Synthesis, optoelectronic properties and applications of halide perovskites

Lata Chouhan a, Sushant Ghimire a, Challapalli Subrahmanyam b, Tsutomu Miyasaka cd and Vasudevanpillai Biju *a
aGraduate School of Environmental Science and Research Institute for Electronic Science, Hokkaido University, Sapporo, Hokkaido 001-0020, Japan. E-mail: biju@es.hokudai.ac.jp
bDepartment of Chemistry, Indian Institute of Technology Hyderabad, Kandi, Sangareddy, Telangana, India
cFaculty of Biomedical Engineering, Toin University of Yokohama, Yokohama, Japan
dResearch Center for Advanced Science & Technology, The University of Tokyo, 4-6-1 Komaba, Meguro-ku, Tokyo 153-8904, Japan

Received 5th December 2019

First published on 27th April 2020


Halide perovskites have emerged as a class of most promising and cost-effective semiconductor materials for next generation photoluminescent, electroluminescent and photovoltaic devices. These perovskites have high optical absorption coefficients and exhibit narrow-band bright photoluminescence, in addition to their halide-dependent tuneable bandgaps, low exciton binding energies, and long-range carrier diffusion. These properties make these perovskites superior to classical semiconductors such as silicon. Most importantly, the simple synthesis of perovskites in the form of high quality films, single crystals, nanocrystals and quantum dots has attracted newcomers to develop novel perovskites with unique optoelectronic properties for optical and photovoltaic applications. Here, we comprehensively review recent advances in the synthesis and optoelectronic properties of films, microcrystals, nanocrystals and quantum dots of lead halide and lead-free halide perovskites. Followed by the classification of synthesis, we address the ensemble and single particle properties of perovskites from the viewpoints of the confinement and transport of charge carriers or excitons. Further, we correlate the charge carrier properties of perovskite films, microcrystals, nanocrystals and quantum dots with the crystal structure and size, halide composition, temperature, and pressure. Finally, we illustrate the emerging applications of perovskites to solar cells, LEDs, and lasers, and discuss the ongoing challenges in the field.


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Lata Chouhan

Lata Chouhan studied chemistry at the Indian Institute of Technology, Bombay. She was awarded a JICA fellowship in 2017 for doctoral research in Hokkaido University. First, she was a JICA research student from October 2017 to March 2018. Since April 2018, she has been a PhD student under the guidance of Vasudevanpillai Biju in Hokkaido University. She focuses her research on single particle microscopy and spectroscopy for the control of charge carrier and exciton recombination in nanocrystals and quantum dots of lead halide perovskites.

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Sushant Ghimire

Sushant Ghimire studied chemistry at Tribhuvan University in Nepal. In 2016, he was awarded a MEXT Scholarship for doctoral research in Hokkaido University. He was a MEXT research student from April 2016 to March 2017 and a PhD student from April 2017 to March 2020 under the guidance of Vasudevanpillai Biju in Hokkaido University. He received his PhD degree in Environmental Materials Science in 2020. His research focuses on the dynamics of charge carriers in perovskite nanocrystals, assemblies, and thin films.

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Challapalli Subrahmanyam

Challapalli Subrahmanyam obtained his PhD degree in Chemistry from the Indian Institute of Technology Madras in 2003. From 2003 to 2007, he was a postdoctoral fellow in the École polytechnique fédérale de Lausanne, Switzerland. He started his academic career as an assistant professor in the National Institute of Technology Trichy, India. In 2009, he moved to the Indian Institute of Technology Hyderabad as an assistant professor, where he is currently a professor of chemistry and the Dean of Academic Affairs. His research focuses on nanomaterials for solar energy harvesting and environmental remediation.

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Tsutomu Miyasaka

Tsutomu Miyasaka obtained his master's (1978) and PhD (1981) degrees from the University of Tokyo (UT). Also, he was a visiting researcher at the University of Quebec. Subsequently, he joined Fuji Photo Film Co., where he was involved in the development of high sensitivity photography films and high capacity battery materials. Since 2001, he has been a professor at the Graduate School of Engineering at Toin University of Yokohama (TUY), where he established Peccell Technologies Inc., serving as CEO. He is presently a project professor at TUY and a Fellow of the Research Centre for Advanced Science and Technology at UT. Since he demonstrated the first halide perovskite solar cell in 2009, he has focused his research on this subject. He was a recipient of the Clarivate Analytics Citation Award in 2017.

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Vasudevanpillai Biju

Vasudevanpillai Biju studied chemistry at CSIR-RRL and the University of Kerala and obtained his PhD in 2000. After postdoctoral research in Japan and the USA, he was a scientist in AIST, Japan. In parallel, he was a visiting scientist at the University of Texas at Austin and a JST PRESTO Researcher. Since February 2016, he has been a Professor at Hokkaido University. His research interests include semiconductor quantum dots, photo-functional molecules, and single-molecule techniques. His research activities are recognized with various awards by the Ozaki Foundation (Gennai Award), the Japanese Photochemistry Association, the Asian and Oceanian Photochemistry Association and the Chemical Society of Japan. Since 2011 he has been a Fellow of the Royal Society of Chemistry.



Key learning points

(1) Synthesis of lead and lead-free halide perovskites: in the form of films, microcrystals, nanocrystals, and quantum dots.

(2) Halide-dependent tunable bandgap and luminescence properties.

(3) Optoelectronic properties governed by excitons and charge carriers.

(4) Photoluminescence blinking in films, microcrystals, nanoparticles, and quantum dots.

(5) Perovskites for high efficiency solar cells, bright LEDs and low threshold lasers.


1. Introduction

Since the discovery of the parental perovskite calcium titanate (CaTiO3) by Gustav Rose in the Ural Mountains, it has taken more than 150 years to optimize the properties of analogous halide perovskites and realize their practical applications. Halide perovskites have become popular due to their cost-effective and simple synthesis, multicolor emission, high photoluminescence quantum yields (PLQYs), and excellent excitonic and charge carrier properties. These properties make halide perovskites suitable for engineering devices such as LEDs,1 lasers,2 photodetectors,3 and solar cells.4 The general formula of halide perovskites is ABX3 (Fig. 1a), where the A-site cation such as Cs+, methylammonium (CH3NH3+, MA) or formamidinium [(CH(NH)2)2+, FA] occupies the voids formed by the corner-sharing BX6 octahedra. Isomorphs of halide perovskites are obtained by varying the A- or B-site cation (Pb2+, Sn2+, Ge2+, Bi3+, In3+ or Sb3+) as well as the halide ion (X = Cl, Br, I). Also, perovskites with mixed A-site and/or B-site cations, such as MAFAPbX3, KCsMAFAPbX3, Cs2AgBiX6, Cs2AgInCl6, and Cs2AgInBiX6, have been developed.
image file: c9cs00848a-f1.tif
Fig. 1 Morphologies and PL properties of halide perovskites. (a) The general structure of ABX3 perovskites. (b) PL images and (c) PL spectra of colloidal solutions of FAPbX3 perovskite nanocrystals. (d) A SEM image of a MAPbI3−xClx perovskite film. (e) A photograph of a FAPbBr3 single crystal. (f and g) TEM images of (f) CsPbBr3 and (g) CsSnI3 nanocrystals. Reproduced with permission from (b and c) ref. 6 (copyright 2017, American Chemical Society), (d) ref. 5 (copyright 2013, American Association for the Advancement of Science), (e) ref. 7 (copyright 2015, The Royal Society of Chemistry), (f) ref. 8 (copyright 2019, American Chemical Society), and (g) ref. 9 (copyright 2016, American Chemical Society).

While the chemical composition dictates the crystal structure and stability of perovskites, the most attractive properties of these materials are halide- and size-dependent bandgap tuning and carrier confinement. The PL color of these perovskites is tuned (Fig. 1b and c) in the visible to near-infrared (NIR) region, which is achieved through the preparation of mixed halide perovskites from mixed halide precursors or post-synthesis halide exchange reactions.5,6 In general, chloride perovskites show violet to blue emission, bromide samples show green emission and iodide samples show deep-red to NIR emission. The PL color can be tuned between violet and green by increasing the proportion of bromide in pure chloride perovskites or chloride in pure bromide perovskites. Similarly, the PL color can be tuned between green and NIR by mixing bromide and iodide ions. This halide-dependent tuning of the emission color or bandgap originates from the modification of band-edge states formed by the hybridized s- and p- orbitals of halides and B-site cations. As the bandgap decreases from chloride to bromide and from bromide to iodide, the effective mass of excitons decreases. Thus, the exciton binding energy decreases from chloride to iodide, which modifies the excitonic and charge carrier lifetimes in perovskites. Other factors affecting the exciton binding energy are A-site cations and crystal size. Images of perovskites with different morphologies, sizes, and chemical compositions are shown in Fig. 1d–g.

In this review, we introduce newcomers to the synthesis, structures, and optical properties of halide perovskites with different sizes, structures, and compositions. Further, we correlate the bandgaps and excitonic and charge carrier properties of these materials with perovskite solar cells, LEDs and lasers.

2. Synthesis

The factors to be considered during the synthesis of halide perovskites are their size, structure, stability, anticipated optical and charge carrier properties, potential device applications, and toxicity. The stability of the ABX3 perovskite structure is determined by the Goldschmidt tolerance factor:
 
image file: c9cs00848a-t1.tif(1)
and the octahedral factor μ = rb/rx, where ra, rb and rx, respectively, are the radii of the A-site cation, B-site cation and halide ion.10 For a stable perovskite structure, the tolerance factor should be between 0.81 and 1.00, and the octahedral factor should be between 0.44 and 0.9, which are determined by the cation and anion sizes. Based on the above stability factors, MA-, FA-, and Cs-based halide perovskites are the most stable structures synthesized in laboratories. While lead perovskites show excellent optical and electronic properties, the environmental costs of lead, particularly its toxicity to aquatic and terrestrial life, have attracted researchers to replace lead by other metal ions such as Sn2+, Ge2+, Cu+, Ag+, Bi3+, Sb3+, and In3+. Among these ions, Sn2+ endows perovskites with properties like Pb2+. However, Sn2+ gets oxidized to Sn4+, and as a result, the geometry of the perovskite distorts and the optical and electronic properties change. Thus, perovskites based on other B-site cations such as Sb3+ are prepared and studied.11 These cations form an A3B2X9 configuration due to the distortion of the BX6 network caused by their small size. Yet another class of lead-free halide perovskites is A2B′B′′X6 double perovskites.12 To fulfil their valency, double perovskites require trivalent and monovalent cations chosen from among Cu+, Ag+, Bi3+, Sb3+, and In3+. While the A- and B-site cations play significant roles in the stability of the perovskites, their optical and charge carrier properties are controlled by halide ions and perovskite crystal size. Fig. 2 summarizes the methods for the synthesis of halide perovskites in the form of thin films, microcrystals, nanocrystals and quantum dots (QDs) with different A- and B-site cations and halide ions.

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Fig. 2 Preparation of perovskite thin films, microcrystals, nanocrystals, and QDs.

2.1. Films

High quality perovskite films are light absorbing layers in perovskite solar cells (PSCs). Such films are fabricated in single4,13–15 step or two steps16 (Fig. 3a) from solutions of their precursor salts (MAX, FAX or CsX, and PbX2) dissolved in solvents such as dimethylformamide (DMF), dimethyl sulfoxide (DMSO), and γ-butyrolactone (GBL). These solvents enable slow crystal nucleation and growth of perovskites, which is due to their low vapor pressures and high boiling temperatures. For instance, a MAPbX3 perovskite film is prepared by simply depositing a precursor solution of MAX and PbX2 in DMF or GBL on a substrate. Antisolvent-assisted deposition, vacuum- and gas-pumping, and hot-casting are other methods for the preparation of uniform perovskite layers in PSCs. Spiccia and co-workers15 used antisolvent-induced fast crystallization of MAPbI3 to form a micro-crystalline film free from grain boundaries. Here, a precursor solution of MAPbI3 in DMF is spin-coated on TiO2, which is followed by exposing the film to the antisolvent dichlorobenzene. Similarly, polycrystalline thin films of MASnX3 perovskites are prepared by spin-coating stoichiometric amounts of MAI and SnI2 dissolved in degassed DMF on TiO2 or Al2O3.13
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Fig. 3 General methods for the synthesis of (a) films and (b and c) crystals of perovskites. (a) Sequential deposition of precursor solution, (b) Inverse temperature crystallization, and (c) antisolvent vapor-assisted crystallization.

In two-step or sequential deposition of perovskite layers (Fig. 3a), heterogenous reaction takes place between the solutions of PbX2 and the organic or inorganic halide salt. For example, a film of mesoporous TiO2 is deposited on a compact thin layer of TiO2, which is followed by the spin-coating of PbI2 dissolved in DMF. After drying the PbI2 solution, MAI is spin-coated and annealed at 100 °C for 20 s to provide a thin film of MAPbI3 cuboids.16

Compositional engineering of perovskite crystals using mixed cations and halide ions in one-step coating provides uniform pinhole-free perovskite films for high efficiency PSCs. Uniform perovskite layers are prepared by spray-coating, soft-cover deposition, brush painting, screen-printing, inkjet printing, blade-coating, and slot-die coating.17 Among these methods, blade-coating is widely used in the fabrication of TiO2 mesoporous thin films for dye-sensitized and QD-sensitized solar cells, and it has gained popularity in the deposition of large area PSCs. In this method, a perovskite precursor solution is swiped on a pre-heated substrate and kept for solvent evaporation. However, the preparation of uniform films over a large coverage area is often challenged by uneven edges of the blade. The speed of blade swiping is another important factor to be considered in the preparation of uniform films, as the speed directs the fluid dynamics and the rate of solvent evaporation. A typical example of blade-coating is the preparation of a MAPbI3 perovskite film on a layer of poly(bis(4-phenyl)(2,4,6-trimethylphenyl)amine) (PTAA).18 First, PTAA is blade-coated on an indium tin oxide (ITO)/glass substrate, which is followed by annealing this layer and blade-coating (50 mm s−1) of a MAPbI3 perovskite solution supplemented with the surfactant L-α-phosphatidylcholine. The use of perovskite films in solar cells is reviewed in Section 6.1.

2.2. Large single crystals and microcrystals

Various processes in the synthesis of single crystals of halide perovskites include inverse temperature crystallization (ITC, Fig. 3b), antisolvent vapor-assisted crystallization (AVC, Fig. 3c), and vapour-phase melt approaches. In ITC, the retrograde solubility of halide perovskites helps in the synthesis of millimetre size crystals. For example, single crystals of FAPbI3 are prepared by controlling the temperature of its precursor (PbI2 and FAI) solution in GBL, which is in the range of 80 to 120 °C. Similarly, large single crystals of FAPbBr3 can be grown by increasing the temperature of a precursor solution composed of PbBr2 and FABr in a mixture (1[thin space (1/6-em)]:[thin space (1/6-em)]1 v/v) of DMF and GBL.7 Here, GBL helps in the crystal growth by lowering the solubility of the precursors. Common antisolvents used in the preparation of perovskite microcrystals are chlorobenzene, ethyl acetate, and dichloromethane. MAPbX3 microcrystals are prepared by the slow diffusion of dichloromethane into a precursor solution containing stoichiometric amounts of MAX and PbX2 in DMF.19 In the vapor-phase melt approach, powders of CsX and PbX2 are kept inside a quartz tube in a furnace where the temperature is slowly increased to 600 °C and kept at this temperature for 15 min to grow the crystals.20 Recently, lead-free double perovskite crystals are prepared by different methods such as hydrothermal crystallization and precipitation. For example, polycrystalline Cs2AgBiX6 double perovskites are prepared by the addition of AgX and BiX3 into a solution of HX and H3PO2 at 120 °C.12

2.3. Nanocrystals and quantum dots

The growth of perovskite crystals can be confined to the nanometer scale by surface capping using long-chain organic ligands which are aliphatic carboxylic acids and amines. Perovskite nanoparticles (or nanocrystals) are routinely prepared by the hot-injection or ligand-assisted reprecipitation (LARP) method shown in Fig. 4. The size and morphology of nanocrystals, such as nanocubes, nanowires, nanoplatelets and QDs, are determined by the chain length of the ligands and the reaction temperature. In typical synthesis of CsPbBr3 perovskite nanocrystals by the hot-injection method, caesium acetate and oleic acid are dissolved in hexadecene at 120 °C, and in parallel, a mixture of PbBr2 and capping ligands such as oleic acid and hexadecyl amine is dissolved in hexadecene at 120 °C. Then, the hot solution of caesium acetate is injected into the PbBr2 solution at 170 °C. Immediately after the injection, the reaction solution is cooled in an ice bath, and nanocrystals are collected by centrifugation.8 Similarly, FAPbBr3 nanocrystals are synthesized at 130 °C by injecting a solution of octadecylammonium bromide in toluene into a solution of lead acetate, formamidinium acetate and oleic acid in octadecene. The reaction is quenched after 10 s, and FAPbBr3 nanocrystals are collected by purification of the reaction mixture using toluene and acetonitrile.21 The LARP method is convenient for the preparation of perovskite nanocrystals at room temperature. MAPbBr3 nanocrystals are prepared by the dropwise addition of a solution of MABr and PbBr2 in DMF into a mixture of octylamine and oleic acid dissolved in toluene and under vigorous stirring.22
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Fig. 4 (a and b) Synthesis of perovskite nanocrystals by the (a) hot-injection method, and (b) LARP method. (c) Ligand-capped nanocrystals of perovskites with different morphologies.

Nanocrystals of lead-free perovskites are synthesized by the hot-injection or LARP method like their lead-based counterparts. For example, CsSnX3 nanocrystals are prepared by suspending Cs2CO3 in a mixture of oleic acid and oleylamine in octadecene, which is followed by injecting SnX2 dissolved in tri-n-octylphosphine and increasing the temperature of the solution to 100–170 °C.9 Here, CsSnX3 nanocrystals precipitate out upon cooling the reaction mixture in an ice-bath. The LARP method is useful for the synthesis of lead-free A3M2X9 perovskite nanocrystals at room temperature. For example, Cs3Sb2Br9 QDs are prepared by mixing CsBr and SbBr3 together in DMF (or DMSO) supplemented with octadecyl amine or n-octylamine, and quickly adding a mixture of octane and oleic acid. Cs3Sb2Br9 QDs are collected from the reaction solution by the addition of acetone or octane and centrifugation. Cs3Sb2Cl9 and Cs3Sb2I9 QDs are prepared by halide exchange in Cs3Sb2Br9 using CsX (X = Cl, I) dissolved in DMSO.11 More recently, 2–6 nm size Cs2AgInxBi1−xCl6 double perovskites are synthesized by antisolvent-assisted recrystallization in a mixture of oleic acid and isopropanol.23 Here, precursor solutions of cations are prepared by dissolving CsCl, AgCl, BiCl3 and InCl3 in DMSO, and the composition of B-site cations is varied by using different molar ratios of BiCl3 to InCl3.

3. Crystal structure

The composition, crystal phase and unit cell parameters of perovskites are characterized by single crystal and powder X-ray diffraction (XRD) analysis. Halide perovskites show different crystal structures and phases such as cubic, orthorhombic, trigonal, monoclinic, pseudocubic, hexagonal and tetragonal (Table 1). XRD patterns of MAPbBr3, CsPbX3 and Sn-based perovskites are shown in Fig. 5. Different crystal phases are assigned according to the Miller indices. The prominent Miller indices of the cubic phase are (001), (011), (002) and (021), which are seen at 2θ ca. 15°, 22°, 30° and 34°, respectively. The difference between the cubic and orthorhombic phases is assigned based on two additional peaks ca. 30°. Organic–inorganic hybrid perovskites form a stable cubic phase at high temperatures, which is due to the reduced degree of rotational movement of the A-site organic cations. When the temperature is decreased, the phase transforms into tetragonal and orthorhombic phases. The Miller indices can be calculated from Bragg's law, = 2dhkl[thin space (1/6-em)]sin[thin space (1/6-em)]θ, where dhkl is the interplanar spacing and λ is the XRD wavelength. The interplanar spacing depends on the lattice constants which are different for different unit cells. Different structural forms of perovskites, such as large single crystals, thin films, nanocrystals, nanocubes, nanoplatelets, nanowires, QDs, quantum cubes, and nanosheets, can be characterized using XRD.
Table 1 Phases and groups of halide perovskites (RT: room temperature)6,9,11,12,19,23–26
Perovskite Halide Phase Group Temperature Ref.
MAPbX3 Br, I Cubic Pm3m RT 80 °C for Br 19
25 and 26
Tetragonal I4/mcm or I4/m 110 °C for I 19 and 25
I Pseudocubic P4mm RT 25
FAPbX3 Cl, Br, I Cubic RT 6
CsPbX3 Cl, Br, I Cubic Pm3m 120–200 °C 24
I Orthorhombic Pnma 315 °C 24
CsSnX3 Br, I Orthorhombic Pnam 170 °C 9
Cl Cubic Pm3m 170 °C 9
Cs3Sb2X9 Br Trigonal P[3 with combining macron]m1 90 °C 11
Cs2AgBiX6 Cl, Br Cubic Fm[3 with combining macron]m 150–210 °C 12
Cs2AgInxBi1−xX6 Cl Cubic Fm[3 with combining macron]m RT 23



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Fig. 5 XRD patterns of (a) MAPbBr3 single crystal, and (b) CsPbX3 and (c) CsSnX3 nanocrystals. Reproduced with permission from (a) ref. 26 (copyright 2015, Nature Publishing Group), (b) ref. 24 (copyright 2015, American Chemical Society), and (c) ref. 9 (copyright 2016, American Chemical Society).

XRD data help one to analyse the crystal phase and chemical composition of perovskites. When compared with single crystals (Fig. 5a), the diffraction peaks of the nanocrystals are slightly shifted to small 2θ angles (Fig. 5b). Further, in single and mixed halide compositions, the diffraction peak shifts from large to small 2θ angles while going from chloride to bromide and bromide to iodide.24 The lattice contraction becomes larger down the halide group, and the corresponding changes can be assigned to the shifting, appearance (ca. 35°, 27°, and 48°) and disappearance (ca. 52°) of some of the diffraction peaks (Fig. 5b and c).9,24 Mixed halide perovskites show slight shift in the diffraction peaks, when compared to single halide perovskites. For example, the diffraction angle increases upon increasing the proportion of chloride in a chloride–bromide mixed perovskite.

4. Energy states

The appealing optical and charge carrier properties of perovskites emerge from the unique distribution of energy states at the band-edge. As halide perovskites are direct bandgap semiconductors (Fig. 6), with the exception of noncentrosymmetric distribution of the conduction band minimum (CBM) and the valence band maximum (VBM) induced by the Rashba effect and Dresselhaus splitting, the band-edge of halide perovskites is composed of antibonding atomic orbitals of [BX6]4− octahedra.10,27 Here, the CBM is constituted by sigma (σ*) and pie (π*) antibonding orbitals, which are formed by the hybridization of 6p-orbitals of lead and s- and p-orbitals of halides, respectively (Fig. 6a). On the other hand, 6s- and 6p-orbitals of lead and s- and p-orbitals of halides contribute to the VBM.
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Fig. 6 (a) Band-edge states of halide perovskites without considering spin–orbit coupling (SOC). (b) Band structure and partial density of states of CsSnBr3 perovskite. (b) is reproduced with permission from ref. 27 (copyright 2013, American Physical Society).

Lead-free perovskites show similar band-edge structures to lead perovskites (Fig. 6b). For instance, in cubic CsSnBr3 perovskite, the VBM antibonding state is a mixture of 4p-orbitals of Br and 5s-orbitals of Sn, whereas the CBM is mainly constituted by 5p-orbitals of Sn.27 Here, the VBM is nondegenerate, while the CBM is three-fold degenerate without any contribution from SOC. When SOC becomes significant, the degeneracy of the CBM is broken, and it splits into spin–orbit split-off doublet states and quadruplet states. Optically allowed transitions from the VBM to the spin–orbit split-off states and quadruplet states impart fine optical absorption features to halide perovskites. Also, an increase in the dispersion of the conduction band through SOC-assisted splitting lowers the optical bandgap and affects their charge carrier and PL properties. In fact, SOC which tunes the optical properties of perovskites is associated with other factors such as the composition, octahedral tilt, and temperature.

5. Optoelectronic properties

5.1. Charge carrier dynamics

The generation, dynamics and recombination of charge carriers and excitons in perovskites largely depend on their composition, crystal structure, and size (Fig. 7). Photogenerated charge carriers in perovskites are either freely diffusing or excitonically bound, depending on the exciton binding energy. The dynamics of free carriers is characterized by the diffusion coefficient (Dn,p) and carrier mobility (μn,p), which are related through the Einstein equation:
 
μn,p = eDn,p/kBT(2)
where e is the electronic charge, kB is the Boltzmann constant and T is the temperature. These parameters are related to the electronic structure through the effective mass (meff) of the carriers. The comparable meff values of electrons and holes in halide perovskites provide this class of materials with balanced electron and hole diffusion lengths. Although different theoretical and experimental methods provide slightly different meff values to electrons and holes, the main factors that affect the masses are SOC, chemical composition, and the crystal size. The diffusion length (LD) of charge carriers is related to their mobility or diffusion coefficient by the equation:29
 
LD = (kBn,pτ/e)1/2 = (Dn,pτ)1/2(3)
where τ is the PL lifetime. The high mobility and long lifetime of charge carriers result in long LD, which is an important requirement for perovskites to be an efficient charge-transport layer in photovoltaics. The long PL lifetimes of single crystal and thin film perovskites show the presence of free carriers formed by the dissociation of photogenerated excitons, when the thermal energy (kBT) overcomes the exciton binding energy.

image file: c9cs00848a-f7.tif
Fig. 7 Charge carrier properties of perovskites. (a) Transient photovoltaic data of MAPbI3 perovskite showing long carrier lifetimes. (b) Diffusion length with respect to charge carrier density. (c) PL decay profiles of MAPbBr3 perovskite with the increasing intensity of excitation light. Reproduced with permission from (a) ref. 29 (copyright 2015, American Association for the Advancement of Science), (b) ref. 31 (copyright 2014, Wiley-VCH Verlag GmbH & Co. KGaA, Weinheim), and (c) ref. 8 (copyright 2019, American Chemical Society).

Charge carrier recombination dynamics in perovskites is explained using the following rate equation:

 
image file: c9cs00848a-t2.tif(4)
where k1, k2, and k3 are respectively the rate constants of monomolecular, bimolecular and Auger recombinations, and n is the total carrier concentration. Monomolecular recombination is a geminate type excitonic recombination or a trap-assisted recombination. On the other hand, bimolecular recombination is a non-geminate recombination of free charge carriers. In Auger recombination, the charge carriers transfer the energy or momentum to a third charge carrier. These recombination mechanisms in perovskites contribute to the total recombination rate and influence the carrier lifetime and diffusion length. The monomolecular recombination in perovskites is mainly governed by the trap-assisted processes, and it depends on the energy, density, and distribution of trap states. Furthermore, the monomolecular recombination rate is influenced by material purity and crystallinity. The trap-assisted recombination rate is in the order of 106 s−1 for single and mixed halide lead perovskites.30,31 On the other hand, for lead-free MASnI3 perovskites, the monomolecular recombination rate (∼109 s−1) is three orders of magnitude greater than that of MAPbI3, which is attributed to the self-doping by Sn4+ ions.13 The activation energy and depth of traps depend on the crystalline phase of perovskites, and thus affect the monomolecular recombination rate. For example, the trap-assisted recombination rate is higher in cubic and tetragonal phases than in an orthorhombic phase, which is due to the difference in the activation energies of trap formation in these phases.30

Bimolecular recombination in perovskites depends on the composition of halide and metal ions. The ratio of the bimolecular recombination rate constant to charge carrier mobility is lower than that predicted by the Langevin theory, according to which bimolecular recombination occurs when electrons and holes move within their coulombically-bound radii. Hence, the deviation from the Langevin theory in perovskites refers to a low bimolecular recombination rate with high carrier mobility. The non-Langevin bimolecular recombination behaviour in perovskites is attributed to the occupation of different molecular orbitals in a perovskite unit cell by an electron and a hole, which are formed by the overlapping of energetically different s- and p-orbitals of halide and metal ions (Pb, Sn). The bimolecular recombination rate largely varies in perovskite single crystals and films. For example, perovskites composed of lighter halides (Br instead of I) and metals (Sn instead of Pb) show larger bimolecular recombination rates. Such rate variations with chemical composition are associated with the modification of band structure and exciton binding energy.

While trap-assisted recombination and bimolecular recombination in perovskites are dominant phenomena under intermediate charge carrier density and one sun illumination, the contribution of Auger-assisted nonradiative recombination processes becomes significantly large at high carrier concentrations. The rate of Auger-assisted recombination is high in the orthorhombic phase at low temperature, whereas it remains low in high temperature tetragonal and cubic phases.30 This phase-specific variation of the rate is due to the varying degrees of SOC and hydrogen bonding. The Auger recombination rate is also influenced by the presence of impurities and phonons. The above carrier recombination dynamics affect the PL lifetime of perovskites. The PL lifetime of nanocrystals varies from picoseconds to several tens of nanoseconds, which largely depends on quantum confinement, whereas it scales up from nanoseconds to microseconds in bulk single crystals. For example, MAPbI3 single crystals show carrier lifetimes as long as 234 μs (Fig. 7a).29 Perovskite thin films also show long carrier lifetimes and diffusion lengths (Fig. 7b and c).5,8,31

5.2. Absorption and photoluminescence

The absorption and photoluminescence of perovskites are tuned in the UV, visible and NIR regions by varying the halide from Cl to I (Fig. 8 and 9) and the external parameters such as the temperature and pressure (Fig. 10). Since the VBM of perovskites is formed by the mixing of halide p-orbitals and metal s-orbitals, the halogen p-character in the VBM can be increased by changing the halide ion from Cl to Br and from Br to I, and as a result SOC increases.10,27 Therefore, by changing the halide ion from Cl to I, the bandgap decreases and the PL spectrum shifts to the red as shown in Fig. 8a.24 The octahedral tilt32 is another factor that modifies the PL properties of perovskites, which occurs due to the decrease of the Pb–X–Pb dihedral angle upon replacement of a large A-site cation (FA+) with a smaller one (Cs+). Bandgap variation with octahedral tilt is shown in Fig. 8b.32 Similarly, the change in crystal structure from cubic to tetragonal or other phases causes an octahedral tilt. Such a tilt in PbX6 octahedra weakens SOC by lowering the orbital overlap between Pb and X, which in turn increases the bandgap and shifts the PL band to the blue. Also, the cation size and hydrogen bonding of MA+ or FA+ ions with [PbX6]4− octahedra influence the electronic and optical properties of perovskites by altering the covalent or ionic character of the Pb–X bond.32 Tetragonal-to-pseudocubic structural change, which is achieved by changing the A-site cation from MA+ to FA+, increases the hydrogen bonding between FA+ and [PbX6]4−. This in turn increases the ionic character of the Pb–X bond and the contribution of the Pb character to the CBM. SOC becomes stronger in the stable pseudocubic structure than the octahedrally tilted tetragonal structure, resulting in a decrease of bandgap. The bandgap energies of different crystalline phases of Cs, MA and FA perovskites are given in Fig. 8c.32
image file: c9cs00848a-f8.tif
Fig. 8 The effects of anions and cations on the optical properties of perovskites. (a) Absorption and PL spectra of CsPbX3 nanoparticles with different halide compositions. (b) Calculated bandgap variations in CsPbI3 with octahedral tilt. (c) Calculated bandgap energies of cubic (c), tetragonal (t) and triclinic (tc) phases of APbI3 (A = Cs/MA/FA) perovskites. Reproduced with permission from (a) ref. 24 (copyright 2015, American Chemical Society), and (b) ref. 32 (copyright 2014, American Chemical Society).

image file: c9cs00848a-f9.tif
Fig. 9 The effects of size and shape on the optical properties of perovskites. (a) Size-dependent absorption and PL spectra of CsPbBr3 perovskites. (b) Experimental and theoretical bandgap values as functions of nanocrystal size. (c and d) Absorption and PL spectra of CsPbBr3 perovskite nanocrystals with different shapes. Reproduced with permission from (a) ref. 33 (copyright 2018, American Chemical Society), (b) ref. 34 (copyright 2016, American Chemical Society), (c) ref. 28 (copyright 2016, American Chemical Society), and (d) ref. 35 (copyright 2016, American Chemical Society).

image file: c9cs00848a-f10.tif
Fig. 10 Pressure-dependent PL of FAPbBr3 perovskites. Reproduced with permission from ref. 36 (copyright 2016, American Chemical Society).

Crystal size is another factor that controls the optoelectronic properties of perovskites. A decrease in size below the exciton Bohr radius results in quantum confinement effects in perovskites, evidenced by the appearance of sharp excitonic peaks, narrow PL bands, and blue-shifted absorption and PL spectra.6,28,33–35 As shown in Fig. 9a, the gradual blue-shift of the absorption onset with decrease in crystal size (from 6.2 to 3.7 nm), which is correlated with the blue-shift of emission peaks, is characteristic of quantum confinement effects in CsPbBr3 perovskites.33 The size and bandgap of MAPbX3 are correlated in Fig. 9b.34 Similarly, as shown in Fig. 9c and d, quantum confinement effects are observed in perovskite nanoplatelets and nanowires. Apart from the size, the type of halogen plays important roles in quantum confinement. For example, the exciton Bohr radius of CsPbX3 perovskites varies as 2.5 nm for CsPbCl3, 3.5 nm for CsPbBr3, and 6 nm for CsPbI3.24

Temperature and pressure also affect the absorption and PL properties of perovskites. The role of temperature is to induce phase transition and vary exciton–phonon interactions by thermal expansion.30 As the temperature of a perovskite crystal is increased, a series of structural changes or phase transitions occur. In lead iodide perovskites, variation of PL properties occurs in three distinct regimes: orthorhombic (low temperature), tetragonal (room temperature), and cubic (high temperature) phases.30 Such a phase transformation affects the PL properties by altering the SOC through octahedral tilt. With increase in temperature, a perovskite crystal tends to stabilize in the cubic phase which renders minimal octahedral tilt and strong SOC. Consequently, the bandgap decreases and the absorption and PL spectra red-shift. Further, the broadening of the absorption and PL spectra observed at higher temperatures is attributed to the enhanced exciton–phonon interactions. Like temperature, pressure modifies the structure as well as electronic and optical properties of perovskites.36 Here, the role of pressure is to alter the SOC through pressure-induced compression of the Pb–X bond, shrinkage of the Pb–X–Pb angle or an increase in octahedral tilt, and amorphization of the crystal lattice. As shown in Fig. 10, pressure-induced changes in the optical properties of FAPbBr3 perovskites are associated with structural changes [cubic (Pm[3 with combining macron]m) → cubic (Im[3 with combining macron]) → orthorhombic (Pnma)].36

6. Photoluminescence blinking

Single-molecule micro-spectroscopy, which is widely used in the field of cadmium and lead chalcogenide QDs, is helpful for gaining insight into the properties of perovskites at the single particle level. Single particle studies reveal the PL characteristics of perovskites in the form of thin films,37 microcrystals,38 nanocrystals,38–40 and QDs.41,42 Blinking is characterized by a sequence of stochastic bright (ON) and dark (OFF) events (Fig. 11–13),38 which are due to the repeated radiative and nonradiative recombinations of photogenerated excitons or charge carriers.
image file: c9cs00848a-f11.tif
Fig. 11 PL blinking in halide perovskites. (Left) PL intensity trajectories in units of brightness and (right) PL images of (a) a MAPbI3 nanocrystal and (b) a localized emitting site located on the top of a large MAPbI3 crystal. Reproduced with permission from ref. 38 (copyright 2015, American Chemical Society).

image file: c9cs00848a-f12.tif
Fig. 12 The origin of PL blinking in perovskites. (a) Quenching and emitting site models of blinking in nano- and micro-crystals. (b) The formation and active–passive state switching of a quenching complex during PL blinking in nanocrystals. (c) A PL intensity trajectory showing two-state blinking in CsPbI3 perovskite QDs. (d) Photoinduced charging–discharging model of two-state blinking in perovskite QDs. Reproduced with permission from (a) ref. 38 (copyright 2015, American Chemical Society), (b) ref. 39 (copyright 2019, Nature Publishing Group), and (c) ref. 41 (copyright 2015, American Chemical Society).

image file: c9cs00848a-f13.tif
Fig. 13 The effect of the local environment on the PL blinking of MAPbI3 single nanocrystals. (a) PL intensity trajectories (i) at an air/glass interface, (ii) in an argon atmosphere, (iii) in a poly(methyl methacrylate) (PMMA) film, and (iv) at an air/glass interface showing recovery of the PL intensity after a long OFF state. (b) A scheme of (upper part) superoxide generation by a MAPbI3 nanocrystal, (middle part) radiative and nonradiative processes associated with various states, and (lower part) superoxide-mediated disintegration of a MAPbI3 nanocrystal. Reproduced with permission from ref. 40 (copyright 2019, Wiley-VCH Verlag GmbH & Co. KGaA, Weinheim).

Halide perovskites show different PL blinking behaviours (Fig. 11 and 12). For instance, an early report on PL blinking in elongated rod-like nanocrystals and microcrystals of MAPbI3 perovskite shows multi-state PL fluctuations with large amplitudes (Fig. 11), which become pronounced at lower excitation intensities.38 The super-resolution localization microscopy technique reveals different intensity levels (or steps) in the multi-state blinking of a rod-like nanocrystal and correlates blinking with different emission localization positions in the crystal. On the other hand, single quantum emitters such as chalcogenide QDs show ON and OFF states. Such blinking in halide perovskites is attributed to the photoinduced activation and deactivation of emitting or quenching sites formed by geometrical distortions or the presence of different chemical structures. The quenching sites are analogous to charge traps or crystal defects. As shown in Fig. 12a, in the ON state, either the whole nanocrystal is emissive with the emission localization position at the centre of the crystal (quenching site model) or PL comes from different positions acting as emitting sites in the crystal (emitting site model). In the quenching site model [Fig. 12a(i)], a photogenerated quencher acts as the charge trap, which is formed at the ends of a rod-like perovskite nanocrystal. This quencher deactivates all the photogenerated electrons and holes. As a result, a decrease in PL intensity and a shift of the emission localization position, marking a ‘step’ with intermediate PL intensity or an OFF state, are observed. On the other hand, in the emitting site model [Fig. 12a(ii)], all the charge carriers in the crystal diffuse efficiently to the emitting sites and recombine radiatively (ON state). If one of the emitting sites is deactivated or quenched, the PL intensity decreases and the localization position shifts (‘steps’ with intermediate PL intensity or OFF state). Alternatively, PL blinking can be assigned to the formation of quenchers near to the emitting sites or by changing the emitting site to a quencher.38 The enhancement of PL intensity and the decrease of amplitude associated with an increase in excitation power density can be attributed to defect- or trap-filling. The high concentration of charge carriers generated at high excitation intensity leads to PL intensity averaging over many emitting sites.

Blinking by photoinduced activation–deactivation of charge traps is supported by temperature-dependent PL measurements of MAPbI3 nanocrystals, correlating blinking with the random switching model of quenchers or switching of charge traps between the active and passive states (Fig. 12b).39 The switching of quenchers is thermally activated, which is supported by suppressed blinking at low temperatures. This blinking suppression is due to the poor diffusion of quenching sites. Although the thermal barrier for the activation of quenching sites is in the order of activation energy for halide migration,39 quenching by the migration of halides or protons alone cannot model PL blinking. Thus, blinking is discussed in terms of a “quenching complex” which is a combination of a shallow mobile (or localized) electron trap present near to the conduction band and a localized (or mobile) hole trap near to the valence band (Fig. 12b). However, such individual shallow defects are unlikely to cause blinking.

PL blinking is observed in perovskite microcrystals and nanocrystals even at low excitation intensities, where the nonradiative Auger process is negligible. This behaviour contradicts with the origin of PL blinking of chalcogenide QDs at higher excitation intensities/energies, which involves photoinduced charging and neutralization. This difference between the blinking behaviours of perovskite nanocrystals and chalcogenide QDs is due to the weakly-confined charge carriers in large perovskite crystals. In contrast, like chalcogenide QDs, single perovskite QDs show two-state PL blinking (Fig. 12c) due to random charging by photoionization and subsequent neutralization.41 This type of blinking is called type-A blinking. At higher excitation intensities, blinking becomes more pronounced with prolonged OFF, which is attributed to the increase in the rates of photoionization and nonradiative Auger recombination processes. As shown in Fig. 12d, the charged excitonic state, called a trion (X*), represents an OFF state and the neutral (X) state represents an ON state. The significant decrease in PL lifetime in the OFF state (ca. 0.78 ns) of CsPbI3 perovskite QDs when compared to the ON state (ca. 13.2 ns) confirms type-A blinking.41 Perovskite QDs also show type-B blinking, which is characterized by the repeated activation and deactivation of traps.42

Apart from charge- and defect-induced blinking, the PL properties of perovskite single particles are affected by oxygen and moisture. Both photobleaching42 and photobrightening37,38 are observed under ambient conditions. For example, a decrease in the PL intensity of MAPbI3 nanocrystals is observed in air than in an argon or polymer environment (Fig. 13a).40 Such a decrease is attributed to the self-sensitized generation of superoxide and the subsequent oxidative disintegration of MAPbI3 nanocrystals (Fig. 13b). Interestingly, the PL intensity recovers after long OFF intervals [Fig. 13a(iv)], showing that ultrafast nonradiative Auger recombination processes in the OFF state prevent the electron transfer to oxygen, generation of superoxide, and oxidation of MAPbI3 nanocrystals. On the other hand, oxidation persists throughout the neutral or ON state, showing that the rate of superoxide generation exceeds radiative relaxation. Single particle studies help one to address the photooxidation and photostability of perovskites.

7. Applications

7.1. Solar cells

Perovskites are promising materials for solar cells, which is due to their high absorption coefficients, long diffusion lengths of charge carriers, defect tolerance, and low exciton binding energies. A PSC was first reported by Miyasaka and coworkers;4 its power conversion efficiency (PCE) was 3.8%. They replaced organic dyes in dye-sensitized solar cells with MAPbI3 perovskite on mesoporous TiO2 electrodes. After a PCE exceeding 10% was achieved using a solid-state MAPbI3 PSC, this research has attracted much attention of physicists and chemists to engineer cost-effective solar cells. For an efficient PSC, the quality of the perovskite layer and a defect-free interfacial structure are very important. Various techniques of device fabrication and optimization of hole and electron transfer processes helped the rapid advancement of PSCs. For example, Spiccia and co-workers15 achieved PCEs as high as 13.9% by using micro-crystalline MAPbI3 perovskite films free from grain boundaries. Subsequently, the PCE was increased to >15% with the use of mixed MAPbI3 and MAPbI3−xClx perovskites,16 >20% with the use of MAPbI3 perovskite and the blade-coating method,18 and >23%14 with the use of mixed cations (MA/FA) and mixed halides (Br/I). More recently, the efficiency of PSCs has increased up to 25.2% (NREL efficiency chart), which, along with the cost-effective device fabrication, has prompted several researchers to commercialize these devices. A comprehensive review on the background of PSCs and the recent progress in PSC fabrication is available elsewhere.17

A typical PSC architecture, as shown in Fig. 14a, has a perovskite layer as the intrinsic light absorbing semiconductor sandwiched in between an electron-transport material (ETM) and a hole-transport material (HTM) on a fluorine-doped tin oxide (FTO) or ITO glass substrate. TiO2 or [6,6]-phenyl-C61-butyric acid methyl ester (PCBM) is a common n-type ETM and 2,2′,7,7′-tetrakis(N,N-dimethoxyphenylamine)-9,9′-spirobifluorene (spiro-OMTAD) or poly(3,4-ethylenedioxythiophene)–polystyrene sulfonate (PEDOT:PSS) is a common p-type HTM. Following the absorption of light by the perovskite layer, photogenerated electrons are transferred to the conduction band of the ETM and subsequently to FTO/ITO. Holes are transferred to the valence band of the HTM and then to the Au/Ag layer. One can design efficient solar cells by selecting perovskites, HTMs and ETMs by referring to the energy states and bandgaps in Fig. 14b. The ETM and HTM in Fig. 14b are selected to facilitate efficient electron and hole-transport from photoexcited perovskites.


image file: c9cs00848a-f14.tif
Fig. 14 Characteristics of perovskite solar cells. (a) Device architecture of a PSC. (b) Energy levels of various electron-transport, hole-transport, perovskite and electrode materials. (c) Current–voltage curves of lead- and tin-based PSCs. (c) is reproduced with permission from ref. 13 (copyright 2014, The Royal Society of Chemistry).

Due to the low exciton binding energy and the ambipolar charge carrier mobility of perovskites, a p–i–n or a n–i–p junction is formed in a PSC. Thus, charge carriers can be extracted even without any HTM or ETM. However, a HTM and an ETM are needed for high efficiency PSCs because the open circuit voltage (VOC), current density (JSC) and fill factor in a solar cell are determined by these materials. As a result, recently, spiro-OMTAD has been replaced by a fluorine-terminated HTM, where the durability and PCE of a (FAPbI3)0.95(MAPbBr3)0.05 PSC are 500 h and >23%, respectively.14 Highly stable PSCs with high PCEs can be fabricated using perovskites with triple/quadruple cations. Also, lead-free PSCs are promising in terms of their performance and efficiency. Although Sn2+ is the most promising substitute of Pb2+, the instability of Sn2+ against oxidation is a challenging issue. Snaith and co-workers reported >6% PCE and >0.88 VOC in a MASnI3-based solar cell (Fig. 14c).13 Recently, Hayase and co-workers stabilized a Sn-based PSC by doping with Ge and achieved a sustained PCE of 6.9%.43 To ensure the stability of PSCs against heat, humidity, and light, studies have focused on the improvement of interfacial structures at grain boundaries and layer junctions.14,17,18,43 The device stability is also highly affected by the quality of HTMs which generally contain diffusible dopants. An important direction of PSCs is to use all inorganic compositions of perovskites as absorbers having high thermal stability.

In the photovoltaic community, perovskites have evolved as promising materials for high efficiency solar cells despite the challenges in their stability. The highly ionic nature of the perovskite lattice results in the degradation of the lattice and leakage of Pb2+ cations in the presence of moisture. Further, the photo-ionic conductivity of halides and metal ions limits the performance of PSCs. Nevertheless, there is enough room to address all these challenges and improve the performance of PSCs to their optimal level.

7.2. LEDs

The tuneable PL and high PLQY of perovskites have attracted people to fabricate multicolor perovskite LEDs (PeLEDs).1,44–46 The low defect density and ambipolar carrier migration in these materials enable efficient injection and transport of charges across the interfaces, ameliorating the device performance. Although the attempts to use halide perovskites in LEDs date back to the early 90s, a room temperature multicolor PeLED was realized only very recently. Recently, the (peak) external quantum efficiencies (EQEs) of green (20.3%)45 and red (21.6%)1 PeLEDs have been considerably improved. Pure chloride and mixed chloride–bromide perovskites are employed in blue PeLEDs. For example, by using CsPb(Br/Cl)3 NCs, Kovalenko and coworkers developed a blue PeLED with 1.2% EQE at 4.4 mA cm−2.44 The structure and electroluminescence (EL) spectra of the blue LED are shown in Fig. 15a and b, respectively. Although the EL wavelength of this LED remained constant at 463 nm (Fig. 15c), the EL intensity monotonically decreased with time. Thus, research on blue PeLEDs is open for improving their EQE and stability.
image file: c9cs00848a-f15.tif
Fig. 15 EL characteristics of a blue PeLED. (a) Device structure of the PeLED, (b) EL spectra of the PeLED at various applied voltages, and (c) EL spectra recorded at 5 V and different times. Reproduced with permission from ref. 44 (copyright 2019, American Chemical Society).

A typical device structure of a PeLED consists of a light-emitting layer of halide perovskites placed in between the charge (electron and hole)-injection (transfer) layers and electrodes. Halide perovskites in their polycrystalline thin film or nanostructure forms (nanocrystals, platelets or QDs) form double heterojunctions with the charge-injection layers. As shown in Fig. 14b and discussed in the ‘Solar cells’ section, the charge-injection layers are large bandgap semiconductors which enable efficient and balanced injection and confinement of charge carriers. PeLED devices are fabricated with an ITO/PEDOT:PSS/halide perovskite/B3PYMPM/LiF/Al configuration as shown in Fig. 15.44 Here, PEDOT:PSS acts as the hole-transport layer, B3PYMPM is the electron-transport layer, LiF is the electron-injection layer, and ITO and Al are the anode and cathode, respectively. A record-breaking EQE (20.3%)45 can be achieved with such a configuration by incorporating a thin layer of insulating PMMA in between the perovskite layer and electron-transport layer.

The intrinsic stability of halide perovskites and the efficiency of PeLEDs are influenced by the crystal structure, temperature, atmosphere, and light. Other factors detrimental to the performance of PeLEDs are unbalanced charge-injection and transport across the perovskite/electron-transport layer or perovskite/hole-transport layer interfaces, the current leakage due to pinholes in the perovskite layer, and nonradiative losses in the perovskite layer. While the balanced charge- injection and transport are achieved by the proper selection of electron- and hole-transport layers and modification of perovskite/charge-injection interfaces, the current leakage can be lifted-up by increasing the surface coverage and decreasing the surface roughness of the perovskite layer. The suppression of nonradiative recombination under an electrical bias is another important factor in achieving high-performance PeLEDs. For example, a PeLED device with record-breaking 20.3% EQE and 80% PLQY is fabricated using a compositionally graded quasi-core/shell CsPbBr3/MABr perovskite layer.45 Here, MABr passivates the nonradiative defect sites, increases the PL lifetime and the surface coverage, and balances charge-injection. In organic–inorganic hybrid perovskites, an interesting strategy of defect passivation is the use of molecules with high affinity to defect sites. For example, a high EQE of 21.6% is achieved for a FAPbI3-based NIR PeLED by passivating defects using 2,2′-[oxybis(ethylenoxy)]diethylamine.1

Besides the passivation of defects, one of the successful approaches to enhance EQE is to use low-dimensional perovskite structures such as nanocrystals, platelets, quasi-2D/3D nanostructures, and QDs. The EQE of PeLEDs constructed using bulk perovskite layers or films is limited by the thermal dissociation of excitons, which suppresses monomolecular or geminate-type excitonic recombination. Further, the bimolecular recombination kinetics in halide perovskites follow a quadratic relation with the charge carrier density, i.e. n2. Therefore, the radiative recombination in such perovskite layers becomes dominant only at higher charge densities. Nanostructure perovskites lift these problems by enhancing the monomolecular radiative recombination through carrier confinement. For example, a red-emitting PeLED with a high EQE (21.3%) is constructed by using anion-exchanged perovskite QDs.46 Despite these approaches to improve the performance of PeLEDs, these devices are still behind commercial LEDs due to the poor stability of perovskites in diverse environments.

7.3. Lasers

The high PLQY, large absorption cross-section and low defect density of perovskites make these materials promising for lasing. After the successful demonstration of low threshold amplified spontaneous emission (ASE) in solution-processed CsPbX3 (X = Cl, Br, and I) nanocrystals,2 investigations of lasing and ASE have been extended to various geometrical shapes of MA-, FA- and Cs-based perovskites.2,20,47–50 With the increase in the intensity of excitation light, the emission spectrum changes from broad, spontaneous emission (SE) to a narrow, red-shifted peak with higher intensity, which is the signature of ASE (Fig. 16a).48 Furthermore, single-mode and multimode lasing with an ultranarrow emission band and a high quality (Q) factor are realized in perovskites (Fig. 16b).20,49 The lasing wavelength in these materials can be tuned over the entire visible to NIR region by controlling the halide composition (Fig. 16c).49 Lasing in all-inorganic, organic–inorganic, and mixed halide perovskites has been demonstrated by the arrangement of these materials to form different cavities (Fig. 17), such as random cavities,2,47 vertical microcavities,50 Whispering Gallery Mode (WGM) cavities,2,49 and Fabry–Pérot cavities.20
image file: c9cs00848a-f16.tif
Fig. 16 ASE and lasing in perovskites. (a) ASE in a FAPbI3 film excited with a 150 fs laser. Inset: Integrated emission intensity and full width at half maximum (FWHM) as functions of pump energy, showing an ASE threshold of 3 μJ cm−2. (b) Triangular rod emission spectra around the lasing threshold. (c) Multicolor, single-mode lasing and the corresponding emission images of CsPbX3 microspheres. Reproduced with permission from (a) ref. 48 (copyright 2017, American Chemical Society), (b) ref. 20 (copyright 2016, American Chemical Society), and (c) ref. 49 (copyright 2017, American Chemical Society).

image file: c9cs00848a-f17.tif
Fig. 17 Lasing cavities in perovskites: (a) random cavity, (b) vertical cavity, (c) WGM cavity, and (d) Fabry–Pérot cavity.

Perovskite thin films exhibit random lasing. The ASE threshold can be lowered, and the optical gain can be improved by improving the quality of the film. For example, the ASE threshold is lowered from 3 μJ cm−2 to 1.6 μJ cm−2 by improving the quality of a FAPbI3 film by treating it with appropriate stoichiometric amounts of MABr and FAI.48 The propagating SE in the film is amplified by stimulated emission when the intensity of the excitation light exceeds a certain threshold, resulting in population inversion. During the propagation of emission in the film, the micro/nanodomains present in it produce coherent backscattering and generate closed-loop cavities (Fig. 17a). When the optical gain exceeds the loss, random lasing occurs at the resonant frequency of the corresponding feedback loop. In a MAPbI3 film, random lasing is observed at a threshold of 102 μJ cm−2 with a high optical gain and strong multiple backscattering provided by the randomly arranged polycrystalline grain boundaries.47

Cavity lasing is achieved by constructing vertical (Fig. 17b), WGM (Fig. 17c) or Fabry–Pérot (Fig. 17d) cavities. A vertical cavity is prepared by sandwiching a perovskite film in between two gallium nitride distributed Bragg's reflector mirrors. The resonators suppress most of the optical states and allow emission from a few states, resulting in mode-controlled lasing. Recently, an ultralow lasing threshold of 0.39 μJ cm−2 was reported for a CsPbBr3 QD vertical cavity.50 Similarly, WGM lasing is attained either in perovskite microspheres49 or by coating perovskite nanocrystals on substrates such as silica microspheres.2 Here, the substrates form circular cavities in which the total internal reflection occurs around the circumference [Fig. 17c(i)]. Ultralow threshold (∼0.42 μJ cm−2) single-mode WGM lasing occurs in a CsPbX3 microsphere cavity with a narrow line width (∼0.09 nm) and a high Q factor (∼6100).49 WGM lasing is observed in various 2D structures as well when the emission is confined within their boundaries [Fig. 17c(ii)]. Fabry–Pérot lasing is observed in perovskite nanocuboids, micro–nanorods, and wires. Such structures act as optical waveguides along the axial or planar direction and create a Fabry–Pérot cavity by confining light via total internal reflection at the two end facets. Multi-mode Fabry–Pérot lasing at room temperature, with low lasing thresholds (∼14.1 μJ cm−2) and high Q-factors (∼3500), is achieved in CsPbX3 rods.20 The lasing threshold in Fabry–Pérot cavities varies largely, depending on the shape, size, end facets, crystalline quality, and halide composition.

Despite the tunable lasing color and low lasing threshold of perovskite lasers, the exact mechanism of optical gain is yet to be rationalized. In organic–inorganic perovskites, the majority of photogenerated charge carriers are free, and the population of bound excitons is negligibly small due to the small exciton binding energy. The concentration of free charge carriers that fill certain states increases as the intensity of excitation light increases, providing a necessary condition for optical gain by population inversion. On the other hand, in all-inorganic perovskites, the optical gain and ASE are assigned to the radiative recombination of biexcitons.

8. Summary and outlook

With straightforward synthesis, halide perovskites have emerged as one of the most promising classes of semiconductor materials for solar cells, lasers, LEDs, and photodetectors. The fascinating optical and electronic properties of these perovskites originate from weakly-bound excitons activated in the band-edge states which are formed by the hybridization of halide orbitals with B-site metal ion orbitals. Thus, the optical and electronic properties of these perovskites strongly depend on halide composition, and their optical absorption and emission color are tuned throughout the visible to near-infrared region by synthesizing pure or mixed halide perovskites. Further, these properties can be modified by varying the temperature and pressure, changing the composition of A-site and B-site cations, and controlling the size and shape of crystals. For example, the emission wavelength and spectral width, charge carrier lifetime, and exciton binding energy can be modified by downsizing perovskites from macroscopic films and microscale crystals to nanocrystals and quantum dots. The excellent abilities of perovskites to absorb light, which is followed by the generation and long-range diffusion of free charge carriers, make them attractive for solar cells. Further, perovskites, owing to their tunable excitonic emission wavelengths and high photoluminescence quantum yields, find applications in high efficiency multicolor LEDs and low threshold multicolor lasers. Although lead perovskites show excellent optoelectronic properties and device performances, the environmental and health costs of lead encourage researchers to develop lead-free perovskites by selecting B-site cations from among Cu+, Ag+, Sn2+, Ge2+, Bi3+, In3+, and Sb3+. Another important challenge in this field is the poor stability of perovskites against moisture and oxygen. Thus, much effort has been dedicated to stabilizing perovskites with shells from materials such as polymers and silica. Commercial devices and biological applications of halide perovskites can be realized with the development of non-toxic perovskites with optimized structures, stability, and optical and charge carrier properties, which cannot be too far.

Conflicts of interest

The authors declare no competing financial interest.

Acknowledgements

VB acknowledges support under the MEXT JSPS Grant-in-Aid for Scientific Research B (19H02550), MEXT JSPS Special Advancement Research Grant (18H05205), and MEXT JSPS Dynamic Alliance for Open Innovation Bridging Human, Environment and Materials. VB and CS acknowledge support under the Scheme for Promotion of Academic and Research Collaboration (SPARC) by the Science and Engineering Research Board of India. CS acknowledges the Japan International Cooperation Agency (JICA) for financial support. LC acknowledges a JICA Scholarship for doctoral research and SG acknowledges a MEXT Scholarship for doctoral research.

References

  1. W. Xu, Q. Hu, S. Bai, C. Bao, Y. Miao, Z. Yuan, T. Borzda, A. J. Barker, E. Tyukalova and Z. Hu, et al. , Nat. Photonics, 2019, 13, 418–424 CrossRef CAS .
  2. S. Yakunin, L. Protesescu, F. Krieg, M. I. Bodnarchuk, G. Nedelcu, M. Humer, G. D. Luca, M. Fiebig, W. Heiss and M. V. Kovalenko, Nat. Commun., 2015, 6, 8056 CrossRef CAS PubMed .
  3. B. Pradhan, G. S. Kumar, S. Sain, A. Dalui, U. K. Ghorai, S. K. Pradhan and S. Acharya, Chem. Mater., 2018, 30, 2135–2142 CrossRef CAS .
  4. A. Kojima, K. Teshima, Y. Shirai and T. Miyasaka, J. Am. Chem. Soc., 2009, 131, 6050–6051 CrossRef CAS PubMed .
  5. S. D. Stranks, G. E. Eperon, G. Grancini, C. Menelaou, M. J. P. Alcocer, T. Leijtens, L. M. Herz, A. Petrozza and H. J. Snaith, Science, 2013, 342, 341–344 CrossRef CAS PubMed .
  6. I. Levchuk, A. Osvet, X. Tang, M. Brandl, J. Perea, F. Hoegl, G. J. Matt, R. Hock, M. Batentschuk and C. J. Brabec, Nano Lett., 2017, 17, 2765–2770 CrossRef CAS PubMed .
  7. M. I. Saidaminov, A. L. Abdelhady, G. Maculana and O. M. Bakr, Chem. Commun., 2015, 51, 17658–17661 RSC .
  8. S. Ghimire, L. Chouhan, Y. Takano, K. Takahashi, T. Nakamura, K. Yuyama and V. Biju, ACS Energy Lett., 2019, 4, 133–141 CrossRef CAS .
  9. T. C. Jellicoe, J. M. Richter, H. F. J. Glass, M. Tabachnyk, R. Brady, S. E. Dutton, A. Rao, R. H. Friend, D. Credgington, N. C. Greenham and M. L. Böhm, J. Am. Chem. Soc., 2016, 138, 2941–2944 CrossRef CAS PubMed .
  10. J. S. Manser, J. A. Christians and P. V. Kamat, Chem. Rev., 2016, 116, 12956–13008 CrossRef CAS PubMed .
  11. J. Zhang, Y. Yang, H. Deng, U. Farooq, X. Yang, J. Khan, J. Tang and H. Song, ACS Nano, 2017, 11, 9294–9302 CrossRef CAS PubMed .
  12. E. T. McClure, M. R. Ball, W. Windl and P. M. Woodward, Chem. Mater., 2016, 28, 1348–1354 CrossRef CAS .
  13. N. K. Noel, S. D. Stranks, A. Abate, C. Wehrenfennig, S. Guarnera, A.-A. Haghighirad, A. Sadhanala, G. E. Eperon, S. K. Pathak and M. B. Johnston, et al. , Energy Environ. Sci., 2014, 7, 3061–3068 RSC .
  14. N. J. Jeon, H. Na, E. H. Jung, T. Yang, Y. G. Lee, G. Kim, H. Shin, S. I. Seok, J. Lee and J. Seo, Nat. Energy, 2018, 3, 682–689 CrossRef CAS .
  15. M. Xiao, F. Huang, W. Huang, Y. Dkhissi, Y. Zhu, J. Etheridge, A. Gray-Weale, U. Bach, Y. Cheng and L. Spiccia, Angew. Chem., 2014, 126, 10056–10061 CrossRef .
  16. J.-H. Im, I. Jang, N. Pellet, M. Grätzel and N. Park, Nat. Nanotechnol., 2014, 9, 927–932 CrossRef CAS PubMed .
  17. A. K. Jena, A. Kulkarni and T. Miyasaka, Chem. Rev., 2019, 119, 3036–3103 CrossRef CAS PubMed .
  18. Y. Deng, X. Zheng, Y. Bai, Q. Wang, J. Zhao and J. Huang, Nat. Energy, 2018, 3, 560–566 CrossRef CAS .
  19. D. Shi, V. Adinolfi, R. Comin, M. Yuan, E. Alarousu, A. Buin, Y. Chen, S. Hoogland, A. Rothenberger and K. Katsiev, et al. , Science, 2015, 347, 519–522 CrossRef CAS PubMed .
  20. H. Zhou, S. Yuan, X. Wang, T. Xu, X. Wang, H. Li, W. Zheng, P. Fan, Y. Li, L. Sun and A. Pan, ACS Nano, 2017, 11, 1189–1195 CrossRef CAS PubMed .
  21. L. Protesescu, S. Yakunin, M. I. Bodnarchuk, F. Bertolotti, N. Masciocchi, A. Guagliardi and M. V. Kovalenko, J. Am. Chem. Soc., 2016, 138, 14202–14205 CrossRef CAS PubMed .
  22. F. Zhang, H. Zhong, C. Chen, X. Wu, X. Hu, H. Huang, J. Han, B. Zou and Y. Dong, ACS Nano, 2015, 9, 4533–4542 CrossRef CAS PubMed .
  23. B. Yang, X. Mao, F. Hong, W. Meng, Y. Tang, X. Xia, S. Yang, W. Deng and K. Han, J. Am. Chem. Soc., 2018, 140, 17001–17006 CrossRef CAS PubMed .
  24. L. Protesescu, S. Yakunin, M. I. Bodnarchuk, F. Krieg, R. Caputo, C. H. Hendon, R. X. Yang, A. Walsh and M. V. Kovalenko, Nano Lett., 2015, 15, 3692–3696 CrossRef CAS PubMed .
  25. L. Wang, N. E. Williams, E. W. Malachosky, J. P. Otto, D. Hayes, R. E. Wood, P. Guyot-Sionnest and G. S. Engel, ACS Nano, 2017, 11, 2689–2696 CrossRef CAS PubMed .
  26. M. I. Saidaminov, A. L. Abdelhady, B. Murali, E. Alarousu, V. M. Burlakov, W. Peng, I. Dursun, L. Wang, Y. He, G. Maculan and A. Goriely, et al. , Nat. Commun., 2015, 6, 7586 CrossRef PubMed .
  27. L. Huang and W. R. L. Lambrecht, Phys. Rev. B: Condens. Matter Mater. Phys., 2013, 88, 165203 CrossRef .
  28. Q. A. Akkerman, S. G. Motti, A. R. S. Kandada, E. Mosconi, V. D’Innocenzo, G. Bertoni, S. Marras, B. A. Kamino, L. Miranda and F. D. Angelis, et al. , J. Am. Chem. Soc., 2016, 138, 1010–1016 CrossRef CAS PubMed .
  29. Q. Dong, Y. Fang, Y. Shao, P. Mulligan, J. Qiu, L. Cao and J. Huang, Science, 2015, 347, 967–970 CrossRef CAS PubMed .
  30. R. L. Milot, G. E. Eperon, H. J. Snaith, M. B. Johnston and L. M. Herz, Adv. Funct. Mater., 2015, 25, 6218–6227 CrossRef CAS .
  31. C. Wehrenfennig, G. E. Eperon, M. B. Johnston, H. J. Snaith and L. M. Herz, Adv. Mater., 2014, 26, 1584–1589 CrossRef CAS PubMed .
  32. A. Amat, E. Mosconi, E. Ronca, C. Quarti, P. Umari, Md. K. Nazeeruddin, M. Grätzel and F. D. Angelis, Nano Lett., 2014, 14, 3608–3616 CrossRef CAS PubMed .
  33. Y. Dong, T. Qiao, D. Kim, D. Parobek, D. Rossi and D. H. Son, Nano Lett., 2018, 18, 3716–3722 CrossRef CAS PubMed .
  34. V. Malgras, S. Tominaka, J. W. Ryan, J. Henzie, T. Takei, K. Ohara and Y. Yamauchi, J. Am. Chem. Soc., 2016, 138, 13874–13881 CrossRef CAS PubMed .
  35. D. Zhang, Y. Yu, Y. Bekenstein, A. B. Wong, A. P. Alivisatos and P. Yang, J. Am. Chem. Soc., 2016, 138, 13155–13158 CrossRef CAS PubMed .
  36. L. Wang, K. Wang and B. Zou, J. Phys. Chem. Lett., 2016, 7, 2556–2562 CrossRef CAS PubMed .
  37. C. Li, Y. Zhong, C. A. M. Luna, T. Unger, K. Deichsel, A. Gräser, J. Köhler, A. Köhler, R. Hildner and S. Huettner, Molecules, 2016, 21, 1081 CrossRef PubMed .
  38. Y. Tian, A. Merdasa, M. Peter, M. Abdellah, K. Zheng, C. S. Ponseca Jr., T. Pullerits, A. Yartsev, V. Sundström and I. G. Scheblykin, Nano Lett., 2015, 15, 1603–1608 CrossRef CAS PubMed .
  39. M. Gerhard, B. Louis, R. Camacho, A. Merdasa, J. Li, A. Kiligaridis, A. Dobrovolsky, J. Hofkens and I. G. Scheblykin, Nat. Commun., 2019, 10, 1698 CrossRef PubMed .
  40. L. Chouhan, S. Ghimire and V. Biju, Angew. Chem., Int. Ed., 2019, 58, 4875–4879 CrossRef CAS PubMed .
  41. Y.-S. Park, S. Guo, N. S. Makarov and V. I. Klimov, ACS Nano, 2015, 9, 10386–10393 CrossRef CAS PubMed .
  42. G. Yuan, C. Ritchie, M. Ritter, S. Murphy, D. E. Gómez and P. Mulvaney, J. Phys. Chem. C, 2018, 122, 13407–13415 CrossRef CAS .
  43. N. Ito, M. A. Kamarudin, D. Hirotani, Y. Zhang, Q. Shen, Y. Ogomi, S. Iikubo, T. Minemoto, K. Yoshino and S. Hayase, J. Phys. Chem. Lett., 2018, 9, 1682–1688 CrossRef CAS PubMed .
  44. S. T. Ochsenbein, F. Krieg, Y. Shynkarenko, G. Rainò and M. V. Kovalenko, ACS Appl. Mater. Interfaces, 2019, 11, 21655–21660 CrossRef CAS PubMed .
  45. K. Lin, J. Xing, L. N. Quan, F. P. G. de Arquer, X. Gong, J. Lu, L. Xie, W. Zhao, D. Zhang and C. Yan, et al. , Nature, 2018, 562, 245–248 CrossRef CAS PubMed .
  46. T. Chiba, Y. Hayashi, H. Ebe, K. Hoshi, J. Sato, S. Sato, Y.-J. Pu, S. Ohisa and J. Kido, Nat. Photonics, 2018, 12, 681–687 CrossRef CAS .
  47. Z.-F. Shi, X.-G. Sun, D. Wu, T.-T. Xu, Y.-T. Tian, Y.-T. Zhang, X.-J. Li and G.-T. Du, J. Mater. Chem. C, 2016, 4, 8373–8379 RSC .
  48. F. Yuan, Z. Wu, H. Dong, J. Xi, K. Xi, G. Divitini, B. Jiao, X. Hou, S. Wang and Q. Gong, J. Phys. Chem. C, 2017, 121, 15318–15325 CrossRef CAS .
  49. B. Tang, H. Dong, L. Sun, W. Zheng, Q. Wang, F. Sun, X. Jiang, A. Pan and L. Zhang, ACS Nano, 2017, 11, 10681–10688 CrossRef CAS PubMed .
  50. C.-Y. Huang, C. Zou, C. Mao, K. L. Corp, Y.-C. Yao, Y.-J. Lee, C. W. Schlenker, A. K. Y. Jen and L. Y. Lin, ACS Photonics, 2017, 4, 2281–2289 CrossRef CAS .

Footnote

These authors equally contributed to this article.

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