Luminescent perovskites: recent advances in theory and experiments

Abstract

Perovskites form an important and enormous class of inorganic compounds. Recently, perovskite materials have attracted extensive research interest owing to their excellent optoelectronic properties. Deep insights into the relationships between the crystal structure, electronic structure and properties play an important role in the development of new functional materials and high-performance devices. In this review, after a brief introduction, we first discuss the crystal structure and crystal chemistry of perovskites according to their three classes: standard perovskites, low-dimensional perovskites and perovskite-like halides. Next, the electronic structure and luminescence from different physical origins are presented. Then, we present a survey on the design, synthesis and luminescence properties of different perovskites, including halide perovskites, oxide perovskites, and lanthanide- or transition metal-doped perovskites, also including dimension-different perovskites (3D, 2D, 1D and quantum dots). We also summarize the strategies for improving the photoluminescence quantum yield (PLQY) and chemical stability, including by surface passivation, encapsulation and doping. Finally, we review their applications and give a brief outlook.

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Zhen Song

Zhen Song (Z. Song) received his Ph.D. degree in Material Science and Engineering from the University of Science and Technology Beijing (USTB) in 2014. He continued to work as a postdoc in Prof. Q. L. Liu's group until Dec. 2016. Since then, he has worked as a lecturer in USTB. From Dec. 2017 to Aug. 2018, he stayed in Prof. Luis Seijo's group at Universidad Autónoma de Madrid (Spain) as a visiting scholar. His current research interests are focused on the theoretical aspects of luminescent materials.

Jing Zhao

Jing Zhao (J. Zhao) received her Ph.D. degree from the Technical Institute of Physics and Chemistry (TIPC) of the Chinese Academy of Sciences (CAS) in 2013. Then, she joined TIPC as a research assistant working on the synthesis of new nonlinear optical materials and crystal growth. She worked as a postdoc in Professor Kanatzidis's group in Northwestern University from Aug. 2015 to Aug. 2017. Currently, she is an associate professor at the University of Science and Technology Beijing (USTB). Her current research interests involve the synthesis of new compounds as photoelectric materials and properties characterizations.

Quanlin Liu

Quanlin Liu (Q. L. Liu) completed his Ph.D. degree in Condensed Matter Physics in 1998 at the Institute of Physics, Chinese Academy of Science (IOP CAS). From 1998–2005, he worked as an assistant and associate professor at IOP CAS, including working as a JSPS fellow at the National Institute for Materials Science, Japan (2001–2003). Since then, he has worked as a full professor in Materials Science at the University of Science and Technology Beijing (USTB) (2005–now). His current research interests mainly concern luminescent materials.

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1. Introduction

Perovskite materials have attracted widespread attention due to their various interesting properties, such as ferroelectricity, photoluminescence, superconductivity, nonlinear optical properties and magnetoresistance.^1,2 Perovskite is a calcium titanium oxide mineral species composed of calcium titanate, with the chemical formula of CaTiO₃. It was discovered by Gustav Rose in 1839 in the Ural Mountains of Russia and is named after the Russian mineralogist L. A. Perovski (1792–1856).³ Recently, much research has been devoted to the luminescence in solid-state perovskites, including in powders, single crystals, thin films and amorphous materials. The earliest written account of a solid-state luminescent material comes from a Chinese text published in the Song dynasty (960–1279 A.D.).⁴ As will be discussed in this review, the luminescent properties of perovskite-type compounds largely depend on the dimensionality of the octahedral network, in either halide, oxide or inorganic–organic hybrid perovskites. It mainly arises from excitons in the inorganic part. For halide and oxide perovskites, luminescence could also be achieved from impurities, such as transition-metal and lanthanide ions. For inorganic–organic hybrid perovskites, unprecedented room-temperature luminescence could be realized due to their large exciton binding energies (>300 meV).

Compared to oxide perovskites, metal halide perovskites possess the advantages of weaker bonding, easier processing and better tunability.⁵ Halide perovskites are solution processable to form photoelectric devices with the characteristics of a long carrier diffusion length, high absorption coefficient, high photoluminescence (PL) quantum efficiency and high defect tolerance. The history of halide perovskites can be traced back to the 19th century, with the first report on CsPbX₃ by Wells.⁶ In the middle of the 20th century, Møller first studied the optical properties of CsPbX₃.⁷ The early studies also included cesium tin(II) trihalides, as reported by Scaife et al.⁸ The most attractive material in solar cells is CH₃NH₃PbI₃ (MAPbI₃), which shows the typical three-dimensional (3D) perovskite structure that was first reported by Weber in 1978.⁹ In 2009, Kojima et al. reported that nanocrystals (NCs) of MAPbX₃ (X = Br, I) attached to a TiO₂ surface in photovoltaic cells showed promise for solar cells.¹⁰ Serving as a prelude to perovskite solar cells (PSCs), research in PSCs began in earnest, and in less than a decade such cells had achieved an efficiency exceeding 23%.¹¹ Compared to their 3D counterparts, 2D perovskites have some advantages, such as large exciton binding energy, low trap density and uniform morphology, which are especially beneficial for obtaining a high photoluminescence quantum yield (PLQY).¹²

Luminescent perovskites have important applications in the fields of nonlinear optical properties,¹³ solar cells,^14–16 scintillators,¹⁷ lighting devices,^18,19 water splitting photocatalysts,^20–22 lasing^23–25 and electronic devices (e.g., capacitors, transducers, actuators).²⁶ Readers are suggested to refer to some excellent review papers for more information.^3,27–32 The most important application of luminescent perovskites is in the fabrication of light emitting diodes (LEDs), or as components used in phosphor-converted LEDs. Another type of structure is the organic–inorganic heterostructure in which an inorganic two-dimensional (2D) semiconductor layer and an organic dielectric layer are alternately piled up, naturally forming a quantum well structure. The research in perovskites used for LEDs has remained a hot topic since 1999, when Mitzi first reviewed the crystal growth and properties of organic–inorganic perovskite structures. This kind of semiconductor has either a single organic layer configuration (H₃N-R-NH₃)MX₄ or a bi-organic layer configuration (R-NH₃)₂MX₄, where R is an organic group, M is a divalent metal (such as Pb²⁺, Sn²⁺, Cu²⁺, Ni²⁺, Mn²⁺, Fe²⁺, Co²⁺, Eu²⁺) and X is a halogen (Cl⁻, Br⁻, I⁻).³³ Excitons are stabilized with large binding energy due to the low dimensionality in the perovskite layer, and exhibit intense exciton absorption and PL from the exciton band even at room temperature. Moreover, the spectral characteristics of the layered perovskites can easily be modified by the replacement of the RNH₃, metal and halide. This feature provides the tunability of the emission colour. In addition, the perovskites possess excellent film processability. By using the conventional spin-coating method, optically high-quality thin films can be easily obtained. From the above-mentioned feature, perovskites are expected to be a promising thin film material for light-emitting devices.

In this review, section 2 includes a detailed description of the standard, low-dimensional perovskites and perovskite-like halide structures. Section 3 covers the electronic structure and luminescence in perovskites. The band structure, luminescence from defects, impurities and excitons, quantum dots and well effects are also discussed. The synthesis and properties of halid perovskite are further discussed in section 4, accompanied by a survey of lanthanide- and transition-metal-doped perovskites. Strategies to improve the luminescecnt efficiency and stability are also included. Section 5 compiles information on the broad application of luminescent perovskites.

2. Crystal structure and materials

The number of perovskite halides has increased rapidly in the past few years, and the definition of perovskites has become blurred.^34–36 In this review, we catalog the perovskites by their crystal structures into: standard-perovskite (with an ABX₃ general formula and 3D structure), low-dimensional (low-D) perovskites and perovskite-like halides. The standard-perovskite and low-D perovskites (2D, 1D, 0D) both contain solely corner-sharing or discrete octahedra, while perovskite-like halides contain solely six coordinated octahedra, but the connection fashion is not limited to edge-sharing or face-sharing octahedra (Fig. 1).

Fig. 1 Standard-perovskites (3D), 2D perovskites with different types of layers (2D, 〈100〉, 〈110〉, 〈111〉), chains of octahedra in 1D perovskites, discrete octahedra separated by organic cations (0D) and perovskite-like halides with moieties of edge- and face-sharing octahedra. Aadapted with permission from ref. 37. Copyright (2018) American Chemical Society. Standard-perovskite structural characteristics.

Space rules need to be followed: The typical perovskite structure is cubic with the space group Pm3m-Oh i.e., SrTiO₃ and CsSnBr₃.³⁸ Take the perovskite-type oxides, ABO₃ for example, where A is the larger cation and B is the smaller cation. In ABO₃ structure, the B cation is 6-fold coordinated and the A cation is 12-fold coordinated with the oxygen atoms. The corner-sharing octahedra form the skeleton with the centre position occupied by the A cation.³⁹ The crystal structure of the perovskite is very flexible, but certain rules must be met to ensure structural stability. The 3D-perovskites have the general formula ABX₃. Goldschmidt's tolerance factor t and the octahedral factor (μ) have been used to screen and discover new halide perovskites,³⁴ as the formulas shown below:

where R_A, R_B, R_X refer to the radii of A, B and X ions, respectively. For a typical 3D halide perovskite, the Goldschmidt factor range is 0.81 ≤ t ≤ 1.0 and the octahedral factor is 0.44 ≤ μ ≤ 0.90.

Electroneutrality: Another condition need to be fulfilled is electroneutrality, i.e. the sum of charges of cations should equal the total charge of oxygen anions. This can be obtained by appropriate charge distribution in the form of A¹⁺B⁵⁺O₃, A²⁺B⁴⁺O₃ or A³⁺B³⁺O₃. In addition the partial substitution of A and B ions by other cations is allowed. However, deficiencies at the A- or B-sites or of oxygen anions are frequent resulting in defective perovskites. WO₃ is a representative of B⁶⁺O₃ type peroskites with no A-site cations.⁴⁰ The nonstoichiometry in perovskites has been widely discussed.³⁹ Oxygen vacancies are more common than those of cations, e.g., Ca₂Fe₂O₅ and La₂Ni₂O₅. Ca₂Fe₂O₅ can be considered as an anion-deficient perovskite with one-sixth of the oxygen ion sites being vacant.

Standard perovskite with single-type B-site cations: In the 3D perovskite structure, the octahedra are connected via corner-sharing to form a 3D network. The chemical formula corresponding to the octahedral unit is MX₃, such as in SrTiO₃, CsPbBr₃ and MAPbX₃ (X = Cl, Br, I)^9,41 According to the tolerance factor, for halide perovskites, the A-site cation is limited to Cs, MA or HC(NH₂)₂⁺ (FA) filling the voids of the BX₆ (X = Cl, Br, I) octahedral network. The B site is normally occupied by a divalent cation, such as Sn²⁺, Pb²⁺, Ge²⁺, Sr²⁺, Ca²⁺, Mg²⁺, Cu²⁺ or Ni²⁺.

Double perovskites: Double perovskites have the formula A₂B′B′′O₆ or A₂B′B′′X₆ (X = Cl, Br, I), where the primes indicate different ions in different oxidation states. Since the B cations generally determine the properties of perovskites, the different kinds of B′ and B′′ ions show a variety of properties. It was reported that the distortion of the double perovskite Sr₂LnRuO₆ from the cubic symmetry is mainly due to the tilting of the octahedra rather than the distortion of the octahedra.⁴² For halide perovskites, one valence Na⁺, Ag⁺, Cu⁺ or Au⁺ cation can take part in mixed occupancy with some trivalent cations, such as Sb³⁺, Bi³⁺ or In³⁺, to form double perovskites A₂BB′X₆.^43–45 For more information, one can refer to the review published by Hoefler et al., in which the possible choices for B-sites are discussed in detail.⁴⁶

2.1. Low-D perovskite structural characteristics

Recently, low-D halide perovskites with ordered vacancies or large separating organic cations have been developed to enhance the stability and/or for a reduced toxicity of the most promising 3D lead-containing halides (e.g. MAPbX₃, FAPbX₃), which showed a high conversion efficiency in solar cells. Low-D perovskites consist of sliced layers, corner-sharing chains, or clusters of the ABX₃ structure.⁴⁷ The low-D structure is less constrained by the tolerance factor than the 3D structure.

2.1.1 Low-D perovskites with ordered vacancies. The standard perovskite structure has been extended to include B-site cations higher than valence 2 and vacancies. Removing 1/3rd of the B cations along a certain crystal direction from the 3D ABX₃ system forms A₃B₂X₉ with 2D structures comprising a corrugated layer,³⁶ which can be denoted as AB_2/3□_1/3X₃ (where □ is a vacancy), such examples include Cs₃Sb₂I₉ and Cs₃Sb₂I₉,⁴⁸ K₃Bi₂I₉, and Rb₃Bi₂I₉.⁴⁹

Such a corrugated layer with ordered vacancies can also be found in organic–inorganic hybrid perovskites (OIHPs). (H₂AEQT)M_2/3I₄ (M = Bi³⁺, Sb³⁺) possesses a 〈100〉-oriented single-layered structure, where the high-valent metal halide inorganic sheet is stabilized by vacancy formations at the metal sites. The inorganic sheet can be written as (M³⁺)_2/3V_1/3X₄ (X = Cl, Br, I), where V represents a vacancy. Given a suitable organic cation layer, this may be further extended to include other higher priced metals. For tetravalent metals (e.g. Sn⁴⁺, Te⁴⁺, Hf⁴⁺), the inorganic anion layer can be represented as (M⁴⁺)_1/2V_1/2X₄²⁻, and for pentavalent metals (e.g. Nb⁵⁺, Ta⁵⁺, Mo⁵⁺) as (M⁵⁺)_2/5V_3/5X₄²⁻ and (Mⁿ⁺)_2/nV_(n−2)/nX₄²⁻ for larger n. However, the vacancy concentration of the perovskite structure may become too high to make the structure stable.

Based on the A₂B′B′′X₆ (X = Cl, Br, I) double perovskite, a B-site cation can be replaced by a vacancy to produce an A₂□BX₆ perovskite. In order to maintain the charge neutrality of the structure, B must be a tetravalent cation. Because the two adjacent octahedra are not connected to each other A₂BX₆ forms 0D structures. A₂BX₆ containing Sn⁴⁺, Te⁴⁺, Pt⁴⁺ and Pd⁴⁺ cations have been reported, e.g. Cs₂SnI₆, Cs₂TeI₆,⁵⁰ Cs₂PdBr₆ ⁵¹ and A₂Pt□I₆ (A = NH⁴⁺; MA⁺; FA⁺; and C(NH₂)³⁺).⁵² Further, to broaden this family, the range of tetravalent cations can also be extended to Mn⁴⁺, Zr⁴⁺, Cr⁴⁺ and Ti⁴⁺, or a combination of these cations.

2.1.2 Low-D perovskites with large A-site cations. 2D perovskites. In 2D perovskite structures, the octahedra are connected by corner-sharing forming infinite layers separated by inorganic/organic units forming quantum wells. In this case, the chemical formula of the octahedral unit is MX₄ for one octahedral layer and M_nX_3n+1 for n layers, such as (C₆H₅C₂H₄NH₃)₂PbI₄, KLaNb₂O₇ and K₂La₂Ti₃O₁₀. In contrast to the 3D structures, the layered 2D systems can accommodate much larger and more complex cations.^53,54

Early in 1978, Arend et al.⁵⁵ reported the synthesis, solubility and crystal growth of layered-structure halide perovskites with unbranched organic chains, such as (C_nH_2n+1NH₃)₂MX₄ and NH₃(CH₂)_mNH₃MX₄ with M = Cd, Cu, Fe, Mn or Pd, X = Br or Cl, n = 1, 2,…, 18, and m = 2, 3,…, 8. In 1985, Day reported layered perovskite halide salts (RNH₃)₂MX₄ (R = organic group; M = Cr, Mn, Cd; X = Cl, Br).⁵⁴ The ammonium groups hydrogen bond to the inorganic sheet by halogens, and the organic tails extend into the space between the layers, holding the structure together via van der Waals interactions. The organic–inorganic perovskite family has yielded a remarkable degree of structural versatility.^56,57

According to the connection of the octahedral layers, the 2D perovskites can be divided into three structural types: (1) the 〈100〉-oriented perovskite; (2) the 〈110〉-oriented perovskite; (3) the 〈111〉-oriented perovskite (Fig. 1).¹² Here, (1) and (2) have the general formula A′₂A_n−1B_nX_3n+1 or A′A_n−1B_nX_3n+1 (A′ = 1+ or 2+; A = 1+ cation; B = Pb²⁺, Sn²⁺, Ge²⁺, Cu²⁺, Cd²⁺, etc.; X = Cl⁻, Br⁻, and I⁻), and (3) has the general formula A′_n+1B_nX_3n+3 (n ≥ 1, where B valence is +3, or a mixed valence averaging +3, e.g. Cu²⁺ and Sb³⁺).⁵⁸ According to a survey of the Cambridge Structural Database, the majority of 2D perovskites possess <100>-oriented layers, with the total number being approximately 250, and there are only a few reports on <110>-oriented structure.^12,59–64

In <100>-oriented perovskites, the thickness of the inorganic layer (corresponding to an n value of A′₂A_n−1B_nX_3n+1) can be controlled by adjusting the ratio between the spacer cations and the smaller cation, with the n value varying from 0 to ∞.^65,66 Similarly, thicker layered 2D perovskites with both <110>- and <111>-oriented inorganic sheets can be obtained (Fig. 2a). We can regard the 3D perovskites as the extreme case of this 2D structure with n = ∞, and 2D perovskites as the case with n = 1.^53,65,67,68 For <110>-oriented perovskites, depending on where the ripple occurs in the corrugated layer, the structures can be defined as 2 × 2, 3 × 3, 4 × 4, “n × n”, where n represents the number of octahedra making up half of the roof (Fig. 2b).⁶³

Fig. 2 (a) The evolution from 2D perovskites to 3D perovskites, with the organic moieties omitted. Reprinted with permission from ref. 58. Copyright (2019) American Chemical Society; (b) corrugated (110)-oriented 2D perovskites with different members of half of the roof octhedra n. Reprinted with permission from ref. 63. Copyright (2017) American Chemical Society.

1D and 0D perovskites. In the 1D perovskite structure, large organic cations separate the infinite chains formed by corner-sharing octahedra, with the unit chemical formula of MX₅. The 0D perovskite structure preserves MX₆ as the unit chemical formula, for which the isolated octahedra are non-interacting and are separated from each other by organic moieties. The 0D OIHP-forming B-site metals include Ti⁴⁺, Hf⁴⁺, Zr⁴⁺, Pd⁴⁺, Pb²⁺, Sn⁴⁺, Te⁴⁺, Sb³⁺, Mn⁴⁺, In³⁺, Bi³⁺ and Cr⁴⁺ or a combination of these cations.^36,69

2.2. Structural characteristics of perovskite-like halides

Perovskite-like halides contain, but are not limited to, octehadra connected by edge-sharing or face-sharing octahedra forming different dimensional structures. This type of structure is weakly associated with perovskites, with the only common feature being that they contain six coordinated metal cations. Therefore, the structure only needs to satisfy the octahedral factor μ. In this review, we discuss the structure and PL properties of this type of perovskite for two main reasons: (1) many compounds with such structures and high PLQY have been reported recently, (2) so far, octahedral connection ways cannot be designed or predicted, and this type of structure occasionally appears when trying to synthesize low-D perovskites.

2.3. Morphology and materials

Luminescent perovskite materials have a variety of morphologies. Zhu et al. systematically synthesized MAPbX₃ (X = I, Br) NCs into dots, rods, plates and sheets by using different solvents and capping ligands.⁷⁰ Morphology evolution is often related to the change in luminescent performance. The luminescence degradation of perovskite NCs mainly originate from the large CsPbBr₃ crystals under illumination.⁷¹ Excitonic emission has only been observed in MAPbI₃ single crystals, but not in MAPbI₃ thin films⁷² The further application of lead halide perovskite materials is hindered by lead toxicity. Different non-lead perovskite NCs can be synthesized by replacing Pb²⁺ with other isoelectronic elements, such as Sn⁴⁺, Sb³⁺ and Bi³⁺.^73,74 New synthesis methods have been developed besides the conventional solution-based colloidal process for perovskite NCs.⁷⁵ Dirin et al. developed a method involving the infiltration of perovskite precursor solutions into the pores of mesoporous silica, followed by drying, leading to the template-assisted formation of perovskite NCs.⁷⁶ Chen et al. reported the grinding synthesis of the whole family of MAPbX₃, FAPbX₃, and CsPbX₃ (X = Cl, Br, I and their mixtures) perovskite NCs, which could be operated at room temperature and in the open atmosphere.⁷⁷

3. Electronic structure and luminescence in perovskites

The luminescent properties of materials are dependent on their electronic structures; while the electronic structure is determined by the crystal structure and chemical composition. To obtain a full understanding of the luminescence in perovskites, it is necessary to study the electronic structure in perovskite materials. The band structure of a perovskite is closely related to its composition. In organic–inorganic hybrid perovskites, the inorganic layer dominates the luminescent properties. Compared with cation A, the octahedra [BX₆] play a more crucial role in the formation of the perovskite electronic structure. For example, in MAPbI₃, the conduction band minimum is mainly affected by the p-orbital electrons of the Pb atom, while the valence band top is mainly composed of the Pb s orbital and the I p orbital. The cation A can only adjust the band structure by changing the bending and stretching between Pb and the halogen atom in the [PbX₆] octahedra.^78,79 For example, the electronic structures between MAPbX₃ (X = I, Br) and CsPbX₃ are the same. The band gap of the hybrid organic–inorganic compounds CH₃NH₃SnI₃ and NH₂CHNH₂SnI₃ is very close to that of a hypothetical CsSnI₃ cubic perovskite with the same cell size.⁸⁰

The luminescence in perovskites originates from the radiative processes, including band-to-band transition, electron–hole recombination and the transitions between emissive sub-band levels. Discussion on the core–valence luminescence in some scintillators, which involves the recombination of a core hole and a valence electron, is not included in this paper.⁸¹ Although some mechanisms have been established in a number of well-developed theories, further studies on the luminescent mechanism are continuing to contribute new insights. For example, there exists a debate about the relative ordering of dark and bright sublevels in halide perovskites. Becker et al. showed that a highly emissive triplet state is the lowest excitonic level in caesium lead halide perovskites (CsPbX₃, with X = Cl, Br or I).⁸² On the contrary, Tamarat et al. proved that the dark singlet exciton state is located several meV below the bright triplet in formamidinium lead bromide (FAPbBr₃) perovskite NCs.⁸³ In Table 1, various luminescent properties are provided, including the mechanism (from impurities or excitons) and emission peak wavelengths.

Table 1 Compilation of luminescent perovskite materials together with the emission peak wavelength

Material name	Luminescence mechanism*	Emission maximum (nm)	Ref.	Material name	Luminescence mechanism*	Emission maximum (nm)	Ref.
*: The letter “I” stands for impurity luminescence and “E” for excitonic luminescence. **: The luminescence in LaInO3 comes from In^3+ as impurities.^84 †: Cerium ion present in Sr2CeO4 is tetra-valent with no radiative emission. The luminescence originates from the host.^98
LaInO3:Bi^3+	I	420	84	Ba2CaTeO6:U^6+	I	500	85
CaZrO3:Bi^3+	I	390	86	KMgF3:Cu^+	I	415	87
LaAlO3:Bi^3+	I	375	84	NaMgF3:Cu^+	I	375	87
CaZrO3:Pb^2+	I	365	88	LiBaF3:Cu^+	I	465	87
LaInO3	I**	515	84	(C10H21NH3)2PbI4	E	516	56
NaMgF3:Eu^2+	I	365	89	YAlO3:Ce^3+	I	370	90
KMgF3:Eu^2+	I	363	89	LuAlO3:Ce^3+	I	365	91
RbMgF3:Eu^2+	I	360	89	BaTiO3	E	485	92
CsMgF3:Eu^2+	I	360	89	(MeNH3)SnI3	E	761	93
KCaF3:Eu^2+	I	485, 520	89	(MeNH3)(C10H21NH3)2Sn2I7	E	733	93
RbCaF3:Eu^2+	I	475	89	(C10H21NH3)2SnI4	E	603	93
CsCaF3:Eu^2+	I	510	89	(MeNH3)SnBr3	E	576	93
RbSrF3:Eu^2+	I	424	89	(MeNH3)(C10H21NH3)2Sn2Br7	E	517	93
CsSrF3:Eu^2+	I	426	89	(MeNH3)PbCl3	E	408	93
Ba5Ta4O15	E	455	94	(C10H21NH3)2PbCl4	E	336	93
Ba5Nb4O15	E	575	94	(C6H5C2H4NH3)2PbI4	E	520	95
KTaO3	E	490	96	(C6H9C2H4NH3)2PbI4	E	510	97
LiTaO3	E	340	96	Sr2CeO4	E^†	485	98
NaTaO3	E	440	96	(C6H5C2H4NH3)2PbBr4	E	406	99
KLaNb2O7	E	590	100	(C6H5C2H4NH3)2PbI4	E	520	99
K2La2Ti3O10	E	475	101	(C4H9NH3)2SnI4	E	616	99
Gd2MgTiO6:Mn^4+	I	681	102	KMgF3:Ce^3+	I	350	103
BaLaLiWO6:U^6+	I	538	104	LiBaF3:Ce^3+	I	325	103
Ba2SrWO6:U^6+	I	512	104	CsPbCl3	I	418	105
Sr2MgWO6:U^6+	I	504	104	(C4H9NH3)2PbBr4	I	412	106
(C4H9NH3)2EuI4	I	460	107	SrSnO3	E	425	108
Ba2MgWO6:U6+	I	510	85	Ba2CaTeO6:U^6+	I	500	85

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3.1. Band gap and affecting factors

The electronic structure plays a vital role in the luminescent properties of perovskites. The forbidden band-gap energy is influenced by both the crystal structures and chemical components. The effect of dimensionality will also be discussed.

Crystal structure factor. For the perovskite structure, the octahedral tilting distortion is the main factor.¹⁰⁹ In the alkaline-earth stannate perovskites (BaSnO₃, SrSnO₃ and CaSnO₃), the conduction bandwidth decreases strongly in response to the octahedral tilting distortion, triggered by the decreasing size of the alkaline-earth cation. This in turn leads to a corresponding increase in the band gap from 3.1 eV in BaSnO₃ to 4.4 eV in CaSnO₃.¹¹⁰ The band gap of CdSnO₃ is relatively small (3.0 eV) considering the large octahedral tilting distortion. The anomaly stems from the mixing between the empty Cd 5s orbitals and the antibonding Sn 5s–O 2p states, which leads to a widening of the conduction band and a corresponding decrease in the band gap. For pyrochlores (Re₂Sn₂O₇, Re = Y, La, Lu), the Sn–O–Sn bonds are highly distorted from the linear geometry in pyrochlore, leading to a relatively narrow conduction band and a wide band gap. In Cd₂Sb₂O₇ and defect pyrochlore oxides Ag₂Sb₂O₆, the Cd²⁺ and Ag⁺ ions exhibit a strong inductive effect, which widens the conduction band and lowers the band gap significantly, very similar to the effect observed in the perovskite form of CdSnO₃.¹¹⁰

Composition factor. Although the variation in chemical composition also induces a change in the crystal structure, the contribution from the chemical component could be prominent in crystal structures with great similarity. The band gaps of AgTaO₃ and AgNbO₃ are 3.4 and 2.8 eV, respectively, being 0.6 eV smaller than the band gaps of NaTaO₃ (4.0 eV) and NaNbO₃, although the crystal structures of AgMO₃ are similar to those of NaMO₃.¹¹¹ Using the plane-wave-based density functional method, it was found that a hybrid orbital of Ag 4d and O 2p forms a valence band at a more negative level than O 2p orbitals, resulting in a decrease in the band gap. A full band gap range of 1.6–2.3 eV could be modulated through MAPbI_3−xBr_x perovskite films.¹⁵ With the change of the mixed-halide, the full-spectrum luminescence (410–700 nm) of CsPbX₃ can be realized.¹¹² Transition metal oxide perovskites usually possess large energy gaps (>3 eV) due to the large energy differences between the transition metal d-orbital conduction band and the O 2p orbital valence band. By replacing O with S or Se in transition metal chalcogenide perovskites, the valence band composed of mainly chalcogen (S, Se) 3p or 4p orbitals could be shifted higher and the band gaps could be decreased to the visible-IR range.¹¹³

Dimensionality factor. The effects of dimensionality on the band gap could be analyzed from the viewpoint of the inductive effect and charge transfer process. Mitzi et al. found that in the layered compound (C₄H₉NH₃)₂EuI₄, the luminescence peak occurs at 460 nm, while for the 3D system, CsEuI₃, the emission peaks at 449 nm.¹⁰⁷ In this case, the 3D crystal structure widens the band and subsequently narrows the band gap. In other words, breaking the corner-sharing octahedra network increases the band gap. Similarly, the band gaps of Sr₂SnO₄, Sr₃Sn₂O₇ and SrSnO₃ decrease as 4.43 (280 nm), 4.13 (300 nm) and 3.88 eV (320 nm), respectively.¹¹⁴ The optical band gaps of A(La_0.98Bi_0.02)Ta₂O₇ (A = Rb, K, and Na) phosphors were measured by their diffuse reflection spectra and estimated to be about 4.10, 3.94 and 3.96 eV, because the 2D perovskite layers are less separated in the sequence Rb, K, Na.¹¹⁵ In the 3D perovskite crystal structure with interconnecting octahedra, the band gap could be widened by smaller A cations with larger electronegativity. The band gap of (Ba_1−xSr_x)₂YSbO₆:0.005Mn⁴⁺ changes from 4.50 to 4.72 eV as the Sr²⁺ content increases from x = 0 to 1.0.¹¹⁶ Through the compositional modulation with increasing Rb, the band gap and emission spectra of Rb_xCs_1−xPbBr₃ are readily tunable over the visible spectral range from 532 to 474 nm.¹¹⁷ The dispersive band edges of CsPbBr₃ do not support self-trapped carriers, which agrees with reports of a weak exciton binding energy and high photocurrent. The larger gap 0D material Cs₄PbBr₆, however, has revealed polaronic and excitonic features.¹¹⁸

The low-Dl structure results in a dielectrically restricted exciton binding energy increase due to the difference in dielectric constant between organic ions and [BX₆] octahedra. Therefore, when gradually increasing the n value, the dimension of the 2D perovskite increases, and the well width of the quantum well increases correspondingly, resulting in weaker exciton binding energy, a reduced band gap and a red-shift of the emission peak. As direct band gap semiconductors, OIHPs have inherent advantages in terms of their conductive and luminous properties, and their band gaps can also be tuned by regulating the inorganic element components.^119,120 Quantum wells are very good fluorescent materials, and their photoluminescence can be rooted to various emission mechanisms, including strongly correlated electron–hole pairs, which are known as free excitons (Fes), permanent lattice defects, transient light-induced defects, like self-trapped excitons (STE), and coordination of the inorganic layer and molecular chromophores.

To obtain a full understanding of the luminescence in perovskites, it is necessary to study the electronic structure in perovskite materials. In OIHPs, the inorganic layer dominates the luminescent properties. For example, the electronic structures between MAPbX₃ (X = I, Br) and CsPbX₃ are the same. Moreover, the material dimensionality also has an effect. Tuning the average crystallite dimension in methylammonium lead trihalide perovskite thin film from tens of nanometres to a few micrometres reveals that larger crystallites present a smaller band gap and longer lifetime.¹²¹

Other factors. The band gaps are also related to the temperature and synthesis conditions. For example, the band gaps of MAPbI₃, MAPbBr₃ and FAPbBr₃ exhibit an unusual blue-shift when raising the temperature from 15 to 300 K, caused by the stabilization of the valence band maximum.¹²² The band gap of Cs₂AgSbCl₆ could be tuned by adding different volumes of HCl during synthesis.¹²³

3.2. Understanding the electronic structure from the viewpoint of charge transfer and the inductive effect

Charge transfer in OIHPs has been addressed as the resonant interaction between Wannier excitons in inorganic materials and Frenkel excitons in adjacent organic layers.^124,125 However, in this paper, charge transfer is mainly focused on the inorganic part, which is commonly regarded as involving a transfer of charge from a ligand anion (oxygen or halide) towards the central metal cation (Ti⁴⁺, Pb²⁺) in the octahedral coordination. In reality, no electron is transferred in the transition, but a considerable reorganization of the charge density distribution around the metal is expected. This reorganization is accompanied by an expansion of the metal–ligand bonds in the excited state, which gives rise to the observation of a large Stokes shift and broad bands.

The absorption edge in perovskites consisting of a highly charged cation with a noble gas configuration and oxygen ions is usually caused by charge-transfer transition. As can be seen in Fig. 3, an electron from the highest filled molecular orbital localized on the oxygen ions is transferred to the lowest empty molecular orbital localized on the highly charged ion, which is mainly 5d(t_2g) for the tungsten ion. Blasse and Corsmit found from the reflection spectra of A₂BWO₆ compounds that the absorption edge is mainly dependent on the A element and weakly influenced by the B element.¹²⁶

Fig. 3 Local anion coordination and schematic shape of the molecular orbitals involved in the charge-transfer transition in the perovskite A₂BWO₆. Redrawn from ref. 108, Copyright (1973), with permission from Elsevier.

With the increase in the ionic radius of A or B, the absorption edge is continuously shifted to longer wavelength. To elucidate the charge-transfer mechanism, the local coordination and the relevant orbitals are redrawn. The absorption corresponds to an electronic transition from the occupied oxygen 2p orbitals to the empty tungsten t_2g orbitals. Since the spectral position of the charge-transfer band depends on the charge and on the radius of the coordinated cations, it is reasonable that the oxygen valence electron feels a weaker field as the A or B ion increases in ionic radius, and therefore less energy is needed to initiate the electron transfer to the highly charged tungsten ion. Consequently, the charge transfer is more sensitive to the four coordinated A ions than the only one coordinated B ion, even though the B ion is located closer to the tungstate group than the A ion.

In this sense, some of the key factors in luminescent perovskites could be understood from the viewpoint of charge transfer and the inductive effect. Exciting one electron from the valence band maximum to the conduction band minimum across the band gap energy could be imagined as a charger transfer from the atomic orbital constituting the valence band to that constituting the conduction band. Meanwhile, since the electron–hole pair of an exciton is created by photon absorption, the process could also be regarded as the charge transfer of one electron from the full valence band to the empty conduction band, with one hole left in the valence band. As a result, the charge-transfer energy could be analyzed by considering the inductive effect.¹²⁷ Take an ABO₃-type perovskite material for example. With a smaller electronegativity of A (larger ionic radius with the same valence state), the electron in ligand O is less attracted by A, and the central B takes less energy to transfer an electron charge from the ligand O. On the other hand, the lower electronegativity of the ligand also favours the charge-transfer process. In CsPbX₃ (X = Cl, Br, I), the excitonic luminescence spectra could be obtained from blue to red as the ligand changes from Cl to I, with decreasing electronegativity. The tuning the light emission wavelength of 2D CsPbBr_xCl_3−x perovskite films from green (504 nm) to blue (470 nm) could be realized through compositional engineering via increasing the content of chloride.¹²⁸

3.3. Interband/exciton luminescence

In semiconductors and insulators, optical transitions across the fundamental band gap excite an electron towards conduction, with a hole created in the valence band. Annihilation of an electron–hole pair leads to interband luminescence, which is also known as radiative electron–hole recombination. However, the band gap energy is usually difficult to determine due to the presence of an exciton. Dorenbos proposed that the band gap energy could be estimated by 1.08 × E_ex, where E_ex is the energy needed to produce an exciton, and the factor of 1.08 means the electron–hole binding energy of the exciton.¹²⁹ So the defined mobility band edge energy is much larger (≈15%) than the fundamental optical absorption band edge energy.

Interband luminescence. The excitonic level is located quite close to the conduction band, such that it is difficult to distinguish interband luminescence from exciton luminescence in some materials. The pronounced conductivity dependence of the emission intensity in SrTiO₃ suggests a direct recombination of the conduction electrons and the oxygen 2p valence band holes.^130,131 Jellicoe et al. observed two luminescent decay channels in CsSnX₃ (X = Cl, Br, I), which were assigned to a fast band-to-band emission and a slow radiative recombination at shallow intrinsic defect sites.⁷³

3.3.1. Exciton luminescence. In semiconductors the absorption grows rapidly above a certain photon energy, which is defined as the fundamental absorption edge. However, this edge usually shows a complicated structure, consisting of a series of lines. These lines correspond to an exciton, which is central to the luminescence in inorganic and organic–inorganic hybrid-perovskites.

Generally, there exist two types of excitons, Frenkel and Wannier excitons. An exciton can be understood as a single excited electron taken out of a band full of electrons according to band theory, which provides a basic description of the electronic states. Band theory claims that all states in the full valence bands correspond to the ground state of the crystal, and in the meantime, all the states in the conduction bands are empty. A hole–electron pair is created in the process of exciting an electron from the valence to the conduction band across the forbidden gap under light absorption. The Coulomb attraction between the electron–hole pair lowers the formation energy of the hydrogen-like state compared to that given by band theory. By solving the Schröndinger equation for a hypothetical 2D hydrogen atom, Shinada and Sugano found that a small peak may appear just above the absorption edge because of the Coulomb interaction between an excited electron and a hole.¹³² The exciton may be described by the effective-mass approximation when the interaction is weak in a medium with a high dielectric constant.¹³³ Frenkel introduced the concept of an exciton as “excitation quanta”.¹³⁴ That is, an excitation wave formed by exciting individual atoms to higher atomic states can be associated with an “excitation quantum” similar to a light quantum. Under interatomic interactions, the motion of this excitation quantum represents the motion of the excitation travelling through the crystal. Its energy is related to the energy difference between the excited and the normal state. In Frenkel's description, in the case of a strong Coulomb interaction, the hole–electron pairs are confined to a single atom, but the excitation state could be cruising from atom to atom.

Wannier found that in the states located near the bottom of the excited Bloch band, the electron cannot escape its hole completely, and no photocurrent can be observed.¹³⁵ Discrete states are included in the lower part, in which the widely spaced lowest states correspond to the excitation of an electron within its cell or to some direct neighbour. The discrete energy spectrum is obtained for bounded excitons. This individual character disappears at the higher states with narrower spacing, in which the electron moves in the Coulomb field of its hole. Furthermore, the continuous Bloch band follows with the electron and hole moving independently, and a current may be observed. Therefore for unbounded excitons, a continuous energy spectrum could be obtained. The free-exciton emission is usually caused by band-gap excitation and has a narrow-band feature.

Exciton luminescence. The luminescence of inorganic–organic hybrid perovskites based on metal halide sheets and optically inert organic cations arises from the exciton states associated with the band gap of the metal halide framework. For example, in the multilayer perovskites (C₄H₉NH₃)₂(CH₃NH₃)_n−1(Ge,Sn,Pb)_nI_3n+1, intense room-temperature photoluminescence has been observed with wavelengths ranging from the ultraviolet through the red spectral region in the germanium(II)-, tin(II)- and lead(II)-based systems.¹⁰⁷ The luminescence originates from the radiative decay of the FEs associated with the 2D inorganic layers in the structure. It should be noted that the luminescence properties of the organic–inorganic hybrid compounds may be governed by the excitonic properties of the inorganic layer.¹³⁶ The emission wavelength is tunable through the choice of metal atom, halogen or the thickness of the perovskite sheets. It should be noted that in the purely inorganic PbI₂, the room-temperature exciton luminescence is quenched because of the small exciton binding energy (approximately 30 meV). Small exciton binding energies of 32 and 41 meV have also been observed for (CH₃NH₃)PbI₃ and MAPbCl₃, respectively, which have 3D perovskite structures.^137,138 On the contrary, the exciton binding energies in lead(II)-based organic–inorganic perovskites approximate 200–400 meV, leading to strong room-temperature excitonic photoluminescence.¹⁰⁷ The increase in binding energy is caused by both the quantum confinement effect (two dimensionality of the structure) and the smaller dielectric constant of the interleaving organic layers sandwiching the metal halide sheets, which enhances the electron–hole Coulomb interaction. For example, in the layer-type perovskite structure of (C₁₀H₂₁NH₃)₂PbI₄, the PbI₄ layers are sandwiched by alkylammonium chains as barrier layers. The perovskite has a large excitonic binding energy of 370 meV, originating both from the 2D characteristic and the small dielectric constant of the barrier layer, with the latter one achieving a much stronger Coulomb interaction between an electron and a hole.¹³⁹

Besides, the excitonic emission wavelength is largely affected by the dimensionality or thickness of the inorganic layer, which could also be explained by the quantum confinement. Generally, quantum confinement leads to a spectral blue-shift towards high energy.¹⁴⁰ In (C₁₀H₂₁NH₃)₂(CH₃NH₃)_n−1Pb_nI3_n+1, as the thickness (n) of the perovskite sheets increases from n = 1 to the 3D n = ∞ compound (i.e. MAPbI₃), the luminescence shows a substantial peak shift towards longer wavelength from 524 nm to 753 nm.¹⁰⁷ This phenomenon in the luminescence spectrum can be understood by a band gap reduction with the increasing perovskite sheet thickness, because the bands are more easily formed as the dimensionality increases. Tabuchi et al. modulated the inorganic layer thickness in the layered perovskites compounds (C_nH_2n+1NH₃)₂(CH₃NH₃)_m−1Pb_mBr_3m+1 by changing the ratio of the two amines (C_nH2_n+1NH₃/CH₃NH₃).¹⁴¹ The strong, clear excitonic absorption peak of the layered perovskite films measured at room temperature was caused by the large exciton binding energy. With increasing the number of inorganic layers from 1 to 3, a red-shift of the excitonic absorption was observed because of the decrease in transfer energy among the inorganic network.

A smaller dielectric constant has the ability to reduce the dielectric screening of the Coulomb interaction between electrons and holes. Compared to MAPbI₃, MABrI₃ has a lower dielectric constant because of the larger band gap energy.¹⁴² This results in a smaller Bohr radius and larger binding energies of the excitons in MAPbBr₃. Meanwhile, the smaller extent of the exciton wave function is reflected by a larger oscillator strength.

3.3.2. Self-trapped exciton. When the photogenerated electron–hole pair strongly couple to the lattice, a geometric lattice relaxation can be expected around the created charge. Consequently, the excitonic emission shows a Stokes shift from the absorption edge.¹⁴³ Those excitons with strong exciton–phonon interactions are called “self-trapped excitons” or STEs for abbreviation.¹⁴⁴ STE luminescence has been widely observed in oxide, organic–inorganic hybrids and all-inorganic halide perovskites.^145–147 Under band-to-band excitation, the observed broad luminescence band in SrTiO₃ at low temperatures was assigned to STEs, taking into account the broad bandwidth (FWHM = 0.5 eV) and the large Stokes shift (0.8 eV).^148–150 The semi-empirical Intermediate Neglect of Differential Overlap (INDO) method was performed with the self-activated crystals PbTiO₃, SrTiO₃, BaTiO₃, KNbO₃ and KTaO₃ by Eglitis et al.^151,152 Universal green luminescence was observed as a result of the radiative recombination of the self-trapped electrons and holes forming the charge-transfer vibronic exciton, which included charge transfer from the O atom onto the nearest Ti/Ta atom. He et al. reported that the photoluminescence in organolead trihalide perovskites is dominated by weakly localized excitons due to the presence of crystal imperfections.¹⁵³ The relationship between material dimensionality and luminescence could also be observed in oxide perovskites. STE luminescence has rarely been detected in single-crystal SrTiO₃ at room temperature due to the small binding energy and the corresponding short radiative lifetime of STEs. When it comes to NCs, an enhanced binding energy and radiative lifetime appear for STEs.¹⁵⁴ The enhanced binding energy originates from both the grain size and the dielectric confinement, leading to room-temperature luminescence. For small-sized SrTiO₃, electric force lines have more opportunities to penetrate into a surrounding medium with a relatively low dielectric constant.¹⁵⁵ Oxygen vacancies may be considered to be highly localized sensitization centres.

White-light (WL) emission could be obtained in thin films of (C₆H₁₁NH₃)₂PbBr₄ resulting from the broad-band, strongly Stokes shifted STE emission.¹⁵⁶ Hu et al. showed that the broad-band Stokes shifted emission in the 2D hybrid perovskite (N-MEDA)[PbBr₄] (N-MEDA = N¹-methylethane-1,2-diammonium) originated from a photogenerated energy distribution of STE states. Almost no potential barrier exists for the transition from FE to STE due to strong electron–phonon coupling, enabling ultrafast formation of the STE states on a femtosecond timescale.¹⁵⁷ Cortecchia et al. conducted a combined, systematic spectroscopic and computational study of the WL emission properties of the layered organic–inorganic perovskites (EDBE)PbCl₄ and (EDBE)PbBr₄. Due to strong Coulomb interactions, the formation of Pb³⁺ and X⁻ (where X = Cl or Br) species were confined within the inorganic perovskite framework, forming self-trapped polaron–excitons.¹⁵⁸ 0D perovskites structurally impose carrier localization and result in the formation of localized Frenkel excitons. In 0D perovskite-derived Cs₄SnBr₆, the substitution of Cs⁺ by Rb⁺ or K⁺ results in a blue-shift of the emission. For 25% substitution, Rb⁺ and K⁺ shift the PL peak from 540 nm to 519 and 500 nm, respectively.⁶⁹ This phenomeon could also be understood from the viewpoint of the inductive effect, which states that cations occupying an A site with a smaller electronegativity are in favour of STE. However, semiconductors generally suffer from severe luminescence quenching due to an insufficient confinement of excitons (bound electron–hole pairs). Sn-Triggered extrinsic self-trapping of excitons in the bulk 2D perovskite crystal PEA₂PbI₄ (PEA = phenylethylammonium) has the ability to improve the luminescence, as reported by Yu et al.¹⁵⁹ However, STE never occurs in the pure state without Sn. The isoelectronic Sn dopants initiate the localization of excitons by inducing a large lattice deformation around the impurities for STE accomodation. The STE luminescence in Sn-doped perovskites generates a broad-band red to near-infrared (NIR) emission at room temperature.

Perovskite-like niobates and tantalates, such as KNbO₃, KTaO₃, Sr₂Nb₂O₇ and Sr₂Ta₂O₇, have corner-sharing NbO₆ or TaO₆ octahedra. They show non-efficient luminescence, which is fully quenched at room temperature. Blasse and Brixner argued that the luminescence originates from self-trapped exciton recombination on NbO₆ or TaO₆ octahedra, because corner-sharing octahedra are favourable for energy-band formation, i.e. electronic delocalization.¹⁶⁰ From the viewpoint of the crystal structure, it is the angles of the M–O–M bonds that are important for delocalization.¹³ Similar results were also observed in niobates MNbO₃ (M = Li, Na, K),¹⁶¹ tantalates MTaO₃ (M = Li, Na, K)⁹⁶ and perovskite-derived Ba₅Ta₄O₁₅ and Ba₅Nb₄O₁₅.⁹⁴ For lanthanide metal ions-doped K₂La₂Ti₃O₁₀, the impurity luminescence could be observed by the host excitation. The energy transfer from the host to the rare-earth ions included both resonant energy transfer and a hole trapping mechanism. Moreover, the persistent luminescence and thermoluminescence observed in Tb³⁺- and Pr³⁺-doped K₂La₂Ti₃O₁₀ indicated a hole trapping process accompanying the valency change of the Tb and Pr ions.¹⁰¹

3.4. Luminescence from defects and impurities

Usually defects and impurities are detrimiental to luminescence as they can act as quenching centres. However, they could serve as luminescent centres by providing sub-band-gap states in the forbidden gap and induce radiative transitions from those levels.

3.4.1. Defect luminescence. First principle DFT calculations have revealed that neutral oxygen vacancies in ABO₃ perovskites are quite shallow defects with an energy level located 0.5–1 eV below the conduction band bottom.¹⁶² On perovskite surfaces, oxygen vacancies become even more shallow defects. The chemical composition of the perovskite affects the defects’ properties. In the zirconates PbZrO₃, oxygen vacancies represent considerably deeper defects, both in the bulk and on the surface. The single-electron centre in SrTiO₃ is a deeper defect compared to neutral vacancies, and reveals a larger segregation energy towards the surface. The strong localization of an electron inside a vacancy, at least in its ground state, is called an F-type centre (colour centre). Trapped F centres are responsible for the luminescence of a large series of perovskite compounds, including NCs.¹⁶³ AT the same time, a hole could also be trapped. Because of the net negative charge of the lattice volume surrounding the activator, there is a large cross-section for the trapping of holes, which may be accompanied by transitions leading to luminescence.¹⁶⁴

Bode and van Oosterhout noticed the defect luminescence in the ordered perovskite A₂BWO₆(Ba₂MgWO₆, Ba₂CaWO₆), which showed two different emission bands.¹⁶⁵ Macke ascribed the two emission bands in the ordered perovskite La₂MgSn_1−xTi_xO₆ to a regular titanate centre and a defect centre.¹⁶⁶ Energy transfer from the regular to the defect centre was also observed. By comparing the luminescent properties between the undiluted titanate and the titanate with tin, it was found that in La₂MgTiO₆ defect luminescence dominated. Kobayashi et al. reported to defect luminescence in CsPbCl₃. They observed two emission peaks, with a fast narrow band at 415 nm close to the band gap and a slower broad one at 600 nm, which suggested a defect origin.¹⁶⁷ An interesting NIR emission peak at 930 nm was observed for Fe doped in SrSnO₃ by Muralidharan et al., originating from the defective states of oxygen vacancies.¹⁶⁸

Chirvony et al. found the dual effects of traps in methylammonium lead bromide perovskite NCs. Although they found that a nonradiative deactivation of the charge carrier occurred at traps, the longer (up to microseconds) luminescent decay components revealed that the traps also acted as a carrier reservoir, resulting from the rapid reversible multiple trapping and detrapping of carriers. The dark states (traps) and bright excitonic states were in dynamic equilibrium, which resulted in long lifetime luminescence.¹⁶⁹

3.4.2. Luminescence from lanthanide ions. Lanthanide ions are widely used as luminescent centres. To deeply understand the optical properties of lanthanide-doped materials, knowledge on their electronic structure is necessary. The luminescent properties of lanthanide ions usually involve two electronic transitions: intra-4fⁿ (f–f) transition and 4fⁿ–4fⁿ⁻¹5d (f–d) transition. Since the 4f shell is shielded from the outer 5s and 5p shells, the excitation/emission wavelength from the f–f transition is relatively insensitive to the host. The f–d transition is a parity-allowed electric dipole transition, and therefore, high quantum efficiencies of absorption and emission could be achieved. The absorption/emission wavelengths can be well tuned by the host due to the strong interaction of the 5d-electron with the ligand anions. The energy levels of the trivalent lanthanide ions in LaCl₃ were determined by Dieke et al.,¹⁷⁰ and we refer to this as a “Dieke diagram”.

Recently, Dorenbos systematically studied how the lanthanide ion levels change with the chemical composition and structure of inorganic compounds.¹²⁹ On the basis of the charge-transfer model and the chemical shift, the host referred binding energy schemes (HRBE) and vacuum referred binding energy schemes (VRBE) can be constructed.

These two schemes are often called “Dorenbos diagrams”, and can well account for the optical properties, such as for lanthanide-doped LaAlO₃ and Pr³⁺-doped (Ca,Ti)_1−x[Na,Nb]_xO₃ perovskite compounds,^171,172 as shown in Fig. 4(a) and (b).

Fig. 4 (a) VRBE scheme for LaAlO₃. Arrow (1) indicates the energy of the CT-band maximum for Ce⁴⁺ and arrow (2) for Eu³⁺. (b) HRBE and VRBE schemes of Ln³⁺-doped (Ca_1−xNa_x) [Ti_1−xNb_x] O₃. The 4f ground states (Ln³⁺:4f) are labelled by the black inverted triangle and connected by solid curves. (E₁: energy of the electron trap depth, E₂: energy of O2–Ti⁴⁺/Nb⁵⁺ CT, E₃: energy of IVCT, E₄: electron transition energy from the top of the valence band to ³H₄, E₅: the band energy. (c) Stacked VRBE schemes for the acceptor levels of Mn^4+/3+, Fe^3+/2+ and Cr^3+/2+ in different phosphors. The valence and conduction bands are represented by the bottom and top bars, respectively. The solid data point is the VRBE of those acceptor levels in a specific compound. Horizontal dashed line denotes the average VRBE for those acceptor levels. (d) VRBE of Cr^3+/2+, Fe^3+/2+ and Mn^4+/3+ as a function of the VBM of different aluminates with octahedral sites. Panels adapted from: a, ref. 171, ©IOP Publishing. Reproduced with permission. All rights reserved; b, reprinted from ref. 172, Copyright (2017), with permission from Elsevier; c and d, reproduced from ref. 173 with permission from The Royal Society of Chemistry.

Perovskite-type compounds can act as hosts to accommodate a large variety of impurities, including transition metal elements, lanthanide (rare-earth) elements, trivalent bismuth,⁸⁴ divalent lead⁸⁸ and hexavalent uranium.¹⁰⁴ An extensive cation substitution is allowed in the high-symmetry perovskite or low-symmetry derived perovskite structures. Luminescence from impurities arises from a more local excitation and is therefore less sensitive to the overall dimensionality.

For example, the luminescent peak of CH₃NH₃EuI₃ has approximately the same wavelength as that of CsEuI₃. Consequently, the emission peaks for the Eu²⁺ family of 3D perovskites and the 2Dlayered system occur at very similar wavelengths.¹⁰⁷ The luminescence could be induced by direct transition from the ground to excited state, and from energy-transfer and charge-transfer processes. The undoped LaInO₃ gives a weak green luminescence resulting from In³⁺ acting as a luminescent centre via charge-transfer transition.⁸⁴ Hair and Blasse found green emission with vibronic lines of the U⁶⁺ ion in the ordered perovskites Ba₂MgWO₆ and Ba₂CaTeO₆.⁸⁵ They ascribed the excitation and emission bands to charge-transfer transitions. (C₄H₉NH₃)₂EuI₄ was the first example of a layered organic–inorganic perovskite with a divalent rare-earth metal in the perovskite sheets.¹⁰⁷ It produces intense blue photoluminescence at room temperature, with a peak wavelength of 460 nm, arising from a more localized excitation between the Eu²⁺ ground state, 4f⁷, and the 4f⁶5d¹ configurations other than the radiative decay of mobile Wannier excitons to produce luminescence. The transitions between the crystal-field levels of transition metal elements display a complex structure composed of zero phonon lines (ZPLs) and broad phonon sidebands caused by electron–phonon interactions.¹⁷⁴ However, Rodriguez et al. argued that it is the vibronic sidebands related to phonons rather than the vibrational local modes of the localized centre that have the most affect.¹⁷⁵ In the emission spectra of KMgF₃:Mn²⁺ and KZnF₃:Mn²⁺, sharp ZPLs were observed on the high energy side at 581 nm and 571 nm, respectively. A nephelauxetic effect results in the centroid shift of excited states for luminescence centres. By increasing the degree of covalency, a spectral shift towards long-wavelength is observed from CaZrO₃:Bi³⁺ to LaInO₃:Bi³⁺.⁸⁴ The choice of compositional component has a large affect on perovskite luminescence. In ABF₃:Eu²⁺ (A = Na, K, Rb, Cs; B = Mg, Ca, Sr), both 5d–4f wide-band emission and 4f–4f sharp-line emission co-exist when B = Mg, but only wide emission exists when B = Ca.⁸⁹ The reason for this lies in the relative energy position of the lowest excited 4f and the lowest 5d levels, which is supposed to be determined by the crystal-field strength. The crystal field of Eu²⁺ sites is very weak in AMgF₃, and consequently the lowest 5d level of Eu²⁺ is located at higher energy, which means 5d–4f band-emission occurs at short wavelengths. This is thus favourable for the occurrence of 4f–4f sharp-line emission.

Charge-transfer transitions have been widely observed in closed-shell transition metal-^176,177 and trivalent lanthanides (Eu³⁺, Sm³⁺, Tm³⁺, Yb³⁺)-doped perovskites as the first absorption band, in contrast to Ce³⁺, Pr³⁺ and Tb³⁺, for which a 4f–5d transition acts as the first band.^178,179 Generally the bandwidth of the charge-transfer band is twice as large as that of the f–d band.¹⁸⁰ Meanwhile, lanthanide luminescence could be implemented by an energy transfer from the charge transfer band, such as the luminescence enhancement of Nd³⁺ or Ho³⁺ by combination with UO₂²⁺.¹⁸¹ The host absorption in the Eu³⁺-doped ionic conductor KGdTiO₄ is mainly ascribed to the charge transition from the O-2p to Ti-3d states.¹⁸² In Yb³⁺ (4f¹³), the excited 4f state, ²F_5/2, is located 10 000 cm⁻¹ above the ground state ²F_7/2. Charge-transfer luminescence is widely reported because of the large energy difference between the charge transfer state and the highest excited 4f state.¹⁸³ When Yb³⁺ is incorporated in a lattice at a larger cationic site, the relaxation in the excited charge-transfer state is larger and therefore the Stokes shift is larger.

The unusual luminescence in Sr₂CeO₄ originates from a ligand-to-metal Ce⁴⁺ charge transfer,⁹⁸ not the isolated valence transitions, since the tetravalent state of cerium usually shows no luminescence. The excitation and emission spectra displayed broad maxima at 310 and 485 nm, respectively, and had a lifetime of 51 μs, which is uncharacteristically long compared to Ce³⁺ excited states.¹⁸⁴ Danielson et al. confirmed by electron spin resonance and magnetic susceptibility that no significant amount of Ce³⁺ was present in the synthesized SrCeO₄. The crystal structure here consisted of linear chains of edge-sharing CeO₆ octahedra parallel to the c axis, and it was this low-D structure with terminal O ligands that was crucial to the observation of luminescence in Sr₂CeO₄. The O atoms in the equatorial plane were shared by two adjacent Ce⁴⁺ centres by edge sharing, with the two remaining terminal O atoms bonded to only one Ce⁴⁺. The highly ionized Ce⁴⁺ and the electron-rich O atom made it possible to facilitate the excited state based on O²⁻ to Ce⁴⁺ ligand-to-metal charge transfer. On the contrary, in the 3D perovskite SrCeO₃, for which a 3D network was formed by corner-sharing octahedra, no luminescence could be observed. It is possible that the low-D structure stabilized the exciton created by the charge-transfer process.

3.4.3. Luminescence from transition-metal ions. Transition metal ions have also played an important role in luminescent materials. They have incomplete 3d shells, and therefore have a number of low-lying energy levels, which may lead to optical transitions. Because 3d electrons are outside the ion core, they strongly interact with neigbouring anions, and consequently, crystal field effects have a key influence on their energy levels. “Tanabe–Sugano diagrams” are usually adopted to describe 3d energy levels.¹⁸⁵

The 4d and 5 d electrons are less tightly bound to the core ion in comparison with the 3d electrons, and therefore charge-transfer transitions take place much easily. This transition is a parity-allowed electric dipole transition, and usually causes strong and broad bands.

Recently Qu and Dorenbos et al. conducted research to predict the location of the defect levels induced by 3d transition metal ions at octahedral sites of aluminate phosphors.¹⁷³ Su et al. gave a brief overview of the crystal field calculations and DFT-based techniques to provide a complementary picture of the electronic structure and optical properties of transition metal- and lanthanide-doped materials, and showed that it is possible to locate the lowest state and all excited state energy levels of an impurity in the host band gap.¹⁸⁶

3.5. Quantum dots and wells

With the decrease in material dimensionality or crystal structure dimensionality, the quantum confinement effect is enhanced, leading to quantum dot and well effects, particularly when at least one dimension is less than about 10a, where a is the Bohr radius of the exciton in the equivalent bulk material. Pronounced high-energy shifts have been observed as a consequence of the spatial confinement of the exciton motion in aggregates, with respect to the direct excitonic transition in the equivalent bulk material.

3.5.1. Quantum dots. Quantum dot NCs have been widely investigated and applied as colour converters to fabricate white LEDs for display backlights.¹⁹⁰ In 1995, Nikl et al. observed an intense narrow emission band peaking at 420 nm with an extremely short lifetime (20–40 ps at 10 K) in CsCl:Pb crystals, which was ascribed to the emission of CsPbCl₃ quantum dots.¹⁰⁵ The CsPbCl₃ phase aggregates were small enough to confine noticeably the motion of the Wannier exciton in the Pb²⁺ sublattice. Then, Nikl et al. reported the emission spectra of CsPbX quantum dots in a CsX host (X = Cl, Br)¹⁹¹ NCs of OIHP MAPbX₃ (X = Br, I, Cl) were synthesized by Zhang et al.¹⁹² Highly improved PLQY originated from the increased exciton binding energy due to the size reduction from the micrometre-sized bulk particles to small-sized MAPbBr₃ quantum dots (average diameter of 3.3 nm). All-inorganic cesium lead halide perovskites (CsPbX₃, X = Cl, Br, and I or mixed halide systems) usually can be synthesized as monodisperse colloidal nanocubes (4–15 nm edge lengths). An excitonic emission peak centred at 2.98 eV was observed for CsPbCl₃, whereas CsPbBr₃ exhibited bound excitonic luminescence peaks located at 2.29 eV.¹⁹³ All-inorganic halide perovskite nanocrystals have great tolerance for defects, for surface dangling bonds and for intrinsic point defects, such as vacancies not forming midgap states, which are known to trap carriers and thereby quench photoluminescence. Ramade et al. investigated the exciton fine structure of CsPbBr₃ nanocrystals by high-resolution photoluminescence spectroscopy. It is interesting to note that the reported spin–orbit coupling term was 1.20 eV, the crystal field term was −0.34 eV and the electron–hole exchange energy was 3 meV.¹⁹⁴ Owing to the quantum confinement effect, the emission could be tuned to cover the whole spectral range from 410–700 nm through compositional modulations,^188,195 as shown in Fig. 5. Through fast internanocrystal anion-exchange, uniform CsPb(Cl/Br)₃ or CsPb(Br/I)₃ compositions could be synthesized simply by mixing appropriate ratios of CsPbCl₃, CsPbBr₃ and CsPbI₃ NCs.¹⁹⁶ However, the internal structure of the anion-exchanged perovskite NCs had significant inhomogeneity in their composition. For instance, the surface of CsPb(Br/I)₃ NC was rich with exchanged iodide ions, whereas the core was rich with native bromide ions. Even after an assumed complete anion-exchange reaction, the obtained CsPbI₃ NCs showed a small amount of bromide ions in the core.¹⁹⁷ The high defects tolerance of halide perovskites allows for enhanced excitonic luminescence.¹⁹⁸ However, these NCs are known to be associated with luminescent blinking, which is assigned to the random charging/discharging driven by photoassisted ionization.¹⁹⁹ The radiative lifetime of CsPbBr₃ NCs is greatly shortened at both room temperature and cryogenic temperature, which is in favour of extremely fast single photon outputs.²⁰⁰ Swarnkar found that the spectral width of a single nanocrystal is the same as that of an ensemble.²⁰¹ Dong et al. reported the precise size control of CsPbX₃ NCs with high ensemble uniformity by varying only the Br : Pb ratio in the reactant and the reaction temperature.²⁰² With the assistance of a fatty acid-capped precursor, He et al. managed to obtain a controlled morphology of CsPbBr₃ NC nanowires to nanoplates, with the thickness varying from 5–7 monolayers.²⁰³

Fig. 5 (a) Transmission electron microscopy (TEM) and high-resolution TEM images of CsPbX₃ perovskite quantum dots. The scale bars represent 100 nm and 5 nm from left to right, respectively. (b) PL spectra (λ = 360 nm) of CsPbX₃ perovskite quantum dots. (c) Size-dependent PL spectra of monodisperse perovskite CsPbBr₃ quantum dots and composition-tunable PL spectra of perovskite CsPbX₃ quantum dots by adding different halides. (d) The corresponding sample of perovskite CsPbX₃ quantum dots. (e) PL spectra of FAPbX₃ NCs prepared by a grinding method and supersaturated recrystallization route. (f) Corresponding luminescence photographs of FAPbX₃ NCs under irradiation of a UV (365 nm) lamp. (g) Schematic illustration of solution-processed perovskite LEDs with a multilayered structure of Al/n-ZnO NPs/CsPbBr₃ QDs/p-NiO/ITO. Coating solutions of C₁₀H₁₄NiO₄ in acetonitrile, CsPbBr₃ QDs in hexane and ZnO NPs in chlorobenzene. EL spectra measured at different voltages, together with a typical emission photograph of the LED with an active area of 2 × 2 mm² (at 5.0 V). (h) Illustration of a multilayer perovskite QLED device. Left: The device structure. Right: Cross-sectional TEM image showing the multiple layers of the material with a distinct contrast. Panels adapted from: a and b, ref. 187; c,d and h, reprinted with permission from ref. 188. Copyright (2015) John Wiley & Sons, Inc.; e and f, reprinted with permission from ref. 77. Copyright (2019) American Chemical Society. g, reprinted with permission from ref. 189. Copyright (2018) American Chemical Society.

The Cs₄PbBr₆ NCs have aroused debate over their luminescent property.²⁰⁴ They have a 0D crystal structure compared to the 3D structure of CsPbBr₃. Their emission colour could also be tuned through the visible-light spectral region through halogen composition modulation. The narrow line-width luminescence originates from exciton recombination confined in the [PbBr₆]⁴⁻ octahedra, with a large exciton binding energy of 222 meV.²⁰⁵ Although both Cs₄PbBr₆ and CsPbBr₃ produce remarkably intense green luminescence, a much longer lifetime is observed in Cs₄PbBr₆.²⁰⁶ Chen reported that the luminescence of CsPbBr₃/Cs₄PbBr₆ composite originates from CsPbBr₃ NCs.²⁰⁷ Lian also ascribed Cs₄PbBr₆ to be optically inactive in a CsPbBr₃/Cs₄PbBr₆ composite.²⁰⁸ Riesen et al. concluded that the green emission from Cs₄PbBr₆ is due to nanocrystalline CsPbBr₃ impurities, as assessed by undertaking cathodoluminescence imaging, which clearly showed the presence of small crystals, with emission peaking at 520 nm, embedded in/or between larger crystallites of Cs₄PbBr₆.²⁰⁹ Zou et al. changed the non-luminescent Cs₄PbBr₆ to blue-emitting NCs by incorporating Sn cations.²¹⁰ On the other hand, some researchers think the photoluminescence in Cs₄PbBr₆ is independent of the presence of CsPbBr₃ NCs.²¹¹ Zhang reported the tunable wavelength from 340 to 378 nm in the 0D perovskite Cs₄PbX₆.²¹² Adhikari confirmed the intrinsic luminescence nature of the Cs₄PbBr₆ crystals by varying the amount of the Cs-oleate precursor to convert CsPbBr₃ with a strong blue emission (462 nm) to lead-depleted Cs₄PbBr₆ crystals with a green (529 nm) emission.²¹³ Yin et al. reported that bromide vacancies in Cs₄PbBr₆ with a low formation energy contribute to a relevant defect level in the midgap radiative state. The purity of the Br-deficient green-emissive Cs₄PbBr₆ NCs was confirmed by atomic-resolution electron imaging, which at the same time excluded the presence of CsPbBr₃.²¹⁴

Lead-free perovskite NCs have been developed with an aim to avoid the toxicity of Pb. Cs₃Bi₂X₉ (X = Cl, Br, I) NCs were synthesized with the emission wavelength ranging from 400 to 560 nm.^215,216 Non-toxic Cs₃Bi₂I₉ and Rb₃Bi₂I₉ were reported by Pa et al.²¹⁷ Cs₃BiBr₆ has a crystal structure of isolated BiBr₆ polyhedra forming a 0D halide perovskite.²¹⁸ Xie et al. synthesized Rb₇Bi₃Cl₁₆ NCs with a bright blue emission peaking at 437 nm.²¹⁹ Men et al. synthesized CsGeX₃ (X = I, Br) perovskite NCs and incorporated 16% Mn²⁺ into the nanosamples.²²⁰ The red luminescence of Cs₂InBr₅·H₂O originates from the self-trapping excitons. It is remarkable that a switchable dual emission is observed during the in situ transformation between hydrated Cs₂InBr₅·H₂O and the dehydrated mixture, which can be exploited as a water-sensor.²²¹ Halide perovskite-derived compounds Rb₂TeX₆ (X = Cl, Br, and I) have also been reported.²²²

3.5.2. Quantum wells. Quantum well structures are 2D systems with unique optical characteristics. Their photoluminescence could be tuned through control of the thickness and metal/halogen content of the perovskite sheets and the chemical/physical properties of the organic cations (e.g. length, shape and polarizability). Their most prominent feature is the exciton stability. An exciton state in a quantum well structure could be clearly resolved from the interband absorption, whereas it is hardly recognizable in bulk crystals.²²³ Meanwhile, the binding energy is increased by four times that in the 3D case.¹³² In accordance with the 2D shrinkage of the exciton wave function in quantum wells, the oscillator strength becomes enhanced.²²⁴ Further, Keldysh²²⁵ pointed out that the exciton binding energy in a thin well layer can be enlarged when surrounded by a medium with a smaller dielectric constant. The reason for this lies in the weakened screening by the Coulomb interaction between the electron and hole by virtue of the smaller dielectric constant, which makes them combine more firmly. This effect is also referred to as “dielectric confinement”, which is fully illustrated in (C_nH_2n+1NH₃)₂MX₄, n = 1,2,…, 18; M = Cd, Mn, Fe,Cu,…; X = Cl; Br with a layer-type perovskite structure.⁵⁶ The MX₄ layer is sandwiched by alkylammonium layers with a wide band gap,⁵⁵ which are natural quantum wells. The electronic transition happens within the inorganic layer, because the alkylamine is transparent in the visible region. Ishihara et al. found that the lowest exciton in Cn–PbI₄ is stable even at room temperature, with a binding energy of about 320 meV due to the dielectric confinement effect in the 2D PbI₄ well layers.⁵⁶ This was further demonstrated by Hong et al. through varying the dielectric constant of the barrier material in PbI₄-based layered structures.²²⁶ They also found from the temperature-dependent photoluminescent spectra that the luminescence efficiency remained constant up to T = 250 K, and dropped approximately as exp(E_a/k_bT), with the thermal activation energy (E_a) in agreement with the exciton binding energy. The lowest exciton energy red-shifts as the exciton is more confined in 3D, 2D (two layers), 2D (one layer) and 0D networks of PbI₆ octahedra, which may be ascribed to a transfer restriction.⁵³ This is another way of describing the quantum confinement effect. The larger exciton binding energy in (C₄H₉NH₃)₂PbBr₄ than that in (C₆H₁₃NH₃)₂PbI₄ could be well explained by the fact that the dielectric constant of the well layer in (C₄H₉NH₃)₂PbBr₄ is smaller than that in (C₆H₁₃NH₃)₂PbI₄.¹⁰⁶ By synthetically manipulating the thickness of the inorganic layer of the 2D (RNH₃)₂(CH₃NH₃)_n−1Pb_nBr_3n+1 structure, different quantum size confinement effects could be realized with a red-shift of the emission from deep blue to bright green.²²⁷ Shang et al. reported a nanocomposite material with homogeneously nanoscale 2D PEA₂Cs_n−1Pb_nBr_3n+1 perovskite quantum wells separated with inorganic crystalline Cs₄PbBr₆. The inclusion of the inorganic matrix improved the stability of an LED device.²²⁸ Gong et al. investigated the effect of electron–phonon interactions on the luminescence of single crystals of 2D perovskites, showing that reducing these interactions can lead to a bright blue emission in such 2D perovskites.²²⁹

3.6. Luminescence efficiency and stability

The luminescence efficiency and stablility of perovskite materials play crucial roles in their applications. Non-radiative transition causes significant energy loss. As the temperature increases, the probility of phonon-assisted non-radiative transitions intensifies through a cross-relaxation of the excited/ground state potential curves (configurational coordinate model), energy migration between defects and luminescent centres and thermal ionization in the conduction band. In addition, some perovskite materials, especially lead-containing halides and quantum dots, are chemically unstable. In section 4.5, we summarize the strategies many researchers adopt to improve the luminescence efficiency and stability.

4. Design, synthesis and properties

Previous works have shown that the PLQYs of 1D perovskites are generally higher than those of 2D perovskites, which in turn are generally higher than those of 3D perovskites (Table 2).^37,230 The light emission of 3D perovskites is the result of FE combination with a small Stokes shift, low FWHM and relatively short nanosecond lifetimes. For 3D materials, nanostructuring the large specific surface of perovskite NC increases the likelihood of surface defects, and the control of the defect density is critical to improve the luminescence properties.²³¹

Table 2 Structure features of standard perovskites, low-dimensional perovskites and perovskite-like halides and their photoluminescence (PL) properties

Compounds	Abbr.*	Structure Feature^†	Dimen.	PL**	Best EX (nm)	FWHM	CRI	PLQY (%)	Ref.
†: FSO = face-sharing octahedra; ESO = edge-sharing octahedra; CSO = corner-sharing octahedral; for <110>-oriented perovskite, depending on where the ripple occurs in the corrugated layer, the structures can be defined as “nד, where n represents the number of octahedra making half of the roof. *: N-MEDA = N1-methylethane-1,2-diammonium; EDBE = 2,2′(ethylenedioxy)bis(ethylammonium); DMEN = 2-(dimethylamino)ethylami; EA = ethyl ammonium; 3Apr = 3aminopyrrolidine; BAPP = 1,4-bis3-aminopropyl; OCTAm = octylammonium; NBT = n-butylammonium; 4amp = 4-(aminomethyl)piperidine; epz = 1-ethylpiperazine; PEA = phenylethylammonium; AMPS = 3,3 0-diaminodiphenyl sulfone; DABCO = 1,4-diazabicyclo[2.2.3]octane; Et = ethyl; tmpa = trimethylphenylammonium; AQ = 3-aminoquinoline; HMTA = hexamethylenetetramine; 2,6-dmpz = 2,6-dimethylpiperazine; hep = heptamethylenimine; mpz = 1-methylpiperazine; 1,4-bbdms = disulfonium cation (CH3)2S(CH2)4S(CH3)2^2+; tms = trimethylsulfonium; ABT = 2-aminobenzothiazole; TDMP = trans-2,5-dimethylpiperazine. **: N stands for narrow-band emission; B stands for broad-band emission.
Standard perovskites
CsZnCl2I	—	Perovskite	3D	432 nm	325	1.12 eV	—	—	234
2D perovskites
(N-MEDA)PbBr4	N-MEDA	<110>, 2 × 2	2D	WL, 420 nm, 558 nm (B)	380	165 nm	82	0.5	235
(EDBE)PbCl4	EDBE	<100>	2D	538 nm (B), 358 nm (N)	310	208 nm	81	2	236
(EDBE)PbBr4-1#	EDBE	<110>, 2 × 2	2D	WL, 573 nm (B), 410 nm (N)	365	215 nm	84	9	236
(EDBE)PbBr4-2#	EDBE	<110>, 2 × 2	2D	523 nm (B)	382	171 nm		18	237
(EDBE)PbI4	EDBE	<110>, 2 × 2	2D	515 nm (N)	400	70 nm		0.5	236
(C6H11NH3)2PbBr4		<100>	2D	WL, 620 nm	325	660 meV			156
α-(DMEN)PbBr4	DMEN	<110>, 3 × 3	2D	WL, 530 nm	355	183 nm	73		63
EA4Pb3Br10−xClx (x = 9.5)	EA	<100>,	2D	WL, 465 nm	355	228 nm	83		238
C4N2H12PbCl4	3Apr	<110>, 2 × 2	2D	WL, 2.01 eV	330	702 meV	85		64
C4N2H12PbBr4	3Apr	<110>, 2 × 2	2D	WL, 2.10 eV	330	743 meV	83		64
C4N2H12PbI4	3Apr	<110>, 2 × 2	2D	WL, 2.29 eV	330	670 meV	77		64
(C6H13N3)PbBr4		<110>, 2 × 2	2D	503 nm (B), 424 nm (N)	360				60
(C6H13N3)PbCl4		<110>, 2 × 2		573 nm(B), 410 nm (N)	355	220 nm	93	<1	239
(C6H5C2H4NH3)2PbCl4		<100>	2D	∼545 nm (B)	340		84	<1	240
(C4H12N)4Pb3I4Br6		<100>, n = 3	2D	Green; 519 nm	508	60 nm	—	—	241
(BAPP)Pb2Br8	BAPP	<110>	2D	WL, 582 nm	367		87	1.5	233
(C6H11NH3)2CdBr4		<100>	2D	2.94 and 2.53 eV	325				242
(OCTAm)2SnBr4	OCTAm	<100>	2D	600 nm (B)	350	136 nm		100	243
(NBT)2PbI4	NBT	<100>	2D	517 nm		25 nm		<1	244
(N-MPDA)PbBr4		<100>	2D	433 nm	410	24 nm			235
(C6H16N2)PbBr4	4amp	<110>	2D	2.38 eV	330	420 meV	76	0.54	37
(C6H16N2)PbBr4	epz	<110>	2D	2.08 eV	330	370 meV	84	0.97	37
(C4H9NH3)2SnBr4		<100>	2D	570 nm (B)	350	0.35–0.5 eV	—	—	245
(C4H9NH3)2PbI4		<100>	2D	525 nm		22 nm			246
(C4H9NH3)2SnI4		<100>	2D	625 nm		38 nm			246
(C4H9NH3)2GeI4		<100>	2D	690 nm		180 nm			246
(CH3(CH2)3NH3)2(MA)Pb2I7		<100>	2D	2.12 eV					65
(CH3(CH2)3NH3)2(MA)2Pb3I10		<100>	2D	2.01 eV					65
(CH3(CH2)3NH3)2(MA)3Pb4I13		<100>	2D	1.90 eV					65
(C8H9NH3)2PbBr4	PEA		2D	410 nm	370	14 nm		10	247
1D and 0D perovskites
(H2O)(C6H8N3)2Pb2Br10	PzPbBr	CSO	1D	WL, 580 nm	365			∼9	248
Cs4SnBr6		Isolated oct.	0D	540 nm	340			15	69
(C12H14N2O2S)[SnCl6]·H2O	AMPS	Isolated oct.	0D	592 nm (B), 482 nm (N)	376	180 nm (592 nm)			249
(CH3NH3)3Bi2I9		Isolated oct.	0D	751 nm (B)	488				250
(C8NH12)4Bi0.57Sb0.43Br7·H2O		Isolated oct.	0D	400–850 nm		—	—	4.5	251
(C8NH12) 4BiBr7·H2O		Isolated oct.	0D	450 nm	400	—	—	0.7	251
(C4N2H14Br)4SnBr6		Isolated oct.	0D	570 nm	355	105 nm	—	95	252
(C4N2H14I)4SnI6		Isolated oct.	0D	620 nm	410	118 nm	—	75	252
(C4N2H14Br)4SnBrxI6−x (x = 3)		Isolated oct.	0D	582 nm	400	126 nm	85	85	253
(C8H12N)2[SnCl6 ]		Isolated oct.	0D	390 nm	273	—	—	—	254
Perovskite-like halides
(H2DABCO)(Pb2Cl6)	DABCO	CSO	3D	455 (N), 585 nm (B)	320		96	2.5	255
(H3O)(Et2-DABCO)8(Pb21Cl59)	DABCO; Et	ESO	3D	420 (N), 690 nm (B)	330		88	1	255
(C9H14N)4Pb3Br10	tmpa	E,CSO	2D	685 nm (B)	375	0.7 eV			256
(C5H14N2)2Pb3Br10	mpz	E,CSO	2D	2.2 eV	330	485 meV	86	0.33	37
(tms)4Pb3Br10	tms	E,CSO	2D	685 nm(B)	350	0.7 eV			256
[(CH3)4N]4Pb3Cl10		F,CSO	2D	402, 496 nm (N), 629 nm (B)	300				257
C4N2H14PbBr4		ESO	1D	WL, 475 nm	379	157 nm		18–20 (bulk)	258
								10–12% (microscale crystals)
(C9H10N2)PbCl4	AQ	ESO	1D	WL, 538 nm, 340 nm (N)					259
C5H14N2PbCl4·H2O		ESO	1D	412 nm (N), 612 nm (B)	330		93.9	1	260
(C6H13N4)3Pb2Br7	HMTA	F,CSO	1D	Yellow-WL, 580 nm	350, 380	158 nm		∼7	261
(C6H16N2)3Pb2Br10	2,6-dmpz	E,CSO	1D	585 nm (B)		325 meV	77	12.24	37
(C7H16N)PbBr3	hep	FSO	1D	1.84 eV	330	285 meV	89	0.63	37
(C6H14N)PbBr3		F,C,ESO	1D	630 nm	375	220 nm			262
(C7H12N2S)2PbBr3	ABT	ESO	1D		394	210 nm			263
((CH3)2S(CH2)4S(CH3)2)3Pb3Br12	1,4- bbdms	FSO	0D	690 nm (B)	328				256
(TDMP)PbBr4	TDMP	ESO	0D	510 nm (B)	330		75	45	233
(C9NH20)6Pb3Br12		FSO	0D	522 nm (B)	371	134 nm		12	264

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Unlike typical 3D perovskites with narrow emission spectra, 2D, 1D and 0D perovskites have larger Stokes shifts and broad-band emissions due to exciton self-trapping. Many low-D perovskites exhibit the coexistence of both FE and STE emissions due to the equilibrium created by thermal activation.²³² Petrozza et al. found that in 2D perovskites, if the inorganic layer is subject to large deformation, the formation of the V_F colour center can be observed, and its radiation attenuation eventually leads to PL broadening.²³³V_F centres here represent electron capture involving halogen vacancies. For corner-sharing PbX₆ octahedra, species such as Pb₂³⁺ or Pb₂²⁺ (reported for PbBr₂) would be difficult to form under excitation because the halogens are located between adjacent Pb. However, distortions in the perovskite layers could assist the creation of such species by shortening the Pb–Pb distances, with increasing the intensity of the resulting WL emission. Unlike 2D perovskites, no observable correlation between structural deformation and the PLQY was found for these 1D wide emissive materials. 1D perovskites exhibit stronger broad-band emission, and their PLQY is generally higher than for 2D perovskites. In 1D systems, the deformation energy is generally low and there are no or only small barriers for the excitons to be self-trapped.

4.1. Luminescence properties of halide perovskites

4.1.1. Standard halide perovskites and quantum dots. Lead-based quantum dots. In 2001, Hayashi et al. first reported that CsPbCl₃ single crystal showed two luminescent peaks:²⁶⁵ a blue luminescent band observed in the temperature range of 4–300 K, due to radiative decay of FE, and a broad red band below ∼180 K, attributed to STE emission. In 2015, Protesescu et al. and Song et al. first developed CsPbX₃ (X = Cl, Br, I) inorganic perovskite quantum dots (QDs), which exhibited an ultrahigh PLQY (∼90%), a tunable luminescence wavelength by anion exchange and a very narrow FWHM of 12–42 nm.^188,266 Since then, for their potential application in illumination and display technology, CsPbX₃ (X = Cl, Br, I) QDs have attracted extensive attention.^{19,192,267–273} A partial or complete exchange of anions in CsPbX₃ can be achieved by adjusting the halide ratio in the colloidal nanocrystal solution,^196,274 resulting in bright PL, which can be adjusted over the entire visible spectrum region (410–700 nm) while maintaining a high PLQY of 80% and narrow emission line width of 10–40 nm. In 2016, Li et al. developed a process for the fast (few seconds) supersaturation recrystallization of CsPbX₃ QDs at RT, and observed red, green and blue emission with very high PLQYs of 80%, 95% and 70%, respectively.²⁷⁵ Very recently, by adopting an in situ PbBr₆⁴⁻ octahedral passivation strategy to reduce the Br vacancies in CsPbBr₃ nanosheets, Wu et al. achieved a PLQY of 96% with a FWHM ∼ 12 nm.²⁷⁶

Doping luminescence in quantum dots. Extensive research has been conducted on Mn²⁺-doped halide perovskite luminescence, which has a very long-lifetime orange emission, originating from its d state transitions. Among halide perovskites, the large band gap material CsPbX₃ (X = Cl or Br) is an ideal host for the efficient transfer of energy.²⁷⁷ Li et al. studied Mn²⁺-doped CsPbCl₃ QDs and found that the emission intensity of Mn²⁺ could be enhanced by controlling the substitution of Zn²⁺ for Mn²⁺.²⁷⁸ In 2019, Du et al. reported that the luminescence of Mn-doped CsPbX₃ (X = Cl or Br) QDs could be tuned from 517 nm to 418 nm by precisely adjusting the ratio of PbBr₂/PbCl₂ and obtained the highest PLQY of 36.7%.²⁷⁹ CsPb_xMn_1−xCl₃ QDs were prepared by a phosphorus-free thermal implantation to replace Pb with Mn. The Mn substitution rate was as high as 46%, and the prepared QDs maintained the tetragonal crystal structure of the CsPbCl₃ host. Significantly, Mn substitution greatly increased the PLQY of CsPbCl₃ from 5% to 54%.²⁸⁰ To name just a few, there have been extensive studies performed on Mn²⁺ doping of halide perovskites in the past few years and the recent breakthroughs were summarized by Adhikari et al.²⁸¹ The photo/electroluminescence (EL) efficiency of CsPbBr₃ NCs can be improved by a simple thermal injection method for doping Ce³⁺ ions. By increasing the doping amount of Ce³⁺ in CsPbBr₃ QDs to 2.88% (where the atomic percentage of Ce is comparable with Pb), the PLQY of the CsPbBr₃ NCs reached 89%. An LED device fabricated by using Ce³⁺-doped CsPbBr₃ NCs as the light-emitting layer showed a significant improvement in EL, with an external quantum efficiency (EQE) of 1.6–4.4%, through Ce³⁺ doping.²⁸² In CsPbCl_1.5Br_1.5:Yb³⁺,Ce³⁺ NCs, a high internal luminescence quantum yield (146%) was observed by Zhou et al.²⁸³ A partial equivalent cation exchange in colloidal CsPbBr₃ NCs, resulting in doped CsPb_1−xM_xBr₃ NCs (M = Sn²⁺, Cd²⁺ and Zn²⁺; 0 < x < = 0.1) could retain the original NC shape. In addition to the small (few %) contraction of unit cells when the guest cation was incorporated, the size of the parent NC remained the same in the product. The portion of Pb′ used for M′ exchange resulted in a blue-shift in the spectrum while maintaining a high PLQY (>50%) and narrow emission.²⁸⁴

Lead-free quantum dots. For environmental reasons,²⁸⁵ researchers have developed Sb, Bi and Sb congeners of lead-containing perovskites with the discovery of CsSnX₃,⁷³ Cs₃Bi₂X₉ ²¹⁵ and Cs₃Sb₂X₉.²⁸⁶ However, their PLQY and stability are not sufficient yet for practical applications.⁷³ The tin(IV)-based Cs₂SnX₆ (X = Cl, Br, I) perovskite is stable to oxygen exposure, but its quantum efficiency is low, with the highest PLQY value of Cs₂SnI₆ QDs being ∼0.48%.²⁸⁷ The mixed halide CsZnCl₂I perovskite shows two emission bands. Together, these two peaks form a very broad-band emission, with the maximum intensity at 2.87 eV with a FWHM of 1.12 eV originating from the mixed halide ions with different energy orbitals.²³⁴

Cs₂AgInCl₆ was reported by Giustino et al. as a promising material, emitting warm-WL with a broad spectrum ranging from 400 to 800 nm.²⁸⁸ Upon 370 nm excitation, Cs₂AgInCl₆ exhibits a distinct red emission peaking at 635 nm due to photo-induced defects, but the PLQY is relatively low ∼6.7%.^289,290 Yang et al. reported bright two-colour luminescence in Cs₂AgBi_xIn_1−xCl₆ double perovskite QDs,²⁹¹ while no luminescence in the bulk Cs₂AgSb_xBi_1−xBr₆ perovskite was observed.²⁹¹ The highest PLQY of ∼86% was obtained by Luo et al. in 0.04% Bi³⁺-doped Cs₂(Ag_0.60Na_0.40)InCl₆. The authors concluded that the reasons for such a high PLQY include: 1. the introduction of Na in Cs₂AgInCl₆ breaks the parity forbidden transition of the electrons, and reduces the electronic dimension, resulting in WL emission originating from STEs; and 2. the addition of Bi³⁺ reduces the defects level, which further improves the PLQY. Furthermore, the Cs₂Ag_0.60Na_0.40InCl₆ powder was directly capsulated in the commercial ultraviolet LED chip (380–410 nm), and the CIE coordinate of the fabricated device was (0.396, 0.448) with a CCTC of 4054 K. The fabricated LED was highly stable in air, with an emission of about 5000 cd m⁻², which lasted for more than 1000 h.²⁹²

Organic–inorganic hybrid perovskite quantum dots. Zhang et al. developed the room temperature reprecipitation and microemulsion preparation method for hybrid perovskite QDs, and obtained MAPbBr₃ QDs with a PLQY up to 70%.¹⁹² Huang et al. reported that the size of MAPbBr₃ perovskite QDs could be controlled by changing the temperature at which precipitation occurs. By changing the synthesis temperature from 0 °C to 60 °C, the resulting QDs exhibited PL from 475 to 520 nm and narrow emission line widths of 28–36 nm and very high PLQYs, ranging from 74% to 93%.²⁹³ The in situ preparation of highly luminescent FAPbBr₃ nanocrystal thin films was carried out by dropping toluene as an anti-solvent during the spin-coating with a perovskite precursor solution using 3,3-diphenylpropylamine bromide (DPPA-Br) as a ligand. The obtained film was homogeneous and consisted of 5–20 nm NCs. The optimized film exhibited strong PL emission at 528 nm with a PLQY up to 78% and an average PL lifetime of 12.7 ns.²⁹⁴

4.1.2. Low-D halide perovskites. Broad-band WL emission has been observed from a variety of low-D metal halides, which tend to have a high colour rendering index (CRI), high thermal stability and low temperature solution processability, showing potential applications in solid-state lighting. 2D perovskite luminescent mechanisms mainly include FE and STE luminescence. The transition between FE and STE can be observed by low temperature spectroscopy.²⁹⁵
4.1.2.1. 2D perovskites for luminescence. A. Self-trapped exciton luminescence

Lead-containing 2D perovskites. WL emission was first observed by Li et al. in <110>-oriented (C₆H₁₃N₃)PbBr₄ (API = N-(3-aminopropyl)imidazole) in 2006 with a relatively low PLQY (<0.5%).⁶⁰ In 2014, Karunadasa's research group observed white emission in the 2D perovskite (N-MEDA)PbBr₄ (N-MEDA = N1-methylethane-1,2-diammonium), which possessed corrugated <110>-oriented inorganic layers, showing a relatively wide band gap of 3.8 eV (Fig. 6).²³⁵ The absorption spectrum of (N-MEDA)PbBr₄ showed a peak exciton band at 395 nm with a shoulder peak at 370 nm. Excited by 380 nm light, it showed wide emission spanning the entire visible spectrum with two peaks: a higher-energy shoulder peak centred at ∼420 nm and a more intense one centred at 558 nm with a broad FWHM of 165 nm. The PLQY of the broad emission ranging from 400 to 700 nm was measured to be ∼0.5%. Through chlorine doping, the bandwidth of this material was further broadened and (N-MEDA)PbBr_3.5Cl_0.5 showed a pure WL emission with CIE coordinates of (0.31, 0.36). By optimizing the chlorine-doping concentration, the highest PLQY of ∼1.5% was obtained with the composition of (N-MEDA)PbBr_2.8Cl_1.2. Furthermore, this sample was stable and even with continuous irradiation for seven days with a 365 nm 4 W lamp, it showed no material degradation or change in PL intensity.

Fig. 6 (a) Structure of <110>-oriented 2D perovskite (N-MEDA)[PbBr₄] (N-MEDA = N1-methylethane-1,2-diammonium). (b) Absorption spectrum (black line) and emission spectrum (red line) for (N-MEDA)[PbBr₄] with excitation at 380 nm; inset shows the luminescence of (N-MEDA)[PbBr₄] powders under 380 nm UV light. Reprinted with permission from ref. 235. Copyright (2014) American Chemical Society.

In the same year, Karunadasa's research group reported another family of WL-emission 2D perovskites (EDBE)PbX₄ (X = Cl, Br, I, EDBE = 2,2′(ethylenedioxy)bis(ethylammonium)). The (EDBE)PbCl₄ structure possess <100>-oriented inorganic sheets. Excited at 310 nm, (EDBE)PbCl₄ showed two peaks. The broad emission spanned the entire visible spectrum, with the maximum at 538 nm and a FWHM of 208 nm, with a less intense shoulder peak centred at 358 nm. The PLQY of (EDBE)PbCl₄ was ∼2%.²³⁶ The Br and I analogues of (EDBE)PbCl₄ also show <110>-oriented inorganic layers. Upon 365 nm excitation, the Br analogue showed a WL emission with a less identifiable shoulder peak centred at 410 nm and the broad one with the maximum at 573 nm and a FWHM of 215 nm with a PLQY of ∼9%, and a CRI of 84. The CIE chromaticity coordinates were (0.39, 0.42), and it possessed a CCT of 3990 K, corresponding to “warm” WL, which makes it suitable for indoor illumination. The origin of the WL emission was the intrinsic emission of the bulk due to STE other than from extrinsic dopants or from surface defect states. The I analogue showed an observable green emission at 515 nm with a FWHM of 70 nm under 400 nm excitation with a PLQY of less than 0.5%. In 2016, Ma's group reported that upon 365 nm UV-LED chip excitation, the (EDBE)PbBr₄ micro-crystal exhibited a PLQY of ∼18%.²³⁷ In 2015, Yangui et al. reported (C₆H₁₁NH₃)₂PbBr₄ with <100>-oriented inorganic layers. The spin-coated thin film, upon 325 nm excitation, showed a very broad WL emission spanning the entire visible light range, with a peak at 2 eV, a FWHM of 660 meV and a large Stokes shift of 1.2 eV.¹⁵⁶

In 2017, Kanatzidis's research group reported three new 2D lead bromide perovskites, and by comparing the PL emission and the structure, they found that there was a correlation between the distortion of the “PbBr₆” octahedra in the 2D layer and the broadening of the PL emission. The most distorted structure with the widest light emission was observed in the most distorted <110>-oriented 3 × 3 2D lead perovskite α-(DMEN)PbBr₄ (DMEN = 2-(dimethylamino)ethylamine). The FWHM was 183 nm and the lifetime was τ_avg = 1.39 ns. Shortly after that, the WL-emitting <100>-oriented 2D perovskite (n = 3) EA₄Pb₃Br_10−xCl_x (EA = ethyl ammonium, x = 0, 2, 4, 6, 8, 9.5 and 10) was reported by the same group (Fig. 7). This had the general formula A′₂A_n−1B_nX_3n+1 (A′ and A = cations; X = halide), with EA⁺ occupying both A′ and A sites in the system. The band gap of EA₄Pb₃Br_10−xCl_x could be monotonously tuned from 3.45 eV (x = 10) to 2.75 eV (x = 0). The light emission was adjustable with the variation in X concentration and a broad-band emission was observed in EA₄Pb₃Cl₁₀, while a narrow blue light emission was observed in EA₄Pb₃Br₁₀, with this difference being related to the distortion levels of EA₄Pb₃Cl₁₀ (large distortion) and EA₄Pb₃Br₁₀ (small distortion). Among all these mixed halide compounds, EA₄Pb₃Br_0.5Cl_9.5 showed the maximum CRI of 83 and EA₄Pb₃Cl₁₀ showed the minimum CRI of 66.²³⁸ Very recently, <110>-oriented 2D perovskites, C₄N₂H₁₂PbX₄ (X = I, Br, Cl), were reported by Kanatzidis and coworkers. The experimental band gaps follow the trend of I < Br < Cl (2.56, 3.29, 3.85 eV, respectively) and all of them showed broad-band emission.⁶⁴

Fig. 7 Structure of three-layered (a) EA₄Pb₃Cl₁₀ and (b) EA₄Pb₃Br₁₀. (c) CIE coordinates of EA₄Pb₃Br_10−xCl_x (x = 0, 2, 4, 6, 8, 9.5 and 10) in 1931 colour space chromaticity diagram. WL emissions observed from EA₄Pb₃Br_10−xCl_x (x = 2, 4, 6, 8, 9.5 and 10) and blue-light emission from EA₄Pb₃Br₁₀ (excited at 315 nm), and (d) corresponding polycrystalline samples under UV light illumination. Reprinted with permission from ref. 238. Copyright (2017) American Chemical Society.

In 2018, Wu et al. synthesized the chlorine analogue²³⁹ of (C₆H₁₃N₃)PbBr₄ (the first compound in which WL emission was observed).²⁷⁵ Upon 355 nm excitation, (C₆H₁₃N₃)PbCl₄ exhibited a broad-band emission with a maximum of 573 nm due to STE emission and a high-energy shoulder peak at 410 nm due to FE emission. The CIE chromaticity coordinates of this material were (0.36, 0.37) and the lifetime was estimated to be 4.416 ns. The CRI of 93 is ultrahigh and is among the highest CRI of broad-band WL-emitting hybrid perovskites. The measured PLQY was <1% and the sample was stable under ambient conditions.²³⁹ Thirumal first reported the WL emission of solution-processed 2D perovskite (C₆H₅C₂H₄NH₃)₂PbCl₄ (with the abbreviation of PEPC) QDs. Upon 340 nm excitation, the PEPC nanoparticle solution, powder, thin film and single crystal samples all showed broad emissions ranging from ∼400 to 900 nm centred at 545 nm. It was demonstrated that the PL spectra of these samples were almost unchanged with the morphology of the materials.²⁴⁰ The PLQY of PEPC was low (<1%). In 2019, Gautier reported (BAPP)Pb₂Br₈ (BAPP = 1,4-bis(3-aminopropyl) piperazine) belonging to the <110>-oriented 2D perovskite class, with a relatively weak WL emission and a PLQY of 1.5%.²³³

Lead-free 2D perovskites. Cd-Based 2D hybrid perovskite with the chemical formula of (C₆H₁₁NH₃)₂CdBr₄ exhibited WL emission under 325 nm UV light illumination, which consisted of a peak at 2.94 eV due to excitons confined in the [CdBr₄]²⁻ inorganic layer and a second peak at 2.53 eV caused by emission from the organic cations. There was a strong correlation between the structural distortion of the CdBr₆ pseudo-octahedra and the broadening characteristics of the WL emission band, and also there was possible energy transfer between the inorganic and organic moieties.²⁴² An intense PL emission was obtained in the 2D <100>-oriented (OCTAm)₂SnBr₄ (OCTAm = octylammonium), which showed a broad emission with a FWHM of ∼136 nm, centred at ∼600 nm and a PLQY of ∼100%.²⁴³ Its large Stokes shift of ∼250 nm and long-lived light emission of ∼3.3 μs were typical for an emission stemming from STE. This material showed high stability with no changes observed under normal humidity over 240 days at RT. By controlling the composition of this material, the luminescence could be tuned from yellow to dark red.²⁴³

B. Free exciton luminescence

Lead-containing 2D perovskites. Dou et al. studied the PL properties of (C₄H₉NH₃)₂PbBr₄ (NBT = n-butylammonium), which had a <100>-oriented layered structure (Fig. 8a). Both the bulk crystal and 2D sheets with different thicknesses (22, 8 and 3 inorganic layers thick) exhibited similar intense violet-blue light emission. The bulk crystal had emission at 411 nm (2.97 eV), while the 2D sheet displayed a slightly blue-shifted peak at ∼406 nm (3.01 eV). The slight increase in the optical band gap of the ultrathin 2D sheets may be caused by lattice expansion, and this was confirmed by theoretical calculations. The PLQY of the 2D sheet was ∼26%, which was much higher than that of the bulk crystal (<1%), indicating the high quality of the single-crystal 2D sheet. The PL lifetime of the 2D sheets displayed a bi-exponential feature with lifetimes of 0.78 ns (67%) and 3.3 ns (33%).¹³⁶ [(C₄H₉NH₃)₂PbI₄] films showed a narrow-band emission at 517 nm with a FWHM of 25 nm, and a small Stokes shift of 4 nm due to FE emission. The absolute PLQY of this material was less than 1%.²⁴⁴ (N-MPDA)PbBr₄ (N-MPDA = N1-methylpropane-1,3-diammonium), which also possesses <100>-oriented layers, also showed a sharp peak at 433 nm due to FE emission (Fig. 8).²³⁵

Fig. 8 (a) Crystal structure of (C₄H₉NH₃)₂PbBr₄ (NBT = n-butylammonium) consisting of <100>-oriented layers. (b) Steady-state absorption (blue) and photoluminescence (green) of (NBT)₂PbI₄. The insets show orange emission from (NBT)₂PbI₄ under UV light illumination. Reprinted with permission from ref. 136. Copyright (2014) American Chemical Society. (c) Crystal structure of the <100>-oriented (N-MPDA)[PbBr₄] (N-MPDA = N-methylpropane-1,3-diammonium). (d) Absorption spectrum (black) and emission spectrum (blue) under excitation at 400 nm for (N-MPDA)[PbBr₄]. The inset shows the luminescence from powders of (N-MPDA)[PbBr₄] under 380 nm irradiation. Reprinted with permission from ref. 235. Copyright (2014) American Chemical Society.

(C₄H₁₂N)₄Pb₃I₄Br₆ is a <100>-oriented 2D perovskite with n = 3. Optical transmission measurements on (C₄H₁₂N)₄Pb₃I₄Br₆ film showed two absorption bands centred at 474 and 508 nm. This compound showed a strong green luminescence emission centred at ∼519 nm, with a FWHM of ∼60 nm. The Stokes shift between the absorption at 508 nm and emission at 519 nm was quite small at ∼11 nm. Thus, the luminescence originated from FE.²⁴¹

Lead-free 2D perovskites. (C₄H₉NH₃)₂EuI₄, reported by Mitzi et al., is the only known example of an f-block metal-based 2D hybrid perovskite, and it displayed strong blue PL centred at 460 nm.

4.1.2.2. 1D and 0D perovskites. Lead-containing 1D perovskites: The structure of (H₂O)(C₆H₈N₃)₂Pb₂Br₁₀ (abbreviated to PzPbBr) consists of twisted corner-sharing metal halide octahedra forming 1D twisted chains two octahedral units wide. It exhibited intense yellow-white emission under 365 nm excitation. The optical characterization of PzPbBr showed a strong Stokes shift and a broad PL emission band peaking at 580 nm with a PLQY of ∼9%.²⁹⁶

Lead-free 0D perovskites: Cs₄SnBr₆ featuring a 0D structure with separated SnBr₆ octahedra exhibits broad PL with a central wavelength of 540 nm and a PLQY of 15.5%, which is attributed to STE emission. The Stokes shift and the STE emission band can be tuned from 500 nm to 620 nm by cationic or anionic mixing following the general formula Cs_4−xA_xSn(Br_1−yI_y)₆ (A = Rb, K; x ≤ 1, y ≤ 1).⁶⁹ The structure of the tin chloride of 3,3 0-diaminodiphenyl sulfone (C₁₂H₁₄N₂O₂S)[SnCl₆]H₂O (abbreviated as (AMPS)[SnCl₆]H₂O) consists of isolated [SnCl₆] octahedra surrounded by organic (AMPS) cations. It showed strong quasi-WL emission that could be observed by the naked eye, even at RT. Wherein (AMPS) molecules act as donors and [SnCl₆] molecules act as receptors, the PL spectrum consists mainly of two bands, a wide and strong band centred at 592 nm and a shoulder peak with the maximum at 482 nm. The strong peak at 592 nm was associated with light-induced exciton formed in the inorganic SnCl₆ octahedra, while the shoulder peak was designated as π–π* transition in the AMPS organic cation.²⁴⁹ The 0D mixed metal halide perovskite (C₈NH₁₂)₄Bi_0.57Sb_0.43Br₇·H₂O contains [BiBr₆]₃ and [SbBr₆]₃ octahedra and showed an experimental band gap of 2.80 eV. The crystal showed an ultra-wide band emission ranging from 400 to 850 nm, with a PLQY of 4.5%, which came from both FE and STE. Meanwhile, (C₈NH₁₂)₄BiBr₇·H₂O showed a narrow emission at 450 nm, which is close to the exciton absorption peak (400 nm), so it belongs to FE emission with a PLQY of 0.7%, demonstrating that the mixed B-site perovskite not only has good ambient stability but also good photostability.²⁵¹ A film of the 0D methylammonium iodonium (CH₃NH₃)₃Bi₂I₉ perovskite exhibited a wide band gap of 2.9 eV, and PL emission was observed at 1.65 eV (751 nm) under 488 nm optical excitation.²⁵⁰ In 2018, Ma's group reported a 0D (C₄N₂H₁₄Br)₄SnBr₆ with a PLQY of ∼95%, and (C₄N₂H₁₄I)₄SnI₆ with a PLQY of ∼75%.²⁵² By mixing with a halide, a yellow emission centered at 582 nm, and a FWHM of 126 nm at RT were observed from (C₄N₂H₁₄Br)₄SnBr_xI_6−x (x = 3). The highest PLQY oberserved was ∼85%. The colour index of a UV-pumped WLED, prepared by using this material and BaMgAl₁₀O₁₇:Eu²⁺ as a blue phosphor, was as high as ∼85.²⁵³ All these values of Sn-based OIHPs reported by Ma's group are among the highest PLQYs of STE luminescence reported so far.

4.1.3. Perovskite-like halides. Perovskite-like halides with 3D structures. The structure of (H₂DABCO)Pb₂Cl₆ (DABCO = 1,4-diazabicyclo[2.2.2]octane) consists of a unique 3D (Pb₂Cl₆)₂ backbone with Pb atoms located in a slightly distorted octahedral coordination environment. This compound showed two emission peaks centred at 455 nm (narrow peak) and 585 nm (broad peak). The 455 nm emission lifetime was 2.89 ns and the 585 nm lifetime was 18.85 ns. Adjusting the excitation from 280 nm to 320 nm resulted in a decrease in the intensity of the emission peak at 455 nm, while the intensity of the emission peak at 585 nm increased. Thus, the emission could be modulated from blue to yellow and an excitation of 300 nm was the best to produce WL emission with a PLQY of 2.5% and CIE of (0.33, 0.34). It is worth noting that the WL emission of this compound showed a CRI value of 96, which is the highest reported value for a single component compound.²⁵⁵

Perovskite-like halides with 1D structures. The structure of (C₉H₁₀N₂)PbCl₄ (3-aminoquinoline abbreviated as AQ) consists of chains of edge-sharing PbCl₆ octahedra extending along the b-axis. This compound showed a strong yellow WL emission, consisting of a yellow broad-band centred at 538 nm and a narrower UV band centred at 340 nm. The wide yellow band was related to the π–π* transition inside the organic molecule, while the UV emission was attributed to the Wannier excitons confined in the inorganic lines.²⁵⁹

Chains of [PbX₄²⁻]_∞ formed by edge-sharing PbX₆ octahedra with organic moieties surrounding show broad emission, such as C₄N₂H₁₄PbBr₄,²⁵⁸ C₅H₁₄N₂PbCl₄·H₂O ²⁶⁰ and (TDMP)PbBr₄ (TDMP = trans-2,5-dimethylpiperazinium).²³³ Among these, (TDMP)PbBr₄ (TDMP = trans-2,5-dimethylpiperazinium) exhibited intense WL emission with a PLQY of 45% and high CRI of 75. The emission originated from STEs through the colour centres, such as Pb₂³⁺, Pb₃⁺, X₂ and X₂⁻ (X = halide).²⁴⁴

The structure of (HMTA)₃Pb₂Br₇ (HMTA = C₆H₁₃N₄⁺) features six face-sharing metal halide dimers (Pb₂Br₉⁵⁻) joined at the corners to form a ring extending in one dimension. The inner and outer surfaces of the tube are surrounded by HMTA (HMTA)₃Pb₂Br₇, which can be excited by ultraviolet light ranging from 250 to 400 nm to produce a broad emission peak ranging from 450 to 750 nm centred at 580 nm, with a large FWHM of 158 nm. The 1D compound (2,6-dmpz)₃Pb₂Br₁₀ (2,6-dmpz = 2,6-dimethylpiperazine) with corner- and edge-sharing octahedra forming chains exhibited a band gap of 2.49 eV, and displayed a broad-band light emission centred at 2.12 eV with a FWHM of 325 meV and average life time of 23.03 ns. Its CIE was (0.44, 0.46) and the CRI was 77. It had the highest PLQY of 12%, owing to its unique structure, which allowed efficient charge carrier relaxation and light emission.³⁷

Perovskite-like halides with 0D structures. The 0D (C₉NH₂₀)₆Pb₃Br₁₂ contains face-sharing PbBr₆ trimer clusters, which are connected by the organic cations.²⁶⁴ A broad-band green PL peaking at about 522 nm with a FWHM of 134 nm was observed. The CIE chromaticity coordinates were determined to be (0.264, 0.392) and the PLQY was around 12%.²⁶⁴

Other compounds. In addition, several sulfonium Pb–Br hybrids were reported showing broad-band visible light emission, such as (tms)₄Pb₃Br₁₀ (tms = trimethylsulfonium; (CH₃)₃S⁺) and (tmpa)₄Pb₃Br₁₀ (tmpa = trimethylphenylammonium).²⁵⁶ Abid et al. reported (C₇H₁₂N₂S)₂PbBr₃ abbreviated as (ABT)₂PbBr₃. The structure consisted of an infinite double-chain constructed by edge-sharing PbBr₆ octahedra. Under UV emission, this material exhibited WL emission with an intensity that can be seen with the naked eye. Its PL spectrum was characterized by a broad emission band covering the visible spectrum, consisting of blue, green, yellow and red components at 450, 475, 530 and 580 nm, respectively. PL measurements with various excitations showed that the WL emission only occurred within a narrow excitation range of about 394 nm.²⁶³ The structure of (1,4-bbdms)₃Pb₃Br₁₂ (1,4-bbdms = (CH₃)₂S(CH₂)₄S(CH₃)₂²⁺) consists of isolated trimers and pentamers of face-sharing octahedra.²⁵⁶ It exhibited broad red PL centered at 690 nm and two higher-energy and lower-intensity PL bands at ca. 375 and 460 nm, which could be attributed to emission from FEs or defects.

4.2. Rare-earth-doped perovskites

Host materials with a perovskite structure have a considerable degree of structural flexibility to accommodate a variety of activating and sensitizing cations. Early in 1967, Yamamoto et al. started to research the luminescence of trivalent rare-earth elements (Eu³⁺, Tb³⁺ and Pr³⁺) in the perovskites SrTiO₃ and BaTiO₃.²⁹⁷ Generally, the characteristics of doped trivalent rare-earth elements are preserved in perovskite structures. For example, upon UV and low voltage cathode-ray excitation, LaInO₃:Sm³⁺, LaInO₃:Pr³⁺ and LaInO₃:Tb³⁺ phosphors prepared by a sol–gel process showed strong orange–red, blue–green and green luminescence, respectively.²⁹⁸ High-throughput combinatorial methods for synthesizing and characterizing oxide perovskites have been developed to identify and optimize interesting new materials.^299,300 The functioning of oxide perovskites can in different material dimensionalities. Meanwhile, it is interesting to note that metal oxide nanosheets could be synthesized by the exfoliation of bulk layered oxides. For example, Ida et al. succeeded in the exfoliation of Bi₂SrTa₂O₉ to obtain a mono-nanosheet of 1.3 nm thickness with blue emission.³⁰¹ By doping Eu³⁺ or Tb³⁺ into perovskite nanosheets prepared by the exfoliation of single- or double-layered perovskite oxides, such as K₂Ln₂Ti₃O₁₀, KLnNb₂O₇ and RbLnTa₂O₇ (Ln: lanthanide ion), intense red and green emission could be observed under UV illumination.³⁰² The visible emission results from energy transfer within the nanosheet, also considering the overlap between the excitation spectrum and the band gap absorbance.

Ce. Allowed 5d–4f transitions of lanthanide ions (e.g. Ce³⁺, Eu²⁺) have short decay times (<1 ms) and have been very successfully utilized for scintillators.³⁰³ Scintillators are used to detect ionizing radiation. Cerium-doped lanthanide perovskites, such as YAlO₃(YAP:Ce³⁺) and LuAlO₃(LuAP:Ce³⁺), are used as scintillators as they exhibit fast scintillation and a high light yield.^90,91,304 The excitation spectrum usually has a maximum intensity at the fundamental absorption edge comparable with that from direct 4f–5d excitation. The free excitions are able to migrate along a long distance and transfer their energy to luminescence centres. This process is very efficient in the excitation energy region, since the capture of a hole by Ce³⁺ is favourable considering the fact that cerium is more stable in the tetravalent state.³⁰⁴ YAP:Ce shows a broad emission band due to the presence of Ce ions.³⁰⁵ Five bands could be clearly distinguished in the excitation spectrum because of the complete splitting of the 5d excited state level at a low Y³⁺ site symmetry.³⁰⁶

Pr. The optical properties of Pr³⁺-doped titanates with a perovskite structure are influenced by the semiconducting nature of the host. In CaTiO₃:Pr³⁺, the excitation of the red luminescence is achieved through the conduction band states and then transferred to the emitting level of Pr³⁺.³⁰⁷ In this process, electron–hole pairs are produced via O(2p)–Ti(3d)–Pr(4f). Meanwhile, the emission enhancement of Pr³⁺ due to charge compensation depends on the compensator types, because the hole capture cross-section is strongly affected by the compensator. For host materials, the intensities of the red luminescence due to the f–f transitions of Pr³⁺ ions increase in the order of cubic SrTiO₃, tetragonal BaTiO₃ and orthorhombic CaTiO₃.³⁰⁸ The increase in f–f transition probability is caused by the point symmetry lowering, which can be controlled by either crystallographic site symmetry or the solid-solution effect at the alkaline-earth site. In the ATiO₃ (A = Pb, Ca, Ba, Sr) perovskite-type amorphous thin films, a tail is observed in the absorbance spectrum curve.³⁰⁹ Usually, the optical properties of amorphous semiconductor compounds are characterized by the presence in the plot of photon energy versus the optical absorption of a tail, in which the optical absorption falls almost asymptotically to zero. This region is normally transparent in crystalline solids.³¹⁰ The tail in the absorbance spectrum curve is caused by the delocalized electronic levels of the fivefold coordination through the displacement of O. The radiationless relaxation pathways lead to the quenching of ³P₀ emission and subsequently to the typical single red luminescence from the ¹D₂ level in CaTiO₃:Pr³⁺.³¹¹ STEs participate in the relaxation process and the Pr³⁺/Ti⁴⁺ ↔ Pr⁴⁺/Ti³⁺ charge-transfer state (CTS) is proposed as the final relaxation channel to the emitting ¹D₂ level. The difference between the luminescence properties of CaTiO₃:Pr³⁺ and CaZrO₃:Pr³⁺ could be explained by the relaxation pathway. The low-lying Pr³⁺/Ti⁴⁺ ↔ Pr⁴⁺/Ti³⁺ charge-transfer state is the final radiationless relaxation pathway to the emitting ¹D₂ level in Pr³⁺-doped CaTiO₃, instead of the low-positioned 4f–5d band. In the zirconate, the relative high stability of the tetravalent state of zirconium is not in favour of a low-lying intervalence charge-transfer state and therefore, when excited in the conduction band at 254 nm, CaZrO₃:Pr³⁺ shows prominent greenish-blue luminescence from the ³P₀ level instead of a single red emission from the ¹D₂ level. The emission profile of praseodymium-doped lanthanum hafnate, La₂Hf₂O₇:Pr³⁺, displayed the involvement of both the ³P₀ and ¹D₂ states, because the Pr³⁺ ions occupy both the Hf⁴⁺ sites and La³⁺ sites.³¹² Pan et al. noticed persistent luminescence in crystalline CaTiO₃:Pr³⁺ nanoparticles prepared from a complex polymer precursor.³¹³ The 612 nm red-emitting persistent luminescence of Ca₃Ti₂O₇:Pr³⁺ can be activated by multiple charge-transfer processes. The red afterglow can last for ∼5 min, with the trap-depth exponentially distributed at 0.69–0.92 eV.³¹⁴

Sm. The reddish-orange emission of Sm³⁺ is caused by transitions from ⁴G_5/2 to ⁶H_J (J = 5/2, 7/2, 9/2, 11/2).³¹⁵ Sm³⁺-doped Sr_n+1Sn_nO_3n+1 (n = 1, 2, ∞) is a good example to elucide the effects of the crystal structure dimensionality on the luminescent properties.¹¹⁴ The mechanoluminescence enhancement process is related to the 2D layered structure and the charge-transfer process. The intensity of mechanoluminescence for Sr₃Sn₂O₇:Sm³⁺ is three orders of magnitude higher than that for SrSnO₃:Sm³⁺ due to the effective confinement of transfer energy in the 2D layer. The band gaps of Sr₂SnO₄, Sr₃Sn₂O₇ and SrSnO₃ decrease as 4.43 eV (280 nm), 4.13 eV (300 nm) and 3.88 eV (320 nm), in accordance with the change in the crystal structure dimensionality. This means that in SrSnO₃ with a 3D connecting octahedral network, the charge-transfer process occurs more easily. The band gap of Sr₃Sn₂O₇:Sm³⁺ could be further tailored by Si or Ge substitution. The more electronegative Si and Ge makes the charge transfer occur more easily and reduces the band gap.¹⁰⁸

Eu. The luminescence of Eu³⁺ originates from the intra-configurational ⁵D₀–⁷F_J transitions. The emission spectra of CaSnO₃:Eu³⁺ and BaSnO₃:Eu³⁺ are dominated by the red ⁵D₀ → ⁷F₂ transition at 614 nm.³¹⁶ The onset of the excitation spectra of MSnO₃ (M = Ca, Sr and Ba) at 77 K occur at almost the same position as for the UV-VIS absorption spectra shown, which indicates that the emission from MSnO₃ is derived from band gap excitation.³¹⁷ The PL spectra were red-shifted from CaSnO₃ to SrSnO₃ to BaSnO₃, as observed in their diffuse reflectance spectra. Eu³⁺-Doped perovskite nanosheets of the form La_0.90Eu_0.05Nb₂O₇ were prepared by the soft chemical exfoliation reaction of K_1−xH_xLa_0.90Eu_0.05Nb₂O₇ with a tetrabuthylammonium hydroxide aqueous solution.³¹⁸ The resulting colloidal La_0.90Eu_0.05Nb₂O₇ nanosheet suspension exhibited a photoluminescence emission from the ⁵D₀ to ⁷F_J manifold transitions of Eu³⁺ by either a direct excitation of Eu³⁺ or by host excitation, whereas no host emission was observed at room temperature. In the case of the bulk precursors K_1−xH_xLa_0.90Eu_0.05Nb₂O₇, the direct excitation yields more intense emission than the host excitation. On the contrary, the most intense emission from the La_0.90Eu_0.05Nb₂O₇ nanosheets was observed by exciting at the broad excitation band maximum (353 nm). The difference in the photoluminescence properties between the La_0.90Eu_0.05Nb₂O₇ nanosheets and their bulk precursors seems to be related to the dimensionality of these host structures and the confinement of the energy-transfer process between the host layer units and the Eu³⁺ activators. Photoluminescence studies into Eu³⁺-doped double perovskites with the formula A₂CaWO₆ (A = Sr, Ba) have revealed that the forced electric dipole (ED) transition is present when Eu³⁺ is substituted at the non-centrosymmetric Sr-site. Substitution at the centrosymmetric Ca-site shows both ED and magnetic dipole (MD) transition.³¹⁹ In A- and B-site substituted double-perovskite Sr₂CaMoO₆ doped by Eu³⁺, the photoluminescence intensity of the B-site substituted Sr₂CaMoO₆ is evidently higher than that of the A-site substituted phosphor.³²⁰ In different hosts, the location of the CT bands of Eu³⁺ is also different, centred at 270, 250 and 263 nm, corresponding to CaZrO₃:1%Eu,SrZrO₃:1%Eu and BaZrO₃:1%Eu.³²¹ Powdered samples of the perovskite BaSnO₃ exhibit strong near-infrared (NIR) luminescence at room temperature following band-gap excitation at 380 nm (3.26 eV).³²² The emission spectrum is characterized by a broad band centered at 905 nm (1.4 eV), tailing on the high-energy side to approximately 760 nm. The luminescence involves a defect state. As the strontium content increases, the excitation maximum and band gap shift further into the UV range, while the intensity of the NIR emission peak decreases and is shifted further into the infrared. This combination leads to an unexpectedly large increase in the Stokes shift. The unusual NIR PL in BaSnO₃ may originate from recombination of a photogenerated valence-band hole and an occupied donor level, probably associated with a Sn²⁺ ion. It was found that the PL intensity of the phosphor NaY_0.7Eu_0.3TiO₄ is about three times higher than that of the phosphor NaGd_0.7Eu_0.3TiO₄ with the optimal composition. This may be due to the larger distortion from the doping of Eu³⁺ in the former compound NaYTiO₄ than that in the latter NaGdTiO₄ lattice, based on the large difference in ionic radii.³²³ Recently, Bala predicted the blue emission (2.48–2.85 eV) of the Eu²⁺-doped CsPbBr₃ perovskite by first-principle calculations. The results showed that Eu²⁺ doping is favourable because of the energetic stablility in the CsPbBr₃ perovskite host with negligible strain. Therefore, further research on Eu²⁺-doped all-inorganic halide perovskites is expected.³²⁴

Tb. The luminescence of Tb³⁺ could be caused by energy transfer. In the Gd³⁺–Tb³⁺-activated LaAlGe₂O₇, the decay time of Tb³⁺ emission under Gd³⁺ excitation at the ⁶I_J energy level was longer than that under direct Tb³⁺ excitation at the ⁵L₁₀ energy level, which reinforces the view of the existence of effective Gd³⁺-to-Tb³⁺ energy transfer.³²⁵ However, the presence of a broad host emission (green, 544 nm) along with strong Tb³⁺ emission (587 and 622 nm) indicates incomplete energy transfer from the host to Tb³⁺ in Tb³⁺-doped SrZrO₃.³²⁶ DFT calculations showed an energy mismatch of the Tb-d states with the Zr-d and O-p states, which explains the difficult energy transfer from the SrZrO₃ host to the Tb³⁺ ion.

Er. The upconversion photoluminescence of the Er³⁺-doped perovskite ABO₃ structure has been widely studied. In Er³⁺-doped BaTiO₃, Er³⁺ occupying the B-site strongly enhances the ⁴S_3/2–⁴I_15/2 emissions due to the thermal quenching and lower symmetry induced by the phase transition. The ²H_11/2 state is thermally quenched to the ⁴S_3/2 state and subsequently contributes to the enhanced population of the ⁴S_3/2 state. The crystal field formed by the octahedral oxygen ions with a lower symmetry than O_h is more suitable for the Er intra-4f transitions.³²⁷

Yb. Yb³⁺-Doped luminescent perovskites have mainly been investigated for the quantum cutting effect, for which two near-infrared photons are emitted for each absorbed visible photon. Besides this, luminescence due to charge transfer is also possible. For YAlO₃ doped with 2% Yb³⁺, the broad emission bands with maxima at 360 and 533 nm are assigned to CT transitions of Yb³⁺, accompanied by a narrow ²F_5/2 → ²F_7/2 emission band peaking at about 1000 nm.¹⁸³ NIR emission has also been observed in Yb³⁺-doped halide perovskite CsPbX₃ NCs³²⁸ Yb³⁺ ions occupy the Pb²⁺ crystallographic sites in CsPbCl₃ NCs.³²⁹ Near-infrared PLQYs of 170% have been measured for Yb³⁺:CsPbCl₃ NCs. Kroupa et al. reported Yb³⁺-doped CsPb(Cl_1–xBr_x)₃ films with extremely high quantum yields reaching over 190%.³³⁰ The extremely efficient sensitization of Yb³⁺ luminescence in CsPbCl₃ NCs is due to the Pb atom with STE, which acts as the energy donor in a quantum cutting process.³³¹ In the Yb³⁺-doped double perovskites Cs₂AgInCl₆ and Cs₂AgBiX₆ (X = Cl⁻, Br⁻), the characteristic f–f transition emissions of Yb³⁺ are also observed due to an energy transfer from the hosts to the ²F_5/2 state of the Yb³⁺ ion.^332,333

4.3. Transition metal-doped perovskites

Cr. Generally, the luminescence of octahedral Cr³⁺ complexes is in the form of sharp lines from the ²E_g state. However, broad-band near-infrared emission from the ⁴T_2g state can be observed, either alone or in combination with the ²E_g lines, in weak crystal fields provided by ligands CI⁻, Br⁻ and I⁻, such as Cs₂NaMX₆ (M = In, Y; X = CI,Br).³³⁴ For Mn⁴⁺, the ground state electronic configuration is 3d³, isoelectronic with that of Cr³⁺. As a result, they share similar luminescence, which consists of a sharp zero-phonon line and an accompanying vibronic sideband emission. The lowest excited state is ²E, and the excitation is usually dominated by O²⁻ to Cr³⁺/Mn⁴⁺ charge-transfer transition. Katayama et al. reported the deep-red persistent luminescence of Cr³⁺. The LaAlO₃:Cr³⁺ sample showed persistent luminescence peaking at a very long wavelength of 734 nm due to a Cr³⁺:²E → ⁴A₂ transition after ultraviolet excitation. The Sm³⁺ ion was found to be a good codopant for increasing the persistent luminescence intensity by more than 35-fold. The Cr³⁺–Sm³⁺-codoped LaAlO₃ sample with peak luminescence at 694 nm is a candidate phosphor for in vivo imaging application.³³⁵

Mn. The photoluminescence properties of Mn⁴⁺-activated perovskites have been reported in a series of oxide phosphors, including germinates, silicates and aluminates. According to Adachi,³³⁶ they could be classified into three groups according to their different PL spectral features. The first type reveals a zero-phonon line (ZPL) emission peak due to the ²E_g → ⁴A_2g transitions in the Mn⁴⁺ ion together with the Stokes and anti-Stokes sideband peaks. However, the second-type of phosphors promise no clear identification of the ZPL emission peaks, even in the PL spectra measured at cryogenic temperatures. The ZPL emission peaks in the third type of phosphors can be tentatively determined from an analysis of the PL spectra using a characteristic Poisson function. The ZPL absorption transition energies in the PL excitation spectra are determined by performing a Franck–Condon analysis within the configurational-coordinate (CC) model. These transition energies and ZPL emission energies are used to obtain the crystal field (Dq) and Racah parameters (B and C) of the Mn⁴⁺ ions in these Mn⁴⁺-activated oxide phosphors. In Gd₂MgTiO₆:Mn⁴⁺, under 315 nm excitation, the sample exhibited a strong zero-phonon line located at 681 nm together with broad sidebands around 700 nm.¹⁰² Brik et al. adopted the crystal structure data and the overlap integrals between the Mn⁴⁺ and O²⁻ ions to demonstrate that, despite the increasing Mn⁴⁺–O²⁻ interatomic distance, the exchange charge contribution to the total values of CFP (which is proportional to the above-mentioned overlap integrals) increased from Y₂Ti₂O₇ to Y₂Sn₂O₇.³³⁷ This increased overlap in Y₂Sn₂O₇ occurs despite the fact that the Mn⁴⁺–O²⁻ bond distance in Y₂Sn₂O₇ is longer than in Y₂Ti₂O₇ and is attributed to a lack of hybridization (covalent bonding) between the filled 2p orbital of the oxygen ion occupying the 48f site of the pyrochlore lattice and the filled Sn⁴⁺ 4d¹⁰ orbital. In the Mn⁴⁺-doped double perovskites La₂LiSbO₆ and La₂MgTiO₆, the ²E_g → ⁴A_2g emission transition was determined by octahedral site distortion. The greater the site distortion, the lower the Mn–O covalent interaction and the higher the energy of the ²E_g → ⁴A_2g emission transition.³³⁸ The double perovstkite-type La₂MgGeO₆:Mn⁴⁺ exhibited deep-red emission peaking at 708 nm under UV irradiation. It also exhibited NIR persistent luminescence in the range from 670 nm to 720 nm, lasting for 60 min.³³⁹ The after-glow behaviour is dependent on the host intrinsic defects, and could be enhanced by the incorporation of Al³⁺ ions. The double-perovskite Ba₂GdSbO₆:Mn⁴⁺ phosphor demonstrated strong red emission, ascribed to a spin-forbidden Mn⁴⁺:²E_g → ⁴A_2g transition in the region of 620–750 nm. Zhong et al. discovered that Li⁺, Mg²⁺, Zn²⁺, Si⁴⁺, Ti⁴⁺ and Ge⁴⁺ dopants are beneficial for enhancing Mn⁴⁺ luminescence.³⁴⁰

Phosphors with Mn²⁺ as the activator usually have a relatively long lifetime in the order of milliseconds due to the spin-forbidden ⁴T₁–⁶A₁ transition. At higher Mn²⁺ concentrations, a red spectral shift and an increase in the oscillator strength have been observed,³⁴¹ which were not simply caused by concentration quenching. Ronda and Amrein proved that the relaxation of the spin-selection rule and the spectral red-shift stem from the spin exchange interaction upon going from isolated Mn²⁺ ions to Mn²⁺ ion pairs. Recently, Pradhan presented a review on a brief history of the development of luminescent Mn-doped NCs over the last 25 years and summarized the important findings and future prospects.³⁴² The typical colour, spectra and energy levels can be found in Fig. 9.

Fig. 9 (a) Photograph of CsPbCl₃ and Mn-doped CsPbCl₃ under 365 nm excitation. (b) Normalized temperature-dependent emission spectra of Y₂MgTiO₆:0.2%Mn⁴⁺ under 355 nm excitation. Crystal structure of the Y₂MgTiO₆:0.2%Mn⁴⁺ double perovskite. Temperature-dependent decay curves of the 698 nm emission in Y₂MgTiO₆:0.2%Mn⁴⁺ under a 355 nm pulsed YAG:Nd laser. (c) Energy level diagram for the Mn²⁺ emitting centre in a free-ion state and in a crystal field of cubic symmetry, energy band structure of CsPb(Cl/Br)₃ PQDs (Br/Cl ratio gradually increases from left to right), and energy transfer mechanisms from PQDs to Mn²⁺ dopants. (d) Multicolour emissions of a series of silica-coated PQD phosphors under the excitation of a UV lamp (from left to right: CsPb(Cl_0.4Br_0.6)₃, CsPbBr₃, CsPb_0.835Mn_0.165Cl₃, CsPb_0.835Mn_0.165(Cl_0.6Br_0.4)₃, CsPb_0.835Mn_0.165(Cl_0.5Br_0.5)₃, CsPb_0.835Mn_0.165(Cl_0.4Br_0.6)₃, CsPb_0.835Mn_0.165(Cl_0.3Br_0.7)₃, and CsPb_0.835Mn_0.165(Cl_0.05Br_0.95)₃). (e) Absorbance (dashed lines) and PL (solid lines) spectra of CsPbCl₃. Panels adapted from: a, b reprinted with permission from ref. 281. Copyright (2019) American Chemical Society; c and d reprinted with permission from ref. 350. Copyright (2017) American Chemical Society; e reprinted with permission from ref. 328. Copyright (2018) American Chemical Society.

Mn²⁺-Doped cesium lead halide (CsPbX₃) perovskite NCs exhibit a broad luminescence, attributed to the spin-forbidden ligand field transition (⁴T₁–⁶A₁) of Mn²⁺ ions at ∼600 nm, together with the exciton luminescence of the host.³⁴³ The d–d transition of Mn²⁺ ions results from exciton-to-Mn energy transfer, indicating a strong exchange coupling between the charge carriers of the host and the dopant d electrons.^281,344 The exciton–dopant energy transfer occurs in the time range of 50–100 ps, slower than the trapping of carriers in the host lattice, which takes 8–10 s. By varying the doping content, the Mn²⁺ luminescence could be tuned from 585–625 nm.³⁴⁵ According to Xu et al., the doping level of Mn²⁺ in CsPbCl₃ could reach 25% of the Pb content, with no substantial change in the morphology of the NCs.³⁴⁶ The emission intensity could be greatly enhanced by the further growth of an undoped shell. A CsMnCl₃ phase with complete Mn dopant substitution by spinodal decomposition was realized by Li et al.³⁴⁷ It had a shorter Mn lifetime, which was consistent with the short Mn–Mn distance within the CsMnCl₃ phase. A single-exponential decay was detected for a low concentration of Mn²⁺, with a concentration-independent lifetime of 1.8 ms. At a high doping level, the shorter and multiexponential decay becomes concentration dependent, reflecting the energy migration from Mn²⁺–Mn²⁺ dimers to traps.³⁴⁸

Mn-Doped CsPbCl₃ could be converted to Mn-doped CsPbBr₃ through an anion exchange reaction, but only weak Mn luminescence was obtained in Mn-doped CsPbBr₃.^278,349 To overcome the reabsorption and anioin-exchange effect in the mixing of multiple perovskite quantum dots, silica-coating was proposed to improve the air stability and suppress the anion-exchange.³⁵⁰ Parobek et al. managed to synthesize Mn-doped CsPbBr₃ NCs with an intermediate structure (L₂[Pb_1–xMn_x]Br₄, L = ligand).³⁵¹ Qiao et al. reported a photoinduced synthesis of Mn-doped CsPbX₃ (X = Cl, Br). The mild nature of the method preserved the size and anisotropic morphology of the NCs.³⁵² The trapped nonradiative energy could be recycled by doped Mn according to Wei et al., who observed the doped Mn snatching energy from the non-radiative trap states rather than from band states.³⁵³ Mir observed that by controlling the size and shape of Mn-doped CsPbBr₃, the intensity of Mn emission decreases with the optical band gap of the host decreasing from 2.92 to 2.53 eV. The quenching of Mn emission was mainly caused by the back energy transfer from Mn to the host.³⁵⁴ Chen et al. introduced a dimethyl sulfoxide (DMSO)–MnBr₂/PbX₂ composite as a precursor for the room-temperature facile synthesis of Mn-doped CsPbX₃ (X = Br, Cl) NCs. By adjusting the PbBr₂/PbCl₂ ratio, the excitonic emission spectra could be tuned from 517 nm to 418 nm.²⁷⁹ The temperature-dependent spectral properties of the pure and Mn-doped CsPbCl₃ NCs revealed that the Mn luminescence was enhanced at 78–270 K, because more carriers located at the excitonic state could transit directly to the thermally excited ⁴T₁ energy level of Mn due to the increased thermal perturbations (k_BT). Higher temperatures reduce the number of carriers located in the excitonic state, and therefore the Mn PL intensity decreases gradually.³⁵⁵ Besides Pb-halide peroskites, non-toxic metal halide double perovskites, such as Cs₂AgInCl₆, have been reported as a host to accommodate Mn²⁺.³⁵⁶

Near-infrared upconversion emission could be observed in high Mn²⁺ doping of KZnF₃:Yb³⁺, Mn²⁺ NCs.³⁵⁷ The 770 nm peaked emission originated from the ⁶A_1g(S)⁴T_1g(G) → ⁶A_1g(S)⁶A_1g(S) transitions of the Mn²⁺–Mn²⁺ dimers.

Ni. Broadband near-infrared photoluminescence could be obtained by exciting KMgF₃:Ni²⁺ using an 808 nm laser diode. The emission peak was centred at 1624 nm with a FWHM larger than 315 nm, originating from the ³T_2g(³F) → ³A_2g(³F) electronic transition of the octahedrally coordinated Ni²⁺. The emission range covered the wide absorption spectrum of typical combustion products, which makes KMgF₃:Ni²⁺ useful for combustion gas sensors³⁵⁸

Both monovalent and divalent Cu cations have been doped into perovskite compounds.⁸⁷ However, the presence of Cu²⁺ just improves the thermal stability and the optical performance of CsPb_1–xCu_x(Br/Cl)₃ NCs, without changing the blue excitonic luminescence.³⁵⁹

4.4. Bi-Doped perovskites

Zhou et al. employed a low-temperature (∼350 °C) ion exchange approach to synthesize (K,Na)La_1–xTa₂O₇:xBi³⁺ with RbLa_1–xTa₂O₇:xBi³⁺ as precursors.¹¹⁵ The 2D-layered perovskite doped with Bi³⁺ exhibited yellow emission due to the transitions from the excited state ³P₁ or ³P₀ to the ground state ¹S₀. Lozhkina et al. argued that although the optical absorption of both the single-crystal and powder samples of CsPbBr₃ doped with Bi presented a red-shift, the band gap of Bi-doped perovskite remained unchanged, according to low-temperature photoluminescence studies.³⁶⁰ The valence band structure was also not changed upon doping Bi³⁺, and only the Fermi level was increased by 0.6 eV, which suggests the formation of electron-type defect states localized closer to the bottom of the conduction band. In double-perovskite Cs₂AgInCl₆, the band gap decreased from 4.25 to 3.28 eV upon 1% Bi²⁺ doping, together with a broad band emission peak at 580 nm.³⁶¹

4.5. Strategies for PLQY and chemical stability improvement

Phosphors for LEDs need to be thermally stable due to the possible high working temperatures of devices. The stability of perovskite QDs need to be further improved because they tend to aggregate and surface oxidation occurs at high temperatures due to incomplete surface passivation.³⁶² The methods for stablity improvement are summarized in this section.

Encapsulation. Wang et al. prevented the anion-exchange effect by mixing perovskite CsPbBr₃ NCs with mesoporous silica particles to form nanocomposites. The green emissive mesoporous silica nanocomposites could be mixed with red perovskite NCs to fabricate an LED without the anion-exchange effect.³⁶³ Embedding CsPbBr₃ NCs into robust and air-stable rhombic prism Cs₄PbBr₆ microcrystals enabled a high emission efficiency in the solid state. The lattice matching contributed to the improved passivation.³⁶⁴ Wei et al. packed CsPbBr₃ NCs into crosslinked polystyrene beads via a simple swelling–shrinking strategy in nonpolar toluene and hexane. The prepared composite beads retained superior water-resistance, still emitting strong luminescence after 9 months water immersion.³⁶⁵ CsPbBr₃ quantum dots incorporated into a silica/alumina monolith exhibited high photostability under the strong illumination of blue light for 300 h due to a robust protective layer of compact SiO₂ and Al₂O₃ against oxygen and moisture.^366,367 Incorporation of perovskite NCs into a polymeric matrix is an effective strategy to improve the water resistance and to prevent anion exchange between different halide NCs in the solid state. This can be performed by mixing pre-synthesized perovskite NCs into a polymer matrix or by the in situ fabrication of perovskite NCs-embedded composite films. Meyns et al. coated CsPbX₃ NCs with poly(maleic anhydride-alt-1-octadecene) (PMA) to improve the stability as well as the processability.³⁶⁸ Zhang et al. embedded CsPbX₃ NCs into microhemispheres of a polystyrene matrix to improve the hydrolysis resistance.³⁶⁹ Antisolvent vapour treatment of CsPbBr₃ embedded in a dielectric polymer matrix of polyethylene oxide (PEO) resulted in a lower trap state density because of the larger crystal size and the fewer grain boundaries, which boosted the luminescent efficiency and the stability of perovskite LEDs.³⁷⁰ Zhu et al. embedded green-light-emitting CsPbBr₃ and red-light-emitting CsPb(Br/I)₃ NCs into carboxyl-containing polymethyl methacrylate (PMMA) by a hot-injection method to improve the efficiency and stability.³⁷¹ By introducing Cs⁺ as a dopant and 3,3-diphenylpropyamine (DPPA) as a capping ligand, the compatibility between NCs and the polymer was improved by decreasing the solubility variance of the precursors. With this method, the PLQY of red-emissive MAPbI₃ was increased from less than 15% to 91%.²¹⁸ Embedding CsPbX₃ (X = Cl, Br, I) perovskite NCs in the cage of zeolite-Y significantly improved its temperature and water resistance.³⁷² Xu et al. formed air-stable CsPbBr₃ nanoplatelets in the matrix of Cs₄PbBr₆ nanosheets by reducing the thickness of Cs₄PbBr₆ to ∼7.6 nm, which is at the scale of the exciton Bohr radius of CsPbBr₃.³⁷³ Hydrophobic solid paraffin could be used to encapsulate all-inorganic NCs to form water-stable composites. Meanwhile, the anioin exchange, which causes undesirable spectral changes, could be simultaneouly inhibited for the isolation of the NCs by a solid paraffin layer.³⁷⁴ He et al. loaded perovskite NCs in a host–guest metal–organic framework (MOF) forming ZJU-28⊃MAPbBr₃, and found this could significantly diminishes the aggregation, provide effective surface passivation and shelter the NCs from the environment.³⁷⁵ Integrating CsPbBr₃ into MOFs also improves the stability of the NCs as the QDs are encapsulated in a porous zeolite matrix. The CsPbBr₃ @Uio-67 composite³⁷⁶ and CsPbX₃-zeolite-Y composite exhibited stable photoluminescence properties under ambient atmospheric conditions.³⁷⁷

Synthesis atmosphere. Motti et al. discovered that the band-to-band radiative emission can be quenched for lead halide perovskites films in an inert environment, independently of the chemical composition. However, this negative effect could be compensated by the presence of oxygen, even in a very small amount.³⁷⁸ Similarly, Brenes et al. investigated the impact of the atmosphere on the local luminescence of MAPbI₃ perovskite grains using confocal photoluminescence measurements.³⁷⁹ The results showed that the emission from each grain depends sensitively on both the environment and the bright/dark nature of the specific grain. In the presence of oxygen and/or water molecules, the dark grains show a substantial improvement in emission, while the bright grain emission stays the same. However, the luminescence detoriates in nitrogen or under vacuum conditions. It is possible that moisture forms a passivating shell on the surfaces of the grains, converting the vacancies on the surfaces to amorphous species. Brenes showed that methylammonium lead iodide (MAPbI₃) polycrystalline perovskite films treated with combined treatments of light and atmosphere exhibit properties comparable with perovskite single crystals.³⁸⁰

Doping engineering. The substitution of Mn²⁺ has been demonstrated to be an effective strategy to fundamentally stabilize the perovskite crystal lattice of CsPbX₃ QDs even at high temperatures of up to 200 °C under ambient air conditions. The first-principles calculations confirmed that the significantly improved thermal stability and optical properties of CsPbX₃:Mn²⁺ QDs were mainly due to the enhanced formation energy originating from the successful doping of Mn²⁺ into CsPbX₃ QD.³⁸¹ α-CsPbI₃ has the most suitable band gap for all-inorganic perovskite solar cell (PSC) application, yet still faces the problem of phase instability at low temperatures in an air atmosphere. However, alloyed CsPb_xMn_1−xI₃ NCs have substantially the same optical characteristics and crystal structure as the parent α-CsPbI₃ system, but they are stable in the film and solution for more than one month.³⁸² Lu et al. chose SrCl₂ as a co-precursor to synthesize CsPbI₃ NCs, which led to an enhanced PLQY due to the simultaneous Sr²⁺ ion doping and surface Cl⁻ ion passivation.³⁸³ The doping of CsPbX₃ NCs with Ni ions could significantly improve the luminescent efficiency by increasing the defect formation energy, which results in a greatly improved short-range order of the perovskite lattice.³⁸⁴ Yb³⁺-Doped CsPbCl₃ NCs emit strong 986 nm NIR light. Yb³⁺/Er³⁺ co-doped CsPbCl₃ QDs emit at 1533 nm. After the incorporation of 2.0% Yb³⁺, the PLQY of CsPbCl₃ QDs was changed from 5.0% to 127.8%. Under continuous ultraviolet (365 nm) illumination, doped CsPbCl₃ NC have better stability than undoped NCs. The PL intensity of the undoped CsPbCl₃ NC decreased to 20% of the initial value within 27 h, while the doped one needed 85 h.³⁸⁵ By doping α-CsPbI₃ with Sb, the phase stability can be enhanced and the film morphology is also improved. It is worth noting that a CsPb_0.96Sb_0.04I₃-based solar cell retained 93% of the initial power conversion efficiency (PCE) after 37 days of storage in an air atmosphere.³⁸⁶ Ding et al. demonstrated that using transition metal halides (FeX₃, CoX₂, NiX₂, CuX₂ and ZnX₂; X = Cl, Br or I) as halide sources could effectively improve the stability of all-inorganic perovskite NCs against heat and moisture.³⁸⁷ The transition metal ions serve as ligand stabilizers, which are doped on the surface of NCs. Lead (Pb) and iodine(I) defects in metal halide perovskite materials can be reduced by doping with europium ions and the Eu³⁺–Eu²⁺ pair acts as a “redox shuttle”, which simultaneously selectively oxidizes Pb and reduces I to improve the long-term durability of the material.³⁸⁸

Surface passivation, defects repair and surface healing. Perovskite quantum dots suffer from trapping defects that give rise to detrimental nonradiative recombination centres. Halide vacancies were found to be the cause of the degradation of halide perovskites, with surface passivation proposed to solve this problem.²⁷⁶ Li et al. designed a recyclable dissolution–recrystallization self-healing strategy to synthesize large-area, crack-free and low-roughness perovskite thin films with improved luminescent performance.³⁸⁹ Tian et al. observed a light-induced photoluminescence enhancement in surface-deposited MAPbI₃ perovskites by time-resolved luminescence microscopy. It is possible that a photochemical reaction involving oxygen has the ability to deactivate the trapping sites where non-radiative charge-recombination occurs. Switching on/off the excitation light or switching the atmosphere between oxygen and nitrogen could thus reverse the enhancement.³⁹⁰ Li et al. performed the surface treatment of CsPbBr₃ NCs by hexane/ethyl acetate to control the ligand density. Through balancing the surface passivation and carrier injection, a 50-fold external quantum efficiency improvement (up to 6.27%) was achieved.²⁷³ Liu et al. introduced the organolead compound trioctylphosphine–PbI₂ (TOP–PbI₂) as the reactive precursor to achieve an almost complete elimination of the trapping defects. The obtained CsPbI₃ perovskite NCs had a high room-temperature PLQY of up to 100% (ref. 391) Koscher performed a postsynthetic modification of CsPbBr₃ NCs by a thiocyanate salt treatment, which significantly improved the quantum yield of both the freshly synthesized and aged NCs. The thiocyanate was able to repair the lead-rich surface, accessing a limited number of surface sites without leading to the destruction of the entire nanoparticles. However, attempts to extend this process to other halide compositions were much less successful, with minor improvements seen for CsPbBr_xCl_3–x compositions, but virtually no change seen for CsPbBr_xI_3–x compositions.³⁹² Nenon et al. introduced a general surface passivation mechanism for CsPbX₃ (X = Cl, Br, I) NCs. Both experimental and theoretical studies confirmed that full trap passivation could be realized by introducing anionic X-type ligands to alter the lead-based defect levels to produce trap-free band gaps.³⁹³ Monodisperse K-modified CsPbBr₃ QDs were synthesized by strictly controlling the amount of K-oleate additive (K/Cs molar ratio = 1.5/1) in the parent solution. Significantly enhanced photoelectric and thermal stabilities were observed with the PLQY increasing from 65% to 83%. The light stability test showed that the film without the K-modifier fell to 50% of its original PL strength after 45 h of irradiation; whereas even after 153 h, the K-modified one retained 100% of its PL.³⁹⁴ Zwitterionic ligands, such as 3-(N,N-dimethyloctadecylamino) propane sulfonate, with multiple anchoring groups can provide effective protection for QDs because zwitterionic ligands exhibit a greater adhesion to QDs surfaces through special chelation.³⁹⁵ In particular, such ligands allow the separation of clean QDs with a high PLQY of more than 90% after four rounds of precipitation/re-dispersion, as well as much higher uniform and colloidally dispersible QDs. Ultralong-term stable cubic CsPbI₃ was synthesized by polymerized polyvinylpyrrolidone (PVP)-induced surface passivation engineering.³⁹⁶ The introduction of trioctylphosphine oxide (TOPO) into a conventional oleic acid/oleylamine system enabled monodisperse CsPbX₃ NCs to be obtained with excellent optoelectronic properties at high temperatures (up to 260 °C). The size of these NCs varies over a relatively wide range. The presence of TOPO could significantly improve the stability of the CsPbX₃ NCs for ethanol treatment. After 100 min of ethanol treatment, the emission intensity of the TOPO-coated sample was decreased by only 5%, while the emission intensity of the non-TOPO-coated NC decreased to86%.³⁹⁷

Lin et al. achieved an external quantum efficiency exceeding 20% in a perovskite LED by managing the compositional distribution in the device. During the mixing of presynthesized CsPbBr₃ and the MABr additive (MA = CH₃NH₃), a CsPbBr₃/MABr quasi-core/shell structure was sequentially crystallized due to the differing solubilities. The MABr shell passivated the nonradiative defects in CsPbBr₃ crystals, thus enabling a balanced charge injection.³⁹⁸ Photoinduced ion segregation leads to band gap instabilities. Abdi-Jalebi mitigated both non-radiative losses and photoinduced ion migration by the decoration of passivating potassium halide layers onto the surfaces and grain boundaries.³⁹⁹ Halogen defects on the surface of CsPbX₃ NCs could be completely removed by post-treatment with a ZnX₂/hexane solution, which simultaneously enhanced the stability and luminescence intensity.⁴⁰⁰ Ahmed et al. boosted the PL of CsPbX₃ (X = Cl, Br, I) perovskite NCs by removing excess lead atoms from the surface using tetrafluoroborate salts.⁴⁰¹ Bodnarchuk et al. proposed a strategy for luminescence recovery by the postsynthesis surface treatment of CsPbX₃ NCs with didodecyldimethylammonium bromide and lead bromide. The function of the core–inner shell–outer shell nanocrystal structure was to heal the surface trap states and improve the colloidal stability.⁴⁰² Bohn et al. added a PbBr₂-ligand solution to repair the surface defects of bromide and lead vacancies in a subensemble of weakly emissive 2D CsPbBr₃ nanoplatelets.⁴⁰³ Li et al. improved the luminescence intensity of CsPbBr₃ NCs by a surface passivation with a silver complex. The Ag⁺ complex had the effect of fixing bromide on the nanocrystal surface and reducing the surface trap density.⁴⁰⁴ Trioctylphosphine (TOP) could instantly recover the luminescence emission and improve the emission intensity of freshly synthesized PQDs, without inducing any detectable structural changes.⁴⁰⁵ CsPbCl₃ and CsPbBr_3−xCl_x synthesized from a halide precursor consisting of copper halide (CuX₂)–oleylamine (OLA) complexes showed improved stability. The origin of their high stability and good crystallinity stemmed from the passivation of defect sites during the recrystallization process with the adsorption of CuCl₂ on the perovskite's surface.⁴⁰⁶

Optimizing the activator concentration. Generally, an excessive doping of luminescent centres devastates the emission intensity remarkably, which is called “concentration quenching”. This phenomenon is caused by the energy loss during the migration of excitation energy between luminescent centres. Layered perovskites with a 2Darrangement of luminescent centres possess high critical concentrations. For example, in Eu³⁺-doped RbLa_1−xEu_xTa₂O₇ and Gd_1−xEu_xTa₃O₉, the critical concentration is x = 0.5 and 0.7, respectively.^407,408 The percolation model has been proposed to account for the quenching behaviour of Eu³⁺, benefiting from the 2Dconversion of Eu³⁺ interactions within the rare-earth elements sublattice.⁴⁰⁹

Photoluminescent blinking. Thin films of the organometal halide perovskite MAPbI₃ exhibited temporally fluctuating PL when observed by fluorescence microscopy, which suggests they could be used as labels in super-resolution optical imaging due to the ultralarge amplitude of photoluminescent blinking.⁴¹⁰ Photoinduced activation or deactivation of the very few emitting or quenching sites per nanocrystal (one site per 104–105 nm³) is the cause of the blinking, including non-radiative channels, which undergo random fluctuations between active and passive states.⁴¹¹In situ analysis of the single-particle photoluminescence imaging of MAPbBr₃ NCs revealed that the photoluminescence quenching and blinking phenomena are most probably caused by charge trapping at surface states, together with the number of 1–4 trapping sites per particle.⁴¹² Fluorescence blinking in the microsecond timescale has also been confirmed to occur for all-inorganic perovskite CsPbBr₃ and CsPbBr₂I NCs. Enhancement of the nonradiative Auger recombination process accelerates a faster blinking at higher excitation power.⁴¹³

Degradation under material characterizations. Scanning electron microscopy or transmission electron microscopy have a potential negative effect on the performance of luminescent materials. Bischak et al. studied the luminescence heterogeneity among different grains of methylammonium lead halide perovskite films using high-resolution cathodoluminescence microscopy.⁴¹⁴ Electron beam-induced luminescence was observed in the films. The variability in intensity characterize the different distributions of the surface and bulk defects. Upon studying the degradation of methylammonium lead iodide perovskite, Yuan et al. proposed reducing ion migration to enhance the stability of perovskite materials. They also noticed that significant perovskite degradation could be readily induced by characterization under scanning electron microscopy or transmission electron microscopy.⁴¹⁵ Meanwhile, photoinduced degradation was observed in methylammonium lead triiodide (MAPbI₃) perovskite NCs under intense light excitation.⁴¹⁶ The degradation was accompanied by an intensity decrease and spatial shifts in the emission localization position.

Crystal morphology control. The crystal morphology of a perovskite layer is correlated to the defect quantity. A large number of grain boundaries and crystal dislocations are expected in a very non-uniform crystal structure with multiple facets, giving rise to increased trap-assisted non-radiative recombination.⁴¹⁷ Yang et al. examined grain boundaries (GBs) with respect to non-GB regions (grain surfaces (GSs) and grain interiors (GIs)) in micrometre-sized perovskite MAPbI₃ thin films. Contrary to previous studies, recombination was reported to happen primarily in the non-GB regions. Further, the lifetimes at the GBs were no worse than those at the GSs/Gis. Those facts suggest that GBs do not dominate non-radiative recombination in MAPbI₃ thin films.⁴¹⁸ By replacing the conventionally used oleic acid with an alkyl phosphinic acid, CsPbI₃ NCs can retain the cubic perovskite phase in solution, avoiding the facile phase transformation to the orthorhombic phase.⁴¹⁹ Quan et al. developed a fabrication strategy to control the different band gaps in PEA₂(MA)_n−1Pb_nBr_3n+1 (phenylethylammonium, PEA) perovskites through composition and solvent engineering during crystallization to obtain highly efficient emission.⁴²⁰

Surface plasmon resonance (SPR). The SPR effect has been applied to metallic nanostructures to enhance the luminescence by directly converting the absorbed photons into electrical energy by generating highly energetic electrons, i.e. hot electrons. Efficient plasmon–hot electron conversion has been reported in Ag–CsPbBr₃ hybrid NCs, which can be ascribed to the increased metal/semiconductor coupling⁴²¹ A small excess of PbI₂ in ABX₃ [A = Cs⁺, MA, or FA; B = Pb or Sn; X = Br, I] suppresses nonradiative charge carrier recombination and enhances luminescence.⁴²²

Lattice-anchoring: Very recently, Wei et al. reported that CsPbX₃ QDs epitaxially synthesized by the surface chemical conversion of Cs₂GeF₆ double perovskites with PbX₂ (X = Cl, Br, I) forming a hybrid structure of CsPbX₃/Cs₂GeF₆. The products has high stability under ambient conditions due to anchoring effects. By halogen substitution, they obtained blue, green and orange-emitting CsPbCl_1.5Br_1.5/Cs₂GeF₆, CsPbBr₃/Cs₂GeF₆ and CsPbBr_1.5I_1.5/Cs₂GeF₆ complexes with quantum efficiencies of 27.3%, 36.4% and 6.2%, respectively.⁴²³ Liu et al. reported a “lattice-anchored” hybrid material that combines CsPbBr_xI_3−x and PbS, where lattice matching between the two materials helped to inhibit the transition of the favourable α-CsPbBr_xI_3−x to the undesired δ-CsPbBr_xI_3−x phase. Compared to the original perovskite, the stability of the PbS-anchored perovskite under ambient conditions was increased by an order of magnitude.⁴²⁴

PLQY and chemical stability improvement for OIHPs. Low-D perovskites with large amounts of inactive ligands provide better stability due to the hydrophobic nature of the organic cations, which prevent direct contact of the water with the perovskite material.⁴²⁵ Smith et al. and Cao et al. reported that the water stability of the perovskite material could be remarkably improved by partially substituting the MA⁺ cation with a long-chain organic cation.^68,426 PEA₂MA₂Pb₃I₁₀ (PEA = C₆H₅(CH₂)₂NH₃⁺) could resist 52% relative humidity, more than twice that of MAPbI₃.⁴²⁶ Photoinduced organometallic halide bond dissociation and reforming play a key role in determining the photostability of OIHPs. Photodissociation of the 1D tin bromide chain followed by structural reorganization led to the formation of a more thermodynamically stable 0D structure. Generally, 0D organometallic halides have a higher stability than 3D organometallic halides. This can be attributed to the fact that the metal halide in the 0D structure is encapsulated and protected by the inactive organic cations from oxygen and moisture. Furthermore, unlike 3D perovskites containing small cations that can migrate in the crystal lattice, 0D perovskites contain large cations that can be relatively more stable. Methods to further improve the stability of the 0D structure may involve the use of larger and more rigid organic components or the post-crosslinking of organic components. It may also be helpful to increase the ionic interaction between the cation and the anion to form a more regular crystal structure.

Coya et al. reported an improvement in the photostability of MAPbI₃ by Bi doping. This increase in stability originated from the strong migration ability of Bi. First, BiI₃ was formed, and then a stable iodonium oxide compound (BiOI) was formed and deposited on the grain surface, which hindered the decomposition of MAPbI₃ into PbI₂ and PbO_x.⁴²⁷

High temperature is another pathway that can lead to the degradation of perovskites, including the effects of thermally induced chemical decomposition and perovskite phase transitions. It is known that stabilizing a perovskite material by substituting its constituent ions is a very promising method. The water stability of the improved 3D perovskite material was demonstrated by controlling the halide composition; whereby substituting I with a bromine anion led to a slight distorion in the PbX₆ octahedra.⁴²⁸ The reduced octahedral tilt and twisted lattice were due to the difference in ionic radii and the hexa-coordination of the I⁻ and Br⁻ ions. The stability of the perovskite material can also be improved by controlling the X-halide and thiocyanate (SCN) composition, as reported by Jiang et al.⁴²⁹ By substituting two I⁻ anions with two SCN⁻ anions, a new perovskite material MAPb(SCN)₂I was obtained. Compared to the pure MAPbI₃ perovskite, MAPb(SCN)₂I decomposed at a slower rate even at 95% relative humidity (RH) (after 4 h of air exposure), with the band gap remaining unchanged. By replacing the organic cation A from MA⁺ into FA⁺ and Cs⁺ in a pure 3D perovskite, the thermal stability of the device could be effectively improved.⁴³⁰ For OIHPs, the decomposition energy and the decomposition temperature are correspondingly low (<300° C), and the moisture/ultraviolet light can further accelerate the decomposition process. A better understanding of the stability limitations is needed.

Improvement of perovskite LED performance. Zou et al. investigated the mechanism of efficiency roll-off in 2D layered perovskite LEDs. By simultaneously measuring EL and PL on the same working device, they found that non-radiative Auger recombination was responsible for the luminescence quenching, which could be suppressed by increasing the width of the quantum wells.⁴³¹ Wang et al. improved the film quality of 2D CsPbBr_xCl_3−x perovskite by designing a NiO_x/LiF hole-transport layer with high affinity to the precursor solution. The quenching effect of the hole-transport layer to the as-prepared perovskite film could be greatly reduced.¹²⁸ Incorporating a Au–Ag alloy nanoparticle in the electronic transport layer of the all-inorganic perovskite LED could increase the luminescence efficiency by 25% through the localized surface plasmon.⁴³² This enhancement could be attributed to the match between the localized surface plasmon resonance wavelength of the Au–Ag nanoparticle and the emission peak of the perovskite LED. Ahn et al. finely controlled the crystallization of MAPbBr₃ by using a polar solvent-soluble self-doped conducting polymer as a hole-injection layer.⁴³³ The induced granular structure makes charge carriers spatially confined more effectively than in a columnar structure. In addition, indium tin oxide (ITO) etching is weakened by reducing the metallic In species released, which are otherwise responsible for exciton quenching.

Luminescence quenching in oxide perovskites. The quenching of In³⁺ luminescence in LaInO₃ comes from the mobile excited state. The corner-sharing InO₆ octahedra construct a 3-dimensional sublattice, from which the energy levels can become broadened into bands easily. This is the reason for the excited state mobility, which accelerates the energy loss when it reaches the quenching centres.⁸⁴ Besides the characteristic luminescence of trivalent rare-earth ions (Pr³⁺, Sm³⁺, Eu³⁺, Tb³⁺, Dy³⁺, and Er³⁺) doped in KLaNb₂O₇, the host luminescence of layered perovskites due to delocalized excited states was observed at 77 K. However, in Pr³⁺-, Sm³⁺-, Eu³⁺- and Tb³⁺-doped KLaNb₂O₇, the host luminescence was completely quenched, probably by an electron- and hole-trapping process occurring at rare earth ions.¹⁰⁰ The addition of Al or Ga is essential for SrTiO₃:Pr³⁺ to achieve a high luminous efficiency of cathodoluminescence and PL. The luminescent enhancement probably arises from the crystallinity improvement, which suppresses the defects and charge compensation by Al³⁺ substituting for Ti⁴⁺.^434,435 This enhancement was also observed for CaTiO₃:Pr³⁺, BaTiO₃:Pr³⁺, BaZrO₃:Pr³⁺ and SrIn₂O₄:Pr³⁺.⁴³⁶ The incorporation of Li ions enhanced the red luminescence of SrTiO₃:Pr³⁺. Tian et al. ascribed the enhancement mechanism to the oxygen vacancy generated by Li doping, which promoted energy transfer from the excited carrier in the lattices to the Pr³⁺ activator ion.⁴³⁷ The introduction of Zn²⁺ ions in BaTiO₃-doping of Er³⁺/Yb³⁺ showed an enhancement in upconversion and cathodoluminescence luminescence, which was caused by the modification of the coordinating environment around the RE³⁺ ions.⁴³⁸

5. Applications

Preparative solid-state chemistry is known for its lack of predictability compared to organic synthesis. However, ongoing research on the structure–property relationship of mixed organic–inorganic compounds is an important step in the predictability of this diverse and interdisciplinary field.

The luminescent perovskites have many important applications in optical materials and devices, see Fig. 10.

Fig. 10 (a) PL spectra of red, green and blue colours on paper. Facial make-up full-colour image and full-colour fluorescent photo of sophisticated patterns. (b) 365 nm UV light photos of red, green and blue quantum dots, and colourful fluorescence cartoon sculptures made from red, green and blue quantum dots. (c) Illustration of full-spectrum persistent luminescence tuning. (d) Emission spectra recorded immediately after exposing the dehydrated (Cs₂InBr₅·H₂O) to air. The three photographs show the colour changes of the dehydrated material under air upon photoexcitation. (e) Visualized dual emission between the hydrated and dehydrated (Cs₂InBr₅·H₂O), fabricated by embedding the materials into an etched butterfly pattern. (f) Sequential optical images and PL emission spectra of MAPbBr₃ NCs@Pb-MOF after one cycle of an impregnation-recovery process. 1, 2 and 3 represent the original, impregnated and recovered powder sample of MAPbBr₃ NCs@Pb-MOF, respectively. (g) Reversible fluorescence switching of the MAPbBr₃ NCs@Pb-MOF pattern in one encryption–decryption cycle (methanol impregnation for encryption and MABr spraying for decryption). Panels adapted from: a, reprinted with permission from ref. 439. Copyright (2019) American Chemical Society; b, reproduced from ref. 374 with permission from The Royal Society of Chemistry; c, reprinted with permission from ref. 440. Copyright (2019) John Wiley & Sons, Inc; d and e, reprinted with permission from ref. 221. Copyright (2019) John Wiley & Sons, Inc; f and g, ref. 441.

5.1. Scintillators

ABX₃-type halide perovskites have a long history as scintillator materials for radiation detection.¹⁰³ Scintillation is the observable luminescence caused by the absorption of high-energy radiation or particles.⁴⁴² Consequently, scintillators have been widely applied in a series of scientific and industrial fields, such as high-energy physics, medical imaging and geophysical exploration. Slow luminescent components constitute the main restrictive factor regarding the usage of scintillators, which relates to the intrinsic recombination luminescence from self-trapped excitons with a lifetime in the order of hundreds of microseconds. To fulfill the application requirements, the contribution of the delayed exciton luminescence should be minimized. A fast scintillation response speed comparable with BaF₂ cross-luminescence has been achieved in Yb-rich YAlO₃ single crystals.⁴⁴³ The direct band-gap semiconductor CsPbBr₃ meets the requirements for radiation detection based on the evidence of its green PL emission when excited by a He–Cd laser, at 325 nm.¹ Hybrid perovskite crystals, such as 3D MAPbI₃, MAPbBr₃ and 2D (EDBE)PbCl₄, also exhibit X-ray scintillator characteristics, owing to the presence of heavy atoms.¹⁷ Compared to 3D perovskites, the thermal effects in 2D (EDBE)PbCl₄ are significantly reduced due to the large exciton binding energy.

5.2. LEDs, NIRs, LEDs, phosphors

The most remarkable application of luminescent perovskites lies in the fabrication of light emitting diodes (LEDs). The OIHPs have attracted much attention due to their tunable excitonic luminescence. In 1994, Era et al. fabricated an organic–inorganic heterostructure electroluminescent (EL) device using the combination of a layered perovskite compound ((C₆H₅C₂H₄NH₃)₂PbI₄, PAPI) and an electron-transporting oxadiazole derivative.⁹⁵ The EL device consisted of an indium–tin–oxid (ITO) anode, a PAPI emitter prepared by a spin-coating method from acetonitrile solution and an electron-transport layer of an oxadiazole derivative and an MgAg cathode. The heterostructure is displayed in Fig. 5(g) and (h). Driven at liquid-nitrogen temperature, an intense green emission peaking at 520 nm was observed, corresponding to the photoluminescent spectrum of the PAPI spin-coated film. However, the EL efficiency dropped dramatically due to the thermal luminescence quenching of the PAPI film, originating from the thermal ionization of the excitons. The light emitting devices typically employ two different morphologies of the layered perovskites: single crystals and spin-coated thin films. Single crystals are usually prepared from a solution phase rather than the melt phase, because the layered perovskites decompose before melting. Thin films can be easily spin-coated from solutions or by a vapour deposition process.⁴⁴⁴ In 2016, Wang et al. reported a multiple quantum wells LED with a very high external quantum efficiency of up to 11.7%.⁴⁴⁵ Kim et al. improved the luminescence efficiency of a formamidinium (FA, CH(NH₂)₂) lead bromide perovskite (FAPbBr₃)-based LED by a ligand engineering method.⁴⁴⁶ Control of the ligand length has the effect of reducing the trap-assisted carrier recombination and improving charge injection and the transport capability in FAPbBr₃ films. Lova et al. reported a potential industrial-scale production route for polymer-perovskite lighting devices by embedding the 2D perovskite (EDBE)PbCl₄ in a flexible polymer, in which emission suppression and enhancement at specific wavelengths could be used to tune the emission colour over the entire visible spectrum without resorting to compositional engineering.⁴⁴⁷ Chin et al. continued to raise the external efficiency of green LEDs above 13% by coupling graded-size formamidinium lead bromide (FAPbBr₃) NCs and microplatelets of octylammonium lead bromide perovskites.⁴⁴⁸ Single-component WL emitters were realized in the Cd-based 2D hybrid perovskite (C₆H₁₁NH₃)₂CdBr₄. Upon 325 nm excitation, a broad-band WL emission was observed, consisting of an excitonic peak at 2.94 eV and a second peak at 2.53 eV resulting from the cyclohexylammonium cations emission.²⁴²

The first use of completely inorganic CsPbBr₃ thin films was achieved in 2015.^18,19 Gangishetty et al. proposed an alternate transport layer structure in a perovskite LED to enable efficient emission across the entire blue–green spectral range.⁴⁴⁹ Khan et al. optimized a CsPbBr₃-based LED device and achieved a significantly reduced turn-on voltage.⁴⁵⁰ Great advances have been achieved in improving the stability of all-inorganic perovskite quantum dots, such as by surface modification or encapsulation in polymer and glass. Shi et al. focused on the emission efficiency and operation stability of perovskite LEDs.¹⁸⁹ CsPbBr₃ quantum dot LEDs using n-ZnO and p-NiO carrier injectors can endure a high humidity (75%, 12 h) and a high working temperature (393 K) even without encapsulation. For highly flexible CsPbI₃ perovskite LEDs fabricated by a photopolymer and ultrasmooth Ag films, good performance was maintained after 1000 times of repeated 180 degree stretch-release bending.⁴⁵¹ Liu et al. presented a simple approach to improve the chemical stability by fabricating CsPbI₃ quantum dots in zinc borosilicate glass through a conventional melting-quenching technique operated at 540 °C.⁴⁵² CsPbBr₃ NCs were embedded in a specially designed TeO₂-based glass matrix, which showed a significant improvement in photon/thermal stability and water resistance.⁴⁵³ CsPbX₃ (X = Br, I) NCs could be embedded into phosphosilicate glasses with low-melting (700 °C) temperature.⁴⁵⁴ The partial replacement of toxic Pb with Mn has also been reported in Mn-doped CsPb(Cl/Br)₃ nanocrystal glass.⁴⁵⁵

Near-infrared (NIR) light-emitting diodes are used in a wide range of applications, including night vision, biomedical imaging, optical communications and computing. By embedding PbS quantum dots in a high-mobility hybrid perovskite matrix MAPbI_xBr_3−x, enhanced radiative recombination in the dots led to them exhibiting 1391 nm emission.⁴⁵⁶ A tunable near-infrared LED was fabricated based on lead-free organo-tin halide perovskite CH₃NH₃SnI₃, achieving a 945 nm near-infrared emission.⁴⁵⁷ Increasing the bromide content led to shorter wavelength emissions, tunable down to 667 nm. Qiu et al. synthesized nanocrystalline methylammonium (MA) lead tin iodide films and used them to fabricate efficient NIR LEDs.⁴⁵⁸ The emission from the mixed lead–tin (Pb–Sn) halide perovskite was tunable from 850 to 950 nm, either by changing the Pb/Sn ratio or by incorporating bromide.

The deep red emission of Mn⁴⁺ (3d³) in the double perovskites La₂LiSbO₆ and La₂MgTiO₆ could be used to develop agricultural (horticultural) applications.³³⁸ Also under 342 nm excitation, NaLaMgWO₆:Mn⁴⁺ and NaLaMgWO₆:Mn⁴⁺ double-perovskite phosphors showed a high-efficiency far-red emission at approximately 700 nm, which suites the requirement for indoor plant growth.⁴⁵⁹

The other application of inorganic perovskite in LED fabrication is to serve as a phosphor in light downconversion and combination. For this, impurity doped phosphors have been widely investigated. Herein, great emphasis has been paid to inorganic halide perovskite nanocrystals. The CsPbBr₃ QDs were chosen as red-emitting components to improve the colour rendering index (CRI) of Ce³⁺:YAG-based white LEDs.⁴²² Owing to the narrow emission band, the colour gamut of CsPbX₃ NCs could cover more than 140% of the NTSC (National Television System Committee) TV colour standard.⁴⁶⁰ CsPbBr₃ QDs exhibited a narrow band with a FWHM at about 20 nm peaking at 534 nm⁴⁶¹ However, the mass production of perovskite phosphors is critical for their wide application.⁴⁶² Perovskite NCs synthesized by spray pyrolysis show an unprecedented stable absolute PLQY of ≈100% in both solution and in the solid-state neat film.⁴⁶³

5.3. Solar cells

Luminescent perovskites play an important role in solar cells. The efficiencies of inorganic–organic lead-halide perovskite solar cells have reached beyond 22% with photovoltages >1.2 V.⁴⁶⁴ Zhou et al. fabricated a novel type of quantum cutting material, CsPbCl_1.5Br_1.5:Yb³⁺,Ce³⁺ NCs, which could realize the emission of multiple near-infrared photons for each ultraviolet/visible photon absorbed, which improved the photoelectric conversion efficiency (PCE) of solar cells.²⁸³ Pazos-Outon et al. showed that methyl ammonium lead iodide (MAPbI₃) likewise has a larger Auger coefficient that expected.⁴⁶⁵ An efficient YCl_3-treated TiO₂ electron-transfer layer was used to fabricate perovskite solar cells with enhanced photovoltaic performance and less hysteresis.⁴⁶⁶ The improved perovskite/TiO₂ interface with YCl₃ treatment was found to separate and extract photogenerated charge rapidly and to effectively suppress recombination. Li et al. provided a facile approach to fabricate core–shell perovskite solar cells with high efficiency and long-term stability against moisture.⁴⁶⁷ The authors introduced gallium(III) acetylacetonate (GaAA3) as a precursor additive to in situ induce a metal–organic-complex monomolecular intermediate ([GaAA3]₄), which allowed realizing Cs_xFA_1−xPbI₃–[GaAA3]₄ (0 < x < 1) hybrid perovskite materials. The formed hybrid perovskites were proven to possess a Cs_xFA_1−xPbI₃ core and [GaAA3]₄ shell. It was reported that surface passivation can be combined with optimized charge carrier selective interfaces to increase the power conversion efficiency of hybrid perovskite (MAPbI₃) thin film solar cells.⁴⁶⁸ Rajagopal et al. incorporated phenylethylammonium (PEA) in a mixed-halide perovskite composition MAPb(I_0.6Br_0.4)₃ to solve the inherent material-level challenges in 1.80–1.85 eV E_g perovskites.⁴⁶⁹ The power conversion efficiency of the lead-free Cs₃Bi₂I₉ mesoscopic solar cells was limited by the poor photocurrent density. Therefore, a continuous network of a 3D crystal structure is more favourable.⁴⁷⁰

5.4. Lasers, non-linear optics, indicators

Deschler et al. constructed and demonstrated the operation of an optically pumped vertical cavity laser comprising a layer of perovskite between a dielectric mirror and evaporated gold top mirrors.⁴⁷¹ The observed broad homogeneous line width in MAPbI_3−xCl_x ought to allow amplification of Fourier-transform-limited pulses with sub-100 fs duration, making these materials interesting for pulsed laser operation.⁴⁷² The optical-pumped rectangular cross-sectional perovskite MAPbI₃ nanowire lasers had a near-infrared wavelength of 777 nm, low threshold of 11 μJ cm⁻² and a quality factor as high as 405.⁴⁷³ Liu et al. systematically investigated the temperature-dependent spontaneous emission and lasing spectra of chemical vapour deposited CsPbBr₃.⁴⁷⁴

Two-photon absorption-induced luminescence is a typical third-order nonlinear optical process. It has been investigated in perovskite CsPbBr₃ quantum dots at a broad temperature range.⁴⁷⁵ The PL spectrum excited by two-photon absorption exhibits different temperature-dependent spectral shift behaviour than general semiconductors, and is probably caused by thermal expansion, electron–phonon interaction and structural phase transition. The inefficient photoluminescence was enhanced by constructing a hybrid dielectric structure, in which a 2D perovskite flake phenylethylamine lead iodide ((PEA)₂PbI₄) was covered by dielectric microspheres (approximately micrometers in diameter).⁴⁷⁶ The emission increased by two orders of magnitude in the hybrid dielectric structure due to the cooperative enhancement of the detection efficiency and quantum efficiency of the 2D perovskite.

Blue-light emission at room temperature from Ar⁺-irradiated SrTiO₃ single crystals could be applied in displays and indicators.⁴⁷⁷ Besides Ar⁺-irradiated oxygen-deficient SrTiO₃, substituting La³⁺ for Sr²⁺ and Nb⁵⁺ for Ti⁴⁺ in SrTiO₃ provided electron carriers in Ti 3d conduction bands, which were responsible for the room-temperature blue-light emission.⁴⁷⁸ The emitting region can be patterned into any size and shape with conventional microscopic fabrication techniques.

5.5. Temperature sensors

BaTiO₃ doped by lanthanide ions (0.3 mol% Er³⁺/3.0 mol% Yb³⁺) can be used as a temperature sensor with a maximum sensitivity of 0.00475 K⁻¹ at 430 K.⁴³⁸ MgTiO₃ nanoparticles doped with Mn⁴⁺ were also synthesized as optical thermal sensors to take advantage of the drastic variations of Mn⁴⁺ luminescence with temperatures.⁴⁷⁹ The temperature evolution of the infrared luminescence of yttrium orthoaluminate nano-perovskite, YAlO₃, doped with 1.0 mol% of Nd³⁺ ions, could be used for both sub-tissue luminescence imaging and luminescence-based thermal sensing.⁴⁸⁰ The Mn⁴⁺-doped Y₂MgTiO₆ phosphors showed far-red emission at ∼715 nm, with a maximum sensor sensitivity as high as 0.00142 K⁻¹ at 153 K.⁴⁸¹ The thermochromism of organic/inorganic halide perovskites Cs₄PbBr₆ has been studied for potential applications as photoluminescence-based temperature sensors.⁴⁸² Pr³⁺-Doped triple-layered perovskite Na₂La₂Ti₃O₁₀ microcrystals promise remarkable performance for temperature sensing over a wide temperature range (125–533 K), with a maximum relative sensitivity of 2.43% K⁻¹ at 423 K.⁴⁸³

5.6. Upconversion, persistent luminescence

Upconversion luminescence could be used to detect IR or NIR light. Usually, nonlinear upconversion luminescence is limited by a small cross-section of multiphoton absorption. Methylammonium lead halide perovskites were combined with erbium silicate nanosheets to give a remarkable spectral response at 1.54 μm. Strong upconversion luminescence from the nanosheets was well confined in cavities, which simultaneously excited the neighbouring perovskite to realize photodetection.⁴⁸⁴ Considering the large absorption coefficient and high PLQY, a convenient approach to realize upconversion luminescence in perovskite quantum dots was reported by Zheng et al. through a sensitization by lanthanide-doped nanoparticles.¹⁸⁷ The sensitization was governed by a radiative energy-transfer upconversion process, and according to Zeng et al., the energy-transfer efficiency between upconversion nanoparticle donors and quantum dot acceptors approached nearly 100%.⁴⁸⁵ Under NIR laser light irradiation, intense green emission from the perovskite CsPbBr₃ quantum dots through BaYF₅:Yb,Ln sensitization was observed even in a liquid suspension. Ma et al. reported WL emission from composites of high-quality CsPbBr₃ quantum dots and β-NaYF₄:30%Yb,0.2%Tm nanoparticles triggered by 980 nm NIR excitation.⁴⁸⁶ The luminescence of the quantum dots was excited by the upconversion luminescence of the nanoparticles. WL emission was achieved from the composite sample with regulation of the concentration. The upconversion luminescence of perovskite quantum dots has received less attention due to the lack of anti-Stokes luminescence properties of the material itself. However, Duan et al. observed upconversion luminescence in CsPb(Br/I)₃ NCs under laser excitation with a wavelength of 635–700 nm. The emission peaks located at 632 nm were consistent with the normal PL emission of CsPb(Br/I)₃ NCs excited by 400 nm UV light.⁴⁸⁷ Therefore, more research is needed to unravel the upconversion luminescent properties in perovskite quantum dots.

Persistent luminescent materials are applied as night or dark-light vision indicators. All-inorganic CsPbX₃ (X = Cl, Br and I) perovskite quantum dots could be combined with the traditional afterglow phosphor CaAl₂O₄:1%Eu²⁺,0.5%Nd³⁺ to produce tunable persistent luminescence covering the full spectrum. The perovskite quantum dots work as light-convertors to absorb the persistently released violet and blue afterglow light and then produce emissions tunable over the entire visible spectral region. In addition, the afterglow decays of different spectral components are highly synchronized due to the same light source.⁴⁴⁰ The other strategy for persistent luminescence depends on the direct energy transfer from hybrid perovskites into the triplet states of organic molecules. Hu et al. reported a 0.2 ms afterglow duration in the 2D perovskite (TTMA)₂PbBr₄(TPB) by performing molecular engineering of the OIHPs. The excitonic emission peak of TPB was located at 420 nm under 390 nm excitation, whereas a weak emission peak at 590 nm corresponded to the phosphorescence from the organic cation TTMABr molecule. Nonradiative recombination was suppressed by the mixed-cation perovskites with a phosphorescence yield of 11.2%. The persistent luminescence colour could be tuned by incorporating different organic cations into the hybrid perovskites. Following this strategy, novel persistent luminescence could be explored in other 2D perovskite systems by changing the dimensionality and the chemical composition.⁴⁸⁸

5.7. Information encryption and decryption

Luminescent perovskites have “smart” functions when applied in confidential information encryption and decryption. In cesium lead halide nanostructures, quantum-confined few-monolayer nanoplatelets are converted to the bulk phase under very low irradiation intensity (≈20 mW cm⁻²).⁴⁸⁹ Benefiting from the remarkable emission colour change during photon-driven transformation, multicolour luminescence photopatterns and facile information photo-encoding can be established. Zhang et al. realized confidential information protection and storage by converting lead-based metal–organic frameworks into luminescent perovskite NCs. By polar solvents impregnation and halide salt conversion, the quenching and recovery of the luminescence lead to reversible on/off switching of the luminescence signal for multiple information encryption and decryption processes.⁴⁴¹ In perovskite NCs CsPbX₃ (X = Cl, Br, and I), the molecular chirality can transfer to the NCs, resulting in circularly polarized luminescence signals.⁴⁹⁰ In addition, confidential information protection can be realized through the anion-exchange reaction of perovskite quantum dots, and emission colour-encoding patterns could be designed by an inkjet-printing technique. The conversion of luminescent 3D from a colourless 0D perovskite is achieved by reaction with invisible halide salts. Therefore, the reversible off/on switching of the luminescence can be used to encrypt and decrypt confidential information.⁴³⁹ Fan et al. fabricated methyl ammonium lead tri-bromide perovskite metasurfaces. Here, the encoded information was visible only when the incident laser was on-resonance, as off-resonance pumping and single-photon excitation just produced a uniform dark background.⁴⁹¹

The use of perovskite NCs as inks was reported by Akkerman et al.⁴⁹² Pure Cs₄PbX₆ (X = Cl, Br, I) and mixed halide compositional NCs with sizes ranging from 9 to 37 nm were synthesized by a colloidal method. They had no visible excitonic emission due to their large band gap. The absorption bands belonging to samples with mixed halide composition were located in the intermediate spectral position between those of the pure halide compounds. Pan et al. synthesized extremely stable CsPbX₃ nanocrystal–polymer composites as solution-processable luminescent inks with remarkable chemical stability towards water.⁴⁹³

5.8. Photosensors

Perovskite materials can be used for photodetection due to their high external quantum efficiency and rapid temporal response. Zheng et al. fabricated a 2D photodetector from high-crystallinity 2D CsPbBr₃ perovskites micro-/nanoflakes⁴⁹⁴ Duan synthesized MAPbBr₃ microdisk arrays with a uniform size distribution and emission spectra and fabricated a perovskite photodetector array with fast rise (<8.3 ms) and decay (<8.3 ms) times.⁴⁸⁷

Room-temperature red luminescence was observed in a 2D layered hybrid lead halide [CH(NH₂)₂][C(NH₂)₃]PbI₄, together with photoconductivity.⁴⁹⁵ The co-existance of the excitonic emission and photoconductivity in hybrid perovskites is rather rare due to the competion between the two processes, as reflected by the low quantum yield of 3.5%. The photoconductivity has a dark specific resistivity value of 1 × 1010 Ω cm, indicating a rather low intrinsic carrier concentration and/or mobility. However, whether trap states or self-trapped excitons cause this photoconductivity remains unclear.

5.9. Stress sensors

Luminescent perovskites could be fabricated into sensors to detect stress or pressure. Sr₃Sn₂O₇:Sm³⁺ has strong reddish–orange emission under compressive load, which is visible to the naked eye.¹¹⁴ Ma et al. observed pressure-induced emission in initially nonfluorescent 0D perovskite quantum dots Cs₄PbBr₆ under a high pressure of 3.01 GPa.⁴⁹⁶ The emission intensified under further compression. This luminescence stems from the enhanced optical activity and the increased binding energy of self-trapped excitons. This phenomenon is ascribed to the large distortion of [PbBr₆]⁴⁻ octahedral motifs resulting from a structural phase transition.

5.10. Ferroelectricity and photochemical reactions

Zhang et al. constructed 3D organic–inorganic perovskite ferroelectrics by the symmetry reduction of spherical moieties⁴⁹⁷ Halide perovskite quantum dots could be used in photochemical conversion (e.g. water splitting or CO₂ reduction). Xu et al. constructed CsPbBr3 quantum dot/graphene oxide composite for photocatalytic CO₂ reduction. Introduction of the conductive graphene oxide caused a photocatalytic enhancement.⁴⁹⁸

6. Conclusion and outlook

Perovskite materials are arousing great research interests due to their excellent luminescence performance. Perovskites have the general formula of ABO₃ and occur both naturally (e.g. CaTiO₃) and synthetically, and can form a wide-range class of derivatives. According to their crystal structure, they can be divided into three sub-classes: standard perovskites, low-D perovskites and perovskite-like halides. We rationalized the luminescence behaviours and various crystal structure types in oxide perovskites and halide perovskites to probe for a potential relationship between the crystal structure, electronic structure and luminescent properties. Perovskite nanomaterials have shown high light emission efficiency with a decrease in their physical dimensionality or crystal structure dimensionality. The electronic migration is more restricted in these cases, leading to a larger exciton binding energy and higher PLQY. Based on the high PLQY of perovskite materials, their use in fabricating illuminating device opens a broad application field. Various material stability and PL efficiency improving methods have been developed by researchers, including surface passivation, encapsulation and doping, to pave the way for practical applications. This review summarized previous research on the structural design, optical property characterizations and property improvements of perovskite-related compounds, but there are still a lot of future research opportunities in this field. At present, the intentional design of high PLQY materials is still challenging, and more efforts are needed to further clarify the relationship between the crystal structure, electronic structure and luminescence properties in the future.

Based on the above, the future development of luminescent perovskites can be focused on the following challenges: 1. gaining a deep understanding of the luminescent properties. For transition-metal and rare-earth elements-doped oxide perovskites, it is important to establish the relationship between the energy level structure and chemical composition/crystal structure by analyzing the crystal-field and nephelauxetic effects. For organic–inorganic hybrid and all-inorganic perovskite nanocrystals, some basic luminescent mechanisms need clarification. For example, there are conflicting reports on the luminescence of CsPbBr₃/Cs₄PbBr₆, and reports where the relative ordering of dark and bright exciton sublevels in the halide perovskites CsPbX₃ and FAPbBr₃ is just the opposite. 2. Improving the surface stability and luminescent efficiency by doping engineering. This is promising to explore the effects of transition-metal and rare-earth elements on surface passivation, and on the defects repair and surface healing of organic–inorganic hybrid and all-inorganic perovskite nanocrystals. 3. The discovery of new Pb-free all-inorganic perovskite nanocrystals. Materials that are less toxic while retaining their luminescence will gain much practical applications in the future.

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

This work is supported by the National Natural Science Foundation of China (No. 51832005, 51602019 and 51702329), Beijing Natural Science Foundation (2182080) and Fundamental Research Funds for the Central Universities (FRF-TP-18-005A1 and FRF-TP-18-002C1).

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