (∼1.59 eV). Another remarkable feature of the photoluminescence spectra is the line shape of the photoluminescence peaks, which is symmetric for the 3D perovskites but not for the 2D analogues. Such a behavior has been recently assigned to trap states in the perovskite framework characteristic of both 2D and 3D perovskites, but the number of trap states has been shown to be higher in the 2D compounds.56,57 Electronic Structure Calculations and Phase Stability. To obtain further insight into the structural character and electronic structure of these 2D perovskite systems, we performed several investigations using DFT calculations. We first computed the relative energies of (BA)2PbI4 (n = 1), (BA)2(MA)2Pb3I10 (n = 3), and (BA)2(MA)3Pb4I13 (n = 4) in both the centrosymmetric and noncentrosymmetric crystal structures to evaluate their differences in stability. In (BA)2PbI4, we found that the Pbc21 and Pcab phases are energetically degenerate within the precision of our calculations, with the Pbc21 structure negligibly lower in energy by 0.58 meV/f.u. Although this could support the classification of (BA)2PbI4 into the polar Pbc21 symmetry, it more accurately indicates that intergrowths of the two phases is likely, granting local variations in noncentrosymmetry. Remarkably, when the number of perovskite layers is increased to n = 3 and n = 4, the noncentrosymmetric phase becomes significantly more stable than the centrosymmetric structure. In (BA)2(MA)2Pb3I10, we find that the C2cb structure is lower in energy than the centric (Acam) phase by 527.9 meV/f.u., while in (BA)2(MA)3Pb4I13, the formation energy of the centric structure (Ccmm) lies much higher than that of the acentric (Cc2m) space group, which is strongly favored. The results are in good agreement with similar calculations on MAPbI3, which predict the acentric (I4cm) structure to lay slightly lower in energy compared to the centric (I4/mcm) structure by 100 meV/f.u.,100 which is very close to the experimentally determined value of 10.4 kJ/mol (= 108 meV/f.u.), obtained earlier from temperature-dependent 1H NMR experiments for the isotropic reorientation of MA ions at room temperature. These large energy differences show that the noncentrosymmetric structures in the 2D perovskite with larger numbers of layers are favorable, despite the weak SHG signal obtained from the compounds experimentally. These initial

In addition to the primary absorption edge, we consistently observe another peak below this region appearing in the 2D perovskites. For clarity, the raw reflectance data clearly illustrating the presence of the second peak are also provided in the Supporting Information (Figure S1). The intensity of the second peak is strongest for the n = 1 compound and progressively subsides as the thickness of the inorganic slabs increases reaching a minimum for n = ∞ where the peak sits nearly on top of the absorption edge (Figure 5a). The presence of the secondary absorption in the optical absorption peak spectra of these compounds, which is similar in nature with that observed for the 3D perovskites CsPbCl3 and CsPbBr3,93 suggests the presence of stable excitons even at room temperature. Because of the presence of excitons, the band gap of the material was estimated by extrapolating the high energy edge (the fundamental edge) to the hypothetical line parallel to the energy axis where the fundamental edge is interrupted by the low energy absorption (the exciton). MAPbI3 in itself was reported to have a relatively large exciton binding energy (Eb) of ∼50 meV,94,95 which becomes smaller in the high-frequency regime due to charge screening from resonant phonons.77 This value has been recently questioned in favor of a negligible binding energy model assuming an exciton Bohr radius of ∼20 nm.96 On the other hand, Eb ≈ 400 meV40 or even 500 meV58 for exfoliated nanosheets has been reported for the n = 1 2D perovskite homologue, whereas the estimated energies from the multiple-layer 2D systems tend to float between these two values.46,54 The effect of higher exciton binding energies for lower dimensional perovskites is typical and can be understood in terms of reduced dielectric screening and stronger quantum confinement effects.50,51 A similar effect was observed in few-layered transition metal dichalcogenide systems.97,98 These excitons are favored by the multiplequantum-well nature of these 2D perovskites. Unlike the artificially constructed quantum well structures of the classical III−V semiconductors (for example, the GaAs/ AlAs/GaAs heterostructures)99 where the excitons are observable only at low temperature, the excitons in BA2(MA)n−1 PbnI3n+1 are stable at room temperature.58 As it has been suggested,41 the stable excitons in the 2D perovskites are due to the Coulombic charge screening effect, which results from the charged nature of the perovskite {MAn−1PbnI3n+1}2− slabs as opposed to the neutral covalent framework of GaAs-based quantum well structures. The consequence of this difference is that upon generation, the exciton is strongly confined within the built-in electric field between the positive organic spacers and the negative perovskite slabs, which acts upon the photogenerated e−/h+ pairs inhibiting exciton splitting into free carriers. This in turn causes strong photoluminescence at room temperature (Figure 5b). We chose to perform the photoluminescence studies using a confocal microscope setup with a steady-state excitation source (473 nm) on oriented crystals (excitation along (010) direction for n = 2, 3, 4, (001) direction for n = 1 and (100) direction for n = ∞ of the (BA)2(MA)n−1PbnI3n+1 compounds. Previous studies have shown both 2D (BA)2PbI4 and 3D MAPbI3 perovskite compounds display photoluminescence at room temperature,31,64 although the photoluminescence intensity and wavelength can vary widely based on the preparation conditions. The photoluminescence emission energy of all the 2D perovskites described here decreases with increasing slab thickness similar to the band gap trend (Figure 5). K


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cannot be expressed simultaneously in centrosymmetric compounds, as described previously.68 Although insightful, this comparison is difficult to achieve in the case of the lead iodide perovskites because of the very strong fluorescence, which masks some of the Raman-active vibrational modes and lowers the quality of the spectra. Therefore, to probe if inversion symmetry is present in the structure, we performed SHG and THG measurements on polycrystalline samples of the perovskites. All of the 2D compounds and the 3D perovskites displayed weak SHG response as well as THG response. Typically, the SHG signal vanishes for a centrosymmetric compound within the dipole approximation, whereas THG signal is insensitive to the crystal symmetry. The criterion for classifying a material as noncentrosymmetric according to the Kurz-Perry method is satisfied if the SHG ratio between the sample and the known α-quartz reference (SiO2: χ(2) = 0.6 pm/V)), defined as I(SHG)/ Iq(SHG), is larger than 10−2,106,107 as indicated by the red dashed line in Figure 8, panel a. The measured SHG ratios for the 2D and 3D perovskite samples are rather close to the borderline for considering them to be noncentrosymmetric (red dashed line), with the n = 1 and n = 2 perovskites exhibiting SHG signals slightly above the red dashed line and the n = 3, n = 4 and n = ∞ lying slightly below. We have confirmed from repeated, averaged out measurements that there is a general trend of an increasing SHG response with a decreasing thickness of perovskite layers (that is, SHG(n = 1) > SHG(n = 2) > SHG(n = 3) > SHG (n = 4) > SHG(n = ∞)), which is unusual since the SHG trend opposes the expected behavior, which dictates that the SHG response should increase following the reduction of the band gap. Since in some instances a centrosymmetric medium may yield a small, nonzero SHG response due to secondary effects such as quadrupole transitions and surface-induced processes,106 the Kurz-Perry method is not an absolute measure but rather a good guideline for the space group assessment. Because of the uncertainty inherent to the Kurz-Perry powder method, we proceeded in evaluating the crystal symmetry of the (BA)2(MA)n−1PbnI3n+1 perovskites by employing an alternative method based on the comparison of the relative intensities between the SHG and the THG signals.108 This method is more reliable because any effects arising from extrinsic sources would cancel out in the “ratio” of the two harmonic generation intensities. In principle, SHG as the lower order NLO term should produce a much stronger response than THG in a noncentrosymmetric medium. Figure 8, panel b shows the experimental SHG/THG ratios of the samples and the reference material α-SiO2. To compensate for the THG absorption effect, which produces photons with higher energy than the band gap, we employed a proper correction factor to each material based on its relative absorbance values at the SHG and THG wavelengths, 900 and 600 nm, respectively (discussed above in the context of Figure 5). As expected, the response for α-SiO2 SHG dominates over THG for α-SiO2 thus confirming that α-SiO2 is indeed noncentrosymmetric. In contrast, all the (BA)2(MA)n−1PbnI3n+1 perovskites exhibit a THG response, which dominates over SHG suggesting that the perovskite samples are more likely to be centrosymmetric and that the SHG response occurs mainly through secondary processes. Nevertheless, it is still remarkable that the SHG response from the bulk material, albeit weak, is not negligible. This finite SHG response seems to arise from local symmetry-breaking effects within the highly polarizable structure of the perovskites. Our analysis therefore implies that the perovskite samples can exhibit large local polar domains with differing polarities, but in the

results reveal that there is a complex interrelationship between the perovskite layer thickness (n) and a yet to be understood correlation length scale required for cooperative lifting of inversion symmetry. Detailed calculations for the structural interrelationships will be explored in a future study. Our DFT calculations reveal a contrasting trend in the stability of the polar structure in the (BA)2(MA)n−1PbnI3n+1 perovskite relative to the experimental SHG results (see Figure 8a below).

Figure 8. (a) I(SHG)/Iq(SHG) (solid bars) for powdered samples of (BA)2(MA)n−1PbnI3n+1 and MAPbI3. The red dashed line is the criterion for the noncentrosymmetry assignment based on the SiO2 method. (b) SHG/THG (solid bars) for the powdered samples of (BA)2(MA)n−1 PbnI3n+1 and MAPbI3.

The DFT calculations show the n = 1 compound to be less susceptible to symmetry-breaking, whereas the compound experimentally exhibits the relatively strongest SHG response. Along the same lines, when the number of layers increases and the material acquires 3D-like characteristics, the SHG response declines reaching a minimum (although nonzero) value for the n = ∞ perovskite despite the DFT calculation results predicting an opposite trend. This behavior could possibly be rationalized by the ability of the perovskite to accommodate defects that could generate individual domains with opposing polarity and lead to the expression of residual local polarization. This is reasonable because the MA ions are small enough to realign inside the perovskite cages at a small energy cost; however, the BA ions are relatively rigid and in that sense they “lock” the crystal structure in a preferred configuration. In other words, MAPbI3 and the 2D slabs favors an acentric configuration that cannot be kinetically maintained because of the MA reorientations, whereas (BA)2PbI4 favors a centric configuration that once perturbed cannot be kinetically restored. These statements can be further supported by a recent report for the chloride analogue of the n = 1 compound where clear ferroelectric behavior was observed,101 whereas the ferroelectric properties of MAPbI3 are still under debate.102 An important aspect of the ferroelectricity discussion that cannot be neglected is the light-enhanced polarization observed in perovskites,103,104 an observation that suggests that the generation and motion of domains in the halide perovskites is in fact light-induced, and therefore the polarization of the halide perovskites should be better described in the context of photoferroics.105 Nonlinear Optical Properties of the 2D Perovskites. Although the compounds presented here were assigned as polar, based on crystallographic criteria and DFT calculation, some ambiguity in the assignment of the space group remains. Indirect evidence of the acentricity of the compounds is also provided by means of Raman and IR spectroscopy through the mutual exclusion principle (see Supporting Information, Figure S2), where the vibrational modes that are both IR and Raman active L


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Chemistry of Materials overall bulk material “matrix” averages out to centrosymmetric. Similar symmetry-breaking phenomena have also been observed for the noncentrosymmetric artificial GaAs/AlAs quantumwells,109 which could be regarded as the III−IV semiconductor equivalents of the natural quantum-well halide Ruddlesden− Popper systems. In view of the NLO results, the 2D perovskites may be best described as noncentrosymmetric (Cc2m for even n and C2cb for odd n) reflecting the presence of local domains generated by local symmetry breaking, whereas the centrosymmetric models (Ccmm for even n and Acam for odd n) refer to the averaged crystal structure resulting from the random orientation of the local polar domains. A similar description may also apply to the MAPbI3 n = ∞ homologue where the local crystal structure is polar and can be described using the noncentrosymmetric (ferroelectric) model (space group I4cm),110,111 whereas the apparent crystal structure of bulk samples is centrosymmetric (space group I4/mcm) due to the canceling out of the polarization from the randomly oriented local polar domains.

semiconducting properties of the perovskites to the observed systematically evolving optoelectronic properties.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.6b00847. Vibrational spectra, detailed crystallographic tables for (BA)2(MA)n−1PbnI3n+1 (n = 2−4) for the noncentrosymmetric (proper assignment) and centrosymmetric (alternative model) refinements (PDF) Crystallographic files (CIF) Crystallographic files (CIF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Author Contributions

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¶

These authors contributed equally to this work.

CONCLUDING REMARKS We have demonstrated the facile, large-scale synthesis of the 2D (BA)2(MA)n−1PbnI3n+1 (n = 1, 2, 3, 4, ∞) perovskites that belong to the Ruddlesden−Popper family. The convenient isolation of the pure materials allowed for the accurate characterization of the crystal structure of the compounds, which are reported here for the first time for n = 2, 3, and 4. The 2D perovskites crystallize in polar space groups reflecting the alignment of the MA cations to generate an uncompensated dipole moment. All 2D homologous compounds have sharp absorption edges in the visible spectral range suggesting a direct band gap, which asymptotically approaches the band gap of the 3D MAPbI3 perovskite. DFT calculations confirm that the 2D perovskites are indeed direct band gap semiconductors with large bandwidths and small effective masses for both electrons and holes. The absorption of the 2D perovskites is accompanied by a strong photoluminescence emission at room temperature, which is characteristically red-shifted with respect to the absorption edge. The optical physics of the 2D perovskites is characterized by the presence of stable excitons, which act like the active chromophores at room temperature, owing to their natural quantum well electronic structure, which is analogous to that of the artificial III−V semiconductor quantum wells. In view of their exciting properties, the 2D perovskites can indeed be a new source of functional, tunable semiconductors expanding on the properties of the 3D species. For example, the intense room-temperature photoluminescence possible from the 2D perovskites points to their potential utility in light-emitting diodes and lasers. On the other hand, the optimal band gaps of higher n members indicate that these compounds can be used as efficient light absorbers for solution-processed solar cells, which offers better solution processability in addition to their superior environmental stability. More complicated concepts like electrooptical modulators can also be envisaged, taking advantage of the quantum well tunable architecture of the 2D perovskites. The great advantage of the 2D perovskites over the 3D ones is that the complex electronic structure of the 2D materials can be chemically tuned through subtle modification of the crystal structure, adjusting the layer-to-layer spacing or the spacerperovskite interactions. Work in progress is directed in synthesizing and structurally characterizing more homologous series to be able to correlate the effects of the chemical composition in the

Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by Grant No. SC0012541 from the U.S. Department of Energy, Office of Science. D.H.C. acknowledges support from the Link Foundation through the Link Foundation Energy Fellowship Program. J.Y. and J.M.R. were supported by the National Science Foundation (NSF) through the Pennsylvania State University MRSEC under Award No. DMR-1420620. DFT calculations were performed on the Extreme Science and Engineering Discovery Environment (XSEDE), which is supported by NSF Grant No. ACI-1053575, and the QUEST high performance computing facility at Northwestern University, which is jointly supported by the Office of the Provost, the Office for Research, and Northwestern University Information Technology. Electron microscopy was done at the Electron Probe Instrumentation Center (EPIC) at Northwestern University. Raman spectroscopy was performed at the Integrated Molecular Structure Education and Research Center (IMSERC) at Northwestern University. Confocal microscope studies were performed at the SPID facility (NUANCE Center-Northwestern University), which has received support from the State of Illinois through the International Institute for Nanotechnology.

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REFERENCES

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