Two-Dimensional Hybrid Halide Perovskites: Principles and Promises

Perspective pubs.acs.org/JACS

Cite This: J. Am. Chem. Soc. 2019, 141, 1171−1190

Two-Dimensional Hybrid Halide Perovskites: Principles and Promises Lingling Mao, Constantinos C. Stoumpos,*,† and Mercouri G. Kanatzidis*

J. Am. Chem. Soc. 2019.141:1171-1190. Downloaded from pubs.acs.org by EASTERN KENTUCKY UNIV on 01/25/19. For personal use only.

Department of Chemistry, Northwestern University, 2145 Sheridan Road, Evanston, Illinois 60208, United States with the works of Wells (1893)2 and Møller (1957−1959)3,4 on CsPbX3 and Auger and Karantassis on CsSnX3 (1925) and CsGeX3 (1935).5,6 Several breakthroughs occurred in 2009 and 2012. In 2009, Miyasaka and co-workers reported liquidbased dye-sensitized solar cells based on nanocrystals of MAPbI3 attached to a TiO2 surface.7 However, the watershed moment occurred in 2012, when (a) the all-solid-state dyesensitized cells based on a solid layer of CsSnI3 perovskite acting as a hole transport layer and light absorber with ∼10% efficiency were reported in 2012 by Kanatzidis and coworkers;8 (b) demonstration of cells based on a solid layer of MAPbI3 devices was reported in 2012 by Miyasaka, Snaith, and co-workers;9 and (c) independent confirmation of the solid layer perovskite of MAPbI3 was reported by Park and coworkers.10 After that, the halide perovskite field really took off, and perovskite solar cells (PSCs) have performed an epic march to >23% efficiencies that compare favorably with those of many established inorganic semiconductors.9−24 Given this history, it is fair to say that the current perovskite solar cell field is more of an intellectual child of the solution-processed dyesensitized cell (Grätzel cell)25 rather than a logical evolution of the more traditional Si, CdTe, and CuInSe2 cells.26 The field has also borrowed heavily from techniques utilized in the organic photovoltaics area. More importantly, the enormous success of PSCs firmly established the emergence of halide perovskites as the next generation of high-performance semiconductors.27−31 Indeed, halide perovskites match, in terms of optoelectronic performance, benchmark inorganic semiconductors in a wide array of applications (LEDs,32−34 FETs,35,36 laser,37−39 hard radiation detectors,40−43 etc.), forecasting a great potential for commercialization. Mechanistically, however, the halide perovskites have distinctive properties that derive from their soft lattices and dynamically disordered crystal structure,44,45 which favorably modifies their charge-carrier recombination lifetimes and renders them nonclassical semiconductor characteristics.46−49 A major obstacle in further development exists in the form of limited composition accessible in the ABX3 (A = Cs+, CH3NH3+, [HC(NH2)2]+; B = Pb2+, Ge2+, Sn2+; X = Cl−, Br−, I−) perovskite structure, with only three A-site cations being able to maintain the 3D corner-sharing inorganic framework,50 in some cases forming only metastable compounds.51 Whereas a solution for stabilizing some metastable phases was found in the form of perovskite nanocrystals,52−54 important relief from the narrow compositional diversity was found in the form of two-dimensional (2D) halide perovskites (Figure 1).

ABSTRACT: Hybrid halide perovskites have become the “next big thing” in emerging semiconductor materials, as the past decade witnessed their successful application in high-performance photovoltaics. This resurgence has encompassed enormous and widespread development of the three-dimensional (3D) perovskites, spearheaded by CH3NH3PbI3. The next generation of halide perovskites, however, is characterized by reduced dimensionality perovskites, emphasizing the two-dimensional (2D) perovskite derivatives which expand the field into a more diverse subgroup of semiconducting hybrids that possesses even higher tunability and excellent photophysical properties. In this Perspective, we begin with a historical flashback to early reports before the “perovskite fever”, and we follow this original work to its fruition in the present day, where 2D halide perovskites are in the spotlight of current research, offering characteristics desirable in high-performance optoelectronics. We approach the evolution of 2D halide perovskites from a structural perspective, providing a way to classify the diverse structure types of the materials, which largely dictate the unusual physical properties observed. We sort the 2D hybrid halide perovskites on the basis of two key components: the inorganic layers and their modification, and the organic cation diversity. As these two heterogeneous components blend, either by synthetic manipulation (shuffling the organic cations or inorganic elements) or by application of external stimuli (temperature and pressure), the modular perovskite structure evolves to construct crystallographically defined quantum wells (QWs). The complex electronic structure that arises is sensitive to the structural features that could be in turn used as a knob to control the dielectric and optical properties the QWs. We conclude this Perspective with the most notable achievements in optoelectronic devices that have been demonstrated to date, with an eye toward future material discovery and potential technological developments.

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INTRODUCTION Hybrid inorganic−organic halide perovskites have made their mark in their scientific world in the past decade, driven by the meteoric rise in the photovoltaic efficiency of the methylammonium lead iodide, CH3NH3PbI3 (or MAPbI3). The discovery of CH3NH3PbI3, a three-dimensional (3D) compound that shares similar structural features with the namesake mineral CaTiO3, can be traced back to Weber’s work in 1978,1 yet the history of the halide perovskites dates long before, that © 2018 American Chemical Society

Received: October 8, 2018 Published: November 6, 2018 1171

DOI: 10.1021/jacs.8b10851 J. Am. Chem. Soc. 2019, 141, 1171−1190

Perspective

Journal of the American Chemical Society

Figure 1. (a) Schematic illustration of the evolution from 2D perovskite to 3D perovskite with key components. (b) 2D quantum well structure illustration.

properties, ascribing physical quantities to structural features and extracting structure−property correlations. We also cover the latest developments in the evolution of these properties as a function of external stimuli, such as temperature and pressure. We conclude with a brief discussion on the applications of 2D halide perovskites, and we provide an look toward promising research directions and offer suggestions for future developments.

Through the chemically accomplished dimensional reduction of the 3D crystal lattice, 2D perovskites, (A′)m(A)n−1BnX3n+1, adopt a new structural and compositional dimension, A′, where monovalent (m = 2) or divalent (m = 1) cations can intercalate between the anions of the 2D perovskite sheets. Besides its apparent simplicity, the concept of dimensional reduction not only grants access to a plethora of hybrid organic−inorganic materials with tailorable functionalities but also gives rise to an unparalleled structural complexity55 which allows for the fine-tuning of the optoelectronic properties. In the 2D hybrid halide perovskites, the organic cations act as insulating barriers that confine the charge carriers in two dimensions, but they also serve as dielectric moderators that determine the electrostatic forces exerted on the photogenerated electron−hole pairs. The specific arrangement of alternating organic−inorganic layers generates a crystallographically ordered 2D multiple-quantum-well (MQW) electronic structure that forms naturally in a bottom-up fashion through self-assembly. The structural coherence of the halide perovskite MQWs resembles that of the artificially synthesized III−V films’ semiconductor heterostructures. In addition, due to the high organic/inorganic dielectric contrast, the perovskite materials generate huge electron−hole binding energies. 2D halide perovskite MQWs can stabilize excitons that can be readily observed even at ambient temperatures. The multitude of possibilities deriving from the MQW structures, already demonstrated in III−V semiconductors,56−59 makes the 2D halide MQWs particularly interesting for the discovery of new fundamental physics and efficient room-temperature optoelectronic applications. In this Perspective, we attempt to categorize the types of 2D halide perovskites, discussing the chemical characteristics that govern the formation of the perovskite layers and how these can alter the semiconducting properties of these materials. As we focus on chemical insights and understanding, we first introduce the general types of 2D perovskites, based on their different octahedral connectivity and layer orientation. We continue the discussion on the structural aspects of the perovskite types based on the (100)-oriented, (110)-oriented, and (111)-oriented 2D perovskites. We then branch out in a more detailed description of the structures of the multilayer (100)-oriented perovskites, specifically the Ruddlesden− Popper (RP) phases,60,61 Dion−Jacobson (DJ) phases,62 and alternating cation in the interlayer space (ACI) phases.63 Based on the crystal structure, we analyze their photophysical

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A BRIEF HISTORY OF 2D HALIDE PEROVSKITES 2D perovskites, especially the transition-metal-based structures (RNH3)2BX4, with B = Cd2+, Mn2+, Fe2+, Cu2+, Hg2+, etc. and X = Cl−, Br−, have been investigated for dielectric and magnetic properties.64−67 Exploration of a lead-based 2D system was reported by Maruyama and co-workers for (CH3(CH2)8NH3)2PbI4 in 1986.68 However, it was mainly Ishihara et al.,66,69−73 Papavassiliou et al.,74−78 and Thorn et al.79 who systematically studied the optical properties of the materials. The latter group introduced the concept of “inbetween phases” (i.e., between the (RNH3)PbI4 and MAPbI3), which was the first conceptual relationship between the standard 2D oxide perovskites and the hybrid 2D halide perovskites. Thorn demonstrated this hypothesis by the successful isolation of whole families of halide perovskites, which instead of discrete compositions could be collectively described as a homologous series (i.e., (CH3(CH2)8NH3)2(CH3NH3)n−1PbnX3n+1, with X = Br−, I−, or (PEA)2(MA)n−1PbnI3n+1, with n = 1−2 and PEA = phenylethylammonium).79 Mitzi and co-workers combined the collective knowledge of that time into a well-defined research field and laid the foundation for the explosive growth of 2D halide perovskite semiconductors. Their report on the evolution of the (C4H9NH3)2(CH3NH3)n−1SnnI3n+1 (n = 1−5) homologous series for many-n members demonstrated how the electrical properties of the compounds may change as a function of n, from the nearly intrinsic n = 1 to the heavily doped, “metallike” n = ∞.80 In addition to the so-called (100)-oriented (reflecting the cleavage plane relative to the ideal cubic 3D perovskite) (C4H9NH3)2(CH3NH3)n−1SnnI3n+1 multilayer series, the (110)-oriented [IC(NH2)2](CH3NH3)mSnmI3m+2 (m = 1−4) series was also described.81 Subsequently, Mitzi and coworkers reported the synthesis and characterization of singlelayer (n = 1) 2D perovskites with different organic cations82−99 and extended the choice of the metal ion to include Ge2+,100 group 15,101 and rare earth elements.102 It was this group’s 1172


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Figure 4’s pool, below. However, on rare occasions the configurational strain that the cations impose to the 2D perovskite breaks the standard connectivity pattern and instead stabilizes edge-sharing or face-sharing networks.133 These networks almost never form with exclusive edge-sharing134 or face-sharing connectivity, since area-filling in two dimensions is difficult to achieve with edge-sharing and altogether improbable with face-sharing, and instead form in combination with partial corner-sharing contacts (Figure 2). For example, (mpz)2Pb3Br10122 (mpz = N-methylpiperazine) is comprised of repeating edge-sharing dimeric units (highlight in red), which are further connected in a corner-sharing fashion to complete a 2D layer. [(H2en)7(C2O4)2]n(Pb4I18)n·4nH2O (en = ethylenediamine) is a rare example of an exclusively cornersharing structure that does not fit (100)-, (110)-, or (111)orientation.135 Similarly, in (tms)4Pb3Br10,136 (tms = trimethylsulfonium), the triads of face-sharing units (highlighted in red) share corners at both terminal apexes to complete a 2D layer resembling the 9R perovskite polytype.137 The way individual octahedra are linked determines the orbital overlap between the metal and the halide, which subsequently influences the band gap of the compound. As a general trend, determined experimentally and also theoretically confirmed, the band gap follows the increasing trend “corner-sharing < edge-sharing < face-sharing”.122,133,138 This is perhaps expected within the special perovskite “inverted” band structure, since corner sharing provides the strongest bonding between the metal and the halogen, which in turn gives rise to a “destabilization” of the ns2 lone pair, pushing it to the top of the valence band.139 As the σ-bonding weakens through the interactions in the edge-sharing and face-sharing motifs, the lone pair is stabilized, localizing deep within the valence band, while the bottom of the conduction band increases in energy, leading to an opening of the bandgap.138 The structural diversity obtained in these 2D perovskites could potentially be exploited to further adjust the optical properties of the hybrid perovskites. The main class of 2D halide perovskite compounds, therefore, consists of corner-sharing octahedra, and these can be further classified into three categories: the (100)-oriented, the (110)-oriented, and the (111)-oriented structures based on the crystallographic layer slicing direction cut along a specific (hkl) plane of the parent 3D structure. The general formula for the (100)-oriented and (110)-oriented 2D perovskites is A′2An−1BnX3n+1 or A′An−1BnX3n+1 (A′ = 1+ or 2+; A = 1+ cation; B = Pb2+, Sn2+, Ge2+, Cu2+, Cd2+, etc.; X = Cl−, Br−, and I−), while for the (111)-oriented 2D perovskite, the general formula is A′n+1BnX3n+3 (n > 1, where B valence is 3+, mixedvalence averaging 3+, or heterometallic, e.g., Cu2+ and Sb3+ in Figure 3).140 Out of the three categories, the (100)-oriented 2D perovskites heavily dominate the classification tree, and they are by far the most commonly obtained 2D perovskite structures. Besides the explosive development in the more standard 2D perovskites which will be discussed in the next section, remarkable progress has been achieved in the unconventional (100)-, (110)-, and (111)-oriented perovskites. In this direction, many new compounds have been discovered, thus broadening the scope of future structural engineering in the halide perovskite family. An interesting extension of the structural variation of the (100)-oriented perovskites comes from blending the dimensional reduction concept141 with the ordered double perovskite structure.142,143 The thus obtained single-layer and bilayer

contribution related to the development of thin-film transistor devices based on 2D halide perovskites that demonstrated their potential as high-performance semiconductors.35 In the current era, 2D halide perovskites provide synthetic opportunities for engineering materials customized to the needs of specific applications in all of the traditional research fields, such as solar cells,62,103−107 light-emitting diodes (LEDs),108−110 and photodetectors.111−113 More exotic forms have also emerged with the rise of 2D halide perovskites such as 2D nanocrystals,114,115 with utility in single-component white-light emission116−123 for solid-state lighting devices,124 strong optical nonlinearity,125 and ferroelectricity.126 2D hybrid perovskite materials not only perform well by themselves but also can be integrated into heterostructures,127,128 enhancing the devices’ functionalities and performance. Because of their fundamental importance, it is therefore crucial to understand at the molecular level the structure−property relationships that govern these materials in order to allow for targeted device engineering and further optimization to take place.129−132

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STRUCTURAL TYPES AND CONNECTIVITY MODES Akin to their 3D congeners, 2D halide perovskites obey stereochemical rules that determine the growth of the perovskite sheets. Three octahedral connectivity modes have been experimentally observed: corner-sharing, edge-sharing, and face-sharing, as can be seen in Figure 2. Different from the

Figure 2. Connectivity modes in 2D perovskite and their representative examples: (a) corner-sharing, (3AMP)PbI4 from ref 62; (b) edge-sharing, (mpz)2Pb3Br10 from ref 122; and (c) facesharing, (tms)4Pb3Br10 from ref 136.

3D structures, where the connectivity depends primarily on the relative size of the A-site cations, 2D perovskites have far fewer size restrictions for the A′-site (A-site still follows the same rules as 3D perovskites), since cations must fit within the interlayer space. Instead, the characteristics that determine the suitability of a cation as a spacer are (i) its net positive charge at the perovskite anchoring position and degree of substitution, i.e., RNH3+ > R2NH2+ > R3NH+ > R4N+ in order of decreasing preference; (ii) its hydrogen-bonding capacity; (iii) its stereochemical configuration, i.e., flexible aliphatic hydrocarbons > rigid aromatic hydrocarbons, including substitution on the protonated site; and (iv) its space-filling ability, i.e., linear interdigitating cations > branched irregular cations. Since there are many cations that fulfill these prerequisites, there are, in fact, many cations that can stabilize the regular (i.e., cornersharing) 2D perovskite structure, which we have included in 1173


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structures both Ag (1+) and Bi (3+) form alternating (MBr6/2)− octahedra, fulfilling both the charge and ionic size requirements of the 2D perovskite layer, while the butylammonium cation bilayer performs its usual role of separating the inorganic layers (Figure 3).60,80 Another interesting 2D perovskite variation is found in the heteroleptic (100)-oriented layered perovskites. The hybrid compound (MA)2PbI2(SCN)2192 as well as the all-inorganic compound Cs2SnI2Cl2193 and Cs2PbI2Cl2194 represent distinctive derivatives of the K2NiF4195 prototype where the lattice positions along the perovskite plane and vertically to it are occupied by different chemical species. In these chemically ordered structures the larger heteroatoms (SCN− in the former, I− in the latter) decorate the surfaces of the layers, whereas the smaller heteroatom is responsible for the corner-sharing expansion of the perovskite along the layer. These three compounds differ significantly from the rest of the hybrid inorganic−organic perovskites as the separation between the layers occurs “naturally”, without the use of bulky organic spacers, and seems to be realized because of the ionic-size mismatch of the two heteroatoms. The corrugated (110)-oriented structures are far less common since there are very few cations that can stabilize them, those being typically small, highly symmetric ones81,196

Figure 3. Examples of the double perovskite (100)-oriented structure (Ag+ and Bi3+ are highlighted in blue and yellow, respectively), the (110)-oriented corrugated structure, and the (111)-oriented perovskite structure (Cu2+ and Sb3+ are highlighted in light green and dark green, respectively).

bromide compounds (BA)4(AgBi)Br8 and (BA)2Cs(AgBi)Br7 represent a very special case of the (100) type arising from this combination with the butylamine (BA+) cation.144 In these

Figure 4. Reported organic cations that form corner-sharing (100)-oriented, (110)-oriented, or (111)-oriented 2D halide perovskite structures (n = layer thickness number): 1, primary alkylammonium (m = 0−1, n = 1;145 m = 1, n = 3;85,121 m = 2, n = 1;146−148 m = 3, n = 1−5;60,61,80,100 m = 4− 5, n = 1;98 m = 6−9, n = 1;149 m = 9−10, n = 1;150 m = 11, 13, 15, 17, n = 1149,151); 2, primary alkyldiammonium (m = 3, n = 1;152 m = 4, n = 1;118,153−156 m = 4−9, n = 1−4;157 m = 5−6, n = 1;97,156,158 m = 8, 10, 12, n = 189,159); 3, m = 2, N1-methylethane-1,2-diammonium (N-MEDA); m = 3, N1-methylpropane-1,3-diammonium (N-MPDA); all n = 1116); 4, m = 2, 2-(dimethylamino)ethylammonium (DMEN); m = 3, 3(dimethylamino)-1-propylammonium (DMAPA); m = 4, 4-(dimethylamino)butylammonium (DMABA); all n = 1120); 5, ammonium 4-butyric acid (n = 283); 6, iodoformamidinium (n = 1−381); 7, guanidinium (GA, n = 1−363,160,161); 8, protonated thiourea cation (n = 1162); 9, 2,2′dithiodiethanammonium (n = 1163,164); 10, 2,2′-(ethylenedioxy)bis(ethylammonium) (EDBE, n = 1117); 11, protonated 2-(aminoethyl)isothiourea (n = 1165); 12, but-3-yn-1-ammonium (BYA, n = 1166); 13, 2-fluoroethylammonium (n = 1167); 14, 2-methylpentane-1,5-diammonium (n = 1152); 15, isobutylammonium (IBA, n = 1168); 16, heteroatom-substituted alkylammonium (n = 1163); 17−20, cyclopropylammonium, cyclobutylammonium, cyclopentylammonium, and cyclohexylammonium (n = 192,169); 21, cyclohexylmethylammonium (n = 1170); 22, 2-(1cyclohexenyl)ethylammonium (n = 1171,172); 23, N-(aminoethyl)piperidinium (n = 197); 24, N-benzylpiperazinium (n = 1173); 25, piperazinium (n = 1173); 26, (carboxy)cyclohexylmethylammonium (n = 197); 27, 3-(aminomethyl)piperidinium (3AMP, n = 1−462); 28, 4-(aminomethyl)piperidinium (4AMP, n = 1−462); 29, cyclohexylphosphonium (n = 1174); 30, 1,4-bis(aminomethyl)cyclohexane (n = 1175); 31, benzylammonium (BZA, n = 195,133,176−179); 32, phenylethylammonium (PEA, n = 1−379,97,104,180,181); 33, m-phenylenediammonium (n = 1182); 34, 4methylbenzylammonium (n = 1−282); 35, 4-fluorophenylethylammonium (n = 1183); 36, 3-iodopyridinium (n = 197; 37, histammonium (n = 195); 38, 2-thienylmethylammonium (n = 284); 39, 2-(2-thienyl)ethanaminium (n = 1184); 40, 2-(ammoniomethyl)pyridinium (n = 1185); 41, 2substituted phenethylammonium (X = F, Cl, Br; n = 1186); 42, N,N-dimethyl-p-phenylenediammonium (n = 1182,187); 43, 4-amidinopyridinium (n = 194); 44, 2-naphthylmethylammonium (n = 199); 45, benzimidazolium (n = 1188); 46, 1,5-diammoniumnaphthalene (n = 189); 47−49, naphthalene-O-ethylammonium, pyrene-O-ethylammonium, perylene-O-ethylammonium (all n = 1189); 50, 5,5′-bis(ammoniumethylsulfanyl)-2,2′bithiophene (BAESBT, n = 1190); 51, 5,5‴-bis(aminoethyl)-2,2′:5′,2″:5″,2‴-quaterthiophene (AEQT, n = 1191). 1174


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Figure 5. Structural comparison among the n = 3 crystal structures of RP phases, DJ phases, and ACI phases: (a,e,i) (PEA)2(MA)2Pb3I10; (b,f,j) (BA)2(MA)2Pb3I10; (c,g,k) (3AMP)2(MA)2Pb3I10; and (d,h,l) (GA)(MA)3Pb3I10.

or flexible ditopic ones.117,120 Until recently, the only building unit known was the most common “2×2” corrugation variation. Nevertheless, new results revealed that the corrugation “length” could be increased to “3×3” 120 and even to a “4×4” 95,197 undulated structure (corrugation length defined by the number of octahedra in one unit, i.e., “2×2” as seen in the structure of (EDBE)PbBr4 in Figure 3) . The (110)-oriented perovskite layers are intrinsically highly distorted, often stabilized by site-specific hydrogen bonding or other secondary bonding interactions. Because of this, (110)-oriented perovskites often exhibit white-light emission at room temperature believed to be caused by the formation of self-trapped excitons (STEs),198 as for example (EDBE)PbBr4 (“2×2”)117 and α-(DMEN)PbBr4 (“3×3”).120 Very recently, the asymmetric “3×2” undulated structure was realized in [CH(NH2)2][C(NH2)3]PbI4,161 which is the first member of the wider class of (210)-oriented 2D perovskites.162 The (111)-oriented 2D perovskites, with a general formula A′n+1BnX3n+3, are in fact a class of defective perovskites.199 Unlike the rest of the layered perovskite configurations, (111)oriented perovskites are M-site-deficient since cleavage along the body diagonal of the 3D cubic cell selectively “eliminates” the metal sites, in contrast to other cleavage planes which eliminate neutral BX2 fragments. As a result, B2+-only halide

(111)-oriented 2D perovskites cannot form, with the structures cleaved along the body diagonal preferring alternative 3D polytype configurations that maintain charge parity.137 Thus, the only homovalent p-block compounds in this class of materials are composed of group 15 B3+ ions (B = Bi, Sb, As) belonging to either the pseudo-layered Cs3Cr2Cl9-type200 0D [B2X9]3− dimers or the Cs3Bi2 Br9 -type 201 2D (B2X9) 3− defective perovskite layers.202−206 Very recently, Solis-Ibarra and co-workers have managed to bypass the charge mismatch obstacle by synthesizing the heterometallic compound Cs4(CuSb2)Cl12, which is the only known example for n > 2.140 In this heterobimetallic perovskite, the Cu2+ (or Mn2+)207 ions can be inserted between the Sb3+ sheets, forming the first example of the n = 3 structure of the general formula. The recent progress in the generally unexplored, “exotic”, 2D perovskite structural types has clearly shown that there is plenty of room for new material discovery. By further elucidation of the synthetic parameters we can rapidly move from the graphical cartoons of the general formulas to the isolation of useful materials on the bench with detailed structural characterization, thus pushing the boundaries of material understanding and exploitation. 1175


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(100)-ORIENTED 2D PEROVSKITES: STRUCTURE DIRECTING EFFECTS OF SPACERS The (100)-oriented 2D perovskites comprise by far the most abundant class, which has been explored for a wide variety of cation templates. The geometrically planar inorganic−organic interface of the (100)-oriented layers not only allows for the incorporation of a plethora of “spacer” cations between the 2D perovskites but also allows for the templated “build-up” of the perovskite itself, which self-assembles in a layer-by-layer fashion between the spacers. These modular multilayer structures were first discovered in the oxide perovskites, where they were used to expand extended homologous series.55,208 Following the oxide perovskite nomenclature, the (100)-oriented oxide perovskites can be further categorized as the Ruddlesden−Popper (RP) phases,209,210 the Dion− Jacobson (DJ) phases,211,212 and the Aurivillius phases,213−215 the general formulas of which can be written as A′2An−1BnO3n+1, A′An−1BnO3n+1, and (Bi2O2)(An−1BnO3n+1), respectively. In the case of halide perovskites, most of these structure types have been realized, with the exception of the Aurivillius phases, which have not yet been reported, possibly due to the reactivity of (Bi2O2)2+ toward the formation of halides. It is entirely possible, however, that such layers can be formed using less reactive rock-salt-type layers, such as (Ba2F2)2+. As in the case of oxide perovskites, also in the case of halide perovskites, the nature of the spacer cation is important. Already from the expression of the chemical formula it is evident that the charge of the spacer cation defines the structural type of 2D perovskites. Nevertheless, both the RP and DJ classes exist for the halide family, having the general formulas A′2An−1BnX3n+1 and A′An−1BnX3n+1 (A′ = interlayer “spacer” cation), respectively, as can be seen in Figure 5. A new structural type, unique to the halide perovskites, is the alternating cation in the interlayer space (ACI) type, where different cations order along the interlayer space.63,160 What distinguishes ACI from the RP and DJ perovskites is reflected directly in the formula A′AnBnX3n+1, as the A-site “perovskitizer” cation (MA+, FA+, or Cs+) not only resides in the perovskite slab but also fills in the interlayer along with the designated spacing cation A. Interestingly, guanidinium (GA+) is the only reported A′ cation that can form the ACI-type structure, which suggests a preference for small cations that are similar in size to the “perovskitizer” cations (MA+, FA+, or Cs+) to form the ACI phases. The crystal structures of all known structural types are shown in Figure 5. The side-by-side comparison of the n = 3 members of each of the RP, DJ, and ACI homologous series illustrates their structural differences and similarities across the three crystallographic directions. For the RP phase (BA)2(MA)2Pb3I1060,105 the layers are offset by one octahedral unit, showing a (1/2,1/2) in-plane displacement. On the other extreme, the layers in the DJ phases (3AMP)(MA)2Pb3I1062 do not exhibit any shift and are stacked perfectly on top of each other (Figure 5k), with (0,0) displacement. The ACI-type structure lies in between and combines the layer stacking characteristics from both RP and DJ structural types (Figure 5l), showing a (1/2,0) displacement. Obviously, the interlayer distance varies with the choice of spacing cation A′, and this turns out to be one of the crucial factors that decide the excitonic PL emission from these quantum well (QW) structures (Figure 5). Typically, the RP phases have larger interlayer distance due to the fact that they

require a bilayer of spacer organic cations, which inevitably will increase the layer separation from approximately 1.5 times the length of the cation in the case of interdigitation (as in the BA case) to more than double the cation’s length in cases where there is no overlap (like in PEA). Thus, in the case of the RP phases the layers are well “isolated”, showing no perovskite− perovskite interactions, thereby producing an excitonic emission that results directly from the spatial electronic confinement (n) of the individual layers. Conversely, the ACI phases, and more so the DJ phases, require the organic cations to snugly fit in the perovskite “pockets” created in the interlayer space by the eclipsed stacking conformation of the perovskite layers. The close proximity of the layers generates a weak perovskite−perovskite interaction which can be readily observed in the position of the exciton, which shows a pronounced redshift. The representative interlayer distances for the RP, DJ, and ACI phases shown in Figure 5 are ∼7, ∼4, and ∼3 Å, respectively. Nevertheless, particularly for the RP phases, the value for BA represents a lower limit, as longer interlayer distances of up to ∼26 Å have been reported, as for example in (C18H40N)2PbI4,151 where the massive alkylammonium cation CH3(CH2)17NH3+ was employed. While the exciton emission trend largely depends on the quantum confinement, there are structural subtleties within a given n member that can also have a significant impact on it. These relate to the distortion of the 2D perovskites lattice, which has the least distorted lattices (in terms of B−X−B angle, see Figure 6a), to have the smallest exciton emission

Figure 6. (a) Definition of the Pb−I−Pb angles. (b) Averaged Pb−I− Pb angles from the n = 1−3 structures of (GA)(MA)nPbnI3n+1, (3AMP)(MA)n−1PbnI3n+1, (BA)2(MA)n−1PbnI3n+1, and (4AMP)(MA)n−1PbnI3n+1. Panel c reproduced with permission from ref 62. Copyright 2018 American Chemical Society.

energy, in complete correspondence with the band gap trends of the 3D perovskites.50 This can be readily observed from the photochromic change in the crystals, moving from a lowsymmetry structure at low temperature to a high-symmetry structure at high temperature in the n = 1 RP series,98,151,216 or from changes in the structure induced by cation templating effects at a given temperature in the DJ series (Figure 6c).62 In selected cases, where the synthesis and characterization of the complete homologous series of the RP, DJ, and ACI phases have been achieved, it is possible to analyze their structure− property correlations. Comparing between n = 1 and 3 for all families, the energy of the excitonic emission (as determined by the PL emission at room temperature) decreases for all the series. However, the relative emission energies are markedly different, following a systematic trend with the emission energy 1176



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(100)-ORIENTED 2D PEROVSKITES: TUNING THE LAYER THICKNESS One of the most attractive features of the 2D halide perovskites is that they can be tailored to the exact n-layer thickness using simple synthetic chemistry, thus allowing for the fine-tuning of the QW electronic structure. The (100)-oriented 2D perovskites are ideal for this “digital” growth due to their favorable geometry that guarantees the “high-fidelity” of the layer thickness. Indeed, careful composition control can generate all the “quantized” RP perovskites up to n = 7 for the (BA)2(MA)n−1PbnI3n+1 series (see Figure 9a, below), which is the thickest 2D perovskite with long-range crystallographic order known to date. Note that higher n members have been observed for oxide-based RP perovskites,217−219 yet morphological effects preclude the formation of the ordered high-n members (n > 5) since these compounds do not have sufficient thermodynamic stabilization.220 The mechanism for this instability is the disproportionation reaction (1):

decreasing in the order 3AMP (DJ) < GA (ACI) < 4AMP (DJ) < BA (RP) (Figure 7a). The origin of this particular trend

Figure 7. (a) Comparison of the PL emission among the n = 1−3 members of (BA)2(MA)n−1PbnI3n+1, (3AMP)(MA)n−1PbnI3n+1, and (GA)(MA)nPbnI3n+1. (b) PL emission energy versus layer thickness (n). (c) Averaged Pb−I−Pb angles versus layer number (n).

disproportionation (RNH3)2 (MA)n − 1PbnI3n + 1 ⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯⎯→

correlates well both with the structural distortion, defined by the octahedron-linking Pb−I−Pb angles, and with the interlayer spacing in the crystal structure. The larger angle represents less distortion, generally leading to a relative narrowing of the PL emission energy among the series. However, the n = 1 is an exception in the trend, with (GA)(MA)PbI4 having a larger average Pb−I−Pb angle but still a slightly higher PL emission energy than (3AMP)PbI4, possibly due to the large out-of-plane distortion of (GA)(MA)PbI4. This deviation indicates that other factors, such as the anisotropy of the lattice118 or bond length variation, also play roles in determining the optical properties of the QW structures. It is important to point out that the differences in the Pb−I− Pb angles become very small as the layers get thicker (higher n members) for all structure types, as can be seen in Figure 6b, suggesting that the distortion of the lattice stops being the determining factor and that the interlayer spacing is becoming the dominant factor. However, in systems where all other things are equal, structural distortion is still important, as in the case of 3AMP and 4AMP DJ perovskites (Figure 6b,c), where the Pb−I−Pb angles are consistently larger for the latter, thus leading to the systematically lower excitonic emission.62

(RNH3)2 (MA)2 Pb3I10 + (n − 3)MAPbI3

(1)

where large perovskite slabs tend to break down into more thermodynamically stable phases such as the thinner 2D perovskite slabs and the 3D perovskite. For instance, the n = 5 member of the BA series undergoes a very slow disproportionation reaction and decomposes into the hydrated form of MAPbI3 (MAPbI3·H2O)221 and the n = 3 2D member over the course of several weeks.61 As the thickness of the perovskite layers increases, the dimensions of the stacking axis also increase, progressively moving from the atomic scale (few Å for n = 1) to the nanoscale (few nm for n = 7). The dimensions of the layers can be even roughly predicted following the general equation (2): d(001) = a*n + x

(2)

where the complete inorganic−organic layer thickness d(001) is determined by the lattice parameter (a*) of the 3D perovskite (∼6.3 Å for MAPbI3), the thickness (x) of a complete organic layer (∼7 Å for the (BA)22+ bilayer; ∼3 Å for the (3AMP)2+ monolayer), and the number (n) of inorganic layers.125 Because of their long-range crystallographic ordering, the

Figure 8. (a) High-resolution PXRD of (BA)2(MA)n−1PbnI3n+1 (n = 1−5). (b) PXRD plotted in d-spacing. (c) Structure of (BA)2(MA)Pb2I7 with the illustration of respective diffraction planes. 1177


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Figure 9. (a) Crystal structures of the RP phase (BA)2(MA)n−1PbnI3n+1 (n = 1−7). Optical absorption spectra (b) and steady-state PL spectra (c) of the series.

tunability offered by the 2D halide provides an excellent playground for designing new materials with unexplored photophysical properties that could be exploited for advanced optoelectronic applications.

phase purity of the bulk materials can be tracked down from their “fingerprint” reflections in X-ray diffraction (XRD) patterns. Thus, the number of so-called basal low-angle Bragg (00l) peaks (when c is the stacking axis) or reflections ((h00) or (0k0) depending on the stacking axis lattice parameter) below the characteristic 3D perovskite reflections (d(110) = 6.26 Å and d(001) = 6.34 Å for tetragonal MAPbI3) precisely defines the number of perovskite layers (Figure 8c). The change in the dimensionality can readily be seen in the evolution of the optical absorption spectra of the QWs. As the 3D perovskite is dimensionally reduced to form the 2D layers, so are the optical absorption edges of the multilayer 2D series, which increase in energy with decreasing n value. Using the (BA)2(MA)n−1PbnI3n+1 series as a representative example, the absorption edge blue shifts from 1.52 eV (n = ∞) to 2.43 eV (n = 1) (Figure 9b). Meanwhile, in addition to the characteristic halide perovskite’s sharp absorption edge, an exciton-like peak starts emerging for the n < 5 members, becoming more pronounced as n decreases further. The PL emission follows the same trend, with the pronounced difference that emission peak maxima now coincide with the cross section of the exciton-like feature and the absorption edge, ranging from 1.6 eV (n = ∞) to 2.35 eV (n = 1) (Figure 9c). The difference between emission and absorption in 2D perovskites is not well understood yet, but it is no different than the corresponding absorption and emission shown in the spectra of the 3D perovskites.222,223 The complexity, however, of the 2D perovskites increases even more, when one considers the magnitude of the exciton binding energy in 2D halide perovskites, which can be as high as Eb ≈ 500 meV for the n = 1 member224 at low temperatures. In order to reconcile the huge exciton binding energies with the observed spectra, alternative electronic structure models have been proposed that include the reduction of the effective dimensionality of the halide perovskites from to 2D to quasi1D models.132,225 Even though theoretical investigations in this direction are still ongoing, it is clear that the wide range of

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(100)-ORIENTED 2D PEROVSKITES: EFFECT OF THE PEROVSKITIZER AND SOLID SOLUTIONS Despite the rapid development of 2D perovskites in the past 5 years, it has been remarkably difficult to isolate homologous series of perovskites that contain A-site cations different than MA+ as “perovskitizers”.137 With the exceptions of ethylammonium (EA+) and isopropylammonium (IPA+), which can produce the line (i.e., not part of a homologous series) 2D compounds (EA)2(EA)2Pb3X10121 (X = Cl, Br) and the metastable (IPA)2(IPA)Sn2I7,137 in which the organic cations occupy both A′- and A-site positions in the same compounds, it has not been possible to expand 2D perovskite homologies using the common Cs+ and/or FA+ perovskitizers. While on paper this would appear to be possible, the intrinsic instability of Cs+ and FA+ toward the formation of the CsPbI3 and FAPbI3 perovskites (both metastable)51 is apparently transferable to the 2D perovskites as well. In contrast to the 3D perovskites, where the thermodynamic phase is a nonperovskitic phase, in 2D perovskites the thermodynamic product appears to be the n = 2 member of the homologous series instead. This has been demonstrated in a number of 2D perovskite families, from RP126 to ACI160 and to double perovskites,144 so that one could tentatively conclude that it is the perovskitizer itself that inhibits the formation of the homologous series, and in spite of the fact that the higher-n members may be possible to form, it is clear that they are not the thermodynamic products of these systems. In order to bypass this obstacle, several researchers have turned to the familiar remedy employed in 3D perovskites:226 the mixing of cations in the A-site in order to tune the tolerance factor.227 Although this approach paid some dividends in terms of improved devices when the mixed compositions were 1178


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Journal of the American Chemical Society employed to cast thin films,107,128,228 the complete lack of structural characterization leaves the exact composition of the films in question. Nevertheless, in a few cases structural characterization has been achieved, showing that incorporation of FA is possible, as in the n = 4 DJ perovskite (3AMP)(MA)2.25(FA)0.75Pb4I13 (unpublished results in our laboratory).229 The mixed-cation, solid solution approach in 2D perovskites is not limited to the A-site cations; it has also been explored with respect to metal alloying and halide mixing. These mixing metals or halides are very appealing since they can substantially change the energy band levels in the perovskite. Conceptually, they can also be envisioned to substitute selectively into different layers. These systems, however, have not yet drawn much attention compared to the 3D perovskite analogues,230−232 and they are limited to a handful of examples. The metal alloying has so far been explored in the single-layer mixed Pb and Sn systems of (HA)Pb1−xSnxI495 and (BZA)2Pb 1−x Sn x I 4 95 (HA = histammonium, BZA = benzylammonium) and in the mixed Sn and Ge system, (PEA)2Ge1−xSnxI4.233 In the former case, the Pb/Sn mixing produces the “typical” anomalous band gap trend for the 3D perovskites, not following Vegard’s law between the two parent compounds.234,235 Particularly in the case of (BZA)2Pb1−xSnxI4, all the intermediate (x = 0.25, 0.5, 0.75) compositions have lower band gaps (1.86, 1.84, and 1.82 eV) than the pure Sn compound (1.89 eV). As expected from the 3D perovskite precedence,236 in the mixed Sn/Ge system the band gap decreases linearly with increasing Sn content, from 2.13 eV for x = 0 to 1.95 eV for x = 0.5. Halide mixing has also seen very little practical study in 2D perovskites (contrary to their extensive exploration in 3D materials237 and 0D nanocrystals32), and it has been mainly employed in Cl/Br systems for optimizing the quantum yield of white-light emission, such as in (EDBE)[PbBr4−xClx],117 EA4Pb3Br10−xClx, and (PEA)2PbBrxCl4−x,238 where a small addition of Cl can significantly broaden the PL emission bandwidth. Mixing Br with I in the (PEA)2SnIxBr4−x series causes shifts of the excitonic emission from 2.66 eV (x = 0) to 1.97 eV (x = 4),239 while in the complete series (PEA)2PbYxZ4−x the whole spectrum between 2.4 eV (Y = I, x = 4) and 3.8 eV (Z = Cl, x = 0) can be accessed by adjusting the halide content between Cl, Br, and I binary mixtures.240

Figure 10. (a) Variation of the interlayer c parameter in (CnH2n+1NH3)2PbI4 (n = 12, 16, 18) compounds as a function of temperature. A small decrease in the interlayer spacing may be seen at the premelting transition, followed by an expansion at the melting transition. (b) Schematic representation of the melting transition (TM) in the (CnH2n+1NH3)2PbI4 compounds. Reproduced with permission from ref 246. Copyright 2003 American Chemical Society.

the organic layers has a dramatic effect on the crystal structure, because it is associated with the interlayer spacing, often exhibiting a large jump in the lattice constant along the stacking direction, as can be seen in Figure 10a for the cases of (CnH2n+1NH3)2PbI4 (n = 12, 16, 18),246,247 and an increase in the symmetry of the inorganic layers, typically from distorted monoclinic to less distorted orthorhombic space groups (Figure 10b).98,151,216 The “melting” temperature is strongly dependent on the length of the alkyl chain in straight alkylammonium spacers, and it drops to just below room temperature (∼5 °C) in the case of the BA cation.98 Naturally, increasing the thickness of perovskite layers to n > 1 increases the phase transition complexity, as the structural changes associated with the MAPbI3 3D perovskite start competing with the structural changes induced by the organic spacers.248 Interestingly, the incremental change in the perovskite thickness does not seem to affect the “melting” of the organic layers, which remains practically constant, yet it has a dramatic effect on the phase transitions that would correspond to the β−γ phase change (in the 3D perovskite), which shifts to higher temperatures with decreasing n. Temperature-induced phase transitions are commonly observed in 2D perovskites, especially in the single-layer systems. For instance, (C7H15NH3)2PbI4 exhibits three phases from 253 to 293 K (Figure 11).216 The band gap of the compound changes according to the level of the structural distortion (i.e., Pb−I−Pb angle),

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2D PEROVSKITES UNDER EXTERNAL INFLUENCES: TEMPERATURE AND PRESSURE EFFECTS Much like the 3D halide perovskites,50,241−244 2D perovskites show great structural versatility when exposed to external influences such as non-ambient temperature and pressure. The presence of the soft organic spacers between the flexible inorganic layers makes the structural response even more complex as the two layers try to adjust independently to the external stimulus. In two dimensions, the inherent liquid-like dynamic disorder of perovskites44,45,245 blends with the weak dispersion forces governing the crystal packing of the organic spacers, producing a truly hybrid material. The temperature-dependent behavior of the 2D perovskites strongly attests to the hybrid nature of the materials, especially for the n = 1 perovskites, which show a characteristic “melting” (i.e., dynamic disorder) of the organic layers at elevated temperatures, whereas the “melting” of the inorganic layers occurs at lower temperatures (Figure 10).246 The “melting” of

Figure 11. (a) Phase transitions and corresponding color change for (C7H15NH3)2PbI4 from (a) monoclinic, P21/a, 253 K, to (b) orthorhombic, Pbca, 278 K, to (c) orthorhombic, Pbca, 293 K. The color change reflects the band gap change induced by the Pb−I−Pb angle evolution with varying temperature. Crystal pictures reprinted with permission from ref 149. Copyright 2012 Royal Society of Chemistry. 1179


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Figure 12. (a) Schematic illustration of Pb−I−Pb bonds in (BA)2PbI4 under compression. Under initial compression from 1 atm to 0.5 GPa, the slight blueshifts of bandgap can be understood from the reduced orbital overlap between Pb and I caused by decreased Pb−I−Pb bond angles. As pressure continues to increase, the perovskite lattice is compressed, with the bond lengths of Pb−I−Pb being unavoidably shortened, leading to the increased electronic wave function overlap between Pb and I, and thereby accounting for the narrowed optical bandgap. (b) Energy variation as a function of bond angle after pressure treatment on (BA)2PbI4 (n = 1) (blue) and (BA)2(MA)2Pb3I10 (n = 3) (yellow). (Inset) Pb−I−Pb angledependent bandgap evolution of (BA)2PbI4 (n = 1). (c) Pressure-dependent bandgap evolutions of (BA)2PbI4 (n = 1) (blue), (BA)2(MA)Pb2I7 (n = 2) (light green), (BA)2(MA)2Pb3I10 (n = 3) (green), (BA)2(MA)3Pb4I13 (n = 4) (orange), and MAPbI3 (n = ∞) (red). Panels a−c reproduced with permission from ref 258. Copyright 2018 National Academy of Sciences. (d) Piezochromism in single crystals of (CH3NH3)2[PbI2(SCN)2]. (e) Lattice-parameter evolution with pressure for (CH3NH3)2[PbI2(SCN)2]. Panels d,e reproduced with permission from ref 260. Copyright 2016 American Chemical Society.

where the larger the angle, the smaller the band gap (darker color of the crystal), as can be seen in Figure 11. Very recently, the application of pressure on the perovskites was shown to provide useful insights into the soft nature of their lattice. Pressure has also been used as an effective tool to manipulate the optical and electrical properties of the 3D perovskites.241,249−252 A panoply of experimental techniques, ranging from the application of hydrostatic pressure242−244 to mechanical indentation,253,254 inelastic scattering,255 and resonance acoustic spectroscopy,256,257 have been brought to bear, aiming to understand the lattice deformation parameters under stress. In this context, studying the compressibility of the 2D perovskites becomes increasingly interesting since the anisotropy of 2D perovskites and the combination of two soft organic and inorganic sublattices may give rise to unusual phenomena. Depending on the direction of the applied pressure, the crystal lattice will distort, causing octahedral tilting and directional bond contraction, while the organic layers can operate as buffers to relieve the excess stress. This in fact turns out to be the case for the (BA)2(MA)n−1PbnI3n+1 homologous series, which exhibit a characteristic structural change upon compression, specifically reflected in the narrowing of the Pb−I−Pb angle and shortening of the Pb− I bond length, as illustrated in Figure 12a.258 The application of isotropic pressure initially shrinks the lattice along the layer stacking axis at relatively lower pressures: ∼7 GPa (n = 1); ∼4 GPa (n = 2); ∼3 GPa (n = 3); ∼2 GPa (n = 4); ∼0.5 GPa (n = ∞). This is followed by the in-plane deformation which was readily observed in the change of the Pb−I−Pb angles of the layers at elevated pressures, prior to amorphization, upon which the materials start behaving nearly isotropically (Figure 12b).258,259

The structural changes are accompanied by changes in the absorption of the materials, showing an initial decrease in the bandgap as the perovskite becomes more symmetric, as can be seen in Figure 12c. When the compressibility of the organic molecules saturates, changes start occurring within the layers, leading to a lowering of the symmetry within the perovskite layers and resulting in the blueshift of the absorption edge. Once this second deformation saturates, 2D perovskites’ absorption starts to decrease again monotonically, similarly to the 3D perovskite.252 The lattice constants and energy bandgaps also shrink with increasing pressure (Figure 12d,e).260 There are pronounced differences in both absorption and emission trends when different organic cations are employed as spacers. For instance, the single-layer n = 1 perovskite (EDBE)CuCl4 undergoes a color change from yellow to orange when the pressure is increased from 0.3 to 4.9 GPa, as it undergoes the α-to-β phase transition, which is followed by a monotonic decrease in the absorption edge at elevated pressures.261 The lack of the initial decrease in absorption in (EDBE)CuCl4 can perhaps be rationalized by the much stronger mechanical coupling between the spacers (2+) and the inorganic layer in the latter case, which can be readily probed by the coherent acoustic phonon oscillations between the n = 1 members at ambient pressure.130 Assessing the properties of 2D halide perovskites upon external stimuli, it is clear that they constitute an unconventional class of “malleable” semiconductors, the properties of which can be drastically modified with small changes in temperature and pressure, clearly providing plenty of room for material design and property engineering. 1180


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2D PEROVSKITES IN OPTOELECTRONICS: SOLAR CELLS, LIGHT-EMITTING DIODES, AND BEYOND Several 2D perovskites have been deployed in actual devices that show superb performance. The unique advantages derive from their nanostructured template and the increased organic component content. Their hybrid inorganic−organic nature allows for the development of multifunctional materials with the added advantage over the 3D perovskites of having enhanced environmental stability,104 reduced ionic migration,262 and easily tunable electrical properties.105 On the other hand, the increased organic content slightly dilutes the functionally active perovskite layers, which as a result decreases somewhat the maximum theoretical light-to-current conversion. The high exciton binding energy limits the diffusion length of the 2D perovskite, where the diffusion length increases with decreasing organic spacer content.20 Though not as high as the 3D counterpart, the carrier mobility in the 2D perovskite can reach about 10 cm2 V−1 s−1 for 50% PEA in the MAn−1PEA2PbnI3n+1 system.20 Despite starting the race with a small disadvantage, 2D perovskites have proven their worth, as the initial successful demonstration of (PEA)2(MA)2Pb3I10,104 which yielded a 4% power conversion efficiency (PCE), was superseded by (BA)2(MA)2Pb3I10, which exhibited a similar PCE.105 However, the massive breakthrough came in the form of (BA)2(MA)3Pb4I13, which exhibited a PCE of 12.52%,263 using the hot-casting method,264 showing impressive long-term stability against both light and moisture. Since the initial demonstration of 2D perovskite solar cells, small improvements were achieved in the BA systems by the mixed A-cation (FA/MA)107 (Figure 13a,b) and Cs/MA228 strategies. Interestingly, higher-n members do not improve the efficiency, leading to an optimized PCE ≈ 10% for the n = 5 member.265 A dramatic improvement of PCE from 1.74% to over 15% can be achieved by addition of MACl into the (ThMA)2(MA)2Pb3I10 system (ThMA = 2-thiophenemethylammonium).266 The PCE could be further improved by employing the so-called “quasi-2D perovskites”, targeting nominal compositions with high layer number (e.g., n = 10, 50, 60). This common strategy afforded solar cells with higher efficiency (∼15%) for various cations and in different 2D systems (n = 60, PEA;103 n = 10, ALA (allylammonium)106). The work by Sargent and coworkers investigated the effects of incorporating different organic spacers (BA, PEA, and ALA) (Figure 13c), where they found that the ALA system, with the highest lifetime, diffusion length, and mobility, reached the highest efficiency compared with BA and PEA (Figure 13c).106 This approach was successfully extended to lead bromidebased systems, yielding a PCE of ∼10% for “(BZA)2(MA)49Pb50Br151”,267 and also in the Sn-based 2D systems, leading to an improved efficiency and stability for “(PEA)2(FA)8Sn9I28” (5.94%).268 It should be noted that the nominal stoichiometry is not equal to the actual layer thickness, where no evidence was found for the existence of a system with n = 50. The resulting experimental input material needs to be validated by XRD, NMR, or other characterization methods. The lack of identification of the material leads to confusion about its corresponding properties and difficulty in comparing materials of the same layer thickness. 2D halide perovskites also perform excellently in the arena of LEDs, in this case displaying an even better performance than the 3D perovskites in certain cases.269 In the (PEA)2(MA)n−1-

Figure 13. (a,b) Solar cell application of the RP phases (BA)2(MA)3Pb4I13 and (BA)2(MA0.8FA0.2)3Pb4I13: device architecture and J−V curve of the device. (c) Transient absorption (TA) spectra for ⟨n⟩ = 10 films of ALA, PEA, and BTA at various delay times. (d) EQE vs applied voltage characteristics and (e) J−V−L characteristics of the n = 3 and 5 (PEA)2(MA)n−1PbnBr3n+1 and MAPbBr3 devices. Reproduced with permission from refs 107, 106, and 109. Copyright 2018, 2018, and 2017, respectively American Chemical Society.

PbnX3n+1 systems,270 the n = 5 member demonstrates a significantly higher external quantum efficiency (EQE) at 8.4% for red light emission for X = I, compared to MAPbI3. The bromide analogue (PEA)2(MA)4Pb5Br16 (nominal composition) also exhibits a high EQE (7.4%) for green light emission and luminescence (8400 cd/m2) at low threshold voltage (Figure 13d,e).109 Interestingly, the BA systems also deliver high quantum yields in the higher “quasi-2D” regime,271 but the efficiency is much lower for registered low-n members.272 The high mobility of the Sn-based 2D halide perovskites has also been utilized in field-effect transistor (FET) devices.273 Early work clearly showed that FET could be realized for the n = 1 (PEA)2SnI4 perovskite derivatives,35,183,274 producing FET mobilities on the order of up to μ ≈ 1.4 cm2 V−1 s−1; this has been recently revisited, showing an impressive enhancement to μ ≈ 15 cm2 V−1 s−1.275 The above qualities of the 2D perovskites make them obviously very promising for implementation in optoelectronic devices, yet a serious setback arises from the fact that, in thinfilm form, the 2D perovskites lose one of their main advantages, that of the controlled tunability of the layer number. This loss of synthetic control affects the excitonic emission wavelength (mostly relevant for Pb-based perovskites) and potentially also the loss of FET mobility (mostly relevant for the Sn-based perovskites). In most cases, targeting a layer thickness in thin films yields a combination of different layer numbers, as seen in the absorption spectra (several slopes) and sometimes low-angle peaks in powder XRD.106,107,265 Despite the loss of structural integrity, the superior performance of the 2D perovskites persists, and this is 1181


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Journal of the American Chemical Society likely attributed to the engineered mixtures of 2D phases that funnel their energy from high band gap to lower band gap phases in a process which surprisingly reduces nonradiative recombination pathways (Figure 13d,e). Among the emerging applications for which 2D perovskites hold great promise due to their strong excitonic emission at room temperature are in the fields of lasing and photodetectors. These device structures are still in their initial steps of development since they require careful engineering of the device and successful mechanical exfoliation of the 2D materials. Nevertheless, the initial results are very promising, as was demonstrated for the n = 1 (PEA)2PbI4 polariton laser276 and the (BA)2(MA)3Pb4I13 engineered photodetectors.111,277 Such new device design principles may even allow for fast hard radiation scintillation detectors for medical applications since the intrinsic excitonic emission in these systems lies in the order of a few picoseconds.278,279 The nonlinear optical responses of the 2D halide perovskites, which tend to amplify upon reduction of dimensionality in excitonic systems,280,281 were also the focus of the early work. The main attention was given to the n = 1 system (C 6H 13 NH3 )2 PbI 4, 282 producing a massive large third harmonic generation coefficient, χ(3) = 1.6 × 10−6 esu, at cryogenic temperatures. The successful isolation of higher-n members in the (BA)2(MA)n−1PbnI3n+1 system allowed for the determination of the χ(3) coefficients for higher-n members as well,125 which for the mid-infrared region at room temperature range between 4.6 × 10−11 and 5.6 × 10−11 esu and could be utilized for medical applications.33 Provided that the 2D perovskites could be integrated in suitable device structures, they should have a bright future in the field of optoelectronics.

ness) is going to be helpful to understand these materials systematically. (3) Controlling the optoelectronics response and quantum ef ficiency. PLQY is an important parameter to understand and optimize if the materials are going to find uses in light-emitting devices. So far little is known regarding the origins of higher PLQY in some 2D systems and its variation with preparation method and material form (films vs crystals). (4) Accessing higher layer thickness (n ≥ 5). So far, most effort has been devoted to synthesis and characterization of simpler forms and lower layer thicknesses (n ≤ 4). The higher layer thicknesses (n ≥ 5) in general are more difficult to synthesize and characterize because of being less thermodynamically favorable. Advancing the synthetic techniques and characterization methods will help unlock a new territory for 2D halide perovskites. Clear identification of the phases and layer thickness is required for studying the properties of the materials, regardless of their form (single-crystal, bulk powder, or thin-film). (5) Phase-selectively depositing 2D perovskites on f ilms. Most of the applications of 2D materials have been reported for cast films incorporated in thin-film devices, yet the fundamental properties of the films are markedly different than those of the pristine crystals, particularly with respect to their QW properties. To be able to understand and control these properties, significant progress will be required in understanding the precise composition and orientation of the thin films.

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We anticipate that addressing these five pillars will provide a huge boost in the potential of these materials in technological applications.

OUTLOOK AND FUTURE DIRECTIONS The 2D halide perovskite family serves as a great platform for examining new physicochemical principles and establishing structure−property relationships. Its modular structural sequence adds additional structural flexibility and allows rational design for future material development. By varying the chemical “pressure” internally or applying stress externally, the optical and electronic properties of the materials can be altered at will, either permanently or temporarily, providing immense opportunities for tailoring the properties of the 2D halide perovskite toward specific applications. As the complexity of these systems incrementally increases, collaborative efforts from multidisciplinary researchers are required in order to understand the physical properties and enhance the performance in optoelectronic devices, so that 2D halide perovskites progress to higher levels. For this to happen, we envision that the following pillars need to be strengthened in the future: (1) Elevating the synthetic chemistry output. The vast library of organic molecules that could intercalate between the perovskite layers provides huge potential for novel hybrid combinations. So far, the properties of the 2D perovskites highly depend on the inorganic counterpart with only a few exceptions. Functionalization on the organic side will open the door for new property discovery. (2) Unveiling the mystery of the excited state. Use of ultrafast spectroscopies to study the properties at the excited state of the different kinds of 2D perovskites (different layer orientation, metal/halide composition, and layer thick-

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AUTHOR INFORMATION

Corresponding Authors

*[email protected] *[email protected] ORCID

Lingling Mao: 0000-0003-3166-8559 Mercouri G. Kanatzidis: 0000-0003-2037-4168 Present Address †

Department of Materials Science and Technology, University of Crete, Vassilika Voutes GR-700 13 Heraklion, Greece. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS Work done by the authors cited in this Perspective article was supported mainly by the U.S. Department of Energy, Office of Science, Basic Energy Sciences, under Grant SC0012541.

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