Two-Dimensional Lead(II) Halide-Based Hybrid Perovskites

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Two-Dimensional Lead(II) Halide-Based Hybrid Perovskites Templated by Acene Alkylamines: Crystal Structures, Optical Properties, and Piezoelectricity Ke-zhao Du,†,⊥ Qing Tu,†,∥,⊥ Xu Zhang,§ Qiwei Han,†,‡ Jie Liu,‡ Stefan Zauscher,†,∥ and David B. Mitzi*,†,‡ †

Department of Mechanical Engineering and Materials Science, ‡Department of Chemistry, and §Department of Biomedical Engineering, Duke University, Durham, North Carolina 27708, United States ∥ NSF Research Triangle Materials Research Science and Engineering Center, Durham, North Carolina 27708, United States S Supporting Information *

ABSTRACT: A series of two-dimensional (2D) hybrid organic−inorganic perovskite (HOIP) crystals, based on acene alkylamine cations (i.e., phenylmethylammonium (PMA), 2-phenylethylammonium (PEA), 1-(2-naphthyl)methanammonium (NMA), and 2-(2-naphthyl)ethanammonium (NEA)) and lead(II) halide (i.e., PbX42−, X = Cl, Br, and I) frameworks, and their corresponding thin films were fabricated and examined for structure−property relationship. Several new or redetermined crystal structures are reported, including those for (NEA)2PbI4, (NEA)2PbBr4, (NMA)2PbBr4, (PMA)2PbBr4, and (PEA)2PbI4. Non-centrosymmetric structures from among these 2D HOIPs were confirmed by piezoresponse force microscopyespecially noteworthy is the structure of (PMA)2PbBr4, which was previously reported as centrosymmetric. Examination of the impact of organic cation and inorganic layer choice on the exciton absorption/ emission properties, among the set of compounds considered, reveals that perovskite layer distortion (i.e., Pb−I−Pb bond angle between adjacent PbI6 octahedra) has a more global effect on the exciton properties than octahedral distortion (i.e., variation of I−Pb−I bond angles and discrepancy among Pb−I bond lengths within each PbI6 octahedron). In addition to the characteristic sharp exciton emission for each perovskite, (PMA)2PbCl4, (PEA)2PbCl4, (NMA)2PbCl4, and (PMA)2PbBr4 exhibit separate, broad “white” emission in the long wavelength range. Piezoelectric compounds identified from these 2D HOIPs may be considered for future piezoresponse-type energy or electronic applications.

H

to hybrid materials with desired properties for transistors,8 electroluminescent devices,9 and for intercalation/chemisorption.10 The wide-ranging selection of the organic cations, with a substantial freedom of motion within the HOIP structure, could also facilitate the design of molecular-based polar materials exhibiting piezoelectricity, ferroelectricity, and second-harmonic generation.11 The piezoelectric effect enables reversible modulation of electric charge with applied mechanical stress, leading to application within sensors, stretchable electronics, actuators, and energy-harvesting devices.12 Currently, most energyharvesting studies that involve HOIPs focus on conversion of solar to electric energy; however, harnessing the piezoelectric effect within HOIPspossibly enabling application of HOIPs for other forms of energy harvestingis still in its infancy, due to the relative scarcity of non-centrosymmetric structures (generally required for piezoelectric activity).13 In current

ybrid organic−inorganic perovskites (HOIPs), especially three-dimensional (3D) AMX3 (A = small organic cation, M = divalent group 14 element, X = halogen) materials, have experienced a surge in popularity within the photovoltaic research community, due to readily achievable high power conversion efficiency (22.1%) and low-cost preparation approaches.1 However, environmental (moisture, oxygen, and UV) instability remains a major drawback for 3D HOIPs.2 Two-dimensional (2D) HOIPs, by contrast, exhibit improved stability and increasingly promising photovoltaic device performance.3 Given the relaxed steric constraints imposed on the organic cations by Goldschmidt’s tolerance factor (i.e., in 3D HOIPs),4 2D HOIPs also provide a much larger compositional space for materials exploration and engineeringthat is, the organic cations do not need to fit into the cage defined by a 3D inorganic framework.5 The 2D structures offer more choices for the metal site (e.g., divalent group 14 element, rare-earth metal, some transition metals, mixed-valent gold)6 as well as for the organic cations (e.g., alkylaryl ammonium, alkyl ammonium, dye cation, mixed cation).7 Such flexibility has led © 2017 American Chemical Society

Received: May 27, 2017 Published: July 27, 2017 9291

DOI: 10.1021/acs.inorgchem.7b01094 Inorg. Chem. 2017, 56, 9291−9302

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Inorganic Chemistry

Figure 1. (a) (PbX4)2− inorganic layer; (b) the molecular structures of the acene alkylamines including PMA, PEA, NMA, and NEA; (c) the PXRD patterns of the 2D HOIP thin films.

leads to the question of where to start in the search for interesting structure−property effects. It is therefore useful to focus on the systematic exploration of related members within a single HOIP structural family, that is, to examine how desired properties change as a function of progressive and independent change to the organic and inorganic components. Among the vast variety of possible organic cation families, acene alkylammoniums have a number of advantages in terms of design flexibility (Figure 1a). On the one hand, one can vary the length of the alkylamine tail tethering the acene group to the inorganic framework. Additionally, the number of fused benzene rings in the acene group can be varied (e.g., phenyl-, naphthyl-, and anthryl-) to impact the size and highest occupied molecular orbital−lowest unoccupied molecular orbital (HOMO−LUMO) separation of the organic cation.22 Finally, substituents or “R” groups (e.g., −H, −F, −Cl, −OH, −CH3) can be placed around the acene or alkyl groups to impact the packing or chiral nature of the organic cation or the steric interaction with the inorganic framework (which in turn can impact the detailed structure and therefore the optoelectronic properties of this layer).7b,21b In the current study, we consider the simplest case of combining lead(II) chloride, bromide, and iodide perovskite layers, with methylamine or ethylamine tethering groups, single (benzyl/phenyl) or double (naphthyl) fused ring acenes, and all possible substituents fixed at R = H (Figure 1). The HOIPs considered for synthesis, crystallization, film growth, and characterization in this study can therefore be denoted as A2PbX4, with A = phenylmethylammonium (PMA; also known as benzylammonium), 2-phenethylammonium (PEA), 1-(2-naphthyl)methanammonium (NMA), and 2-(2naphthyl)ethanammonium (NEA), and with X = Cl, Br, or I. We develop a detailed understanding of the relationship between the crystal structures and optical/piezoelectric properties of these materials by subjecting them to powder X-ray diffraction (PXRD), single-crystal X-ray diffraction (SCXRD), photoluminescence (PL), fluorescence microscopy, optical absorption, and piezoresponse force microscopy (PFM). Among the 10 2D HOIPs we synthesized, five systemsthat

piezoelectric applications, Complementary Metal-Oxide-Semiconductor (CMOS)-compatible Micro-Electro-Mechanical Systems (MEMS) processing is a key challenge, due to a high crystallization temperature for many currently available piezoelectrics.14 For example, the classical piezoelectric material, Lead Zirconate Titanate (PZT), requires temperatures in the range of 600−700 °C for proper film crystallization.15 Furthermore, while piezoelectric polymers, such as polyvinylidene difluoride (PVDF), can be used for CMOS compatible energy harvesting, they require mechanical stretching to convert to the piezoelectric phase and can therefore only be employed on flexible substrates.16 Some piezoelectric thin films, for example, ZnO and AlN, have been fabricated using sputter deposition at relatively low temperature.17 However, noncentrosymmetric HOIPs, for which films can readily be solution-processed at near-ambient temperature on varied substrates,18 provide an even more compelling opportunity for low-cost piezoelectric application. Among HOIPs, CH3NH3PbI3 is currently the dominant material for piezoelectricity research.19 Its piezoelectric performance, however, may be limited, because the apparent polarization likely arises from the rotations of dipolar CH3NH3+ groups (with a slower depolarization than polarization rate) instead of global structural noncentrosymmetry.19a This shortcoming motivates the search for and development of stable and structurally tunable 2D HOIPs with non-centrosymmetric crystal structures for piezoelectric applications. Here, the broad range of available organic cations for 2D HOIPs provides a convenient means to alter and tune the optoelectronic/piezoelectric properties. Specifically tailored organic cations may not only introduce independent optical character to the hybrid (e.g., luminescence7a,20) but may also modify the structure of the inorganic framework and therefore tune the electrical/optical/polar character of this component of the structure.21 For example, the variation of the Sn−I−Sn bond angle in Sn(II)-based 2D HOIPs is sensitive to the specific hydrogen bonding interaction with the organic cation, and it can change the bandgap by as much as 1 eV.21b While 2D HOIPs provide a high degree of tunability, this flexibility also 9292

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Inorganic Chemistry Table 1. Selected Room-Temperature Crystal Unit Cell Parameters of the 2D HOIPs Considered in this Worka unit cell parameters, with the error noted within parentheses (Å for length, Å3 for volume, and deg for angle) (PMA)2PbI426

(PMA)2PbBr4

(PMA)2PbCl422c

(PEA)2PbI4

(PEA)2PbBr427

(PEA)2PbCl428

(NMA)2PbBr4

(NMA)2PbCl422b

(NEA)2PbI4

(NEA)2PbBr4

a = 8.689(5) b = 28.78(1) c = 9.162(4) a = 33.3529(7) b = 8.1466(2) c = 8.1232(2) a = 33.619(7) b = 7.820(1) c = 7.728(1) a = 8.7389(2) b = 8.7403(2) c = 32.9952(6) a = 11.6150(4) b = 11.6275(5) c = 17.5751(6) a = 11.1463(3) b = 11.2181(3) c = 17.6966(5) a = 41.331(1) b = 8.0801(2) c = 8.0836(2) a = 7.748(2) b = 7.790(2) c = 20.859(3) a = 12.266(5) b = 12.276(5) c = 38.65(2) a = 8.2989(4) b = 8.2986(4) c = 19.8768(9)

α = 90 β = 90 γ = 90 α = 90 β = 90 γ = 90 α = 90 β = 90 γ = 90 α = 84.646(1) β = 84.657(1) γ = 89.643(1) α = 99.547(1) β = 105.724(1) γ = 89.977(1) α = 99.173(1) β = 104.634(1) γ = 89.999(1) α = 90 β = 90 γ = 90 α = 90 β = 90 γ = 90 α = 90 β = 98.87(1) γ = 90 α = 92.827(1) β = 99.124(1) γ = 90.142(1)

R_gt/wR_ref

V = 2291(2) Pbca

0.0573/ 0.1030

V = 2207.18(9) Cmc21 Flack = 0.01(2) V = 2031.6(6) Cmc21

0.0303/ 0.0785

V = 2498.29(9) P1̅

0.0503/ 0.1205

V = 2250.6(2) P1̅

0.042/ 0.106

V = 2111.8(1) P1̅

0.031/ 0.044

V = 2699.6(1) Cmc21 Flack = 0.015(9) V = 1259.07(2) Pbam

0.0225/ 0.0418

V = 5751(4) Pn Flack = 0.001(3) V = 1349.8(1) P1 Flack = 0.076(7)

0.021/ 0.054

0.048/ 0.111 0.0615/ 0.1507 0.0346/ 0.0843

a

In cases where the structure comes from the literature, the reference to the structure is given next to the compound name. Structures with bold name are either new to this work or redetermined.

study, with structure refinement details given in Table S1. For (PMA)2PbBr4, we determined the space group as Cmc21 rather than the previously reported Cmca option.11a Furthermore, a distinct room-temperature structure is presented for (PEA)2PbI4, which was previously reported at 203 K.24 Finally, among all the hybrids considered, (NEA)2PbI4, (NEA)2PbBr4, (NMA)2PbBr4, (PMA)2PbBr4, and (PMA)2PbCl4 crystallize in non-centrosymmetric space groups, whereas all the others adopted centrosymmetric alternatives. To explore the origin of the non-centrosymmetric nature within the above 2D HOIPs, the Platon program was used to check the crystalline symmetry of the original structures, as well as the associated structures with organic cations and the inorganic anions independently removed.25 As seen in Table S3, each of the non-centrosymmetric structures maintains the originally assigned symmetry without the inorganic anion, while the removal of organic cation in some cases induces a new suggested space group (i.e., often a centrosymmetric space group). These results therefore suggest that the organic cation conformation/positioning play a key role in establishing the acentric nature in the 2D HOIP structures (as previously noted in several earlier studies on related compounds).11a,c Spincoated thin films of the non-centrosymmetric hybrid structures showed characteristic piezoelectric behavior (see Piezoelectricity section), while the structural agreement of these films

is, (PMA)2PbBr4, (PMA)2PbCl4, (NMA)2PbBr4, (NEA)2PbI4, and (NEA)2PbBr4are identified as piezoelectrics.

■

RESULT AND DISCUSSION The bulk crystals for each HOIP were synthesized using a solution-based growth method. After crystal growth, for each unknown crystal structure, SCXRD data were collected. Subsequently, the crystals were dissolved in dimethylformamide (DMF) and spin-coated on different substrates to obtain crystalline thin films for the various characterization studies. See the Experimental Section for more detail. Crystal Structures. The inorganic layers in each of the target compounds consist of single ⟨100⟩-oriented (PbX4)2− layers,23 constructed from corner-sharing PbX6 octahedra (Figure 1a). All of the organic cations are asymmetric, with a single tethering ammonium group (Figure 1b), and form organic cation bilayers that alternate with the inorganic anion layers. A convenient and general formula, (B)n-(CH2)m-NH3+, is introduced to represent the different acene alkylammonium moieties in this work. Here, (B)n describes the acene chain with “n” fused benzene rings, (CH2)m designates the carbon chain with “m” carbons, linking the aromatic body and ammonium group. The selected unit-cell parameters of all the 2D HOIPs in this work appear in Table 1. Structures with bold names in the table are either new (i.e., as for (NEA)2PbI4, (NEA)2PbBr4, and (NMA)2PbBr4) or redetermined within this 9293

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Figure 2. Definition of organic cation penetration (a); schematic of perovskite layer distortion (b); in-/out-of-plane distortion (c).

Figure 3a, the Pb−X−Pb bond angle varies directly with NH3 penetration (which ranges from 0.35 to 0.7 Å). From Figure 3b,

with the single crystal structures was confirmed by thin-film PXRD data (Figure 1c). The conformation for each HOIP organic cation is described in the Supporting Information (Figure S1). Different cation penetration, defined as the average distance between the nitrogen (N) atoms of the organic cation and the plane of the axial halogen atoms of the perovskite sheet (Figure 2a), induces different distortions in the inorganic framework,5c which in turn can be divided into two parts. One part is the distortion of the PbX6 octahedra themselves. To quantify PbX6 distortion in the current set of compounds, we calculate their bond length quadratic elongation (⟨λ⟩) and their bond angle variance (σ2) using VESTA software:29 λ

1 = 6

⎛ di ⎞2 ∑⎜ ⎟ d i=1 ⎝ o ⎠ 6

(1)

where di are the Pb−X bond lengths, do is the Pb−X bond length for a regular octahedron of the same volume (⟨λ⟩ is dimensionless and independent of the effective size of the octahedron); and σ2 =

1 11

Figure 3. Pb−X−Pb bond angle (a), in-plane and out-of-plane distortion (b), and octahedral distortion (d) of the compounds in this work as a function of the organic cation penetration. Summary of Pb− X−Pb and Sn−I−Sn bond angle dependence on the organic cation penetration for literature-reported structures (c).7f,21b,31 A negative penetration value means that the N atom is above the plane of axial X atoms, whereas a positive value means that the N is below the plane. The black, red, green, and blue symbols in (a) correspond to data for the HOIPs with PMA, PEA, NMA, and NEA cations, respectively.

12

∑ (αi − 90)2 i=1

(2)

where αi are the X−Pb−X bond angles for the octahedra. The other important structural distortion is the more global perovskite layer distortion between adjacent octahedra, which can be determined from the Pb−X−Pb bond angle(s) (Figure 2b). A larger deviation of Pb−X−Pb bond angle away from the ideal 180° indicates a higher perovskite layer distortion. As previously reported, these two distortion types in the inorganic sheet have a dominant effect on the exciton absorption and PL for semiconducting 2D HOIPs based on germanium(II), tin(II), and lead(II) halide frameworks.21b,30 For example, increased perovskite layer distortion decreases the bandwidth leading to a larger bandgap,21b while increased octahedral distortion lowers the energy barrier for some interface defect states leading to a broader PL emission.30 The in-plane Pb−X− Pb bond angle is determined by the projection of the bridging X atoms onto the plane of Pb atoms. The in-plane and out-ofplane distortions contributing to the overall perovskite layer distortion can be defined by the supplementary angle to the inplane Pb−X−Pb bond angle and the shift of the axial Pb−X bonds away from the normal to the plane of Pb atoms, respectively (Figure 2c).21b The parameters ⟨λ⟩ and σ2, together with the Pb−X length, Pb−X−Pb angle, and the in/ out-of-plane distortion, are provided in Table S2, for each of the compounds examined in this study. In the current study, we compare the effects of organic cation penetration on octahedral and layer distortion. As shown in

we see that the in-plane distortion has the strongest correlation with cation penetration. This correlation agrees with analogous tin-iodide-based layered perovskite data extracted from the literature (Figure S2).21b However, in all of these compounds mentioned above, the NH3 unit is more than 0.35 Å below the plane of the axial halogen atoms. To give a more complete picture of how Pb−X−Pb bond angle depends on organic cation penetration, we extract additional structural data from the literature,7f,21b,31 with each structure based on a single-layer hybrid Pb/Sn-based perovskite templated by organic cations with a single tethering ammonium group. As shown in Figure 3c, Pb−X−Pb bond angle decreases for penetration in the range from −0.7 Å (above the plane of axial halogen atoms) to 0.4 Å (below the plane of axial halogen atoms) and then increases in the range from 0.4 to 0.9 Å. The Sn−I−Sn bond angle shows a similar trend as a function of NH3 penetration in the range from 0.30 to 0.7 Å. As mentioned above, bulky organic cations that hydrogen bond with terminal halogens of the metal halide octahedra5a,28 likely affect the Pb(Sn)−X−Pb(Sn) bond angle. The hydrogenbonding interaction is operative within a certain distance range between the NH3 and the axial halogen. Taking the penetration of ∼0.4 Å as a boundary from the experimental data (Figure 3c), the interaction becomes stronger (i.e., the Pb(Sn)−X− 9294

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distortion decreases the width of both the valence and conduction bands, thereby increasing the bandgaps and also the associated exciton energies (Figure 4b).21b The dependence of exciton energies on the equatorial (or bridging) Pb−X bond length, which can perturb the top of the valence band,21b shows a similar trend (Figure S4a). However, no clear relationship can be discerned between the exciton absorption energies and octahedral distortion (Figure S4b). All compounds studied exhibit a strong and sharp exciton emission, with the peak wavelength determined predominantly by the selection of the inorganic layer (i.e., X = Cl, Br, or I; Figure 5a; see also Figure S3), where the iodide emits in the green (∼520 nm), the bromide emits in the violet (∼400 nm), and the chloride emits in the ultraviolet (∼350 nm). Fluorescence microscope images in Figures 5b,c directly show the emission for (PEA)2PbI4 and (PEA)2PbBr4, as examples. For compounds sharing the same inorganic framework, materials with the same (CH2)m have more similar Pb−X− Pb bond angles, which in turn implies closer exciton emission peak positions (Figure 5d and Table S3). Because of its closer position to the terminal ammonium group, the (CH2)m unit plays a more important role than (B)n when the ammonium groups interact with the inorganic frameworks. The structural basis for the specific exciton emission peak positioning therefore likely arises predominantly from the Pb−X−Pb bond angle, which is known to play a major role in establishing the specific band gap for a given lead/tin(II) halide 2D HOIP.21b Similar to the role in exciton absorption, the equatorial (or bridging) Pb−X bond length has a secondary effect on the exciton emission peak position (Figure S4c). Given these results, it is clear that we can fine-tune the emission of the 2D HOIPs based on “m” in the organic cation (B)n(CH2)m-NH3. Both (PMA)2PbI4 and (PEA)2PbBr4 have been used to fabricate green and violet light-emitting diodes, respectively.9d,e With regard to luminescence, the broader family of 2D HOIPs discussed here may therefore prove useful for application as green, violet, and ultraviolet light-emitting diodes, with the latter ultraviolet light emission being of particular interest.34 Notably, (PMA)2PbCl4, (NMA)2PbCl4, and (PMA)2PbBr4 also exhibit separate, broad “white” emission in the long wavelength range (Figure S3), similar to the case previously reported for (PEA)2PbCl4 and attributed to selftrapped excitons involving the organic cations.35 Besides exciton emission peak positioning, it was recently proposed that a broadened PL peak may correlate with the radiative decay from halogen vacancy defects, induced by highly distorted octahedra within the inorganic sheets.30,36 According to this earlier study, which considered (NBT)2PbI4 and (EDBE)2PbI4 (NBT = n-butylammonium and EDBE = 2,2(ethylenedioxy)bis(ethylammonium)), increased distortion of the octahedra lowers the energy barrier for hole self-trapping, thereby broadening the PL line width. In the current study, which considers a broader array of compounds, we also find the narrowest PL peaks ((PEA)2PbI4 and (NMA)2PbCl4) in the phase space of low octahedral distortion, as measured by λ and σ2 (Figure S5a). However, the correlation between the octahedral distortion and PL line width is not obvious (e.g., (PMA)2PbBr4 has small octahedral distortion but relatively wide PL line width). The possible effects of perovskite layer distortion (Figure S5b) and film crystallinity (Figure S6a−c) on the PL line width were also considered. Neither factor appears to play a significant role in determining the PL line width within the current set of samples. Further study is therefore

Pb(Sn) bond angle becomes smaller) as the organic cation penetrates down nearer to the boundary from above the axial halogen planethat is, for penetration in the range from −0.7 to 0.4 Å. If the NH3 penetrates down further, the distance between the NH3 and the axial halogen likely increases, thereby weakening this hydrogen-bonding interaction and producing a larger Pb(Sn)−X−Pb(Sn) bond angle for penetration in the range from 0.4 to 0.9 Å. Given the directional nature of hydrogen bonding and prospects for steric interaction between the “R” group and the apical halides, the detailed orientation of the R−NH3+ group should also play a role, which may explain the scatter of the data in Figure 3c. Importantly, the Pb(Sn)− X−Pb(Sn) bond angle induced in the metal halide framework may affect the structure symmetry and may also impact optoelectronic properties. In contrast to the effect of layer distortion, there is no clear correlation between octahedral distortion and sheet penetration for the set of structures considered (Figure 3d). Optical Properties. The optical absorption of the HOIP thin films was collected by UV−visible spectroscopy (Figure S3) and plotted (Figure 4a) using the direct bandgap Tauc

Figure 4. (a) UV/vis absorption spectra for the 2D HOIP films; (b) dependence of the exciton energy on Pb−X−Pb bond angle; (c) Stokes shift plotted against the fwhm of the exciton absorption peak. The black, red, green, and blue symbols in (b, c) represent different organic cations, corresponding to the data for the compounds with PMA, PEA, NMA, and NEA cations, respectively. The black and red dash lines in (c) are the calculated Stokes shift according to the topographical and quasiequilibrium models, respectively.

formalism (details can be found in the Supporting Information), yielding bandgaps ranging from 2.7 to 4.5 eV. Quantum and dielectric confinement effects give rise to narrow and intense ambient-temperature exciton absorption and emission, which is a specific feature of the 2D HOIPs.32 The exciton absorption mainly arises from electronic transitions within the inorganic layer, for example, from predominantly Pb(6s)−I(5p) mixed states to predominantly Pb(6p) states for the lead(II) iodide inorganic layer, with substitution of Br(4p) and Cl(3p) for I(5p) in the cases of the lead(II) bromide and chloride, respectively.33 For the same inorganic layer type (i.e., the same halogen), the exciton absorption peak most strongly correlates with the variation of the perovskite layer distortion. This 9295

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Figure 5. (a) The PL spectra of 2D HOIPs in this work; (b, c) fluorescence images of (PEA)2PbI4 and (PEA)2PbBr4 thin films; (d) the exciton emission peak positions of the 2D HOIPs showing that, for a given lead(II) halide framework, exciton emission peak position correlates most strongly with the length of the alkyl tethering group (rather than the number of rings in the acene component of the organic cation).

limiting growth (i.e., each compound considered consists of inorganic layers that are exactly one perovskite layer thick, and charge balance enforces this alternation in organic barrier and inorganic well layer). Accordingly, a well-defined linear or quadratic relation between the Stokes shift and the exciton absorption peak line width is not observed for the HOIP quantum wells (Figure 4c). Another consideration for 2D HOIPs, however, is that different organic cations (e.g., as examined in this study) induce different structural distortions (e.g., NH3 penetration effect on the perovskite layer distortion) for the same type of MX42− perovskite layer, which has great impact on the exciton properties.30 The vibrations of organic and inorganic components together with the lattice thermal energy can also affect the exciton absorption peak line width and Stokes shift.39 These extra aspects, arising from the distinct properties of organic−inorganic hybrids, may provide a basis for the non-monotonic dependence of Stokes shift on exciton absorption peak line width (Figure 4c). Notably, we do not find obvious dependence between the exciton absorption peak line width and the film crystallinity/homogeneity (Figure S6d−f), although such dependences cannot be completely ruled out. Piezoelectricity. PFM measurements were performed on all 2D HOIP films listed in Table 1. Given the possibility of artifacts in the PFM data,40 stiff conductive AFM cantilevers (spring constant 2−5 N/m) were used to minimize the

needed to unravel the factors, beyond octahedral distortion, impacting PL line width in 2D HOIPs. The energy difference between exciton absorption and emission defines the Stokes shift for a given film, which in turn typically quantifies the exciton energy loss through vibrational relaxation. 2D HOIPs can be regarded as quantum well structures, most often with the inorganic layer framework serving as wells and the organic cations acting as barrier layers.5c,9b,20b,37 The origin and nature of the Stokes shift in classical semiconductor quantum wells (e.g., InxGa1−xAs/GaAs) is still debated and has been the focus of recent studies, including a proposed topographical model and quasiequilibrium model (see more detailed discussion in the Supporting Information).38 In the topographical model, quantum well thickness fluctuation determines the exciton energy distribution. The exciton relaxing into a local minima in the exciton energy distribution before exciton emission represents the Stokes shift of the quantum well, which has a linear relationship with the exciton absorption peak line width.38f In the quasiequilibrium model, in which thermal quasiequilibrium become an emphasis, the Stokes shift follows a quadratic dependence on the exciton absorption peak line width.38e Well width variation, which plays an important role in the exciton properties in classical quantum wells,38 is less likely to impact single-layer 2D HOIPs, given that these structures exhibit self9296

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Figure 6. AFM topographic images of a selected region on the films of (PMA)2PbBr4 (a), (PEA)2PbI4 (d), and (PMA)2PbCl4 (g). Piezoresponse plotted against the driving frequency for films of (PMA)2PbBr4 (b), (PEA)2PbI4 (e), and (PMA)2PbCl4 (h), fitted using a DSHO function. The AC driving voltage is 3 V. Driving voltage-dependent piezoresponse for (PMA)2PbBr4 (c), (PEA)2PbI4 (f), and (PMA)2PbCl4 (i) films.

driving AC electrical field.43,47 Other mechanisms that can give rise to PFM amplitude have either frequency-dependent background (ion-migration induced Vegard strain) or a nonlinear piezoresponse-AC field relation (electrostriction or induced polarization).41a,43,48 The non-centrosymmetric crystal structures of the 2D HOIPs were confirmed by their piezoelectricity. (PMA)2PbBr4 crystal unit-cell parameters from the current study resemble those from the previous report, in which its space group was determined as centrosymmetric Cmca at room temperature.11a The argument for this assignment revolves around the fact that (PMA)2PbBr4 crystals did not exhibit optical second-harmonic generation (SHG). However, the absence of a detectable SHG signal does not prove the existence of a centrosymmetric space group. 4 9 Through PFM experiments, we find that (PMA)2PbBr4 exhibits piezoelectricity (Figure 6a−c), which arises from the polarization from the non-centrosymmetric structure, leading to the conclusion that (PMA)2PbBr4 actually crystallizes in the non-centrosymmetric space group Cmc21. For the (PMA)2PbBr4 films, the PFM amplitude shows frequencyindependent piezoresponse for frequencies well-separated from the resonance, and the resonance peak can be well fit to the Damped Simple Harmonic Oscillator (DSHO) mode function (Figure 6b). The extracted piezoresponse signal increases linearly with increasing driving voltage (Figure 6c). As control experiments, the PFM measurements on centrosymmetric (PEA)2PbI4 and ferroelectric (PMA)2PbCl4 were also performed. In contrast to the case for (PMA)2PbBr4, the extracted piezoresponse signal of the (PEA)2PbI4 film is highly nonlinear with respect to the driving AC voltage (Figure 6d−f), consistent with a centrosymmetric space group. For the ferroelectric (PMA)2PbCl4,11a the relationship between the

electrostatic contribution to the over signal.41 AFM cantilever resonance was further utilized to improve the electromechanical piezoresponse signal42 and to suppress the electrostatic contribution via dynamic stiffening.41a The PFM amplitude was recorded while sweeping a frequency range that contains the first contact resonance of the AFM cantilever. The spectra were then fitted to a Damped Simple Harmonic Oscillator (DSHO) to extract the piezoresponse signal.42,43 A(ω) =

APR ω0 2 (ω 2 − ω0 2)2 − (ωω0 /Q )2

(3)

where A, ω, ω0, Q, and APR are the amplitude, the driving frequency, the contact resonance frequency of the AFM cantilever, the quality factor of the resonance peak, and the piezoresponse signal, respectively. For piezoelectric materials, due to the converse piezoelectric effect, the field-induced strain sj is given by44

sj = dijEi

(4)

where dij are components of the piezoelectric tensor (in Voigt notation), and Ei is the applied electric field. The AFM measured field-induced displacement will be APR = deff VAC

(5)

where deff is the effective converse piezoelectric coefficient, which depends on the piezoelectric tensor dij, and VAC is the applied alternating current (AC) voltage.44,45 Thus, piezoelectricity induced by polarization from a non-centrosymmetric structure will result in a frequency-independent PFM amplitude when it is away from the cantilever resonance,46 and the extracted piezoresponse signal should linearly depend on the 9297

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(PMA)2PbCl4. PbCl2 (38.1 mg) was dissolved in 0.5 mL of HCl (37%). CH3OH (1 ml) was placed on the top of the PbCl2 solution. Then 0.030 mL of PMA liquid was added into the CH3OH layer. The colorless and laminar crystals (yield ∼62% according to the PbCl2) formed in the solution overnight. (PEA)2PbI4. PbI2 (54.6 mg) was dissolved in 0.5 mL of HI (57%). CH3OH (1 ml) was placed on the top of the PbI2 solution. Then 0.030 mL of PEA liquid was added into the CH3OH layer. The red and laminar crystals (yield ∼76% according to the PbI2) formed in the solution overnight. (PEA)2PbBr4. PbBr2 (43.5 mg) was dissolved in 0.5 mL of HBr (48%). CH3OH (1 ml) was placed on the top of the PbBr2 solution. Then 0.030 mL of PEA liquid was added into the CH3OH layer. The colorless and laminar crystals (yield ∼74% according to the PbBr2) formed in the solution overnight. (PEA)2PbCl4. PbCl2 (32.9 mg) was dissolved in 0.5 mL of HCl (37%). CH3OH (1 ml) was placed on the top of the PbCl2 solution. Then 0.030 mL of PEA liquid was added into the CH3OH layer. The colorless and laminar crystals (yield ∼60% according to the PbCl2) formed in the solution overnight. (NMA)2PbBr4. NMA·HBr was prepared by adding a stoichiometric amount of 48% concentrated HBr into the NMA methanol solution. After the evaporation of methanol, colorless NMA·HBr crystals formed and were separated. Then, NMA·HBr (9.5 mg) and PbBr2 (7.3 mg) were dissolved into DMF (0.2 mL). The colorless and laminar crystals (yield ∼78% according to the PbBr2) were separated after the partial evaporation of DMF. (NMA)2PbCl4. A stoichiometric ratio of NMA·HCl (9.5 mg) and PbCl2 (7 mg) was dissolved into the mixture of DMF (0.3 mL) and HCl (37%, 0.040 mL). The solution was partially evaporated, and the target colorless crystals formed (yield ∼71% according to the PbCl2). (NEA)2PbI4. PbI2 (27.0 mg) was dissolved in 0.5 mL of HI (57%). CH3OH (1 ml) was placed on the top of the PbI2 solution. Then 0.020 mL of NEA liquid was added into the CH3OH layer. The red and laminar crystals (yield ∼70% according to the PbI2) formed in the solution overnight. (NEA)2PbBr4. PbBr2 (21.5 mg) was dissolved in 0.5 mL of HBr (48%). CH3OH (1 ml) was placed on the top of the PbBr2 solution. Then 0.020 mL of NEA liquid was added into the CH3OH layer. The colorless and laminar crystals (yield ∼78% according to the PbBr2) formed in the solution overnight. (NMA)2PbI4 and (NEA)2PbCl4 did not form using the procedures used for the other compounds. Thin-Film Fabrication. All 2D HOIP crystals in this study effectively dissolved into DMF solvent. The concentration of the solution was appoximately 6%∼10% relative to the total weight. For the optical absorption studies, the chloride compounds were spincoated at 3000 rpm for 30 s on quartz substrates, while the bromide and iodide compounds were spin-coated on glass substrates at 3000 rpm for 30 s. For the piezoresponse force microscopy measurements, all of the compounds were spin-coated at 6000 rpm for 30 s on fluorine-doped tin oxide (FTO) on glass substrates. All of the films were annealed at 100 °C in air for 10 min before measurement. Characterization. Single-crystal X-ray diffraction data were collected using a Bruker D8 ADVANCE Series II at room temperature. The crystal structures were solved and refined by Shelxl and Olex software (more details are given in the Supporting Information).50 Powder X-ray diffraction measurements were performed on a PANalytical Empyrean powder X-ray diffractometer using Cu Kα radiation. Optical absorption spectra were obtained using a Shimadzu UV-3600 spectrophotometer. The photoluminescence spectra were measured using a Horiba-Jobi-Yvon LabRAM ARAMIS system, with a 325 nm He−Cd laser excitation. The laser beam was collimated and focused through a 40× UV objective onto the sample surface at room temperature. The photoluminescence peaks and exciton absorption peaks were fitted using a Gaussian function. The blue and green fluorescence images were captured with a Nikon Eclipse TE2000-U Microscope using 405 and 488 nm excitation, respectively. The piezoresponse force microscopy was conducted using an Asylum MFP-

piezoresponse signal and driving AC voltage is again highly linear (as for (PMA)2PbBr4; Figure 6g−i). The PFM results of (PEA)2PbI4 and (PMA)2PbCl4 agree with their expected symmetries. We also performed PFM on films of the other materials (Figure S7), finding that (PMA)2PbI4, (PEA)2PbBr4, (PEA)2PbCl4, and (NMA)2PbCl4 are non-piezoelectrics, while (NMA)2PbBr4, (NEA)2PbI4, and (NEA)2PbBr4 are piezoelectrics. Future work will explore in more detail the piezoelectric or ferroelectric properties of these latter compounds.

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CONCLUSIONS In conclusion, we have synthesized and characterized the structures of a series of 2D HOIPs containing acene alkylamines, including several new or redetermined roomtemperature structuresthat is, (PMA)2PbBr4, (PEA)2PbI4, (NMA)2PbBr4, (NEA)2PbI4, and (NEA)2PbBr4. The noncentrosymmetric space groups of several of these 2D HOIPs have been confirmed by PFM detection of piezoelectricity, including (PMA)2PbBr4, which was previously listed as centrosymmetric. The relationship between the structure and exciton absorption/emission has been studied. The perovskite layer distortion, which plays a dominant role in tailoring the optical properties of exciton absorption/emission in these systems, has a direct relationship with organic cation penetration depth. Besides the sharp exciton emission, (PMA) 2 PbCl 4 , (NMA) 2 PbCl 4 , (PMA) 2 PbBr 4 , and (PEA)2PbCl4 exhibit separate, broad “white” emission in the long-wavelength range. The relationship between the Stokes shift and the full width at half-maximum (fwhm) of the sharp exciton absorption peak in these 2D HOIP quantum wells exhibits a different functional dependence compared with classical inorganic quantum wells, for which well width fluctuation plays an important role. Furthermore, we have screened several new piezoelectric materials from among these 2D HOIPs through the use of PFM measurements. The frequency-independent background of the PFM amplitude and the linear piezoresponse-AC field relation confirm that the piezoresponse arises from their non-centrosymmetric structures. These results provide opportunities for HOIPs as future energy-harvesting/scavenging media (conversion of mechanical to electrical energy).

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EXPERIMENTAL SECTION

Chemical. PbI2 (99.999% trace metal basis), PbBr2 (99.999% trace metal basis), phenylmethylamine (>99.0%), 2-phenylethylamine (99%), HI (57 wt % in H2O, with hypophosphorous acid as stabilizer, assay 99.95%), HBr (48 wt % in H2O, assay >99.99%), HCl (ACS reagent, 37%), and CH3OH (>99.9%) were purchased from SigmaAldrich company. PbCl2 (99.999% metal basis) was purchased from Alfa Aesar company. 1-(2-Naphthyl)methanamine hydrochloride (95%) and 2-(2-naphthyl)ethanamine (95%) were purchased from Enamine company. 1-(2-Naphthyl)methanamine (95+%) was purchased from Ark Pharm company. Crystal Synthesis. (PMA)2PbI4. A stoichiometric mixture of PMA (0.060 mL) and PbI2 (126.4 mg) was heated to 100 °C in 1 mL of the HI (57%) solvent. After several hours, a white solid (which appeared after adding the PMA) in the solution completed transformation into red and laminar crystals (yield ∼77% according to the PbI2). (PMA)2PbBr4. PbBr2 (50.3 mg) was dissolved in 0.5 mL of HBr (48%). CH3OH (1 ml) was placed on the top of the PbBr2 solution. Then 0.030 mL of PMA liquid was added into the CH3OH layer. Colorless and laminar crystals (yield ∼73% according to the PbBr2) of the product formed in the solution overnight. 9298

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Inorganic Chemistry 3D AFM (Oxford Instrument, CA). AFM cantilevers employed tips coated by Ir or Pt (Asyelec-01 or HQ: NSC14/Pt).

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T. Organometal Halide Perovskites as Visible-Light Sensitizers for Photovoltaic Cells. J. Am. Chem. Soc. 2009, 131, 6050−6051. (c) Green, M. A.; Emery, K.; Hishikawa, Y.; Warta, W.; Dunlop, E. D. Solar cell efficiency tables (version 48). Prog. Photovoltaics 2016, 24, 905−913. (2) (a) Leijtens, T.; Eperon, G. E.; Noel, N. K.; Habisreutinger, S. N.; Petrozza, A.; Snaith, H. J. Stability of Metal Halide Perovskite Solar Cells. Adv. Mater. 2015, 5, 1500963. (b) Xiao, Z.; Du, K. Z.; Meng, W.; Wang, J.; Mitzi, D. B.; Yan, Y. Intrinsic Instability of Cs2In(I)M(III)X6 (M = Bi, Sb; X = Halogen) Double Perovskites: A Combined Density Functional Theory and Experimental Study. J. Am. Chem. Soc. 2017, 139, 6054−6057. (3) (a) Xiao, Z.; Meng, W.; Wang, J.; Mitzi, D. B.; Yan, Y. Searching for promising new perovskite-based photovoltaic absorbers: the importance of electronic dimensionality. Mater. Horiz. 2017, 4, 206− 216. (b) Stoumpos, C. C.; Cao, D. H.; Clark, D. J.; Young, J.; Rondinelli, J. M.; Jang, J. I.; Hupp, J. T.; Kanatzidis, M. G. RuddlesdenPopper Hybrid Lead Iodide Perovskite 2D Homologous Semiconductors. Chem. Mater. 2016, 28, 2852−2867. (c) Tsai, H.; Nie, W.; Blancon, J. C.; Stoumpos, C. C.; Asadpour, R.; Harutyunyan, B.; Neukirch, A. J.; Verduzco, R.; Crochet, J. J.; Tretiak, S.; Pedesseau, L.; Even, J.; Alam, M. A.; Gupta, G.; Lou, J.; Ajayan, P. M.; Bedzyk, M. J.; Kanatzidis, M. G.; et al. High-efficiency two-dimensional RuddlesdenPopper perovskite solar cells. Nature 2016, 536, 312−316. (4) Goldschmidt, V. M. Die gesetze der krystallochemie. Naturwissenschaften 1926, 14, 477−485. (5) (a) Mitzi, D. B. Synthesis, Structure, and Properties of OrganicInorganic Perovskites and Related Materials. In Progress in Inorganic Chemistry; John Wiley & Sons, Inc., 2007; pp 1−121. (b) Pedesseau, L.; Sapori, D.; Traore, B.; Robles, R.; Fang, H. H.; Loi, M. A.; Tsai, H.; Nie, W.; Blancon, J. C.; Neukirch, A.; Tretiak, S.; Mohite, A. D.; Katan, C.; Even, J.; Kepenekian, M. Advances and Promises of Layered Halide Hybrid Perovskite Semiconductors. ACS Nano 2016, 10, 9776−9786. (c) Saparov, B.; Mitzi, D. B. Organic-Inorganic Perovskites: Structural Versatility for Functional Materials Design. Chem. Rev. 2016, 116, 4558−4596. (6) (a) Mitzi, D. B. Synthesis, crystal structure, and optical and thermal properties of (C4H9NH3)2MI4 (M= Ge, Sn, Pb). Chem. Mater. 1996, 8, 791−800. (b) Mitzi, D. B.; Liang, K. Preparation and Properties of (C4H9NH3)2EuI4: A Luminescent Organic−Inorganic Perovskite with a Divalent Rare-Earth Metal Halide Framework. Chem. Mater. 1997, 9, 2990−2995. (c) Svensson, P. H.; Kloo, L. Metal Iodides in Polyiodide Networks: Synthesis and Structure of Binary Metal Iodide−Iodine Compounds Stable under Ambient Conditions. Inorg. Chem. 1999, 38, 3390−3393. (d) De Jongh, L. J.; Bloembergen, P.; Colpa, J. H. P. Transition temperatures of weakly anisotropic layertype magnets. Physica 1972, 58, 305−314. (e) Castro-Castro, L. M.; Guloy, A. M. Organic-based layered perovskites of mixed-valent gold(I)/gold(III) iodides. Angew. Chem., Int. Ed. 2003, 42, 2771− 2774. (7) (a) Mitzi, D. B.; Chondroudis, K.; Kagan, C. R. Design, Structure, and Optical Properties of Organic−Inorganic Perovskites Containing an Oligothiophene Chromophore. Inorg. Chem. 1999, 38, 6246−6256. (b) Xu, Z. T.; Mitzi, D. B. SnI42--based hybrid perovskites templated by multiple organic cations: Combining organic functionalities through noncovalent interactions. Chem. Mater. 2003, 15, 3632−3637. (c) Mitzi, D. B. Organic-inorganic perovskites containing trivalent metal halide layers: The templating influence of the organic cation layer. Inorg. Chem. 2000, 39, 6107−6113. (d) Naik, V. V.; Vasudevan, S. Melting of an Anchored Bilayer: Molecular Dynamics Simulations of the Structural Transition in (CnH2n+1NH3)2PbI4 (n = 12, 14, 16, 18). J. Phys. Chem. C 2010, 114, 4536−4543. (e) Papavassiliou, G. C.; Koutselas, I. B.; Terzis, A.; Whangbo, M. H. Structural and electronic properties of the natural quantum-well system (C6H5CH2CH2NH3)2SnI4. Solid State Commun. 1994, 91, 695−698. (f) Billing, D. G.; Lemmerer, A. Synthesis, characterization and phase transitions of the inorganic-organic layered perovskite-type hybrids [(CnH2n+1NH3)2PbI4] (n = 12, 14, 16 and 18). New J. Chem. 2008, 32, 1736−1746.

ASSOCIATED CONTENT

S Supporting Information *

The cif files for each of the new/redetermined structures, The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.inorgchem.7b01094. Details of general molecule structure, crystal data and structure refinement parameters, the conformation and orientation of organic cations, raw data of optical absorption, more structure−property relationship figures and piezoresponse as a function of the driving frequency for the compounds not listed in the main text. The CCDC Nos. for (PMA) 2 PbBr 4 , (PEA) 2 PbI 4 , (NMA)2PbBr4, (NEA)2PbI4, and (NEA)2PbBr4 are 1542460, 1542461, 1542462, 1542463, and 1542464, respectively. (PDF) Accession Codes

CCDC 1542460−1542464 contain the supplementary crystallographic data for this paper. These data can be obtained free of charge via www.ccdc.cam.ac.uk/data_request/cif, or by emailing [email protected], or by contacting The Cambridge Crystallographic Data Centre, 12 Union Road, Cambridge CB2 1EZ, UK; fax: +44 1223 336033.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Jie Liu: 0000-0003-0451-6111 Stefan Zauscher: 0000-0002-2290-7178 David B. Mitzi: 0000-0001-5189-4612 Author Contributions ⊥

These authors contributed equally. K.D. conducted the synthesis and the structural and optical characterizations; Q.T. and K.D. conducted the piezoresponse force microscopy. The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript.

Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This material is based upon work financially supported by the NSF through the Research Triangle MRSEC (DMR-1121107). The work was performed in part at the Duke University Shared Materials Instrumentation Facility (SMIF), a member of the North Carolina Research Triangle Nanotechnology Network (RTNN), which is supported by the National Science Foundation (Grant No. ECCS-1542015) as part of the National Nanotechnology Coordinated Infrastructure (NNCI). All opinions expressed in this paper are the authors’ and do not necessarily reflect the policies and views of the NSF.

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REFERENCES

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