Recent progress in Ce³⁺/Eu²⁺-activated LEDs and persistent phosphors: focusing on the local structure and the electronic structure

Abstract

Ce³⁺/Eu²⁺ activated luminescent materials offer a versatile platform for precise emission light manipulation through structure control on the basis of the composition–structure–property correlations. To date, Ce³⁺/Eu²⁺ activated phosphors have been well developed as an indispensable component in the lighting industry and display systems due to their superior performance. Meanwhile, many persistent phosphors contain Ce³⁺/Eu²⁺ together with other lanthanide or transition-metal co-dopants. It is therefore of great importance to focus on their similarities and gain insight into the interplay effect of the local structure and electronic structure on emission peak modulation and persistent duration elongation. Here, we review the theoretical and experimental progress in the discovery and optimization of Ce³⁺/Eu²⁺ activated LEDs and persistent phosphors. The Dorenbos model on f–d transitions and the latest developments in the correlation of the local structure and luminescence characteristics are elaborated to give an overall vision on the composition–structure–property correlations in Ce³⁺/Eu²⁺-activated phosphors. Particular attention is devoted to highlighting the critical role of the electronic structure in tuning the properties of phosphors. The development and optimization routines of some typical phosphors are expounded, with an emphasis on phosphor design principles, aiming at providing inspirations for tailoring and optimizing the properties of Ce³⁺/Eu²⁺-activated phosphors toward specific applications. Finally, we propose an outlook toward potential theory developments and future material discovery.

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Shuxin Wang

Shuxin Wang (S. X. Wang) received his PhD degree in Materials Science and Engineering from the University of Science and Technology Beijing (USTB) in 2021. From Dec. 2019 to Jan. 2021, she studied in Prof. Seshadri Ram's group at the University of California, Santa Barbara, as a visiting scholar. Now, she is a postdoc in Prof. Wai-yeung Wong's group in the Department of Applied Biology and Chemical Technology at The Hong Kong Polytechnic University. Her current research interests are focused on luminescent materials.

Zhen Song

Zhen Song (Z. Song) received his PhD degree in Materials 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. Now, he is an associate professor 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 theoretical aspects of luminescent materials.

Quanlin Liu

Quanlin Liu (Q. L. Liu) completed his PhD 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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Introduction

Phosphor-converted light emitting diodes (pc-LEDs) have made their mark in the lighting industry and display systems as an indispensable solid-state light source driven by their unique properties including high luminescence efficiency, small volume, long lifetime, fast switching, and excellent durability.^1–6 The concept of pc-LEDs, assembling blue/violet LED chips with phosphors to partially down-convert some of the blue/violet emission to longer wavelengths corresponding to colors such as green, yellow, and red, can be traced back to Bando's work in 1996,⁷ shortly after candela-class blue LEDs were first reported.⁸ The first commercially available pc-LED device was invented by Nichia Corporation, utilizing yellow emitting YAG:Ce (a cerium-doped yttrium aluminum garnet; Y₃Al₅O₁₂:Ce³⁺) overcoated onto a blue InGaN LED chip.⁹ Since then, the pc-LED field really took off, and various fabrication strategies for creating white light were proposed.¹⁰ Meanwhile, tremendous attention has been focused on exploring high performance phosphors for white LED (w-LED) and LED backlights. Phosphors typically use optically active lanthanide or transition-metal ions as emission centers embedded in an inorganic matrix. The general requirements for appropriate LED phosphors include the following: (i) a broad excitation spectrum matching well with the emission spectrum of a LED chip; (ii) a suitable emission spectrum; (iii) high quantum efficiency; (iv) small thermal luminescence quenching; and (v) high chemical stability.^11–13 Many of these requirements have been met by the discovery and development of new phosphors over the past decades including a variety of Ce³⁺/Eu²⁺-activated silicate,^14,15 aluminate,^16,17 nitride etc.,^18,19 leading to commercial LEDs that cover a full range of white CCTs (correlated color temperatures) (Fig. 1a). The parity-allowed electric dipole f–d transitions of Ce³⁺ and Eu²⁺ are of particular interest due to their unique luminescence properties including a broad emission band and high efficiency.^20–22 Besides its high performance, the strong interaction of the 5d-electron with the neighboring anion ligands in the host lattice not only grants access to tune the luminescence properties by structural modulation but also provides feasibility to design new phosphors based on composition–structure–property correlations, which forecasts great potential for more “smart” pc-LED devices.

Fig. 1 (a) Commercial LED phosphors for w-LEDs using blue chips. (b) Important persistent phosphors with promising or potential PersL properties.

On the basis of LED phosphors, by further introducing appropriate trap centers that can capture excited electrons and then release them by thermal assistance, mysterious persistent luminescence (PersL) can be achieved.²³ Compared with LED phosphors, persistent phosphors have distinctive properties which can give continuous luminescence for seconds, minutes, hours, or even days after ceasing the stimulations.²⁴ Owing to their unique optical performance, persistent phosphors are emerging as promising materials for safety signage, in vivo bioimaging, and energy storage applications in the past decades.^25–28 The PersL materials have a rich and longstanding history that can be traced back to as early as 1000 years ago, when they were first used as special “night-vision” inks for ancient Chinese paintings.²⁹ The landmark breakthrough occurred in 1996, when Murayama et al. reported the green persistent phosphor SrAl₂O₄:Eu²⁺,Dy³⁺ with extremely bright and long duration PersL in the dark (over 30 h before the emission intensity drops to 0.32 mcd m⁻², 100 times higher than the sensitivity of the dark-adapted human eyes).³⁰ Since then, research on persistent phosphors has become increasingly popular. Extensive studies have been devoted to discover new PersL materials, tune their properties, and explore the fundamental mechanism, aiming at broadening the application scope and realizing the ultimate “dream” of night lighting. To date, a variety of persistent phosphors have been discovered, and some important ones with promising or potential PersL properties are presented in Fig. 1b. It is worth noting that the majority of the persistent phosphors are originally from conventional phosphors. By controlling defect concentration and intentionally introducing aliovalent or isovalent co-dopants as trap centers, their intensities and afterglow durations can be effectively tuned, such as Y₃Al₂Ga₃O₅:Ce³⁺,Yb³⁺ and M₂Si₅N₈:Eu²⁺,Tm³⁺ (M = Ca, Sr).^31–33 The potential PersL possibilities of LED phosphors make them particularly interesting for the discovery of new promising PersL materials.

Generally, luminescence is determined by the average and local crystal structures of phosphors. Unlike band-to-band optical transitions, f–d transitions from isolated luminescent centers such as Ce³⁺/Eu²⁺ are strongly influenced by the coordinated crystalline environment.²⁰ Within the local structure, energy levels of Ce³⁺/Eu²⁺ may be split and shifted, resulting in peak shifts observed in both photoluminescence (PL) and photoluminescence excitation (PLE) spectra.^34,35 So, it is necessary to deeply understand the relationship between the local structure and luminescence. Recently, Xia's group developed several Eu²⁺-doped oxide phosphors, such as NaLi₃SiO₄:Eu²⁺, KSrScSi₂O₇:Eu²⁺, and Sr₂Li(Al,Ga)O₄:Eu²⁺. Through gaining insight into the local structure, the luminescence mechanism was well explained, and the luminescence properties were successfully engineered at will via local structure modification.^36–46 On the other hand, the luminescent properties of materials are strongly dependent on their electronic structures; while the electronic structure is mainly determined by the crystal structure. To fully understand the luminescence, especially the thermal luminescence quenching behavior, it is necessary to investigate the electronic structure of rare-earth-doped luminescent materials. Recently, Dorenbos has proposed a method to construct the host referred binding energy (HRBE) and vacuum referred binding energy (VRBE) schemes, showing all the information regarding lanthanide energy levels and host bands.^21,47,48 The electronic structure based on VRBE schemes is of vital importance to deeply understand Ce³⁺/Eu²⁺-activated LEDs and persistent phosphors.

The main object of this review is to provide a detailed and deep understanding of well-known Ce³⁺/Eu²⁺-activated LEDs and persistent phosphors, from the view point of the local structure and the electronic structure, to highlight the important design considerations for new high-performance phosphors and PersL materials. As we focus on the composition–structure–property correlations, we first introduce basic theory on f–d transitions to specify how the chemistry and structure affect the luminescence properties. Then the PersL mechanism for persistent phosphors is illustrated, which is followed by demonstrating the electronic structure diagram as a powerful design tool towards promising properties. We continue the discussion on the widely used LED phosphors. Based on the crystal structure, we analyze their luminescence properties, ascribe physical quantities to structural features (crystal and electronic structure) and extract the composition–structure–property correlations. A parallel focus is placed on the progress of converting LED phosphors into persistent phosphors by structure modification by virtue of electronic structure engineering. We conclude with a brief discussion on the prerequisites for LED and persistent phosphors and provide an outlook toward promising research directions and suggestions for future developments.

Basic theory on Ce³⁺/Eu²⁺-activated phosphors and persistent luminescence

Structural features are critical to physical properties. Understanding and developing the composition–structure–property correlations are the core of materials science, which lead mankind to materials revolution from time to time. In the past decade, scientific advances in experiments have enabled abundant theoretical insights into the rational design of high performance Ce³⁺/Eu²⁺-activated inorganic luminescent materials. Experimental techniques providing information about the local coordination environment have been developed, such as high-resolution transmission microscopy (FRTEM),⁴⁹ extended X-ray absorption fine structure (EXAFS),⁵⁰ solid-state nuclear magnetic resonance⁵¹ and Mössbauer spectroscopy.⁵² It has been widely acknowledged that luminescence properties are determined by the average and local structures of phosphors. Unlike band-to-band optical transitions, f–d transitions from the isolated luminescent center Ce³⁺/Eu²⁺ are strongly influenced by the coordinated crystalline environment, under which energy levels of Ce³⁺/Eu²⁺ may be split and shifted, resulting in their highly tunable luminescence properties. On the other hand, the electronic structures, deriving from the crystal structure, with information regarding dopant energy levels and host bands have superiority in gaining insight into the thermal luminescence stability and PersL properties. We discuss here the intrinsic properties of Ce³⁺/Eu²⁺-activated LED and persistent phosphors, and the effects of the host crystal and the dopant ions on phosphor optical properties.

Dorenbos model on f–d transition

In general, the luminescence process of Ce³⁺/Eu²⁺-doped phosphor can be divided into energy absorption, non-radiative transition and photon emission.^53,54 Upon excitation, absorption occurs due to the transitions from the 4f ground state to the 5d excited states. Then, due to the coupling of the 5d electron with lattice phonons, partial energy is lost by non-radiative transition. Finally, emission occurs by the transition from the lowest 5d excited state to the 4f ground state. Therefore, the relative energy difference between the 4f and 5d levels dictates the excitation energy. Different from the invariable energy levels of 4f orbitals shielded by the filled 5s²5p⁶ sub-shells, the energy levels of unshielded 5d electron can be tuned by tens of thousands of wave numbers from one compound to another benefiting from their subtle response to the host lattice, resulting in their crucial effect on excitation properties.²⁰ The effect of the host lattice on 5d energy levels is typically characterized by two parameters, i.e., the centroid shift ε_c and the crystal field splitting ε_cfs, as illustrated in Fig. 2. Please note that Fig. 2 is a simplified schematic plot. Actually, for Eu²⁺ with seven 4f electrons, the situation is more complicated due to its more excited states. The details could be found in ref. 55. In the host, the 5d levels shift toward lower energy due to a decrease in the interelectron repulsion. Moreover, the degenerate 5d levels split into at most five different 5d states depending on the site symmetry for the activator ion. The downshift of the average energy of the five 5d levels is the centroid shift ε_c and the energy difference between the lowest 5d₁ and the highest 5d₅ energy levels defines the crystal field splitting ε_cfs.^56,57 The overall effect of centroid shift ε_c and crystal field splitting ε_cfs leads to a decrease in energy difference between the 5d and 4f levels, directly determining the excitation energy. Based on the excitation energy, the emission energy is further determined by the Stokes shift ΔS, defined as the energy difference between the absorption and emission maxima of the same electronic transition, resulting from the surrounding lattice relaxation due to the coupling of the 5d electron with the lattice phonons upon the excitation of the activator from the 4f to the 5d configuration.^58,59 In the past few decades, a series of sound theories have been established to gain insight into the composition–structure–property correlations and account for the physical nature behind them. With the rapid development of first principles calculation, some theoretical models have been established to finely understand the luminescence properties. In addition, as an increasing number of phosphors were developed, the effects of host crystal on the luminescence properties have been further qualitatively or quantitatively demonstrated by analyzing the f–d luminescence properties in more than 300 Ce³⁺/Eu²⁺-activated phosphors by virtue of the phenomenological approach proposed by Dorenbos.²⁰ Detailed analysis and results are summarized below.

Fig. 2 A schematic energy level diagram for Ce³⁺/Eu²⁺ shows the effects of the host crystal, including the centroid shift, crystal field splitting and Stokes shift. Such effects lead to an overall decrease in the energy difference between the 5d and 4f levels. Reproduced with permission from ref. 60, copyright 2017 American Chemical Society.

Centroid shift. Once Ce³⁺/Eu²⁺ ions are introduced into a crystalline host, the average position of the 5d levels of Ce³⁺/Eu²⁺, i.e., the centroid, is lowered relative to that of the free ion. The downward shift of the centroid position of the five 5d-levels relative to that of the free ion is defined as the centroid shift ε_c. The centroid shift is commonly associated with the expansion of the charge cloud, i.e., the nephelauxetic effect, which is often attributed to the covalency between the 5d orbital and the p-orbitals of the anions. The covalency effect results in a larger average distance and hence a reduced Coulomb repulsion between the electrons of the lanthanide cation, giving rise to the centroid shift phenomenon.⁶¹ In 1980 Morrison first suggested a physical origin for the centroid shift involving ligand polarization, addressing the vital contribution of the correlated motion of 5d electron and anion ligand electrons to the centroid shift.⁶² The instantaneous position of the metal electron polarizes the surrounding ligands, which reacts back on the metal electron itself. Essentially a self-induced potential is generated that reduces the interelectron Coulomb repulsion between the metal electrons. It leads, like the covalency effect, to a lowering of the centroid energy. Models on covalency predict in the first order that a centroid shift is proportional to the S ≡ 〈ϕ_M|ψ_L〉 square of the overlap integral between the metal (M) and ligand (L) orbital as shown in the following formula:^61,63where a_i and b_i are appropriate constants. From the perspective of ligand polarization model, based on the model suggested by Morrison⁶² and Aull and Jenssen,⁶¹ centroid shift ε_c can in first approximation be written aswhere r represents the radial position of the electron in either 5d or 4f orbital, and 〈r²〉 is the expectation value of r², α_i is the polarizability of the ligand i located at a distance R_i from the metal ion, e is the elementary charge, and ε₀ is the permittivity of vacuum. The summation is over all the N nearest coordinating anion ligands. As shown in Fig. 3a, it is experimentally observed that the centroid shift tends to increase following the nephelauxetic sequence of the anion and to decrease following the nephelauxetic sequence of the most electronegative cation (the neighboring small cation of the lanthanide ion) in the compounds as shown below, proposed by Dorenbos⁶⁴

Fig. 3 (a) The centroid shift for Ce³⁺ 5d configuration in halides and oxide compounds. Reproduced with permission from ref. 64, copyright 2012 The Electrochemical Society. (b) Spectroscopic polarizability against the inverse square of the average cation electronegativity in Ce³⁺ doped nitrides, oxides, and fluorides. Reproduced with permission from ref. 65, copyright 2015 Elsevier B.V. (c) The revised correlation between the spectroscopic polarizability and the average cation electronegativity in Ce³⁺ doped nitrides. (d) The relative energy of the 5d level centroid and the lowest and highest 5d levels in Ce³⁺ doped nitrides. Panels c and d reproduced with permission from ref. 66, copyright 2018 Elsevier B.V.

It can be understood from the viewpoint of charge cloud attraction, which can be comprehended as a form of bonding that increases the binding energy of the anion electrons. A weaker nephelauxetic effect of the anion ligand leads to the stronger binding with the lanthanide ion, which implies a larger oscillation force constant and therewith smaller anion polarizability, resulting in a smaller centroid shift according to eqn (2). A stronger nephelauxetic effect of the neighboring cation (the second nearest atom) leads to a stronger binding with the anion ligand, which also reduces the covalency between the ligand charge cloud and the 5d orbital of lanthanide ions behaving as a more ionic surrounding. In this case eqn (1) predicts a smaller centroid shift. Overall, strong binding lowers the electron donating power of the anion ligand toward the lanthanide ion, which is the original interpretation for not only the centroid shift but also the nephelauxetic effect.

The centroid shift has been investigated by Dorenbos using a phenomenologically parameterized model with the inclusion of spectroscopic polarizability, α_sp, which merges together the effect of (1) ligand polarization, (2) covalency, and (3) possible charge cloud expansion.³⁵ On the basis of the assumption that (1) the total centroid shift is the result of the added contribution of each coordinating anion individually; (2) only the nearest-neighbor anions give a significant contribution to the centroid shift; and (3) polarizability of all the ligands is identical, eqn (2) can be re-written as⁶⁷

where R_j (pm) is the distance from Ce³⁺ to the anion j in the unrelaxed lattice. The summation is overall N coordination anions of Ce³⁺. ΔR ≡ R_M − R_Ln, where R_M is the ionic radius of the cation which is replaced by Ce³⁺ with the ionic radius R_Ln. 0.6ΔR is an estimation of the bond length relaxation. The most appealing aspect of this model is that it is a zero-parameter model, and it provides a direct quantitative and qualitative physical interpretation of the centroid shift. Pearson's work reported a proportional relationship between the force constant K of the metal–hydride bond in diatomic MH molecules and the value for the electronegativity χ of the metal atom.⁶⁸ Since in atomic physics the polarizability α of an atom is related to the binding force constants K_i of its electrons as , the correlation between spectroscopic polarizability α_sp and the average cation electronegativity χ_av was constructed as follows:³⁵where α₀ is the limiting polarizability in the case of very large χ, and b represents the susceptibility of the anion to change its polarizability due to the bonding with the metal ion. The average cation electronegativity χ_av is defined as follows:where n_i is the number of the cation i with charge z_i and electronegativity χ_i in the chemical formula.

On the basis of the theory above, employing the phenomenological approach, Dorenbos obtained the quantitative correlations between α_sp and χ_av as shown in Fig. 3b by analyzing the spectral and structural data in Ce³⁺-activated oxide and fluoride compounds.^57,64,69 Over the past few decades, Ce³⁺/Eu²⁺-activated nitride and oxonitride phosphors have attracted tremendous attention due to their unique luminescence properties and potential application in pc-LEDs.^70–73 The bonding in nitrides is significantly more covalent than that in oxides, because of the higher formation energy of N³⁻ from atomic N (+2300 kJ mol⁻¹) than that of O²⁻ from atomic O (+700 kJ mol⁻¹).⁷⁴ Compared with oxide hosts, nitride and oxonitride hosts enable a more downward shift of 5d energy levels, giving rise to a large centroid shift ε_c. In this way Ce³⁺/Eu²⁺-activated nitride phosphors can absorb blue LED radiation and then emit visible light effectively, which are promising for pc-LED applications. In addition, nitridosilicates and their derivatives, as an important class of nitrides, typically consist of SiN₄ tetrahedra, in which a partial nitrogen/silicon can be substituted by oxygen/aluminum to form Si/Al[O/N]₄ tetrahedra.⁷⁵ These tetrahedron units stack together by corner sharing to form the condensed framework, generally contributing to the chemical and thermal luminescence stability of nitridosilicates.⁷⁶ Furthermore, these tetrahedron units take various stacking topologies providing Ce³⁺/Eu²⁺ with abundant coordination environments, allowing the Ce³⁺/Eu²⁺-activated nitridosilicate phosphors to be highly tunable. With the aim of approaching a guideline for further development of Ce³⁺/Eu²⁺-activated nitridosilicate phosphors, our group works on exploring the α_sp and χ_av correlations in Ce³⁺-activated nitrides based on Dorenbos’ work.⁶⁵ As expected from the nephelauxetic effect, the tendency of b^F < b^O < b^N can be observed as demonstrated in Fig. 3b, which means that α_sp is increasingly sensitive to χ_av in the sequence of fluorides, oxides and nitrides. This sequence provides a guideline for phosphor design that the larger ε_c and thereby the redshift of the spectrum can be realized by the substitution of the anion with a stronger nephelauxetic effect. The structural coordination data, as the critical structure parameter, play a vital role in studying the composition–structure–property correlations, such as the ε_c analysis. However, the determination of coordination number for Wyckoff sites in the crystal structure lacks explicit criteria, which results in that the coordination number of the cations in some compounds is ambiguous and sometimes there are conflicts for coordination number in different reports regarding certain compound.^{18,19,77–80} In order to clear it up, our group further proposed a uniform standard to determine the coordination number of cations in nitrides according to the bond valence theory and the requirement for stability of the coordination polyhedron.⁶⁶ In general, the nitrogen anions which contribute more than 4% to the bond valence of the central cation are considered as the coordination atom. The threshold value can be adjusted on a small scale in order to obtain a higher stability of the coordination polyhedron. Based on the revised structural data, supplementing data in the latest nitride phosphors, our group re-investigated ε_c of the 5d configuration of Ce³⁺ in the nitride compounds. The updated correlation between α_sp and χ_av is displayed in Fig. 3c, the credibility of which is ensured by the crystal field splitting ε_cfs data as shown in Fig. 3d. With the above results, one can predict the 5d centroid shift of Ce³⁺ in phosphors based on the composition and structure of the compounds.

Nevertheless, we would like to remind the readers of the overestimation of the centroid shift in garnet series, as pointed out by Seijo and Barandiarán.⁸¹ They performed spin–orbit coupling, relativistic, embedded cluster, wave function-based ab initio calculations on the (CeO8)¹³⁻ cluster under the effects of the embedding potentials of garnets and obtained significantly smaller centroid-shift values with regard to those estimated using the Dorenbos model. Since the ab initio results are more advocated by the garnet family, the completeness of the Dorenbos’ model needs further improvement by elucidating some coefficients.

It is worth noting that the research on ε_c of the 5d configuration of Eu²⁺ is missing, which is caused by the difficulty in the identification of the transition energy from the 4f⁷ ground state to the five 4f⁶5d excited states. In the f–d excitation spectra of Eu²⁺ always about 0.8 eV-wide bands appear because the 6 electrons remaining in the 4f-shell may occupy one of the seven ⁷F_J states that spread about 0.6 eV in energy.^82,83 They overlap with the transitions to the 4f⁶5d₂ and higher 4f⁶5d_i states, smearing out into a featureless 1 eV broad band, which makes it difficult to identify the energy of the five f–d transitions for developing ε_c theory in Eu²⁺-activated phosphors. In a recent paper, Joos et al. applied state-of-the-art multiconfigurational ab initio embedded-cluster methods to investigate the staircase structure in more detail.⁵⁵ Other than the decoupled scheme, which attributes the fine structure to the J = 0–6 levels of the 4f⁶ (⁷F_J) subshell, it is the interplay between the spin–orbit coupling within the 4f⁶ (⁷F_J) subshell and the 4f–5d exchange splitting that presents the fine structure.

Crystal field splitting. Ce³⁺/Eu²⁺-activated phosphors have the advantage of parity-allowed electric dipole f–d transition, giving rise to their broad and highly efficient emission. The crystal field splitting ε_cfs of 5d level further broadens the excitation spectra, which allows that Ce³⁺/Eu²⁺-activated phosphors generally could be excited by lower energy irradiation such as blue light. Crystal field splitting is caused by Coulomb interactions between the 5d electron and anion ligands. By analyzing the luminescence and structural data of more than 60 Ce³⁺/Eu²⁺-activated phosphors including halide, oxide, sulfide, and selenide compounds, Dorenbos revealed that the crystal field splitting ε_cfs is related to the polyhedral shape and size, while irrelevant to the types of the anions, as shown in Fig. 4a and b.³⁴ For the compounds with the same shape of coordination polyhedron around Ce³⁺/Eu²⁺, the crystal field splitting appears to behave as follows:

ε_cfs = β^Q_polyR_av⁻², (7)

where β^Q_poly is a constant determined by the type of the coordination polyhedron. R_av is defined aswhere R_i is the individual bond length of the N coordinating anions in the unrelaxed lattice. ΔR ≡ R_M − R_Ln, where R_M is the ionic radius of the cation which is replaced by the lanthanide Ln with an ionic radius R_Ln. 0.6ΔR is an estimation of the bond length relaxation. By fitting the data in Ce³⁺/Eu²⁺-activated phosphors with 6-fold octahedral (octa), 8-fold cubal (cubal), and 12-fold cuboctahedral (cubo) coordination through eqn (7), β_poly was obtained. As an example, β_octa is 1.36 × 10⁵ eV pm² for Eu²⁺, and β_octa : β_cubal : β_cubo equals 1 : 0.89 : 0.44 for both Ce³⁺ and Eu²⁺. It is revealed that higher coordination number tends to reduce the crystal field splitting, while the type of the anion, whether it is F, Cl, Br, I, O, S or Se, does not matter at all.³⁴ Furthermore, comparing the Eu²⁺ε_cfs data with the Ce³⁺ data shows that⁸⁴

β_poly²⁺(Eu²⁺) = 0.81β_poly³⁺(Ce³⁺). (9)

Fig. 4 The crystal field splitting of the 5d configuration of (a) Ce³⁺ and (b) Eu²⁺ and (c) their correlation in phosphors including halide, oxide, sulfide, and selenide compounds. Panels a–c reproduced with permission from ref. 34, 64, 84, copyright 2002 Elsevier, 2012 The Electrochemical Society, and 2003 Institute of Physics Publishing. The crystal field splitting of the 5d configuration of (d) Ce³⁺ and (e) Eu²⁺, and (f) their correlation in nitride phosphors. Panels d–f reproduced with permission from ref. 85, copyright 2017 Elsevier B.V.

Taking the 12 pm larger size of Eu²⁺ into account, the correlation between Ce³⁺ and Eu²⁺ crystal field splitting ε_cfs was further achieved (Fig. 4c)⁸⁴

ε_cfs(Eu²⁺) = 0.77ε_cfs(Ce³⁺). (10)

As the emerging of nitride phosphors, our group then branched out to explore the structure-crystal field splitting ε_cfs correlations in Ce³⁺/Eu²⁺-activated nitrides. By analyzing the luminescence and structural data of tens of Ce³⁺/Eu²⁺-doped nitrides, the relationship between the crystal field splitting ε_cfs and the shape and size of the coordination polyhedron in nitride compounds is obtained as shown in Fig. 4d and e, which is consistent with the results of halide, oxide, sulfide, and selenide compounds, further confirming that ε_cfs is irrelevant of anion types.⁸⁵ The crystal field splitting ε_cfs behaves as shown in eqn (7), and decreases with the coordination number increasing. A linear relationship as shown in eqn (10) also exists in nitride phosphors with a similar multiplication factor of 0.76 (Fig. 4f). In summary, with these quantitative correlations, the crystal field splitting ε_cfs in Ce³⁺/Eu²⁺-activated inorganic phosphors can be theoretically predicted based on the crystal structure.

Electronic structure. Both the luminescence and persistent properties of lanthanide-doped phosphors are mainly determined by the electronic structure, specifically the energy of the lanthanide excited and ground states with respect to the band states of the host.^86–88 With the aim of illustrating the electronic structure, Dorenbos proposed the concept of host referred binding energy (HRBE) and vacuum referred binding energy (VRBE) schemes, which provide energy information of all the lanthanide levels and host bands with sufficient accuracy.^21,47,48 With the assumption that the valence band (VB) maximum energy is 0 eV, HRBE schemes provide the relative energy of 4f and 5d levels of divalent and trivalent lanthanides with respect to the band states of the host. On the basis of HRBE schemes, Dorenbos further proposed the chemical shift model, which relates the host referred binding energy with the vacuum referred binding energy, giving rise to the VRBE schemes representing the absolute energy of 4f and 5d levels of divalent and trivalent lanthanides as well as the band states of the host.^89,90 Detailed construction instructions of HRBE and VRBE schemes have been elaborated in ref. 60, 86, 91. As an example, Fig. 5 shows the HRBE and VRBE schemes for Y₃Al₅O₁₂ (YAG), in which the red and blue zigzag curves represent the energy for the 4f ground states of divalent and trivalent lanthanide ions, respectively, and the other red and blue curves represent the energy for the lowest 5d states of divalent and trivalent lanthanides, respectively. The HRBE and VRBE schemes enable us to interpret, predict, and engineer the properties of phosphors including but not limited to the thermal luminescence stability, trap depth introduced by lanthanides, the preferred valence of the lanthanides, and the presence or absence of emission.^86,92,93 On the basis of the relative energy between the 5d states of the Ce³⁺/Eu²⁺ emission center and the CB edge of the host, one can rationally select a suitable host lattice with appropriate ionization energy for LED or persistent phosphors. In addition, the energy of the 5d states of divalent lanthanides gives information on the depth of the traps introduced by trivalent lanthanides, which provides a feasibility to rationally choose the trap center for persistent phosphors.^94–96 As a summary, HRBE and VRBE schemes make it possible to systematically manipulate the LED and persistent phosphor properties by tuning the energy of the CB edge via cation substitution in the host and co-doping trivalent lanthanides as trap centers, which are the so-called electronic structure engineering. The examples for the excellent application of electronic structure engineering will be presented in detail in Section 3.

Fig. 5 The HRBE and VRBE schemes of Y₃Al₅O₁₂. Red and blue ‘zigzag’ curves connect the energy for electrons in the 4f ground state of divalent and trivalent lanthanides, respectively; The other red and blue curves connect the energy for the lowest 5d states of divalent and trivalent lanthanides, respectively. Reproduced with permission from ref. 86, copyright 2012 Royal Society of Chemistry.

It is worth noting that even though the Dorenbos model is of vital importance for illustrating, predicting, and tuning phosphor properties, it has limitations. It rather parameterizes the trends than explaining the mechanism, and it fails to explain individual phosphors in detail. Therefore, further exploration of the correlation between the Dorenbos model and the luminescence theory is required to optimize the Dorenbos model and gain insight into its physical origins. Nevertheless, the empirical/phenomenological approaches bear the advantages of the simple form and easy-to-use style. Both beginners and veterans could gain useful insights from those simple models involving bond lengths, bond angles and constitutional components. Yet it can be only applied in a limited region, and fails to respond to the structure–property relationships across different material systems. On the contrary, the advanced first principles/ab initio approaches could provide more information on the dopant energy levels and electronic structure of the host. Besides, the orbital analysis and defect level calculation are unique and give more fundamental insights into the luminescence behavior. Although those advanced calculations are prevailing in recent publications, the huge consumption of computational resources still remains a barrier for most researchers. Meanwhile, some user-friendly automatic DFT software may undermine the expertise needed for the calculation process.

Effects of the local structure on d–f luminescence

Crystal field. The emission and absorption of light by a phosphor originate from the electronic transitions between the discrete energy levels of the luminescent center embedded in the host. The energy level distribution is determined by the coulomb interactions between nuclei/electron and electron/electron, together with the Stark effect of the valence electron exerted by the surrounding ligands. Rare-earth-doped phosphors employing f–d electronic transitions exhibit intensified oscillator strength from luminescent center, such as Ce³⁺ and Eu²⁺, compared to those of f–f or d–d transitions. For rare-earth ions, 4f electrons are screened by the outer-shell 5s²5p⁶ electrons. Therefore, the energy levels are insensitive to the subtle change in the ligand environment. Unlike 4f electrons, the 5d electron is directly exposed to the crystalline environment and its energy levels could be tuned by the modifications of the surrounding ligands. As a result, the versatile host structures have a significant impact on 5d energy levels, and the emission spectrum of f–d transitions could reproduce the whole spectral region from ultraviolet to far red. The effect of crystalline ligand environment on luminescence tuning constitutes the main topic of crystal-field theory.

Crystal field theory has successful explanations in high-symmetry ligand polyhedra, such as ideal octahedron (6-coordination) and cube (8-coordination), both of which share the point symmetry O_h.^97,98 The detailed analytical description is derived in the framework of perturbation theory. First, the crystal-field potentials of one d electron exerted by the surrounding ligands of octahedron and cube are⁹⁹

respectively, with the substitution . In the above formula, (x,y,z) stands for the spatial coordinates of the d electron, r for the distance between electron and the nuclei, and a for the center–ligand distance. In order to calculate the perturbation matrix elements , the basis function could be selected as one d electron wavefunction, with an angular part |2,m_l〉, where m_l stands for the magnetic quantum number (±2, ±1 and 0 in the case of l = 2). With further substitution of , in which R_nl stands for the radial part of the d electron wavefunction, the perturbation matrix could be constructed in the form of D and q product, and the diagonalization result gives the splitting of 5-fold degenerate energy levels by the Stark effect due to the electrostatic interaction with ligand ions. It is interesting to note that the energy level splitting schemes are −4Dq (triple) and 6Dq (double) for octahedron and −16/3Dq(double) and 32/9Dq(triple) for the cube. Octahedron has the largest crystal-field splitting of 90/9Dq, whereas cube has a smaller one of 80/9Dq. In addition, in the case of ideal cube two orbitals are positioned in the lower energy scale in comparison to the three degenerate orbitals in the octahedron. Since both are proportional to D, then a reciprocal 5th-power relationship could be obtained between the total crystal-field splitting and the center–ligand distance. This simple rule paves the way for explaining the energy level shift from the viewpoint of the ligand environment and serves as a bridge between the crystal structure and the luminescence properties.

The 5th-power rule seems to provide an ultimate answer to the structure–property problem in luminescence. Under this simple and clear guidance, only the 1D bond length magnitude instead of the 3D complexity of crystal structure needs to be considered. This rule, however, meets astonishing contradictory cases in a ligand polyhedron with a local symmetry lower than that of O_h. The structure–property relationship itself proves to be not only a mere bond-length comparison game, as will be demonstrated in the garnet-type phosphors.

Garnet-type phosphors constitute a large family of commercial phosphors. It has a remarkable capability to accommodate various luminescent dopants including both rare-earth and transition-metal elements, such as Ce³⁺, Eu³⁺, Tb³⁺, Cr³⁺, etc. Among them, Ce³⁺ doped garnet phosphors are most widely used in the field of wLEDs because of the relatively cheap cost, chemical stability, and convenience for mass production. The luminescence of Ce³⁺ in the garnet host originates from the parity-allowed d–f transitions, which was first reported by Blasse et al. in 1960s.¹⁰⁰ In the garnet structure, the 8-ligand coordinated dodecahedral (A) site with point symmetry D₂ accommodates Ce³⁺. In this low-symmetry crystal-field potential, the 5-fold degenerate 5d energy levels will be totally split. The energy level scheme could be understood as first split by a T_d crystal-field potential, with ²E (2-fold, low-lying) and ²T₂ (3-fold, higher) levels, and then further split by a D₂ crystal-field potential.^101–103 By composition substitution, the coordination environment of the dodecahedral site could be modulated and the luminescence of the Ce³⁺ doped garnet could be tuned. However, inspection of excitation spectra of Ce³⁺ in the garnet series of Y₃Al₅O₁₂(YAG), Y₃Ga₅O₁₂(YGG), Gd₃Al₅O₁₂(GdAG), Gd₃Ga₅O₁₂(GdGG), Lu₃Al₅O12 (LuAG), Lu₃Ga₅O₁₂(LuGG) and their solid solutions presents a contradictory relationship between the crystal structure and crystal-field splitting.^104–107

Composition substitution takes effect due to the difference in the ionic radius. For the above-mentioned garnet series, it has been confirmed that the cell parameters, volumes of dodecahedron (A) and mean bond-lengths have a consistent trend with the ionic radius of the substituted elements.¹⁰⁸ However, the crystal-field splitting demonstrates contradictory trends with the ionic radius. Here, for convenience the crystal-field splitting is represented by the energy difference between the two lowest 5d levels, which could be easily read out as two excitation peaks. When the cell is expanding from YAG to YGG due to the larger Ga, the crystal-field splitting is decreasing,¹⁰⁹ in accordance with the 5th-power rule. Yet from YAG to LuAG the crystal-field splitting is reduced with cell shrinkage.^110,111 The deviation from the 5th-power rule in garnet phosphors is named the reverse garnet effect.^5,112 It is surprising that the normal and reverse phenomena co-exist in the same material family.

Many efforts have been devoted to resolve this self-contradictory phenomenon. For example, the self-adjustment of the local structure on doping is considered as one of the reasons. DFT investigations on the structural relaxation of Ce³⁺ in the lutetium aluminum garnet shows longer Ce–O bonds, which could explain the smaller crystal-field splitting according to the 5th-power rule.¹¹³ Yet, the dilemma lies in that the 5th-power rule itself is derived without the consideration of local structure relaxation. In fact, the simple 5th-power rule only constitutes a small part of the crystal-field theory. It is derived under the highest-symmetry point group, i.e., O_h, which will eliminate all the 2nd rank crystal-field parameters due to the symmetry limitation in the crystal-field potential.^114,115 This special constraint leaves only one 4th rank crystal-field parameter, i.e. B₄₀, in the expression of the O_h crystal-field potential for d electron. Under the framework of the point-charge electrostatic model (PCEM), B₄₀ has the form of a 5th power exponential of bond length in the denominator. But further symmetry degradation will set the 2nd crystal-field parameters nonzero, which will introduce 3rd power exponential of bond length together with that of the 5th power. In the real D₂ site symmetry, the crystal-field potential could be expressed as¹¹⁴

The final analytical expression of the split energy levels will inevitably include the 2nd rank crystal-field parameters, and thus break the monotonous relationship between crystal-field splitting and center–vertex distance. With the D₂ crystal-field potential, the crystal-field analysis has been performed assuming the crystal-field strength in the sequence of cubic field ≫ the spin–orbit interaction ≫D₂ crystal-field,¹⁰³ to calculate the energy level splitting of Ce³⁺ doped garnets.^108,109 The results provide convincing explanation for the above-mentioned self-contradictory phenomena. For the aluminum and gallium series, the crystal-field splitting increases as Lu < Y < Gd. For LnAG/LnGG paris, i.e. LuAG/LuGG, YAG/YGG and GdAG/GdGG, it drops as Ga substitutes for Al. It can be concluded that the contradiction vanishes because both the normal and reverse phenomena originate from the crystal-field effect.

Although there have been similar responses of cell and polyhedral volumes on larger ion substitution at different cationic sites, subtle differences exist in how the oxygen ligands respond to different site substitutions. The ligand movement has two distinct patterns with regards to site substitution, as illustrated by the example of evolvements from LuAG to YAG and LuAG to LuGG in Fig. 6a. Further investigation shows that these two patterns are closely related to the squeeze or stretch of the octahedron (16a site), and the interplanar distance of the octahedron is selected to characterize the multi-site substitution. Meanwhile, the distance has a simple expression of , in which L is the cell parameter and (x, y, z) stands for the oxygen coordinates resembling the form of (−0.0306, 0.0506, 0.1493).¹¹⁶ This proves to be an effective indicator of the crystal-field splitting of Ce³⁺ in the garnet structure, as shown in Fig. 6b.

Fig. 6 (a) Left: ligand movement pattern for LuAG to YAG; right: ligand movement pattern for LuAG to LuGG. The arrows indicate the moving direction of ligands when the garnet compound changes. The magnitude of the arrow has been magnified 30 times for better visualization. Reproduced with permission from ref. 108, copyright 2018 American Chemical Society. (b) Crystal-field splitting as a function of octahedral interplanar distance for different garnet series. The inserted graph shows octahedron and its ligands numbered from 1 to 6, with the fractional coordinates aside. The centers of the two regular triangles are denoted as small purple dots. A dotted purple line connecting the two centers represents the interplanar distance. Reproduced with permission from ref. 116, copyright 2019 American Chemical Society.

The validity of the octahedral interplanar distance as the structural indicator stems from its relation to the tetragonal distortion of the dodecahedron. The tetragonal distortion was first systematically studied by Seijo and Barandiarán, who performed quantum chemistry calculations of Ce³⁺ doped garnets based on AIMPs (ab initio model potentials).¹¹⁸ They establish the 8-coordinated distorted cube by transforming a reference perfect cube through a series of SBT (stretching, bending and twisting) operations. They find that only the symmetric stretching (S1) and symmetric bending (S3) modes have important impacts on the lowest 4f–5d transition. In addition, the symmetric bending (S3) also leads to tetragonal distortion of a perfect cube, which induces elongation or compression of the cube, resulting in a cuboid. Early in 2007, Wu et al.¹¹² have found the dodecahedral edge-ratio as an indicator of tetragonal distortion. It also presents a monotonous relationship of excitation properties of Ce³⁺ in garnets. A simplified cuboid model is set up to provide a deeper understanding of the effect of tetragonal distortion on crystal-field splitting. The cuboid degrades the point symmetry from O_h of a perfect cube to D_4h, and the energy level splitting scheme also differs as depicted in Fig. 7. Numerical simulation of the crystal-field analysis shows that in certain geometrical configurations, the tetragonal distortion induces strengthening of the crystal-field splitting even with a larger cell size and a longer center–vertex distance.¹¹⁷

Fig. 7 Cuboid model with Cartesian coordinates of vertex and energy level splitting scheme. Reproduced with permission from ref. 117, copyright 2019 American Chemical Society.

Since tetragonal distortion has such a profound impact on the complex crystal-field splitting variation of Ce³⁺ in garnets, it is necessary to find their crystallographic origin. The above-mentioned numerical simulation shows that on isotropic expansion or shrinkage of the ligand polyhedron, the 5th-power rule will be obeyed. But the tetragonal distortion is anisotropic, caused by the neighboring polyhedron competition that has the effect of compromising the volumetric change of all the polyhedra. In the real garnet structure, polyhedra are linked by shared edges, and polyhedron deformations induced by compositional substitution behave in different manners with or without the neighboring polyhedra. The dodecahedral deformation is far beyond isotropic because of the competitions between dodecahedron and neighboring polyhedra, i.e., suppression by the octahedron/tetrahedron grid. Therefore, the complexity in the structure–property relationship is principally portrayed by the subtle local crystal structure changes.

Inductive effect from second nearest atoms

The structure and composition are decisive for the physical and chemical properties. Typically, 5d energy level of Eu²⁺/Ce³⁺, acting as a crucial factor for luminescence characteristics, is mainly determined by the structure and composition of the host, as illustrated in Fig. 8a. The effect of average compositions and structure on the energy of 5d level is represented as the nephelauxetic effect, which determines the centroid shift of the 5d level. Besides, the local structure also plays a critical role in 5d level determination. The effect of the nearest coordination anion on the energy of the 5d level is expressed by crystal field splitting, as a well-known parameter widely used for predicting and understanding the luminescence properties. Likewise, the second nearest cation also has a significantly influence on the energy of 5d level, which can be conveyed by the inductive effect.

Fig. 8 (a) Schematic diagram of structures on 5d levels. (b) Schematic diagram of the inductive effect of M. The thickness of arrow represents the degree of ability of donating electrons. (c) Centroid shift of the 5d levels of Ce³⁺versus the inductive factor in (qua)ternary nitridosilicates. Panels (b and c) reproduced with permission from ref. 119, copyright 2018 American Chemical Society.

Since the inductive effect was first proposed to explain the subtle changes of Si–O bond lengths in silicates by Noll in 1963,¹²⁰ it has been widely used in understanding the physical and chemical properties of materials including inorganic and organic compounds. As illustrated in Fig. 8b, for a ternary compounds T_xM_yX_z, where X is an anion and T and M are cations, if T is less electronegative than M, it will tend to act as an electron donor. Then the anion X prefers to get the electron density from T instead of M, causing the changes of the M–X bond. This is called the “inductive effect”.¹²¹ In 2017, it was introduced into luminescent materials and accounts for the tunable luminescence properties in K_2(0.995−x)-Na_2xAl₂B₂O₇:0.01Eu²⁺ (KAB:2xNa).¹²² The excitation spectra analysis shows that the centroid shift decreases with x increasing, and then the inductive effect is employed to understand this phenomenon. Given that Eu occupies the K site and its neighboring cation K is increasingly replaced by Na as x increases, the covalency of the Eu–O bonds would be reduced since Na is more electronegative than K, which leads to a smaller centroid shift of the 5d levels. After this, our group further provided insight into the influence of the inductive effect on structural chemistry and luminescence properties in nitridosilicates and oxysilicates.¹¹⁹ A new parameter, the inductive factor, was proposed to relate the difference of electronegativity between metal elements and silicon. Through collecting and analyzing the structural and luminescence data of more than one hundred compounds, a linear relationship between the average length of the Si–N/Si–O bonds and the inductive factor is revealed. And centroid shift ε_c shows obvious positive correlation with the inductive factor as shown in Fig. 8c. In addition, the inductive effect also accounts well for the charge transfer (CT) process. CT transition is another kind of electric dipole allowed transition of the electron from ligands to lanthanide ions, which may give out intense luminescence, such as the CT luminescence of Yb³⁺ in Yb³⁺-containing compounds, and the CT state sensitized luminescence in red phosphor Y₂O₃:Eu³⁺.¹²³ By introducing electronegativity factor and ionic radius factor, semiquantitative models are proposed that CT energy has a strong positive correlation with the electronegativity factor and a strong negative correlation with the ionic radius. These works highlight the key role of the neighboring cations in structural chemistry and luminescence properties, which provide us an interesting perspective for phosphor design.

Stokes shift. The Stokes shift ΔS is defined as the energy difference between the excitation and emission maxima of the same electronic transition. As the excitation properties can be predicted based on the crystal field splitting ε_cfs and the centroid shift ε_c, combining with the Stokes shift ΔS, the emission properties can be further forecasted based on the composition and structure of the compounds. It is fair to say that gaining the knowledge of the effect of the composition and structure on the Stokes shift ΔS is of guiding importance for the development of Ce³⁺/Eu²⁺-activated inorganic phosphors.

For Ce³⁺/Eu²⁺-activated inorganic compounds, the degenerate 5d-level splits into at most five 5d states due to the interaction of the 5d-electron with the neighboring anion ligands. Excitation occurs by the transitions from the 4f ground state to all the 5d states. At room temperature, the decay from the higher 5d states to the lowest 5d state is dominantly non-radiative and emission occurs by the transition from the lowest 5d state to the ground state, as shown in Fig. 9a. Therefore, more specifically, the Stokes shift is the energy difference between excitation and emission maxima of transition between the lowest 5d and the 4f ground states.^20,58,59 Based on the assumptions (1) Franck–Condon approximation, implicating that an optical transition on the dopant ion occurs so quickly that the lattice arrangement does not change during the transition, and (2) considering only one mode of lattice vibration, i.e., the breathing mode, which states that the surrounding lattice pulsates in and out about the dopant ion. The electronic and lattice states can be represented in a single energy level diagram called the configurational coordinate diagram.¹²⁵ In Fig. 9a, the processes of the excitation and emission transitions are shown in the configurational coordinate diagram represented by the ground and excited state potential energy curves with vibrational levels depicted as horizontal lines, and the Stokes shift also being marked in the picture. For simplicity, in the analysis, the harmonic approximation and the assumption that the vibrational frequencies are the same in both ground and excited states are generally adopted. The Stokes shift ΔS has the relation of¹²⁵

ΔS = (2S − 1)ħω, (14)

where ħ is the reduced Planck constant, ω is the angular frequency of vibration, and ħω is the energy difference between the consecutive vibrational states (phonon energy), as shown in Fig. 9a. Dimensionless Huang–Rhys factor S gives information about the coupling of electronic and vibrational states (i.e., electron–phonon coupling). The relation ΔS = (2S − 1) ħω has been used in many reports, while the alternative relation ΔS = 2Sħω has also been used in some studies.^126–132 Jong et al. resolved this ambiguity and pointed out that both the expressions are only approximately correct and the accurate Stokes shift is always between them.⁵⁸ Blasse classified the cases of weak (S < 1), intermediate (1 < S < 5) and strong (S > 5) electron–phonon energy.¹³³ Huang–Rhys factor S is given bywhere α = Mω/ħ and M is the reduced mass of the vibrating system,¹²⁵ and ΔQ_e is the equilibrium position offset, as shown in Fig. 9a. The existence of such a displacement corresponds to different equilibrium geometries in the electronic ground and excited states, which is caused by the change in the bonding between ions or atoms in the ground state and excited states as a result of the change in the electronic configuration.⁵⁸ It is affected by the cation substituted by lanthanide, coordination anions and local structure in the compounds. Blasse et al. reported that the larger size of the cation substituted by lanthanide ion leads to a larger Stokes shift in three orthoborates.¹³⁴ In addition, the phonon energy ħω is linearly related to the Stokes shift, as shown in eqn (14). It was reported that phonon energy decreases in going through the halogenide series from fluoride (∼350 cm⁻¹) via chloride (∼260 cm⁻¹) to bromide (∼200 cm⁻¹) compounds.¹³⁵ It also has been affirmed that oxides (∼600 cm⁻¹) have a larger phonon energy than sulfides (∼270 cm⁻¹), and the selenides (150 cm⁻¹) are the smallest among them.¹³⁶ This means that vibrational frequency ω decreases from oxides via sulfides to selenides, which originates from the increasing atomic weight, resulting in the increasing atomic radius, decreasing electronegativity and the weaker bond strength. And so does the situation in halogenides.

Fig. 9 (a) Energy level diagram showing the crystal field splitting ε_cfs, centroid shift ε_c and the Stokes shift ΔS. Some vibrational states and phonon energy are indicated in the configurational coordinate diagram. Reproduced with permission from ref. 124, copyright 2019 Elsevier B.V. (b) A summary scheme representing the correlations between the Stokes shift, phonon energy, Huang–Rhys factor, chemical composition and structure of the compounds.

Through statistical analysis of the structural and luminescence data of more than 60 Ce³⁺/Eu²⁺-activated inorganic phosphors, our group delved into the effect of the chemical composition and structure on the Stokes shift, phonon energy, and Huang–Rhys factor.¹²⁴ It has been demonstrated that both the cation substituted by Ce³⁺/Eu²⁺ and the coordination anion have crucial effects on the Stokes shift, phonon energy and Huang–Rhys factor. The Stokes shift ΔS has a positive correlation to the effective average coordination bond length R_av. It decreases in going through the halogenide series from fluorides to bromides, and the chalcogenides from oxides to selenides, respectively, while it has little difference in the compounds with the anions of the same period, which can be well digested using the effective charge concept. By replacing the formal charge Z_fc of the anion (1 for fluorides, 2 for oxides and 3 for nitrides) with an effective charge Z_ec = Q/N (Q is 3 for Ce³⁺ and 2 for Eu²⁺, and N represents the number of coordination anions), the electric field around the Ce³⁺/Eu²⁺ is approximately independent of the anion type in the same period, which results in no obvious influence of the ligand charge on the bonding strength. It also reveals that the Stokes shift ΔS has no obvious correlation with the coordination number of cation. Besides, the phonon energy is negatively correlated to the average bond length R, while the Huang–Rhys factor S is positively related to R. These composition–structure–Stokes shift correlations provide a powerful tool to predict the emission spectrum of Ce³⁺/Eu²⁺-activated inorganic phosphors.

Quantum efficiency and luminescence quenching. Luminescence of rare-earth elements originates from the electronic transition between separated energy levels. For a ground-state electron, it jumps to the high-energy level under light activation and stays at a meta-stable excitation state after the cross-relaxation process by emitting phonons. Next, those excited electrons will return to the low-energy ground state by emitting photons, which is called radiative transitions. However, this process occurs non-simultaneously for all the excited electrons. Those electrons could be thought as waiting in a sequence characterized by the decay lifetime. From the beginning to the lifetime, which has the magnitude of nanoseconds and microseconds for Ce³⁺ and Eu²⁺, respectively, about 63.2% of the excited electron ensemble will undergo the radiative transitions. Therefore, there is enough time for the excited electron to take part in other non-radiative transitions. When the energy-curve cross point between excited/ground states lies closer to the meta-stable excitation state, the excited electrons have the chance to return to the ground state with the help of thermal activation or by the tunneling process.¹³⁷ In this case, the energy is dissipated within the original activator by emitting phonons. In other cases, the excited energy would be lost to “killer” centers such as defects, vacancies, etc., during the energy transfer between activators or through the conduction band. In order to characterize the luminescence efficiency, quantum efficiency (QE) is usually measured for the luminescent material.

Apart from suitable excitation and emission spectra, high (QE) and good thermal luminescence stability also plays key roles in the practical applications of LED phosphors. QE of phosphors is defined as the ratio between the number of emitted and absorbed photons.¹³⁸ When an electron is excited from the ground state to the excited states, it will return to the ground state via two paths, either radiative transition or non-radiative transition. Radiative transition refers to the emission of lights, while in the non-radiative transition the energy is dissipated in some other way except for luminescence, which is generally called luminescence quenching.¹³⁹ Thus, the QE of phosphors fundamentally depends on the radiative recombination efficiency of the electrons, thereby it is of vital importance to have the knowledge of the nature of the non-radiative channels that compete with radiative emission.

Till now much endeavors have been devoted to explore the luminescence quenching process, and three quenching mechanisms are widely recognized including configurational coordinate model, ionization model, and impurities and lattice defects model.^140,141 For the configurational coordinate model, sufficient thermal energy could assist the excited electron to an upper vibrational level in the excited state. If the upper vibrational level is higher than the intersection of the excited and ground potential energy curves, the excited electron can return to the ground state by losing energy to the host lattice in the form of phonons (lattice vibration), which leads to luminescence quenching.¹⁴² The excited electron also has chance to return to the ground state by the tunneling process.¹³⁷ In this model, the luminescence efficiency is mainly determined by the phonon energy ħω and the Huang–Rhys factor S of the LED phosphors. In the ionization model, the 5d₁ excited electron can be photoionized or thermally ionized into CB. Subsequently, the ionized electron becomes delocalized and travels across the whole crystal, which increases the possibility to be trapped by a defect trap and lose excitation energy, resulting in luminescence quenching.^143,144 In this case, the luminescence efficiency depends on the ionization energy of the 5d excited electron, which is the energy difference between the CB bottom edge of the host and the 5d¹ excited state of the emission center. For the impurities and lattice defects model, energy can be transferred to defects or impurities, which have additional electronic levels to which excitation energy can be transferred. Such transitions are limited to the cases where the impurities or defects are sufficiently close to the activation ions, or the concentration of the activation ions is high enough so that the impurities or defects can be reached via energy migration over multiple activation ions, which is well known as concentration quenching. What is noteworthy is that these quenching processes related to impurities and defects are not intrinsic to the material and can be avoided by optimizing synthesis conditions and dopant concentrations. As for the first two quenching processes, while different, they are the intrinsic properties, related to the crossover point of the 4f and 5d potential wells and the relative energy position of the activation ions with respect to the host crystal bandgap, which can be adjusted by choosing a suitable host material or modifying the existed phosphors. Electronic structure schemes with the energy information on the activation ions and the host crystals are recently used as powerful tools for optimizing the luminescence efficiency and thermal luminescence stability as discussed above.

It is worth noting that luminescence quenching will be aggravated as the operating temperature increases since the quenching processes in configurational coordinate, ionization model and energy transfer to defects are thermally activated. While functioning, LED chips produce heat and transfer it to the coating phosphors. This heat activates more phonon modes of the phosphor, thereby decreasing the efficiency and luminous intensity.^5,145 This phenomenon is typically onset between 100 and 200 °C, and is described as thermal luminescence quenching, which not only lowers the overall efficiency but also causes the undesirable change of the color rendering of the device. The preservation of phosphor emission intensity and quality with an increase in temperature is referred as the thermal luminescence stability. Therefore, taking thermal and concentration quenching into account, luminescent materials with optimized dopant concentration, extraordinary thermal stability, and fine structural rigidity are necessary to fulfil the requirements for LED applications.

Persistent luminescence mechanism. Persistent phosphors generally consist of the host lattice and two kinds of activation centers including emission centers and trap centers. Emission centers are mostly Eu²⁺ or Ce³⁺ that give out luminescence. Trap centers are typically lattice/intrinsic defects (such as oxygen vacancies), or intentionally introduced aliovalent or isovalent co-dopants (such as Dy³⁺ or Cr³⁺), which can capture the excitation electron and then release it by thermal assistance, causing the luminescence delay or continuous luminescence as in PersL.²⁴ Since the green persistent phosphor SrAl₂O₄:Eu²⁺,Dy³⁺ with excellent properties was discovered,³⁰ tremendous research studies have been devoted to explore its underlying mechanisms. Although several models were proposed in order to give convincible explanations for this attractive phenomenon,¹⁴⁶ we will limit ourselves to the most widely used electron trapping–detrapping mechanism for Ce³⁺/Eu²⁺ doped persistent phosphors.²³ As illustrated in Fig. 10, the PersL process can be typically divided into four phases: (a) delocalization of electron. Upon effective optical charging by ultraviolent (UV) or visible (VIS) light irradiation, there are two pathways to get delocalized electrons: (1) the 4f electrons are excited to 5d excited states and thus thermally ionized to the conduction band (CB); and (2) 4f electrons are excited directly to CB. (b) Storage of electrons (charging): the delocalized electrons are captured by the traps. (c) Release of captured electrons: after ceasing the excitation, the captured electrons are thermally released from the traps via the CB to 5d excited states of Ce³⁺/Eu²⁺. (d) Recombination process: the released 5d excited electrons move back to the 4f ground state yielding delayed luminescence, which is called PersL. Consequently, the emitting properties of PersL including the emission wavelength and the bandwidth are determined by the emission center, while persistent properties including intensity and duration are determined by the characteristic of the trap center, such as trap depth and density. Trap depth is the energy gap (ΔE) between the bottom of CB and the trap level. Generally, electrons in shallow traps with a small ΔE will be released fast at the operating temperature, which contributes to a strong initial intensity but a short PersL duration. In contrast, electrons in deep traps with a large ΔE will be gradually or even hardly released at the operating temperature, which results in a long PersL duration but weak initial intensity. Thus, an ideal ΔE needs to match the operating temperature so that the captured electrons can be slowly released by the relevant assisted energy, giving rise to optimized PersL properties. In addition, an appropriate trap density generally contributes to a strong PersL intensity, which can be achieved by tuning the trap center concentration or introducing a co-dopant. What is more, emission center is also critical, since it not only determines the luminescence properties but also plays a vital role in achieving delocalized electrons (thermal ionization of 5d excited electrons). Therefore, two essentials for promising persistent phosphors can be summarized: (1) an appropriate host crystal with an emission center; and (2) trap center with a suitable trap density and depth for specific applications. Accordingly, the PersL performance can be optimized following two strategies: (1) manipulating the bandgap structure of the host especially the CB minimum energy by structure modulation; and (2) introducing rare earth or transition metal ions or deficiencies as trap centers. Both strategies can be approached by virtue of electronic structure schemes described above.

Fig. 10 Illustration for the PersL mechanism.

Well known Ce³⁺/Eu²⁺-activated LED and persistent phosphors: focusing on the local structure and the electronic structure

The past several decades have witnessed the rapid development of LED phosphors and the increasing interest in strategies to improve their properties toward specific applications. In this section, we present several widely used promising phosphors, focusing on their crystal structure, local structure, electronic structure, and luminescence properties. Meanwhile, with the aim at illustrating the strategies to screen phosphors with potential PersL properties, the LED phosphors, in which decent PersL has been achieved successfully, get the priority here. With the emphasis on structure design principles and routines, here we discuss the composition–structure–property correlations and how the local structure and the electronic structure can alter the luminescence properties including emission characteristics, luminescence efficiency, thermal luminescence stability, and PersL properties. The electronic structure, luminescence and PersL properties of these materials are summarized in Tables 1–4 with the prospect of motivating inspirations of interested readers for further development of luminescent materials and other functional materials.

Table 1 Luminescence properties and band gap of garnet phosphors. λ_ex,max and λ_em,max are the wavelengths corresponding to the peak of excitation and emission bands, respectively; QY is the quantum yield; I_150 °C/I_RT represents the relative emission intensity at 150 °C to that at room temperature

Chemical composition	λ ex,max (nm)	λ em,max (nm)	QY (%)	Band gap (eV)	I 150 °C/IRT (%)	Ref.
Y3Al5O12:7%Ce^3+	460	555	80	6.4	93	144, 145 and 154
(Y2.2Gd0.8)Al5O12:6%Ce^3+	460	568	—	—	—	153 and 162
Gd3Al5O12:1%Ce^3+	470	600	—	6.3	32	156 and 163
Gd3GaAl4O12:1%Ce^3+	458	587	—	6.3	37	163
Gd3Ga2Al3O12:1%Ce^3+	448	574	—	6.3	39	163
Gd3Ga3Al2O12:1%Ce^3+	442	566	58–82	6.1	11	163 and 164
Gd3Ga4AlO12:1%Ce^3+	437	561	—	5.9	—	163
Gd3Ga5O12:1%Ce^3+	428	490	—	5.4	—	163
Lu3Al5O12:7%Ce^3+	450	540	84	7	96	154
Ca3Sc2Si3O12:Ce^3+	455	505	—	7.35	90	161 and 165
Tb3Al5O12:1%Ce^3+	470	552	—	—	56	156 and 166–168
Y3Sb2Al3O12:4%Ce^3+	465	528	—	—	—	169
Y3Al4.9Si0.1O11.9N0.1:3%Ce^3+	470	595	68	—	53	170 and 171
Y3Mg2AlSi2O12:1.5%Ce^3+	475	600	45	—	44	172 and 173
CaLu2Mg2Si3O12:Ce^3+	470	605	80	—	72	174
CaLu2Al4SiO12:1%Ce^3+	450	520	77	5.77	85	175
Mg3Y2Ge3O12:3%Ce^3+	460	535, 573	4.5	5.47	—	176
MgY2Al4SiO12:6%Ce^3+	452	568	—	—	—	177
Ca2GdZr2(AlO4)3:2%Ce^3+	417	500	40	—	33	178
Ca2LuZr2(AlO4)3:2%Ce^3+	411	480	50	—	60	179
Ca2YZr2(AlO4)3:2%Ce^3+	412	481	43	—	51	178
CaGd2ZrSc(AlO4)3:2%Ce^3+	448	545	13	—	18	178
Ca2LaZr2Ga3O12:6%Ce^3+	430	515	35	4.54	32	159
Ca3Hf2SiAl2O12:0.5%Ce^3+	400	457, 482	75	—	73	180
Tb2.2Lu0.8Al5O12:0.5%Ce^3+	430	550	—	—	—	181
Mg3Y2(Ge0.3Si0.7)3O12:Ce^3+	460	604	54.3	5.82	—	176
Y3AlGa4O12:Ce^3+	422	504	17	—	13	160 and 182
Y3Al2Ga3O12:Ce^3+	430	512	47	5.9	21	144 and 182
Y3Al3Ga2O12:Ce^3+	440	523	83	—	85	182 and 183
Y3Al4GaO12:Ce^3+	447	537	86	—	94	182
Mg3YGdGe3O12:Ce^3+	467	575	—	—	—	160
Mg3Y0.25Gd0.75Ge3O12:Ce^3+	465	590	—	—	—	160
Mg3Gd2Ge2SiO12:Ce^3+	467	606	—	—	—	160
Mg3Gd2Ge3O12:2%Ce^3+	464	600	—	—	—	160
Lu3Al4.8Si0.2O11.8N0.2:3%Ce^3+	470	587	—	—	71	170
Y3Al3MgSiO12:0.25%Ce^3+	450	577	68	—	56	184
Y2LuAl3MgSiO12:0.25%Ce^3+	450	570	71	—	62	185
YLu2Al3MgSiO12:0.25%Ce^3+	450	564	78	—	70	185
Lu3Al3MgSiO12:0.25%Ce^3+	450	552	86	—	82	185
Lu3(Al1.5Mg0.5)(Al2.5Si0.5)O12:3%Ce^3+	447	551	85	—	83	186
Lu3(AlMg)(Al2Si)O12:3%Ce^3+	448	562	81	—	—	186
Lu3(Al0.5Mg1.5)(Al1.5Si1.5)O12:3%Ce^3+	455	571	71	—	—	186
Lu3Mg2Si2O12:3%Ce^3+	460	571	66	—	—	186
Y3Sc2Al3O12:5%Ce^3+	460	520	74	—	—	141
Y3Sc2Al2GaO12:5%Ce^3+	445	510	65	—	—	141
Y3Sc2AlGa2O12:5%Ce^3+	430	500	63	—	—	141
Y3Sc2Ga3O12:5%Ce^3+	420	496	37	—	—	141
(Lu2Mg)(Al4Si)O12:4%Ce^3+	434	533	47	—	76	187
(Lu2Ca)(Al4Si)O12:4%Ce^3+	438	527	68	—	85	187
(Lu2Sr)(Al4Si)O12:4%Ce^3+	445	511	78	—	92	187 and 188
(Lu2Ba)(Al4Si)O12:4%Ce^3+	445	511	81	—	93	187
(Y0.95La0.05)3Al4.2(Mg,Si)0.4O12:Ce^3+	456	555	—	—	65	162
Sr3Y2Ge3O12:Ce^3+	427	513	—	6.37	0	165
Y2LuCaGaAl2SiO12:6%Ce^3+	435	505	—	—	—	189
Ca2YHf2Ga3O12:4%Ce^3+	398	479	26	—	—	190
Ca3Sc2Ge3O12:8%Ce^3+	450	530	13	5.45	90	191
Ca2LaHf2Al3O12:5%Ce^3+	408	515	47	—	—	192
Lu2Mg2Al2Si2O12:8%Ce^3+	436	575	73	4.83	75	193
Y2Mg2Al2Si2O12:6%Ce^3+	460	599	57	—	∼85	194

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Table 2 The electronic structure and PersL properties of garnet persistent phosphors. λ_ex is the excitation wavelength. λ_em,max is the wavelength corresponding to the peak of emission band

Chemical composition	λ ex (nm)	λ em,max (nm)	PersL duration	Trap depth (eV)	Ref.
Y3Al2Ga3O12:1%Ce^3+	430	512	∼20 min	0.73, 0.86	183
Y3Al2Ga3O12:1%Ce^3+,1%Pr^3+	430	512	∼1 h	0.69, 0.71	183
Y3Al2Ga3O12:0.5%Ce^3+,0.5%Cr^3+	460	505, 700	265 min	0.81	182, 196 and 204
Y3Al2Ga3O12:0.2%Ce^3+,0.1%Yb^3+	455	520	138.8 h	1.01	31 and 197
Lu2CaMg2SiGe2O12:1%Ce^3+	460	568	∼1 s	0	202
Gd3Al4GaO12:0.2%Ce^3+,0.1%Cr^3+	450	537	22 min	0.56	205
Gd3Al3Ga2O12:0.2%Ce^3+,0.1%Cr^3+	450	528	126 min	0.47	205
Gd3Al2.5Ga2.5O12:0.2%Ce^3+,0.1%Cr^3+	450	515	405 min	0.45	205
Gd3Al2Ga3O12:0.2%Ce^3+,0.1%Cr^3+	450	510	401 min	0.37	205 and 206
Gd3AlGa4O12:0.2%Ce^3+,0.1%Cr^3+	450	511	8 min	0.29	205
Gd3Sc2Ga3O12:Cr^3+,Eu^3+/Yb^3+	254	784	—	—	198
Gd3Sc2Ga3O12:Cr^3+,Yb^3+	254	784	—	—	198
Y3Sc2Ga3O12:0.5%Ce^3+	440	500	1 h	—	203
Y3Sc2Ga2.5Al0.5O12:4%Ce^3+	365	503	∼500 s	0.62, 0.79	201

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Table 3 Luminescence properties and band gap of commercial LED phosphors reviewed in this work. λ_ex,max and λ_em,max are the wavelength corresponding to the peak of excitation and emission band, respectively; QY is the quantum yield; I_150 °C/I_RT represents the relative emission intensity at 150 °C to that at room temperature

Chemical composition	λ ex,max (nm)	λ em,max (nm)	QY (%)	Bandgap (eV)	I 150 °C/IRT (%)	Ref.
CaAlSiN3:Eu^2+	455	650	87	5.5	90	18, 79, 225 and 228–230
Sr0.8Ca0.2AlSiN3:0.8%Eu^2+	455	628	83	—	79	209 and 231
SrAlSiN3:0.8%Eu^2+	455	610	—	—	81	209
CaAlSi1+2xN3+2xOx:2%Eu^2+	460	605	—	—	80	223
Ca2Si5N8:Eu^2+	395	605	∼35	5	40	232 and 233
Sr2Si5N8:4%Eu^2+	450	626	79	4.9	86	232–234
Ba2Si5N8:Eu^2+	395	574	∼75	4.9	83	232 and 233
CaSrSi5N8:2%Eu^2+	455	632	—	4.41	82	235 and 236
Sr0.5(Ca0.55Ba0.45)1.5Si5N8:2%Eu^2+	460	654	—	—	96	237
La3Si6N11:6%Ce^3+	460	535	100	4.0	∼100	238 and 239
La2.9Al0.3Si5.7N10.9:1%Ce^3+	530	535, 665	13		55	239
CaSi2O2N2:2%Eu^2+	355	560	76	6.6c	83	240
SrSi2O2N2:2%Eu^2+	360	538	91	6.67	99	240 and 241
BaSi2O2N2:2%Eu^2+	380	494	71	6.6c	83	240
(Sr0.5Ba0.5Eu0.02)Si2O2N2:2%Eu^2+	450	548	92	—	94	240
β-SiAlON (z = 0.18):0.02%Eu^2+	325	540	82	7.2	97	242 and 243
Ca2SiO4:Eu^2+	365	502	—	5.2	60	244–248
Sr2SiO4:Eu^2+	350	574	—	5.29	∼80	142 and 249–251
Ba2SiO4:Eu^2+	370	505	—	5.5	57	142, 252 and 253
Ca1.2Eu0.8SiO4	450	653	50	—	43	14
Sr0.9Ba1.1SiO4:1%Eu^2+	394	530	90	—	∼75	51
(Ba0.68Ca0.32)2SiO4:Eu^2+	340	484	—	—	79	244
CaSrSiO4:0.25%Eu^2+	450	620	68	—	—	254
Sr3SiO5:5%Eu^2+	460	581	68	6.05	85	255–257
Ba3SiO5:1%Eu^2+	350	504, 566	43	5.42	77	257 and 258
Sr2.6Ba0.4SiO5:0.02Eu^2+	417	598	76	—	—	259
Sr3SiO5: Ce^3+,Li^+	415	540	79	6.05	∼50	5 and 259–262

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Table 4 The electronic structure and PersL properties of persistent phosphors reviewed in this work. λ_ex is the excitation wavelength. λ_em,max is the wavelength corresponding to the peak of emission band

Chemical composition	λ ex (nm)	λ em,max (nm)	PersL duration	Trap depth (eV)	Ref.
Sr0.8Ca0.2AlSiN3:0.15%Eu^2+	254	628	9600 s	0.86	94
Ca2Si5N8:1%Eu^2+	365	610	150 s	0.87	263
Sr2Si5N8:1%Eu^2+	365	620	80 s	0.66	263 and 264
Ba2Si5N8:1%Eu^2+	365	580	400 s	0.71	263
Ca2Si5N8:1%Eu^2+,Tm^3+	365	610	2500 s	0.91	263
Sr2Si5N8:Eu^2+,Tm^3+	254	610	600 s	0.65	264 and 265
Sr2Si5N8:0.5%Eu^2+,1%H3BO3	254	610	—	0.69	264
SrCaSi5N8:2%Eu^2+,2%Tm^3+	365	647	∼30 min	—	266
SrSi2O2N2:Eu^2+,Ho^3+	254	540	360 min	1.05	241
SrSi2O2N2:Eu^2+,Dy^3+	254	540	50 min	1.18	241
β-Sialon (z = 0.5):1.5%Eu2+	254	537	400 s	0.27–0.59	243
Sr2SiO4:Eu^2+,Dy^3+	313	550	∼5 min	0.5–0.66	267 and 268
Sr2SiO4:1%Eu^2+,0.5%Tm^3+	365	540	—	1.35	269
Ba2SiO4:0.5%Eu^2+,1%Ho^3+	254	504	24 h	0.71–0.86	253
Ba2SiO4:0.2%Eu^2+,2%Er^3+	360	510	402 min	0.59, 0.78	270
Sr3SiO5: Eu^2+,Dy^3+	370	570	6 h	0.66, 1.05, 1.21	271
Sr3SiO5:0.1%Eu^2+,0.5%Nd^3+	254	580	14 h	0.79–1.05	272
Sr3SiO5:0.01%Eu^2+,0.5%Ge^3+	365	580	7000 s	0.73, 0.97	273
(Sr0.88Ba0.12)3SiO5:Eu^2+,Nb^5+	254	590	17 h	0.48–0.71	274

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Ce³⁺-activated garnets

The Ce³⁺-activated garnets, as a class of highly efficient luminescent materials, have been widely used in various applications ranging from lighting and displays to persistent illumination fields.¹⁴⁰ The garnet structures crystallize in the cubic space group Ia3̄d, and the crystal chemical formula is (A)₃{B}₂[C]₃O₁₂ where (), {}, and [] represent the dodecahedral, octahedral and tetrahedral sites, respectively.¹⁴⁷ As shown in Fig. 11a, tetrahedral C and octahedral B form a corner-connected 3D network, with the 8-coordinating A atoms filling the remaining space. Owing to the unique crystal structure, garnet based compounds have a large variety of compositional derivatives. Recently, Song et al. borrowed theory from the perovskite structure, introducing the tolerance factor into the garnet structure for interpretation and prediction of the phase stability with the garnet structures.¹⁴⁸ A survey of the tolerance factor over 130 garnet compounds reveals the correlation between the tolerance factor and the garnet-phase stability, demonstrating the critical role of the tolerance factor not only in understanding the crystal chemistry of garnet compounds but also in predicting the stability of the garnet phase. The various cation sites and the rich structural diversity allow Ce³⁺-activated garnets to be remarkably flexible in tuning and optimizing luminescence properties for specific applications. The Ce³⁺-activated yttrium aluminum garnet (YAG:Ce) with its blue to yellow down-conversion luminescence was first reported by Blasse & Bril in 1967 as a new phosphor for flying-spot cathode-ray tubes for color television.¹⁴⁹ After the bright blue InGaN LED chip became available, YAG:Ce was widely used as the down-converting phosphors fabricated with the blue chips to produce white light.⁹ Taking the advantage of the parity-allowed electric dipole f–d transition of Ce³⁺, the emission of YAG:Ce features with a broad emission band ranging from 500–650 nm, as shown in Fig. 11b.^11,150 Owing to the strong crystal field splitting of the 5d energy level of Ce³⁺, the excitation from the ²F_5/2 term to the lowest-energy 5d orbital of Ce³⁺ occurs in a broad range of 440–480 nm, which matches well with the commercial blue LED chip, leading to its successful applications in general lighting and backlights. The improved thermal stability was successfully realized by coating an ultrathin layer of Al₂O₃.¹⁵¹

Fig. 11 (a) The crystal structure of garnet phosphor Y₃Al₅O₁₂:Ce³⁺ (YAG:Ce). Reproduced with permission from ref. 152, copyright 2013 American Chemical Society. (b) Typical excitation and emission spectra of YAG:2%Ce. Reproduced with permission from ref. 11, copyright 2020 Elsevier Ltd. (c) Emission spectra (λ_ex = 460 nm) of (Y_2.94−xCe_0.06Gd_x)Al₅O₁₂ phosphors for different Gd contents. Reproduced with permission from ref. 153, copyright 2011 Elsevier GmbH. (d) Excitation and emission spectra of Y_2.93−xLu_xAl₅O₁₂:Ce_0.07 phosphors. Reproduced with permission from ref. 154, copyright 2011 Elsevier B.V. (e) Schematic energy diagram of Ce³⁺ in YAG, TAG, and GAG and the emission process of Ce³⁺. Note that in this schematic picture the variation of the bandgap with composition is not taken into account. Reproduced with permission from ref. 155, copyright 2007 The Electrochemical Society. (f) Excitation and emission spectra of Y₃Al_5−xGa_xO₁₂:0.5%Ce³⁺. Reproduced with permission from ref. 109, copyright 2017 Elsevier B.V.

However, pc-LEDs fabricated with sole YAG:Ce phosphor have drawbacks. One deficiency is that they are limited to cold white light due to the lack of a red spectral component. Many endeavors have been devoted to improving the emitting color of garnet phosphors by cation or anion substitution. Through replacing the dodecahedral Y³⁺ ions with larger ions such as Gd³⁺ and Tb³⁺,^153,156 the red shifted emission can be achieved, while replacing Y³⁺ with smaller Lu³⁺ ions leads to blue shifted emission with respect to the emission of YAG:Ce.¹⁵⁴ The emission spectra for a series of (Y_2.94−xCe_0.06Gd_x)Al₅O₁₂ are presented in Fig. 11c as an example. As the Gd³⁺ concentration increases, the emission peak wavelength shifts from 532 to 568 nm, which enables (Y_3−xGd_x)Al₅O₁₂:Ce³⁺ to be applied for w-LEDs with a warmer white light. Of the various garnets, LuAG has a more condensed structure and a high effective atomic number, and thereby a high stopping power for ionizing radiation, which allows the successful application of LuAG:Ce as scintillator materials in medical imaging and security.^157,158 As shown in Fig. 11d, the photoluminescence excitation (PLE) spectra of Y_2.93−xLu_xAl₅O₁₂:Ce³⁺ comprise two excitation bands including a low-energy band peaking at ∼455 nm and high-energy band peaking at ∼345 nm, which are assigned to the transitions from the 4f ground state to the 5d₁ and 5d₂ excited states, respectively. With increasing Lu content, the peak of the low-energy band shifts toward a shorter wavelength while that of a high-energy band shifts toward a longer wavelength, indicating that the 5d₁–5d₂ crystal field splitting decreases upon substitution of Y³⁺ by Lu³⁺. A schematic energy diagram for Ce³⁺ in LnAG (Ln = Y, Lu, Gd) is summarized in Fig. 11e. It is intriguing to find that the crystal field splitting increases as the size of the cations increases, which contradicts with the crystal field theory revealing that a larger polyhedron leads to a smaller crystal-field splitting. Extensive research has been devoted to gain insight into this so-called reverse garnet effect phenomenon, and it has been demonstrated that, besides the shape and size of the polyhedron, polyhedron distortion also has an effect on crystal field splitting and in this case it is dominant.^108,116,117 In addition, the dodecahedral A site can also accommodate Ca, La, Mg, Sr and Ba atoms, giving rise to the abundant structural diversity and tunable luminescence properties of garnet phosphor.^159–161 The electronic structure and luminescence properties are summarized in Table 1.

Cation substitution at the tetrahedral and octahedral Al sites in Ln₃Al₅O₁₂:Ce³⁺ also helps in tuning their luminescence properties. Since Al³⁺ ions in garnets occupy two Wyckoff sites, the garnet structure has different cationic substitution modes. Previous research studies have demonstrated that the octahedral Al site can be replaced by Ga³⁺, Sc³⁺, Mg²⁺, Sb³⁺, Zr³⁺, and Hf³⁺ and the tetrahedral Al site can be replaced by Ga³⁺, Si⁴⁺, and Ge⁴⁺ while maintaining the garnet-type structure.¹⁴⁰ For example, Ga can occupy both tetrahedral and octahedral sites. Replacing partial Al with Ga leads to a blue shift in both the lowest energy 4f–5d excitation and 5d–4f emission of Ce³⁺ (Fig. 11f), which can be understood from the view of the inductive effect.¹⁹⁵ The emission of Y₃Al_5−xGa_xO₁₂:Ce³⁺ can be tuned gradually from yellow to green, which enables it to be fabricated together with narrow-band red-emitting phosphors to realize a wide gamut of liquid crystal display (LCD). It is noteworthy that although these substitutions improve the red component of the emission, they reduce the thermal luminescence stability (Table 1) since the structure rigidity and the bandgap structure can be changed by the cation or anion substitution. By means of temperature dependence of photoconductivity measurement, it is verified that the thermal ionization process takes the responsibility for luminescence quenching in Y₃Al_5−xGa_xO₁₂:Ce³⁺.¹⁴⁴ Inspired by the empirical theory that the lower electronegativity of N compared with O generally leads to a larger 5d centroid shift and thereby lower energy emission, Setlur et al. successfully prepared garnet phosphors with a low energy emission by incorporating Si⁴⁺-N³⁻ into Ce³⁺ doped RE₃Al₅O₁₂:Ce³⁺ (RE = Lu, Y, and Tb).¹⁷⁰

What is surprising is that garnet compounds also serve as superior PersL materials. Originally, Holloway and Kestigian first observed several seconds of PersL in Y₃Al₂Ga₃O₁₂:Ce³⁺ and Y₃Al_1.5Ga_3.5O₁₂:Ce³⁺ after removal of UV excitation.¹⁹⁹ About 40 years later, Kanai et al. reported another garnet phosphor Gd_3+d(Al, Ga)_5−dO₁₂:Ce³⁺ which exhibited PersL for hundreds of milliseconds using X-ray excitation.²⁰⁰ Since then, more and more Ce³⁺ singly doped garnets with PersL performance were reported successively such as Y₃Al_5−xGa_xO₁₂:Ce³⁺, Y₃Sc₂Al_3−xGa_xO₁₂:Ce³⁺ and Lu₂CaMg₂(Si_1−xGe_x)₃O₁₂:Ce³⁺, but the brightness and duration were still rather limited.^{183,201–203} A breakthrough was made by Ueda et al. in 2014.¹⁹⁶ By introducing Cr³⁺ into Y₃Al₂Ga₃O₁₂:Ce³⁺ as a trap center, the green PersL of Ce³⁺ in Y₃Al_5−xGa_xO₁₂ was dramatically (∼3900 times) improved, and the duration can be reached over 5 hours after blue light excitation. They also took a quick look into the PersL of Lu₃Al₂Ga₃O₁₂:Ce³⁺,Cr³⁺ and Gd₃Al₂Ga₃O₁₂:Ce³⁺,Cr³⁺, which gave out PersL with tunable emission colors from greenish-blue to yellow, as shown in Fig. 12a. The electronic structure and PersL properties of garnet persistent phosphors are summarized in Table 2.

Fig. 12 (a) Image of Ln₃Al₂Ga₃O₁₂:Ce³⁺ (Ln = Y, Lu, Gd) and SrAl₂O₄:Eu²⁺, Dy³⁺ at 5 min after ceasing the excitation. Reproduced with permission from ref. 196, copyright 2014 American Institute of Physics. (b) VRBE scheme of Y₃Al_5−xGa_xO₁₂:Ce³⁺, Cr³⁺. Reproduced with permission from ref. 182, copyright 2015 Royal Society of Chemistry. (c) PersL decay curves of Y₃Al₂Ga₃O₁₂:0.2%Ce³⁺, 0.1%Yb³⁺ transparent ceramics, Y₃Al₂Ga₃O₁₂:0.2%Ce³⁺, 0.1%Yb³⁺ normal ceramics, and SrAl₂O₄:Eu²⁺, Dy³⁺ powder after blue light charging. Reproduced with permission from ref. 31, copyright 2018 American Chemical Society. (d) VRBE schemes revealing the PersL mechanism for Y₃Al₂Ga₃O₁₂:Ce³⁺, Yb³⁺, B³⁺ phosphors with and without Si⁴⁺. Reproduced with permission from ref. 197, copyright 2020 Elsevier B.V. (e) HRBE scheme of Gd₃Sc₂Ga₃O₁₂. Reproduced with permission from ref. 198, copyright 2017 Wiley. (f) Correlation curves between thermal quenching performance and the longest tested duration of (Y_1−xCe_x)₃(Al_1−xGa_x)₅O₁₂:Ce³⁺ samples. The time for samples with a 40% Ga content is set to zero. Reproduced with permission from ref. 195, copyright 2017 Elsevier B.V.

As the concept of electronic structure schemes matures, the HRBE and VRBE schemes of garnets are applied for interpreting and predicting PersL properties, and designing new potential persistent phosphors.⁹² Since PersL behavior is realized by the electron transfer between luminescent centers and trap centers via CB, the control of bandgap structure and trap depth are two effective strategies towards promising persistent phosphors, which are named bandgap engineering and trap depth engineering, respectively. By means of bandgap engineering, the bandgap and thereby trap depth have been systematically tuned by varying the Ga content in Y₃Al_5−xGa_xO₁₂:Ce³⁺,Cr³⁺.¹⁸² By constructing the VRBE scheme with energy level of the 3d electron of Cr²⁺ (Fig. 12b), it is revealed that, with increasing Ga content, the energy level of the CB bottom decreases while the energy level of the 3d electron of Cr²⁺ remains almost the same, giving rise to a smaller trap depth and thereby different PersL properties for various applications, such as phosphors with deeper traps Y₃Al_5−xGa_xO₁₂:Ce³⁺,Cr³⁺, x = 0–1 for energy storage and those with shallower traps Y₃Al_5−xGa_xO₁₂:Ce³⁺,Cr³⁺, x = 2.5–3.5 for emergency signage at room temperature. Since the ionization energy for luminescence quenching originating from thermal ionization directly affects the charging efficiency for PersL at room temperature, by comprehensively analyzing luminescence quenching and PersL properties, the optimum energy gap between the 5d₁ level and the CB bottom for efficient charging is estimated to be around 0.6 eV, which can improve the PersL properties at room temperature without compromising the blue light charging efficiency and quantum efficiency. In order to extend the PersL duration of Y₃Al₂Ga₃O₁₂:Ce³⁺,Cr³⁺, the trap center with a deeper depth relative to Cr³⁺ is required. Employing trap depth engineering, Yb³⁺ with a lower energy level (−3.52 eV) with respect to Cr²⁺ (−1.99 eV) in VRBE schemes was introduced into Y₃Al₂Ga₃O₁₂:Ce³⁺ as a trap center.³¹ As expected the Y₃Al₂Ga₃O₁₂:Ce³⁺,Yb³⁺ as-made transparent ceramic phosphor shows super-long persistent luminescence for over 138.8 h after blue light charging, but its initial intensity is relatively lower compared to that of Y₃Al₂Ga₃O₁₂:Ce³⁺,Cr³⁺ (Fig. 12c), which is common in the case of deep traps. Recently, Zhou et al. proposed a multi-doping strategy towards outstanding PersL performance combining strong intensity and long duration through introducing wide trap distribution ranging from shallow traps for strong initial intensity to deep traps for long duration.^197,207 By multi-doping B³⁺ and Si⁴⁺ into Y₃Al₂Ga₃O₁₂:Ce³⁺,Yb³⁺, the wide trap distribution from 0.47 to 1.11 eV is achieved (Fig. 12d), and the PersL performance including intensity and duration is improved substantially. As another example for trap depth engineering, Xu et al. prepared Gd₃Sc₂Ga₃O₁₂:Cr³⁺,Ln³⁺ garnet persistent phosphors with Cr³⁺ as the emission center and Ln³⁺ as the trap center. The lanthanide ions were selected to be Sm³⁺, Eu³⁺, Tm³⁺, and Yb³⁺ with 4f energy level of the relevant divalent ions located below the CB bottom according to the HRBE scheme as shown in Fig. 12e.^198,208 Through analysis of the thermoluminescence glow curves, it was confirmed that Eu³⁺ and Yb³⁺ can introduce trap centers with different depths while Sm³⁺ and Tm³⁺ cannot, which suggests that a relatively larger energy gap between the 4f energy level of divalent lanthanide ions and the CB bottom is required for efficient electron storage. Considering that both luminescence quenching and PersL properties are related to the energy gap between the 5d₁ level and the CB bottom, Song et al. investigated their correlations in the (Y_1−xCe_x)₃(Al_1−xGa_x)₅O₁₂ system.¹⁹⁵ As expected, a negative correlation between luminescence quenching and PersL properties is revealed (Fig. 12f), which provides us a clue that the luminescence quenching properties of LED phosphor can be regarded as a criterion to evaluate the possibility of LED phosphors possessing potential PersL performance. The development of Ce³⁺-activated garnets with promising properties has been established as an inspiring template for further design and modulation of novel outstanding LED and persistent phosphors.

Eu²⁺-activated MAlSiN₃ (M = Ca, Sr)

As discussed above, conventional w-LEDs composed of blue InGaN LED chip and yellow-emitting YAG:Ce phosphor have a low color rendering index because they lack a red component. Therefore, stable red-emitting phosphors, which can be efficiently excited by blue light, are essential to compensate the red color deficiency and thereby obtain w-LEDs with warm white light. For developing red-emitting phosphors, three approaches have been recognized as promising: (1) Eu²⁺-activated nitrides and oxynitrides;²⁰⁹ (2) Eu³⁺-activated phosphors;²¹⁰ and (3) Mn⁴⁺-activated phosphors.²¹¹ Since the line-shaped excitation band corresponding to the forbidden f–f transition of Eu³⁺ is not suitable for solid-state lighting for which a broad excitation band is generally required,²¹² and the valence state of Mn in host materials is hard to control,²¹³ Eu²⁺-activated nitrides and oxynitrides are regarded as the most potential materials to be efficient red emitters.

CaAlSiN₃:Eu²⁺ is one of the widely used red-emitting phosphors in w-LEDs due to its high quantum efficiency and outstanding thermal luminescence stability. The crystal structure of CaAlSiN₃:Eu²⁺ was first reported by Ottinger in 2004.²¹⁴ It has an orthorhombic structure with the space group of Cmc2₁. As shown in Fig. 13a, alternating [SiN₄] and [AlN₄] tetrahedra form conner-connected 3D network, with the 5-coordinated Ca²⁺ ions in the cavities. It has a rigid crystal structure consisting of a highly dense and corner-sharing MN₄ (M = Al³⁺, Si⁴⁺) tetrahedral network. There are two sheets of six-membered MN₄ tetrahedra rings as shown in Fig. 13b. The two sheets are 180° rotated and overlaid to form a rigid three-dimensional framework. The crystal chemical formula of CaAlSiN₃ could be written as Ca(Al/Si)₂N^I₂N^II,²¹⁵ in which Ca atoms occupy the 4a site, Al/Si atoms take up 8b site randomly, and N^I, N^II are accommodated in two different crystallographic sites of 8b and 4a, respectively. Many nitrides are crystalized in the CaAlSiN₃-type structure with the general formula of AB₂X₂X′, such as LiSi₂N₃ and NaSi₂N₃ (with A = Li or Na, B = Si and X = X′ =N), Si₂N₂O (A = vacancy, B = Si, X = N, X′ = O), and SrAlSiN₃ and MgAlSiN₃ (A = Sr/Mg, B = Al/Si). Generally, A site could accommodate larger cations such as Li, Na, Mg, Ca and Sr ions, or vacancies. B site has various Al/Si ratios. In X′-sites O could substitute for N.

Fig. 13 (a) Crystal structure of CaAlSiN₃:Eu²⁺ with the AlN₄ and SiN₄ tetrahedra. Reproduced with permission from ref. 222, copyright 2018 American Institute of Physics. (b) Two 180° rotated sheets consisting of six-membered MN₄ tetrahedra rings. They are overlaid to form a rigid three-dimensional framework. (c) PLE and PL spectra of Ca_1−xAlSiN₃:xEu²⁺ (x = 0.01, 0.1, 0.2, 0.3, and 0.5%). Reproduced with permission from ref. 216, copyright 2016 American Chemical Society. (d) Emission spectra of Sr_xCa_1−xAlSiN₃:Eu²⁺. The figures indicate Sr occupancy. Reproduced with permission from ref. 209, copyright 2008 Elsevier B.V. (e) Fitting curves by bi-Gaussian method of Ca_0.625[Al_0.75Si_1.25]N₂ (N_0.85O_0.15):0.02Eu. The inset figures reveal the EuN^I₂N^II₃ and EuN^I₂N^II₂O models corresponding to P1 (1.910 eV) and P2 (2.115 eV), respectively. (f) Fitting curves by the bi-Gaussian method of Ca(Al/Si)₂N₂(N_0.8O_0.2):0.02Eu. The figures indicate different excitation wavelengths. The arrows show that the relative intensity of fitting curves change with increasing excitation wavelength. With longer wavelength excitation, P1 intensities increase and P2 intensities decrease. Panels (e and f): reproduced with permission from ref. 215, copyright 2013 Elsevier B.V. The normalized emission spectra of (Ca_1−xLi_x)_0.98(Al_1−xSi_1+x)N₃:0.02Eu²⁺ samples for different ranges: (g) x = 0–0.2, (h) x = 0.3–0.9, (i) x = 0.8–1. (j) Proposed model on the chemical unit co-substitution strategy represented by structural evolution of (Ca_1−xLi_x)_0.98(Al_1−xSi_1+x)N₃:0.02Eu²⁺, from CaAlSiN₃ (x = 0) to LiSi₂N₃ (x = 1) via the (LiSi)⁵⁺ substitution for (CaAl)⁵⁺ couple. It corresponds to three stages in the evolution of photoluminescence. Panels (g–j): reproduced with permission from ref. 219, copyright 2016 American Chemical Society. (k) Photographs of the persistent luminescence phenomenon in Sr_yCa_1−yAlSiN₃:0.05%Eu²⁺ (y = 0–0.95) (up) under UV light excitation (254 nm), and (middle) 15 s and (down) 60 s after the removal of the excitation light. (l) PLE, TLE and DR spectra of Sr_0.8Ca_0.2AlSiN₃:0.15% Eu²⁺. The monitoring emission wavelength is 628 nm for the PLE spectrum. The TLE spectral intensity is obtained by integration of the TL glow between 30 and 400 °C. The DR spectrum is for the Sr_0.8Ca_0.2AlSiN₃ host. (m) Persistent luminescence mechanism for Sr_0.8Ca_0.2AlSiN₃:xEu²⁺. Panels (k–m): reproduced with permission from ref. 94, copyright 2019 Elsevier B.V.

In 2006 Uheda et al. reported the luminescence properties of CaAlSiN₃:Eu²⁺ synthesized under high temperature and pressure, and pointed out its great potential as a red-emitting phosphor for practical application in w-LEDs with a high color rendering index.¹⁸ As displayed in Fig. 13c, CaAlSiN₃:Eu²⁺ exhibits a broad excitation band ranging from UV to blue light, and under 455 nm excitation it gives out a deep-red emission peaking at ∼650 nm.²¹⁶ In order to save energy and reduce cost, self-propagating high-temperature and low-temperature alloy-nitridation synthesis methods were successively developed to prepare high-performance CaAlSiN₃:Eu²⁺ phosphors.^79,217,218 For lighting applications with the aim of both high color rendering and efficiency, deep-red emission is desired to be shifted to shorter wavelengths. For this purpose, Watanable et al. replaced partial Ca in CaAlSiN₃:Eu²⁺ by larger Sr ions to weaken the crystal field strength around Eu²⁺ and thus tune the emission color.²⁰⁹ The solid solutions Eu²⁺-doped Sr_xCa_1−xAlSiN₃ (0 ≤ x ≤ 1) with an isotopic structure of CaAlSiN₃ were synthesized under high pressure and temperature. As x increases, the Eu²⁺ emission gets a blue shift from 650 nm to 610 nm due to the decrease in the crystal field strength caused by unit-cell expansion (Fig. 13d). What is surprising is that the x = 0.8 sample Sr_0.8Ca_0.2AlSiN₃ with the optimum Eu²⁺ concentration has high internal quantum efficiency (80–83%) and external quantum efficiency (70%), which are comparable to that of the x = 0 sample CaAlSiN₃. In addition, by replacing partial N with O, the blue shift of the emission wavelength can also be achieved in Ca(Al/Si)₂N₂(N_1−xO_x):Eu²⁺.²¹⁵ Fitting of the emission spectra into two sub-Gaussian peaks P1 and P2 (Fig. 13e) suggests two types of Eu²⁺ luminescent centers with different ligand environments: EuN^I₂N^II₃ and EuN^I₂N^II₂O tetrahedra.^18,79,215 By increasing the O concentration, P1 is intensified while P2 weakens (Fig. 13f), indicating that oxygen atoms only substitute for one kind of N atoms, i.e., N^II at the 4a crystallographic site.²¹⁵ Co-substitution of the (LiSi)⁵⁺ pair for the (CaAl)⁵⁺ pair in the CaAlSiN₃ host forms a completely miscible solid solution, i.e., (Ca_1−xLi_x)_0.98(Al_1−xSi_1+x)N₃. The Eu²⁺-doped samples exhibit dark red emission with a broad band peaking at about 650–700 nm apart from the compound with x = 1, i.e., LiSi₂N₃:Eu²⁺, as shown in Fig. 13g–i. The 3-stage spectral evolution depends on chemical compositions, which could be explained by local structural evolution of the (Ca_1−xLi_x)_0.98(Al_1−xSi_1+x)N₃:0.02Eu²⁺, from CaAlSiN₃ (x = 0) to LiSi₂N₃ (x = 1) via the (LiSi)⁵⁺ substitution for (CaAl)⁵⁺ couple as shown in Fig. 13j.²¹⁹

It is worth mentioning that there are some well-known Eu²⁺-doped aluminate phosphors, such as BaMgAl₁₀O₁₇:Eu²⁺ and BaAl₂O₄:Eu²⁺, showing potential for LED applications.^220,221 Fine-tuning of the luminescence properties may be achieved by modification of the phosphors as a result of the incorporation of nitrogen.

Compared with the well commercialized green-emitting and blue-emitting persistent phosphors, such as SrAl₂O₄:Eu²⁺,Dy³⁺, CaAl₂O₄:Eu²⁺,Nd³⁺, Sr₂MgSi₂O₇:Eu²⁺,Dy³⁺, and Sr₄Al₁₄O₂₅:Eu²⁺,Dy³⁺, red-emitting persistent phosphors with efficient PersL properties are relatively deficient. Although red PersL phenomena have been observed in some oxides, sulfides, and oxysulfides, such as Y₂O₂S:Eu³⁺,Mg²⁺,Ti³⁺, CaS:Eu²⁺,Ti³⁺,Ce³⁺, Ca₃Ti₂O₇:Pr³⁺, Sr₂ZnSi₂O₇:Sm³⁺, and β-Zn₃(PO₄)₂:Mn²⁺,Sm³⁺/Al³⁺/Ga³⁺, their applications are highly restricted due to either low chemical stability or poor PersL performance.^223–227 Hence, developing efficient red-emitting PersL phosphors has attracted continuous attention. Note that a lot of persistent phosphors are original LED phosphors, and their PersL properties can be effectively improved by controlling defect concentration and introducing trap centers. Therefore, exploring the potential PersL possibilities of red-emitting LED phosphors has been focused on as an effective approach for the discovery of new promising red PersL materials. Initially, the red PersL phenomenon in CaAlSiN₃:Eu²⁺ was accidently observed by Ueda et al. in 2018 when they studied its luminescence quenching mechanism.²²⁸ Shortly after, while Suda et al. studied the origin of the abnormal blue emission of the red phosphor CaAlSiN₃:Eu²⁺ at low Eu²⁺ concentration, the red PersL phenomenon was observed again, and it is proposed that the nitrogen vacancies which can capture electrons take the responsibility for both the blue emission and the red PersL from Eu²⁺.²²² In 2019, Liu et al. systemically investigated the PersL properties of Sr_yCa_1−yAlSiN₃ (0 ≤ y ≤ 0.95) doped with various Eu²⁺ concentrations under UV light excitation, and found that the y = 0.8 sample Sr_0.8Ca_0.2AlSiN₃ doped with 0.15% Eu²⁺ shows the best red PersL performance with a emission peak at ∼628 nm and a duration of 9600 s at the 0.32 mcd m⁻² threshold value (Fig. 13k).⁹⁴ By comparing the thermoluminescence excitation spectrum (TLES) to the photoluminescence excitation spectrum (PLES) (Fig. 13l), a new PersL mechanism different from that of SrAl₂O₄:Eu²⁺,Dy³⁺ was proposed. In the trapping process, the electrons captured by the traps are first excited from the 4f states to the 5d states of Eu²⁺ and then thermally stimulated to CB in SrAl₂O₄:Eu²⁺,Dy³⁺, while in Sr_0.8Ca_0.2AlSiN₃:Eu²⁺ the electrons are directly excited from the 4f states or VB to CB and then captured by the traps (Fig. 13m). Looking forward, future work on exploring the appropriate rare-earth or transition metal co-dopant as the trap center would improve the PersL performance, providing opportunities for Sr_0.8Ca_0.2AlSiN₃:Eu²⁺ toward commercial PersL applications.

Eu²⁺-activated M₂Si₅N₈ (M = Ca, Sr, Ba)

Eu²⁺-activated M₂Si₅N₈ (M = Ca, Sr, Ba), as a class of high-performance orange-red emitting phosphors, have been successfully commercialized in w-LED applications due to their high efficiency and high color purity under near-UV to blue light excitation and excellent thermal luminescence stability up to 200 °C. The structures of M₂Si₅N₈ (M = Ca, Sr, Ba) were first reported by Schnick et al. in 1995.^275,276 M₂Si₅N₈ (M = Sr, Ba) phosphors share the same orthorhombic structure in the space group of Pmn2₁, while Ca₂Si₅N₈ crystalizes in the monoclinic structure with a space group of Cc. The structures of all these three compounds are constructed by [SiN₄] tetrahedra through corner sharing, as shown in Fig. 14a. What is different is that Sr²⁺ and Ba²⁺ ions are located in cavities formed by Si₆N₆ rings parallel along the [100] direction whereas Ca²⁺ ions are accommodated in the structure along the channels of the [010] direction. In 2000, van Krevel first reported the orange-red Eu²⁺ emission of M₂Si₅N₈ (M = Ca, Sr, Ba) under blue light excitation.²⁷⁷ The character of blue light excitation can be attributed to the large covalency of nitrogen and thereby a larger centroid shift of the 5d energy levels of Eu²⁺. Later, Höppe et al. studied the luminescence properties of Ba₂Si₅N₈ with various Eu²⁺ concentrations and ascribed the red shift of the emission maxima with increasing Eu²⁺ concentration to the reabsorption process of Eu²⁺.⁷⁸ After that, dozens of research studies came up, systematically investigating the effect of M cation and Eu²⁺ concentration on the luminescence properties of M₂Si₅N₈ (M = Ca, Sr, Ba):Eu²⁺ (Fig. 14b).^232,235 The maximum solubility of Eu²⁺ in Ca₂Si₅N₈ is about 7 mol%, whereas M cation in M₂Si₅N₈ (M = Sr, Ba) can be completely substituted by Eu²⁺ due to the isostructural character of M₂Si₅N₈ (M = Sr, Ba) and Eu₂Si₅N₈. By varying the composition of different M cations and Eu²⁺ concentrations, the emission wavelength of M₂Si₅N₈ (M = Ca, Sr, Ba):Eu²⁺ can be tuned within a wide range from 570 nm to 680 nm. It was proposed that the emission band shift caused by both different M cation compositions and Eu concentrations originates from the difference in the Stokes shift, which is generally affected by the cation size and the structural rigidity. Among them, Sr₂Si₅N₈:Eu²⁺ exhibits the highest quantum yield of 79% at room temperature and the best luminescence thermal stability with only a few percent of quantum efficiency loss up to 150 °C (Table 3).^234,278 In 2017, Lin et al. further studied the effect of the cation disorder on the luminescence properties in Sr_1.98−x(Ca_0.55Ba_0.45)_xSi₅N₈:2% Eu (x = 0–1.5) systems, demonstrating that replacing Sr by the (Ca_0.55Ba_0.45) mixture with the same average radius as Sr but increased cation size variance results in the increased quantum efficiency and thermal stability of Eu²⁺ emission.²³⁷ It was proposed that cation disorder can suppress nonradiative processes through disruption of lattice vibrations and create deep traps that release electrons to compensate for thermal quenching, which provides a fresh perspective for performance optimization of luminescent materials. Since high chemical stability and minimal thermal degradation are the prerequisites for practical applications, it is revealed that a higher degree of cross-linking between Si(O,N)₄ tetrahedral enhances the chemical stability of M–Si–N compounds and the oxidation resistance of Ln–Si–O–N and Ce–Si–O compounds.^279,280 It has been demonstrated that the degradation stability of Sr₂Si₅N₈:Eu²⁺ can be enhanced by deposition of an ultra-thin Al₂O₃ coating using atomic layer deposition in a fluidized bed reactor.²⁸¹

Fig. 14 (a) Crystal structure of M₂Si₅N₈ (M = Ca, Sr). (b) Emission spectra and excitation spectra of Sr_2−xSi₅N₈:Eu_x with varied contents from x = 0.02–0.40 and Ba_2−xSi₅N₈:Eu_x (x = 0.10). Reproduced with permission from ref. 282, copyright 2012 American Chemical Society. (c) Afterglow decay curves of the Ca₂Si₅N₈:2% Eu²⁺, Ca₂Si₅N₈:2% Eu²⁺, 2% Tm³⁺ and Ca₂Si₅N₈:0.5% Eu²⁺, 0.5% Tm³⁺ phosphor λ_ex = 400 nm and λ_em = 600 nm. Reproduced with permission from ref. 32, copyright 2010 The Electrochemical Society. (d) Afterglow intensity of Ca₂Si₅N₈:Eu,R 1 min after excitation with 400 nm light, as a function of the co-dopant. Reproduced with permission from ref. 263, copyright 2011 Multidisciplinary Digital Publishing Institute. (e) Energy level scheme for M₂Si₅N₈:Eu persistent phosphors. Reproduced with permission from ref. 96, copyright 2011 Elsevier B.V.

The PersL performance of M₂Si₅N₈ (M = Ca, Sr, Ba):Eu²⁺ was first discovered in Ba₂Si₅N₈:Eu²⁺, but its PersL duration is pretty short of about 15 min under the naked eye.⁷⁸ Until 2009, Miyamoto et al.³³ and Eeckhout et al.²⁸³ independently reported PersL in Ca₂Si₅N₈:Eu²⁺, and found that its PersL intensity and duration can be enhanced by co-doping with Tm³⁺ as shown in Fig. 14c. Although the PersL duration of Ca₂Si₅N₈ with a single Eu²⁺ dopant is only 150 s under 400 nm excitation taking 0.32 mcd m⁻² as the threshold; it can be increased to 1 h upon co-doping with Tm³⁺, which is comparable with the conventional red-emitting persistent phosphor Y₂O₂S:Eu³⁺,Mg²⁺,Ti³⁺ but with better luminescence and chemical stability.³² After that, Eeckhout et al. further investigated the PersL behavior in the M₂Si₅N₈ (M = Ca, Sr, Ba):Eu²⁺ family and thoroughly studied the effect of the rare-earth co-dopant, dopant concentration, and starting materials on their PersL properties.²⁶³ All of the samples show PersL after 300 nm irradiation, but it is weak for all the M = Sr samples. Ca₂Si₅N₈:Eu²⁺,Tm³⁺ shows the brightest and longest PersL duration while Dy³⁺, Nd³⁺, Ho³⁺ or Sm³⁺ co-dopant scarcely improves the PersL performance of Ca₂Si₅N₈:Eu²⁺ (Fig. 14d), and thermoluminescence measurements indicate that the addition of Tm³⁺ in Ca₂Si₅N₈:Eu²⁺ deepens the relevant charge trap to a more suitable depth for improved PersL performance while others do not. It was verified by the HRBE scheme of Ca₂Si₅N₈ constructed soon after.²³³ The energy gap between the 4f level of Tm³⁺ and the CB bottom is well consistent with the trap depth introduced by Tm³⁺ of about 0.91 eV extracted from thermoluminescence spectra, which is suitable for PersL at room temperature. Whereas, trap depth introduced by Dy³⁺, Nd³⁺, Ho³⁺ or Sm³⁺ inferred from the HRBE scheme is too shallow to contribute to PersL at room temperature. It was also suggested that the dopant and co-dopant in their fluoride form with concentrations not higher than 1 mol% work better for PersL performance. In addition, Teng et al. reported that the PersL properties of Sr₂Si₅N₈:Eu²⁺ can also be enhanced by co-doping Tm³⁺.²⁶⁵ Inspired by the excellent PersL performance of SrAl₂O₄:Eu²⁺,Dy³⁺,B³⁺, our group further improved the PersL performance of Sr₂Si₅N₈:Eu²⁺,Tm³⁺ by co-doping B³⁺ and Zhang et al. improved that of Ca₂Si₅N₈:Eu²⁺,Tm³⁺ by co-doping Dy³⁺.^264,284 Apart from emergency signage, Ca₂Si₅N₈:Eu²⁺,Tm³⁺ also shows potential for application in in vivo imaging, by following the PersL of nano-sized particles which were excited just before injection.²⁸⁵

Combining thermoluminescence excitation spectroscopy and a study on the influence of temperature during excitation, Smet et al. studied the trap filling behavior in M₂Si₅N₈:Eu²⁺ (Tm³⁺) in depth and addressed the crucial role of the thermal barrier between the 5d excited state of Eu²⁺ and the CB bottom in PersL performance.⁹⁶ They showed that the traps relevant for PersL can be filled after electrons are excited into the 5d energy levels of Eu²⁺, and for those into lower 5d levels, a thermal barrier is present before trapping can occur. By virtue of HRBE schemes, it is demonstrated that Ca₂Si₅N₈:Eu²⁺,Tm³⁺ being the phosphor with the best PersL performance upon excitation with visible light has the smallest thermal barrier (Fig. 14e), which makes it easy for electrons to be captured by the trap and then released to give out PersL. Note that this thermal barrier also takes responsibility for thermal luminescence stability. Hence, the large thermal barrier endows Sr₂Si₅N₈:Eu²⁺ with the best thermal luminescence stability, enabling it to be developed as a high-performance LED phosphor, while Ca₂Si₅N₈:Eu²⁺ with a small thermal barrier exhibits the best PersL performance. This work provides a clue that the quest for Eu²⁺-based persistent phosphors needs to focus on those compounds with a relatively low thermal quenching temperature, as this makes the trap filling process easier.

Ce³⁺-activated La₃Si₆N₁₁

Yellow emitting La₃Si₆N₁₁:Ce³⁺ is an important commercial phosphor for high-power w-LEDs due to its excellent thermal luminescence stability (almost no emission loss at 150 °C) and sufficient emission efficiency (nearly 100%).^238,286 The crystal structure of La₃Si₆N₁₁ was first established in 1983 in the compound Sm₃Si₆N₁₁.²⁸⁷ It crystallizes in the tetragonal space group P4bm. The structure consists of La layers and SiN₄ tetrahedra. As shown in Fig. 15a, the SiN₄ groups are connected via corner sharing, forming a three-dimensional network, in which the La atoms occupy two different cavities both coordinated by eight nitrogen neighbors. As the problem of underestimation of the band gap for first principles density functional theory (DFT) calculations has been reduced by more recent first principles approaches, it is revealed that for M–Si–N compounds the top of the VB band is composed of N atoms with the lowest coordination number with Si, implying that the band gap of the M–Si–N compounds becomes higher for increasing degree of cross-linking between SiN₄ tetrahedra.^288–291 It is also demonstrated that a higher Si/N ratio generally leads to a higher Eu²⁺ 4f–5d absorption energy, a higher 5d–4f emission energy and a larger Stokes shift.²⁹² Under 460 nm excitation, La₃Si₆N₁₁:Ce³⁺ shows a broad-band emission peaking at 535 nm with an FWHM of 112 nm (Fig. 15b).²⁹³ With regard to the luminescence mechanism, several different viewpoints have been put forward. George et al. believed that Ce³⁺ on different La³⁺ sites exhibit similar emission properties for their identical La/Ce–N coordination numbers and similar bond lengths,²³⁸ whereas Jia et al. proposed that the yellow emission could be ascribed to Ce_La(2), and no luminescence in Ce_La(1) for its 5d excited states immersed into the conduction band (CB) based on DFT calculations.²⁹⁴ Coincidentally, in 2015, Park et al. prepared La₃Si₆N₁₁:Ce³⁺ with a low Ce³⁺ concentration to study its decay behavior and observed an extra blue emission ranging from 400 to 475 nm under 355 nm excitation.²⁹⁵ Based on this phenomenon, combining with both theoretical and experimental methods, Zhong et al. ascribed the yellow emission to Ce_La(2) and the blue emission to Ce_La(1).²⁹⁶

Fig. 15 (a) Crystal structure and local structure around La sites of La₃Si₆N₁₁. (b) Normalized emission spectra of La_2.86−xY_xSi₆N₁₁:0.14Ce³⁺ (x = 0–0.9) phosphors under 460 nm excitation. (c) Emission intensity and peak location of La_3−zSi₆N₁₁:zCe³⁺ with varying Ce³⁺ concentrations (z = 0–0.2). Panels (a–c): reproduced with permission from ref. 293, copyright 2016 Royal Society of Chemistry. (d) Excitation spectra monitored at 580, 597, 612, 627, 643 and 664 nm, respectively, and emission spectra under 450 nm excitation of La_2.9Si_6−xAl_xN_11−x/3:0.1Ce³⁺ (x = 0–0.3). Reproduced with permission from ref. 239, copyright 2020 American Chemical Society. (e) Temperature-dependent normalized integrated PL intensities of La_2.9Si_6−xAl_xN_11−x/3:0.1Ce³⁺ (x = 0–0.3). Reproduced with permission from ref. 297, copyright 2017 Royal Society of Chemistry.

Compared with the YAG:Ce³⁺ phosphor, La₃Si₆N₁₁:Ce³⁺ exhibits superior thermal luminescence stability and high quantum efficiency due to the high rigidity of the host.^238,293 However, it suffers from the low color rendering and high color temperature of the fabricated devices resulting from its greenish yellow emission (peaking at 535 nm), and also the pure phase requires harsh synthesis conditions.²⁹³ Therefore, extensive research studies have been devoted to optimizing the synthesis process and tuning the luminescence properties. In 2017, Chen et al.²⁹⁸ and Suehiro et al.²⁹⁹ reported that incorporating M (M = Ca, Sr, Ba) at La sites in La₃Si₆N₁₁:Ce³⁺ could not only improve the sample purity but also lead to the slight red shift of emission with an increased external quantum efficiency up to ∼42%. Du et al. also made a great effort to shift the emission of La₃Si₆N₁₁:Ce³⁺ toward a longer wavelength, including varying Ce³⁺ concentrations, Y³⁺ substitution for La³⁺, and Al³⁺ substitution for Si⁴⁺, but unfortunately the effect was fairly insignificant either way.^293,297 As shown in Fig. 15b and c, the emission spectra of La_3−z−xY_xSi_6−yAl_yN_11−y/3:zCe³⁺ (z = 0.14, y = 0) show a very small red shift from 535 nm to 552 nm as the substituting Y³⁺ concentration x increases from 0 to 0.9, and no obvious change in the emission peak could be observed with either different Ce³⁺ concentrations (0 ≤ z ≤ 0.2) or Al substitution (0 ≤ y ≤ 0.1). Recently, almost at the same time, You et al.²³⁹ and Nitta et al.³⁰⁰ reported that an efficient red shift can be achieved by increasing the substituting Al³⁺ concentration (Fig. 15d). Under 530 nm excitation, La_2.9Si_6−xAl_xN_11−x/3:0.1Ce³⁺ (x = 0.3) shows red emission peaking at ∼650 nm, but the thermal luminescence stability decreases as Al³⁺ concentration increases (Fig. 15e). Through DFT calculation, You et al.²³⁹ believed that this unusual red emission band originates from the interstitial Ce³⁺ site occupying the [Si₈N₈] void in the c ∼ 1/2 layer, created as a charge compensator for the aliovalent Al³⁺ substitution for Si⁴⁺. Therefore, the interstitial site engineering was proposed, for the first time, as an efficient strategy for phosphor design, and the aliovalent ion co-doping was proved to be an effective way to create interstitial sites.

Another characteristic of the Ce³⁺ luminescence in La₃Si₆N₁₁ is the neutralization of cationic substitution on the excitation peak shift. As reported by Du et al., both yttrium and cerium substitution for lanthanum almost have no effect on the peak shift of excitation spectra, as shown in Fig. 16a and b.^286,293 In contrast, the lattice parameter and bond lengths demonstrate significant changes depending on cationic substitution contents. This anomaly has been investigated by crystal structure modelling and crystal-field calculation.³⁰¹ The real coordination polyhedron is simplified to a truncated square pyramid with point symmetry C_4v. The different sizes of the parallel top and down squares, together with the pyramid height serve as the varying coefficient in expressing the split crystal-field levels. From the simulation results, it is found that specific geometrical shapes have the effect of freezing the crystal-field levels on cationic substitution, regardless of the significant bond-length variance. Further investigation of the orbital shape in Fig. 16c reveals possible explanation of the crystal-field passivation, i.e., the diminished influence of the expanding square on z-oriented orbitals d_z², d_xz and d_yz.

Fig. 16 (a) PLE spectra of La_2.86−xY_xSi₆N₁₁:0.14Ce³⁺ with different Y contents. Reproduced with permission from ref. 293, copyright 2016 Royal Society of Chemistry. (b) PLE spectra of La_3−xSi₆N₁₁:xCe³⁺ with different cerium contents. Reproduced with permission from ref. 286, copyright 2017 Elsevier B. V. (c) Schematic orbital shape and the passivation of crystal-field effect. Reproduced with permission from ref. 301, copyright 2020 Royal Society of Chemistry.

Eu²⁺-activated UCr₄C₄-type compounds

UCr₄C₄-type compounds have been investigated for narrow band emitting phosphors due to their highly condensed, rigid framework structure. The general formula of UCr₄C₄-type structures is written as M(A,B)₄X₄. [AX₄] and [BX₄] tetrahedra connect to each other by vertex- and edge-sharing to form the highly condensed three-dimensional network with the degree of condensation k = 1 (k equals to the atomic ratio (A,B):X). M cations (alkali/alkali-earth metal ions) fill in a highly symmetric cube-like site within vierer ring channels along the [001] direction.¹¹ The narrow band emitting phosphors based on the UCr₄C₄ structure were first discovered by Schnik's group. They reported a series of Eu²⁺-doped nitridoaluminates which give out narrow band red emitting characteristics, such as M[LiAl₃N₄]:Eu²⁺ (M = Ca, Sr),^1,302 M[Mg₂Al₂N₄]:Eu²⁺ (M = Ca, Sr, Ba),³⁰³ and Sr[Mg₃SiN₄]:Eu²⁺.³⁰⁴ Among them, SrLiAl₃N₄:Eu²⁺ shows the most superior luminescence properties. As shown in Fig. 17a, the ordered edge- and corner-sharing AlN₄ and LiN₄ tetrahedra with channels of vierer rings along [011] form the highly condensed and rigid framework of Sr[LiAl₃N₄].¹ There are two crystallographic Sr sites, each coordinated by eight N atoms in a highly symmetric cuboid-like environment. The Eu²⁺ dopant is expected to replace Sr ions.³⁰⁵ Sr[LiAl₃N₄]:Eu²⁺, which can be efficiently excited by GaN-based blue LEDs, shows highly efficient red emission peaking at ∼650 nm with a full-width at half-maximum of ∼50 nm and pretty low thermal quenching (>95% relative to the quantum efficiency at 200 °C) (Fig. 17b). The emission peak position of Sr[LiAl₃N₄]:Eu²⁺ is similar to that of commercial CaAlSiN₃:Eu²⁺, but the significantly narrowed bandwidth can improve the color gamut and concentrate the emission within the visible spectral region. The color gamut of LED fabricated by β-SiAlON:Eu²⁺, Sr[LiAl₃N₄]:Eu²⁺ with a blue LED chip reaches up to 84.2% of NTSC in CIE 1931 after passing through the color filter.³⁰⁶ By DFT calculations, the electronic structure of Sr[LiAl₃N₄] was investigated and it was revealed that the DOS of the two non-equivalent Sr sites shows a high degree of similarity, contributing to the narrow band emission of Sr[LiAl₃N₄]:Eu²⁺ (Fig. 17c).^60,307 The PersL properties of Sr[LiAl₃N₄]:Eu²⁺ were also explored. After 254 nm light pre-irradiation, Sr[LiAl₃N₄]:0.1%Eu²⁺ phosphor shows 395 s red PersL at a 0.32 mcd m⁻² threshold value as shown in Fig. 17d.³⁰⁸

Fig. 17 (a) Crystal structure of SrLiAl₃N₄. Yellow spheres are Sr; blue tetrahedra are AlN₄; red polyhedral are LiN₄; and blue spheres are N. (b) Excitation (Sr[LiAl₃N₄]:Eu²⁺, blue; CaAlSiN₃:Eu²⁺, light grey) and emission spectra for λ_exc = 440 m of Sr[LiAl₃N₄]:Eu²⁺ (pink) and CaAlSiN₃:Eu²⁺ (dark grey). The dotted curve indicates the upper limit of sensitivity of the human eye. (c) Calculated DOS in the VB (upper half of left panel) and CB (lower half of left panel) and the band structure (right panel) of Sr[LiAl₃N₄]. (d) The persistent decay curves (log–log plot) of Sr[LiAl₃N₄]:Eu²⁺ after 254 nm light irradiation for 5 min at RT. (e) Crystal structures of RbLi(Li₃SiO₄)₂:Eu²⁺ and RbNa(Li₃SiO₄)₂:Eu²⁺. (f) Excitation and emission spectra of RbLi(Li₃SiO₄)₂:Eu²⁺ (RLSO:Eu²⁺), RbNa(Li₃SiO₄)₂:Eu²⁺ (RN:Eu²⁺), and β-SiAlON:Eu². Panels a and b: reproduced with permission from ref. 1, copyright 2014 Macmillan Publishers Limited. Panel c: reproduced from ref. 307, copyright 2015 Willy-VCH. Panel (d): reproduced from ref. 308, copyright The Royal Society of Chemistry 2020. Panels e and f: reproduced from ref. 312, copyright 2019 Willy-VCH.

Based on the insight into nitridoaluminates with the UCr₄C₄-type structure, Huppertz and co-workers together with OSRAM Opto Semidonductors issued the patent WO 2018/029304 A1 on the alkali lithosilicate phosphors in Germany, where several narrow band emitting representatives with the UCr₄C₄-type structure were described.³⁰⁹ They also reported that alkali metal silicates such as Na[Li₃SiO₄]:Eu²⁺ and K[Li₃SiO₄]:Eu²⁺ show efficient narrow band emission.³¹⁰ Almost at the same time, Xia et al. independently reported a series of Eu²⁺-doped oxidesilicate phosphors with the UCr₄C₄-type structure which exhibit narrow band emitting characteristics, such as RbLi(Li₃SiO₄)₂:Eu²⁺,³¹¹ RbNa(Li₃SiO₄)₂:Eu²⁺,³¹² RbNa₃(Li₃SiO₄)₄:Eu²⁺,³¹³ and Na_0.5K_0.5Li₃SiO₄:Eu²⁺.³¹⁴ RbLi(Li₃SiO₄)₂:Eu²⁺ exhibits a narrow band green emission peaking at 530 nm with the full width of highest maximum of 42 nm. Later, by cation substitution strategy, RbNa(Li₃SiO₄)₂:Eu²⁺ was designed via partially substituting Li by Na in RbLi(Li₃SiO₄)₂. It shows similar luminescence properties to RbLi(Li₃SiO₄)₂:Eu²⁺ with emission peaking at 523 nm and the full width of highest maximum of 41 nm. Both phosphors share a similar crystal structure with the monoclinic system and a space group of C2/m (Fig. 17e). Compared with the commercial phosphor β-SiAlON:Eu²⁺, the PL spectra of RbLi(Li₃SiO₄)₂:Eu²⁺ and RbNa(Li₃SiO₄)₂:Eu²⁺ show narrower emission band (Fig. 17f), and thus the color gamut of RbLi(Li₃SiO₄)₂:Eu²⁺ and RbNa(Li₃SiO₄)₂:Eu²⁺ based LED is 107% and 113% of NTSC, respectively, which is larger than that of β-SiAlON:Eu²⁺ based LED. These inspiring breakthroughs in narrow band emitting phosphors provide more opportunities for wide color gamut LCD technology. However, these Eu²⁺-doped UCr₄C₄-type phosphors suffer from poor chemical stability. Even though endeavors have been devoted to improve the chemical stability by hydrophobic modification, its commercial application is still restricted.^315,316 Further work for improving the chemical stability of Eu²⁺-doped UCr₄C₄-type phosphors is highly required to realize their practical applications.

Eu²⁺-activated MSi₂O₂N₂ (M = Ca, Sr, Ba)

Oxynitridosilicates, also known as siones, represent a class of solid compounds, which can be formally derived from oxysilicates by partial substitution of oxygen by nitrogen.⁷⁰ These materials combine attractive material properties like high mechanical hardness and strength, and great thermal and chemical stability. Meanwhile, higher condensed siones are known to be highly covalent and stable towards oxidation and hydrolysis. Among them, MSi₂O₂N₂ (M = Ca, Sr, Ba) have been well developed as famous phosphors applied in LED applications by doping with Eu²⁺. This family of compounds is constructed with silicate layers separated by M²⁺ layers.^317–319 As an example, Fig. 18a shows the triclinic structure of SrSi₂O₂N₂ that crystallizes in space group P1. The CaSi₂O₂N₂ crystallizes in the monoclinic space group P2₁, whereas BaSi₂O₂N₂ crystallizes in the orthorhombic space group Cmcm. Each of the four M sites in MSi₂O₂N₂ (M = Ca, Sr) is coordinated with six oxygen ions and one nitrogen ion. When doped with Eu²⁺, these compounds exhibit a broad excitation spectrum covering the UV-blue range (250–460 nm), and yield efficient emission in the blue-green to yellow region. As shown in Fig. 18b, CaSi₂O₂N₂:Eu²⁺ shows yellow emission with a maximum at 560 nm, and green-yellow emission over 530–570 nm is observed in SrSi₂O₂N₂:Eu²⁺ depending on the O/N ratio and Eu²⁺ concentration.^240,320 Remarkably, BaSi₂O₂N₂:Eu²⁺ shows blue-green emission with a narrow emission band at 494 nm, which indicates its potential application in LED backlights. The high quantum yield (90%) and high quenching temperature (>500 k) of SrSi₂O₂N₂:Eu²⁺ were also verified, demonstrating that it is a promising material for pc-LED application.

Fig. 18 (a) Crystal structure of SrSi₂O₂N₂. Reproduced with permission from ref. 241, copyright 2017 American Chemical Society. (b) PLE and PL spectra of 2% Eu²⁺ doped CaSi₂O₂N₂, SrSi₂O₂N₂, and BaSi₂O₂N₂ phosphors, as recorded at 4 K; PL spectra were taken upon excitation at 450 nm; PLE spectra were obtained at the corresponding emission maxima. (c) Temperature dependence of the integrated emission intensity Sr_0.98−xBa_xSi₂O₂N₂:Eu 2% (x = 0, 0.25, 0.5). Panels (b and c): reproduced with permission from ref. 240, copyright 2009 American Chemical Society. (d) Normalized PL emission spectra (λ_ex = 405 nm): (i) digital phosphor images upon irradiation using a 365 nm UV lamp of the typical (Sr_0.98−xBa_xEu_0.02)Si₂O₂N₂ phosphors; (ii) x = 0–0.63; (iii) x = 0.65–0.77; and (iv) x = 0.78–0.98. Reproduced with permission from ref. 321, copyright 2014 American Chemical Society.

To improve the color rendering index, correlated color temperature, and thermal luminescence stability of MSi₂O₂N₂ (M = Ca, Sr, Ba):Eu²⁺ phosphors, cation substitutions and Eu²⁺ concentration variation have been used to adjust their luminescent properties.^240,321 As Eu²⁺ concentration increases upon 2%, the emission gets a red shift, which can be attributed to the energy migration and energy transfer to Eu²⁺ ions emitting at longer wavelengths. Unfortunately, the lowering of quantum efficiency and thermal quenching temperature due to concentration quenching occurs along with the red shift. On the other hand, partial substitution of Sr²⁺ by either Ca²⁺ or Ba²⁺ also results in a red shifted Eu²⁺ emission. Similar to the case of increasing Eu²⁺ concentration, the quenching temperature drops upon substitution of Sr by Ca. What is exciting is that the quantum efficiency and quenching temperature remain high for Ba substitution, as shown in Fig. 18c and Table 3. For Ca²⁺ substitution, the red shift is expected and can be understood by an increased crystal field splitting for Eu²⁺ on the smaller Ca²⁺ cation site. Whereas, the red shift from substitution of Sr²⁺ with Ba²⁺ cation is intriguing since a blue shift of the f–d emission of Eu²⁺ is generally observed in case of substitution of Sr²⁺ by the larger Ba²⁺ cation. In order to explore the underlying mechanism of this abnormal red shift, Li et al. conducted a detailed work on the phase structure and tunable luminescence of (Sr_1−xBa_x)Si₂O₂N₂:Eu²⁺ (0 ≤ x ≤ 1). The multicolor emission of (Sr_0.98−xBa_xEu_0.02)Si₂O₂N₂ (0 ≤ x ≤ 1) phosphors can be clearly observed upon irradiation using a 365 nm UV lamp as displayed in Fig. 18d–i. By virtue of X-ray diffraction measurements and first-principles calculations, it is revealed that the crystal structures of (Sr_1−xBa_x)Si₂O₂N₂:Eu²⁺ (0 ≤ x ≤ 1) can be nominally divided into three sections, namely, Phase 1 (0 ≤ x ≤ 0.65), Phase 2 (0.65 < x < 0.80), and Phase 3 (0.80 ≤ x ≤ 1). Through extended X-ray absorption fine structure analysis, it is proposed that the abnormal red shift in Phase 1 with gradual replacement of Sr²⁺ by Ba²⁺ (Fig. 18d-ii) is caused by the Sr(Eu)-O/N bond length shrinkage in the local structure. The controllable luminescence in Phase 2 from blue to white (Fig. 18d-iii) and Phase 3 from cyan to yellowish green (Fig. 18d-iv) is attributed to phase segregation based on the high-resolution transmission electron microscopy and selected area electron diffraction analysis. This work serves as a guide in developing oxynitride luminescent materials with controllable optical properties based on variations in local coordination environments through cation substitutions.

The green-yellow PersL in SrSi₂O₂N₂:Eu²⁺ was first reported by Zhang et al. in 2013.³²² It was in 2018 that Zhuang et al. investigated its potential for information storage applications.²⁴¹ Different from room temperature applications, the information storage applications require PersL materials with a much larger trap depth, typically in the range of 0.8–1.6 eV, to retain the recorded information at room temperature for a sufficiently long duration. On the basis of HRBE and VRBE schemes, led by the strategy of trap depth engineering, Dy³⁺, Ho³⁺, and Er³⁺ with a 0.84–1.16 eV energy gap between their 4f energy levels and the CB bottom are introduced into SrSi₂O₂N₂:Eu²⁺ as the trap center (Fig. 19a). After irradiated by a 254 nm lamp, SrSi₂O₂N₂:Eu²⁺,Ho³⁺ shows the highest intensity and the longest duration (360 min) at room temperature, while SrSi₂O₂N₂:Eu²⁺,Dy³⁺ shows the weakest intensity and shortest duration as shown in Fig. 19b and c. The capability of these materials for information storage application is demonstrated in Fig. 19d. After removing the UV light, the samples were kept at room temperature for 60 min, and finally heated to 100 °C. The Er³⁺ co-doped sample gives the most intense emission in the first several minutes, whereas the luminous pattern from the Dy co-doped samples is almost unrecognizable in the first 60 min. However, when the sample is heated to 100 °C, the emission is greatly enhanced in all samples, and the emission intensity increases in the sequence of Er, Ho, and Dy. This is exactly opposite to the intensity that is detected at RT, which can be explained by the increased trap depth in the sequence of Er, Ho, and Dy from 0.8 to 1.2 eV derived from thermal luminescence spectra shown in Fig. 19e–g and Table 4. The captured electrons in the shallow traps can be released by thermal assistance at room temperature, giving rise to the most intense initial PersL intensity of SrSi₂O₂N₂:Eu²⁺,Er³⁺, while the sufficiently deep traps introduced by Dy³⁺ have the ability to effectively restrain the detrapping at room temperature and thus ensure intense emissions as optical signals for the information readout. This work not only validates the capability of deep-trap PersL materials like SrSi₂O₂N₂:Eu²⁺,Dy³⁺ in the information storage applications, but also provides inspirations to trigger more explorations toward new materials for applications in anticounterfeiting and advanced displays.

Fig. 19 (a) HRBE (left ordinate) and VRBE (right ordinate) schemes of SrSi₂O₂N₂ (SSON). The 4f ground states (Ln²⁺:4f) and the lowest 5d excited states (Ln²⁺:5d) of Ln²⁺ are labeled by black dots and hollow circles and connected by solid curves. The 4f ground states of Ln³⁺ (Ln³⁺:4f) are connected by a dashed curve. (b) PersL decay curves in SSON:Eu,Ln (Ln = Dy, Ho, and Er). The samples are irradiated by a Hg lamp dominantly at 254 nm for 20 s before recording. The insets show enlarged parts in the beginning stage. (c) Photographic images of the pellet samples under natural light (i); under UV excitation (ii); and 10 s (iii), 1 min (iv), 5 min (v), and 20 min (vi) after removal of the excitation source at RT. The samples for 1, 2, 3 are SSON:Eu,Ln, Ln = Dy, Ho, and Er, respectively. The UV excitation source is 254 nm. (d) Schematic illustration of the application in information storage using deep-trap persistent luminescent materials. (e) Decay curves of the luminous patterns on SSON:Eu,Ln. The samples were heated from RT to 100 °C at 60 min. (f) Photographic images of the pellets at RT and 100 °C. Exposure times of the digital camera are 0.5 s. The photos marked by RT and 100 °C were taken at approximately 1 and 61 min after removal of the UV light. The scale bar is 5 mm. (g) TL glow curves of SSON:Eu,Ln. Panels a-g: reproduced with permission from ref. 241, copyright 2017 American Chemical Society.

Eu²⁺-activated sialon

Sialons are solid solutions of the M–Si–Al–O–N system derived from silicon nitride Si₃N₄,³²³ which has two phases denoted as α and β. They could be described as different stackings of Si–N layers in ABCD… and ABAB… sequences, respectively. α-Si₃N₄ has two large closed interstitial sites per unit cell and β-Si₃N₄ has long continuous channels.³²⁴ As shown in Fig. 20, α-Si₃N₄ has two Si sites, and β-Si₃N₄ has only one Si site. The composition of α-sialon could be expressed by M_xSi_12−(m+n)Al_(m+n)O_nN_16−n, where m (Al–N) pairs and n (Al–O) pairs replace m + n (Si–N) pairs in each unit cell. The molar fraction x is obtained by m divided by the valence of M cation, which includes Li, Mg, Ca, Y or heavier lanthanide ions. Since the two interstitial sites exist in every unit cell, the value of x representing the M content has a limit of 2, and m needs to be adjusted to meet the electroneutrality requirement.

Fig. 20 Crystal structures of Si₃N₄. (a) α-Si₃N₄ has two Si sites as represented by the blue and orange SiN4 tetrahedra. (b) β-Si₃N₄ only has one type of Si site. The different colors of the SiN4 tetrahedra reveal the regularly placed 6 unit ring channel.

In 2002, van Krevel et al. and Xie et al. reported the luminescence properties of Eu²⁺-doped Ca-α-sialon.^277,325,326 It can be excited by UV to blue light (300–470 nm) and give yellow emission (570–600 nm). The emission could be tuned by composition modulation, such as changing the m and n values, substituting Ca for other metal ions such as Li, Mg and Y, adjusting the O/N content and varying the Eu concentration.^327,328 Yang et al. synthesized Eu²⁺-doped pure-nitride α-sialons with nominal formula (Ca_0.995Eu_0.005)_m/2Si_12−mAl_mN₁₆ using the alloy-nitridation method.³²⁹ Samples with 2.4 ≤ m ≤ 4.0 have the strongest emission, and exhibit a red-shift with increasing m values as well as Eu concentrations.

The PersL of Eu²⁺-doped Ca-α-sialon has been reported in CaSi_10−nAl_2+nO_nN_16−n:xEu²⁺(m = 2,n = 0–1, x = 0.1–8%).³³⁰ It has a maximum persistent duration of about 60 min for the CaSi₁₀A1₂N₁₆:0.5% Eu²⁺ sample. DFT calculations prove that the charge carriers are electrons which are excited from the 4f ground state of Eu to the conduction band directly in the charging process, and the trap levels are introduced by oxygen/nitrogen defects. It has a potential application in anti-counterfeit and information storage.

As a currently commercial narrow-band green phosphor, β-sialon:Eu²⁺ is widely applied to fabricate wide color gamut backlights owing to its outstanding luminescence properties with a narrow emission band (FWHM ∼ 55 nm), superior thermal luminescence stability (∼10% emission loss at 150 °C), and high quantum efficiency.^242,331 The solid solution β-sialon with formula Si_6−zAl_zO_zN_8−z, 0 < z ≤ 4.2, structurally derived from β-Si₃N₄via simultaneous equivalent substitution of Al–O for Si–N, was discovered by Oyama and Kamigaito in 1971.³³² It crystallizes in a hexagonal structure with the space group of P6₃/m. As shown in Fig. 21a, the structure is built up by corner sharing (Si/Al)(N/O)₄. About 30 years later, Hirosaki et al. reported for the first time the green photoluminescence of Eu²⁺-doped β-sialon with z = 0.17, which exhibits a narrow-band green emission centering at 535 nm with an FWHM of 55 nm under near UV or blue light excitation.³³³ Then, Xie et al. demonstrated that its emission wavelength can be tuned from 528 to 550 nm with an increasing z value.³³⁴ The typical PLE and PL spectra for β-sialon:Eu²⁺ with the nominal composition of Eu_0.015Si_5.5Al_0.485O_0.515N_7.485 are displayed in Fig. 21b as an example. In 2008, Li. et al. gave the formula Eu_xSi_6−zAl_z−xO_z+xN_8−z−x for Eu²⁺-doped β-sialon, and proposed that Eu²⁺ ions are located in the large channels along the c-axis, which are formed by six (Si/Al)(O/N)₄ tetrahedra.³³⁵ Later, it was confirmed by Kimoto et al. using scanning transmission electron microscopy (STEM).⁴⁹ What is more, many endeavours have been devoted to investigating the origin of the narrow band emission, and it is revealed that the interstitial site in the highly symmetric hexagonal channel accommodating Eu²⁺ plays a crucial role in the narrow band emission.³³¹ Since the unusual occupation of Eu²⁺ in the host lattice and some Eu³⁺ ions existing in the channel of β-sialon instead of Eu²⁺ generally result in decreased luminescence efficiency, Li et al. promote the reduction of Eu³⁺ to Eu²⁺ in β-sialon by post-annealing in an N₂–H₂ reducing atmosphere and thereby improve the luminescence efficiency of β-sialon:Eu²⁺.³³⁶ Recently by virtue of density functional theory (DFT) calculations, Wang et al. revealed that Eu²⁺ is first coordinated by nine nitrogen atoms, forming a highly symmetric EuN₉ polyhedron in β-sialon:Eu²⁺.³³⁷ They also attributed its excellent thermal luminescence stability to the large energy gap between the 5d levels and the conduction band (CB) of the host (∼0.6 eV) calculated by DFT calculations.

β-sialon:Eu²⁺ also exhibits PersL behaviour. After 254 nm irradiation, green PersL is observed for 400 s with the threshold value of 0.32 mcd m⁻² (Fig. 21c).²⁴³ By virtue of density functional theory (DFT) calculations, it is revealed that the trap levels responsible for PersL are impurity levels induced by Si–O bonds located below the bottom of the conduction band (CB) on random substitution of Al–O for Si–N in β-Si₃N₄, as shown in Fig. 21d. Through comparing thermoluminescence excitation (TLE) and photoluminescence excitation (PLE) spectra in Fig. 21e, it is confirmed that, in the charging process for PersL, electrons are directly excited from the 4f⁷ ground state of Eu²⁺ to the CB, which can be explained by the large barrier for thermal ionization of β-sialon:Eu²⁺ and is well consistent with its excellent thermal luminescence stability. Based on the TLE spectrum, a novel strategy is proposed to construct HRBE and VRBE schemes, i.e., using the onset energy of the TLE spectrum as the energy difference between the 4f ground state and the bottom of CB to pinpoint the 4f energy level location of Eu²⁺ (Fig. 21f). Hopefully, through co-dopant introduction and ion substitution led by electronic structure engineering, the PersL properties of β-sialon:Eu²⁺ could be further improved toward practical applications in emergency signage and optical information storage.

Fig. 21 (a) Crystal structure of β-sialon:Eu²⁺. (b) PLE (λ_em = 537 nm) and PL (λ_ex = 302 nm) spectra of Eu_0.015Si_5.5Al_0.485O_0.515N_7.485 (ESAON). (c) PersL decay time of ESAON after 254 nm irradiation for 10 min at room temperature. (d) Computed electronic band dispersion, DOS, and PDOS of the perfect compound β-sialon (z = 0.5). The impurity levels of the Si–O bond are denoted by orange lines. (e) PLE and TLE spectra of ESAON. (f) The HRBE and VRBE schemes for ESAON. Red and blue ‘zigzag’ curves connect the energy for electrons in the 4f ground state of divalent and trivalent lanthanides, respectively. The other red and blue curves connect the energy for the lowest 5d states of divalent and trivalent lanthanides, respectively. Panels a-f: reproduced with permission from ref. 243, 2019 Royal Society of Chemistry.

Eu²⁺-activated M₂SiO₄ (M = Ca, Sr, Ba)

Although nitride and oxynitride phosphors, taking advantage of their excellent luminescent properties, low thermal luminescence quenching, and high chemical stability due to their high covalencies, are the most qualified candidates for LED applications, their synthesis with refractory Si₃N₄/AlN and unstable alkaline-earth nitrides as raw materials, generally requiring high temperature (>1500 °C) and high pressure (>0.1 Mpa), is rigorous, energy intensive and expensive. By contrast, the production of oxide phosphors is more convenient and energy-efficient as their raw materials are non-refractory and air-stable.³³⁸ Therefore, research studies on orthosilicate phosphors for LED applications have been the hotspot in luminescent materials science. Among them, M₂SiO₄ (M = Ca, Sr, Ba):Eu²⁺ are a classical class of phosphors widely used for LED applications like daily lighting and traffic lights. M₂SiO₄ (M = Ca, Sr, Ba) adopt slightly different structures, depending on the M cations, but all look similar to Ba₂SiO₄ shown in Fig. 22a. Ba compound only exists in an orthorhombic structure with a space group of Pmcn.³³⁹ Two types of cations exist in the crystal structure, and thus the crystal chemical formula could be written as MM′SiO4. The M and M′ sites differ in the coordination number and polyhedral volume. For Sr compound, at 85 °C the displacive phase transition from α′ to β occurs.³⁴⁰ β-Sr₂SiO₄ is monoclinic with space group P2₁/n. α′-Sr₂SiO₄ is isostructural with Ba₂SiO₄ and can be stabilized at room temperature by adding a small amount of Ba.³⁴⁰ The solid solution MM’SiO4 with M, M′ = Ba, Sr presents preferential site occupation, i.e., Sr prefers the site with a small volumetric size while Ba prefers the larger one.¹⁴² The stable Ca₂SiO₄ polymorph at room temperature is the γ type, which has an orthorhombic olivine structure with space group Pbnm. Besides, it also possesses a metastable β phase, which is an isotype with the β-Sr₂SiO₄ phase.³⁴¹

Fig. 22 (a) Crystal structures of Ba₂SiO₄. Reproduced with permission from ref. 5, copyright 2013 Annual Reviews. (b) Relative emission spectra of a white-emitting InGaN-based YAG:Ce LED and GaN-based Sr₂SiO₄:Eu LED under a 20 mA drive current. Reproduced with permission from ref. 342, copyright 2003 American Institute of Physics. (c) Emission spectra (λ_ex = 365 nm) of Ca_2−xEu_xSiO₄ with the Eu²⁺ content (x) equal to 0.2, 0.4, 0.6, and 0.8. Inset: Photographs of these samples upon photo-excitation at 365 nm. Reproduced with permission from ref. 14, copyright 2014 Wiley. (d) Emission spectra collected at room temperature using the optimal excitation wavelength for each sample show broad emission across the green to yellow region of the visible spectrum. Reproduced with permission from ref. 51, copyright 2014 American Chemical Society. (e) Photoluminescence spectra excited at 370 nm of Sr₂SiO₄:Eu²⁺ phosphor at various temperatures. (f) Photoluminescence spectra excited at 370 nm of Ba₂SiO₄:Eu²⁺ phosphor at various temperatures. Panels e and f: reproduced with permission from ref. 343, copyright 2004 Elsevier Ltd.

The first report of phosphors within this system was provided as a brief description by Jenkins and McKeag in 1950,³⁴⁴ and a full description of their luminescence properties was made by Barry³⁴⁵ and Blasse³⁴⁶ in the 1960s. From an efficiency viewpoint, they found that phosphors based on solid solution between Ba₂SiO₄ and Sr₂SiO₄ with the orthorhombic phase have higher efficiencies versus that based upon monoclinic β-Ca₂SiO₄. Then it was in 1990s, since the blue LED chips were discovered successfully, the research on luminescence properties of M₂SiO₄ (M = Ca, Sr, Ba):Eu²⁺ came again under the spotlight. All of them can be efficiently excited with both blue and near-UV light and give out emissions within the visible region, making these materials ideal for LED applications. The yellow-emitting Sr₂SiO₄:Eu²⁺ phosphors fabricated with blue GaN chips (400 nm) exhibit high-quality white light with a better luminous efficiency than that of the InGaN chips (460 nm) coated with YAG:Ce phosphors as shown in Fig. 22b.³⁴² The applications in traffic lights and automotive displays of green-emitting Ba₂SiO₄:Eu²⁺ phosphors combined with 395 nm emitting InGaN chips were also demonstrated by Zhang et al.²⁵² What is intriguing is that the emission of Ca₂SiO₄:Eu²⁺ varies from green-yellow to deep red upon different Eu²⁺ concentrations (Fig. 22c), which was suggested to be resulting from the different Eu²⁺ substitutions on two types of Ca sites in α_L′-Ca_2−xEu_xSiO₄ (0.2 ≤ x ≤ 0.8) including 10-coordinated Ca₁ sites and 8-coordianted Ca₂ sites.¹⁴ Following this work, Sato et al. further studied the effect of Eu²⁺ concentrations on the photoluminescence properties of α′-Sr₂SiO₄ and α′-Ba₂SiO₄.³⁴⁷ In the case of Sr_2−xEu_xSiO₄ the emission wavelength was shifted from 585 nm to 611 nm with increasing Eu²⁺ concentration (x) from 0.1 to o.8, and a similar trend occurs in Ba_2−xEu_xSiO₄ in which the emission was shifted from 513–545 nm. Also the low energy edges of the excitation spectra for both samples were significantly shifted to a longer wavelength by 40–50 nm. The induction of such large redshifts in the emission and excitation spectra of both systems was attributed to the increasing Eu²⁺ substitution in Sr2 and Ba2 sites, the coordinated polyhedra of which are smaller and more distorted than that of Sr1 and Ba1 sites.

Besides, cation or anion substitution also plays an important role in the modification of the luminescence properties in the M₂SiO₄ (M = Ca, Sr, Ba):Eu²⁺ system. As an example, Denault et al. studied the luminescence properties of Sr_xBa_2−xSiO₄:Eu²⁺ and explored the structure–property correlations.⁵¹ With the increasing Sr content, the excitation spectra exhibit a slight red-shift while the emission spectra show a large red-shift of about 65 nm (Fig. 22d), which were originating from the decreased polyhedral volume and increased distortion as Sr content increases. Surprisingly, the intermediate composition with 46% Sr has the highest thermal luminescence stability, remaining stable up to 413 K, and maintaining an emission efficiency of 75%. Structural studies using high resolution X-ray diffraction and total scattering show that the intermediate compositions possess optimal bonding that creates a more rigid crystal structure compared to the end-member compositions. The increased structural rigidity of Sr_0.9Ba_1.1SiO₄ is confirmed by its highest Debye temperature determined by the low temperature heat capacity. The high Debye temperature limits access to phonon modes which contribute to luminescence thermal quenching. This work provides an impetus for searching novel efficient phosphor hosts that possess highly connected structures and optimized bonding interactions. Furthermore, He et al. found that the composition dependence of quenching temperatures coincides with that of Sr/Ba atomic ordering degree in the host lattice.¹⁴² Therefore, the improved thermal luminescence quenching resistance of the intermediate compositions in (Ba_1−xSr_x)₂SiO₄:Eu²⁺ phosphors originates from the ordered cationic distribution, which could be represented by the degree of ordering η = Occ_Sr2 − Occ_Sr1 = Occ_Ba1 − Occ_Ba2. The volumetric difference of the two cationic sites in MM’SiO₄ suggests that BaSrSiO₄ could be thought as the end-member compound since the size-mismatch between Ba and Sr ions is annihilated. From this point of view, Ba₂SiO₄ and Sr₂SiO₄ have the intermediate compositions of which the size-mismatch will degrade structural rigidity and deteriorate the thermal luminescence quenching resistance. In addition, the temperature-dependent emission spectra of M₂SiO₄ (M = Ca, Sr, Ba):Eu²⁺ were investigated by Kim et al.³⁴³ It was demonstrated that two emission bands of Sr₂SiO₄:Eu²⁺ show normal redshift with increasing bandwidth and decreasing intensity as temperature increases (Fig. 22e), which can be explained by the interaction between the electron and thermally active phonon. Different from Sr₂SiO₄:Eu²⁺, the emission bands of M₂SiO₄ (M = Ca, Ba):Eu²⁺ show an anomalous blueshift (Fig. 22f), which can be described in terms of back tunneling from the excited states of low-energy emission band to that of high-energy emission band with the assistance of thermally active phonon.

The optimal bonding has also been reported in the T phase of the (Ba,Ca)₂SiO₄ system. Fukuda et al. found that the disordered crystal structure contains many split crystallographic sites.³⁴⁸ He et al. carefully evaluated the coordination environments of all the cation sites of the representative T-phase (Ba_0.65Ca_0.35)₂SiO₄, to understand the origin of the atom site splitting phenomenon, as shown in Fig. 23a and b.²⁴⁴ By analyzing the cation site coordination environments by split- and unsplit-atom-site models, it is found that the three cation sites in the split-atom-site structure are optimally bonded with the ligand O atoms. The coordination environments of the five cation sites together with the corresponding average bond lengths are shown in Fig. 23c. Due to the split of the O2 site, the unsplit M5 site has the average bond length and coordination number modified from 3.063–2.815 Å and 12–8, respectively. Meanwhile, the M1 and M2 sites are split for optimal bonding due to the co-occupation of Ba²⁺/Ca²⁺ ions, and two sets of coordination environments are formed for larger Ba²⁺ ions and smaller Ca²⁺ ions, leading to a more rigid structure and better thermal luminescence stability compared to end-member Ba₂SiO₄ and Ca₂SiO₄ crystals.

Fig. 23 Crystal structures of the T-phase (Ba,Ca)₂SiO₄ for (a) the spilt-atom-site and (b) unsplit-atom-site models. (c) Average bond lengths and the coordination numbers of the different Ca²⁺/Ba²⁺ sites for the split-atom-site model (⊗) and unsplit-atom-site model (♦) for T-phase (Ba,Ca)SiO₄, respectively. Panels (a–c): reproduced with permission from ref. 244, copyright 2018 Royal Society of Chemistry.

M₂SiO₄ (M = Ca, Sr, Ba):Eu²⁺ with Ln³⁺ co-dopant also exhibit PersL performance. In 2008, Lakshminarasimhan et al. first reported that Sr₂SiO₄:Eu²⁺,Dy³⁺ shows yellow emitting PersL for about 5 min after 313 nm irradiation as displayed in Fig. 24a.³⁴⁹ After that, Dutczak et al. further studied the effect of the excitation energy on the PersL properties of Sr₂SiO₄:Eu²⁺,Dy³⁺, and found that two times stronger PersL can be achieved after supra-bandgap vacuum ultra violet (VUV) excitation than after sub-bandgap near-UV or visible excitation, which indicates that the direct creation of free charge carriers under supra-bandgap excitation is followed by trapping and slow release of charge carriers resulting in a strong yellow PersL.²⁶⁷ Then in 2014, Zhang et al. doped Tm³⁺ into Sr₂SiO₄:Eu²⁺ to introduce deep traps for optical storage applications.²⁶⁹ As shown in Fig. 24b, Sr₂SiO₄:Eu²⁺,Tm³⁺ shows green photostimulated luminescence (PSL) upon infrared stimulation at 980 nm after being pre-irradiated by UV light. In 2018, the outstanding PersL performance of Ba₂SiO₄:Eu²⁺,Ho³⁺ was demonstrated by Wang et al.²⁵³ It shows supper-long green PersL which remains above 2.97 mcd m⁻² even 24 h after the removal of UV irradiation as shown in Fig. 24c. What is more, green luminescence by 980 nm stimulation is still observable even 15 days after the UV irradiation (Fig. 24d). It is proposed that the outstanding PersL and PSL performance is mainly derived from the tailoring of trap distribution via Ho³⁺ co-doping (Fig. 24e). On the basis of HBRE and VRBE schemes of Ba₂SiO₄, Er³⁺ with a suitable energy gap between the CB bottom and the 4f ground state, 0.64 eV was selected as an efficient co-dopant, and 7 h PersL duration was achieved for Ba₂SiO₄:Eu²⁺,Er³⁺.²⁷⁰ All of these studies show the great potential of the M₂SiO₄ (M = Ca, Sr, Ba):Eu²⁺ family in PersL applications. With the advantage of bandgap engineering and trap depth engineering based on HRBE and VRBE schemes, further research studies toward improved PersL in solid solution or with multi-doping are expected to give a bright future of M₂SiO₄ (M = Ca, Sr, Ba):Eu²⁺ in the practical applications of emergency signage and optical information storage.

Fig. 24 (a) Decay curves of Eu²⁺ emission in Sr_1.98SiO₄:0.01Eu²⁺,0.01RE³⁺ [RE = Ce, Nd, Sm and Dy]; λ_ex = 313 nm. The inset shows the fitting of the decay curve of Sr_1.98SiO₄:0.01Eu²⁺,0.01Dy³⁺. Reproduced with permission from ref. 349, copyright 2007 Elsevier Ltd. (b) The PSL emission spectrum of β-Sr_1.997−xSiO₄:0.003Eu²⁺,xTm³⁺ (x = 0–0.004). The photos present the corresponding brightness green PSL of serial samples under 980 nm stimulation after preirradiation with UV excitation. Reproduced with permission from ref. 269, copyright 2014 American Ceramic Society. (c) PersL decay curves of Ba₂SiO₄:0.005Eu²⁺,yHo³⁺ (y = 0.002–0.03). (d) The write-in using UV (254 nm) irradiation and read-out (980 nm laser) of the rabbit pattern on the flexible phosphor film. (i) Phosphor film and photomask (rabbit pattern). (ii) Phosphor film is covered by photomask using UV 254 nm irradiation to write in the patterns. (iii) The phosphor film was kept at RT in the dark for 15 days, after removing the UV light and photomasks. (iv) The phosphor film was stimulated by a 980 nm laser to read out the stored patterns. (e) Schematic illustration of the PersL and PSL mechanisms of the Ba₂SiO₄:Eu²⁺,Ho³⁺ phosphor. Panels (c–e): reproduced with permission from ref. 253, copyright 2018 Royal Society of Chemistry.

Ce³⁺/Eu²⁺-activated M₃SiO₅ (M = Sr, Ba)

Within the MO–SiO₂ (M = Ca, Sr, Ba) phase diagram, another compound with M₃SiO₅ stoichiometry, exhibiting efficient yellow-orange emission under near-UV to blue excitation when doped with either Eu²⁺ or Ce³⁺, has been successfully commercialized for applications in warm w-LEDs. The symmetry of these materials is different for each M cation, generally upgrading in the symmetry with the size of the alkaline earth cation. For the Ca, Sr, Ba variants, the crystal structures are monoclinic with a space group of Cm, tetragonal with a space group of P4/ncc, and tetragonal with a space group of I4/mcm, respectively.^350,351 The structure of Sr₃SiO₅ is displayed in Fig. 25a as an example. It has two sites of strontium ions in the unit cell, i.e. 8f and 4c sites. The latter one has a larger polyhedral volume. Its crystal chemical formula could be thus written as (Sr)₂[Sr]SiO₅. Blasse et al. first reported these compounds as phosphors in 1968, doping them with Eu²⁺.³⁴⁶ These researchers showed that these phosphors can be excited efficiently by near-UV or blue light, and give out a broad emission peaking at approximately 500–600 nm (Fig. 25b), which can be systematically tuned toward the lower-energy part of the spectrum by varying the M cation from smaller to larger. In 2004, Park et al. demonstrated that InGaN (460 nm)-based Sr₃SiO₅:Eu²⁺ exhibits a better luminous efficiency than InGaN (460 nm)-based YAG:Ce, but a lower color rendering index (CRI, 64) due to the lack of red component.²⁵⁵ Then, in order to improve the quality of Sr₃SiO₅:Eu²⁺-based w-LEDs, the emission of Sr₃SiO₅:Eu²⁺ was shifted to a longer wavelength by adding the co-dopant Ba²⁺.³⁵² Through integration of the InGaN (460 nm) with two phosphor blends (Sr₂SiO₄:Eu²⁺ yellow phosphors and a Ba²⁺ co-doped Sr₃SiO₅:Eu²⁺ yellow-orange phosphor) into a single package, CRI of the as-made w-LEDs was improved to over 85. It is worth noting that the red shift leading by the larger cation does not conform with the typical positive relationship between size of cation occupied by activator ions and emission energy. And this abnormal red shift may result from the different structures of Sr₃SiO₅ and Ba₃SiO₅, and the different occupancy modes of Eu²⁺ in two M²⁺ sites of M₃SiO₅ (M = Sr, Ba). The temperature-dependent PL intensity in Sr_3−xEu_xSiO₅ (x = 0.05) is shown in Fig. 25c. Different with the normal case like in YAG:Ce phosphors that the emission intensity continuously drops as the temperature increases since the possibility of non-radiative process increases with thermal assistance, in Sr_3−xEu_xSiO₅, the emission intensity slightly increases at 50 °C and then decreases with increasing temperature.²⁵⁶ This abnormal emission intensity increase under the elevated temperature originates from the traps in the host, which can capture the electrons and then release them at elevated temperatures. These traps also contribute to the PersL performance of Sr₃SiO₅:Eu²⁺, which will be discussed below. This phenomenon indicates that the traps with an appropriate trap depth not only contribute to the PersL, but also could give rise to improved thermal luminescence stability.

Fig. 25 (a) Crystal structure of Sr₃SiO₅. Reproduced with permission from ref. 3, copyright 2016 Elsevier Ltd. (b) Excitation and emission spectra of Sr_3−xEu_xSiO₅ (x = 0.05) phosphors annealed at 1600 °C. (c) Temperature dependence of the emission intensities of commercial YAG:Ce³⁺ phosphor and Sr_3−xEu_xSiO₅ (x = 0.05) phosphor synthesized via the Pechini process. (Annealing temperature = 1600 °C, λ_ex = 460 nm). Panels (b and c) reproduced with permission from ref. 256, copyright 2011 The electrochemical Society. (d) Room temperature EPR spectra of Sr₃SiO₅:Eu²⁺ (red and blue lines) and Sr₃SiO₅:Eu²⁺,Nb (green and black lines). Magnetic field from 3420 to 3650 Gauss before and after 100 s UV radiation. The sharp signal at about 3530 G corresponds to oxygen vacancies. (e) Calculated defect formation energies for oxygen vacancies in Sr₃SiO₅ with incorporation of a most probable 2Nb_Si–V_Sr configurations (as shown pictorially in the inset). (f) A schematic representation of the yellow-orange PersL mechanisms of (Sr_1−xBa_x)₃SiO₅:Eu²⁺,Nb⁵⁺. (g) Photographic images of ‘‘Pisces’’ patterns for different commercialized SrAl₂O₄:Eu²⁺,Dy³⁺ persistent phosphor and Sr₃SiO₅:Eu²⁺,Nb phosphor ratios (denoted by numbers 1–5) under natural light (NL), 254 nm UV excitation, and 5 s after shutting down the UV light (PersL), persistent emission spectra of ‘‘Pisces’’ patterns, and CIE coordinates of ‘‘Pisces’’ patterns. (h) Warm-light PersL images of (Sr, Ba)₃SiO₅:Eu²⁺,Nb⁵⁺ obtained at different PersL decay times after 254 nm UV-lamp excitation for 5 min (upper image) and after 254 nm + 365 nm UV-lamp excitation for 5 min (lower image), and TL glow curves of Sr₃SiO₅:Eu²⁺,Nb⁵⁺, Y₂O₂S:Eu³⁺,Mg²⁺,Ti⁴⁺, (Sr_0.88Ba_0.12)₃SiO₅:Eu²⁺,Nb⁵⁺ and (Ca,Sr)S:Eu²⁺. Panels (d, e, and g): reproduced with permission from ref. 272, copyright 2020 Royal Society of Chemistry. Panels (f and h): reproduced with permission from ref. 274, copyright 2021 Royal Society of Chemistry.

Apart from Eu²⁺ activator, the green-yellow phosphors for LED applications can be created when the M₃SiO₅ host is doped with Ce³⁺. As an example, Sr₃SiO₅:Ce³⁺,Li⁺ phosphor exhibits a very broad emission peaking at ∼540 nm under near-UV or blue light excitation.²⁶⁰ Thereinto, Li⁺ was added for charge compensation since Sr²⁺ was replaced by aliovalent Ce³⁺. The concentration quenching happens at Ce³⁺ levels of ∼0.75%, which results from the energy migration over multiple activation ions and then impurities or defects in phosphor can be reached out. Besides single phosphor LEDs, Sr₃SiO₅:Ce³⁺,Li⁺ phosphors were also used in combination with red-emitting CdSe quantum dots to give white light with a high CRI. Unfortunately, Ce³⁺-activated M₃SiO₅ phosphors suffer from severe thermal luminescence quenching at T > 100 °C. At 170 °C, Sr₃SiO₅:Ce³⁺,Li⁺ loses about 50% of its room temperature intensity, which to a large extent limits its applications. Hopefully, by exerting bandgap engineering based on VRBE schemes, the thermal luminescence stability of M₃SiO₅:Ce³⁺ could be improved by manipulating the energy gap between the 5d energy level and the CB bottom by cation or anion substitution, giving them a bright prospect for LED applications.

The yellow-orange PersL and photostimulated luminescence (PSL) of Sr₃SiO₅:Eu²⁺ were first discovered by Sun et al. in 2009, and they found that the introduction of Dy³⁺ can enhance PSL and prolong the PersL duration to 6 h after UV irradiation.³⁵³ Subsequently, the trap depth distribution and retrapping process in Sr₃SiO₅:Eu²⁺,Dy³⁺ were further studied. By analyzing the result of time response in PSL, it is revealed that both shallow and deep traps can retrap electrons during the PersL and PSL processes. Then Dong et al. reported that yellow-orange PersL of Sr₃SiO₅:Eu²⁺,Dy³⁺ can be further improved by adding alkaline earth metal fluorides MF₂ (M = Ca, Sr, Ba).³⁵⁴ In particular, the PersL of Sr_2.95SiO₅,0.05BaF₂:Eu²⁺,Dy³⁺ can last for 12 h. In 2013, Li et al. systemically studied the effect of rare earth codopant on PersL of Sr₃SiO₅:Eu²⁺, and found that, after 365 nm irradiation, Sr₃SiO₅:Eu²⁺,Dy³⁺ has the longest PersL duration while Sr₃SiO₅:Eu²⁺,Nd³⁺ exhibits the highest initial PersL intensity.³⁵⁵ Recently, Wang et al. further investigated the mechanism of PersL enhancement by Nb-incorporation. By virtue of electron paramagnetic resonance (EPR) and density functional theory (DFT) (Fig. 25d and e), it is revealed that Nb-incorporation reduces the formation energy of oxygen vacancies and thereby increases the density of oxygen vacancies, which act as electron trapping centers with appropriate trap depths (Fig. 25f), giving rise to the enhanced PersL performance of Sr₃SiO₅:Eu²⁺,Nb³⁺. It is noteworthy that Sr₃SiO₅:Eu²⁺,Nb³⁺ can also be activated by sunlight (Fig. 25g), which indicates its potential in specific applications, such as emergency signage and night lighting. These researchers also reported that PersL of Sr₃SiO₅:Eu²⁺ can also be enhanced through Ge incorporation with a similar mechanism as Nd incorporation, by inducing increased oxygen vacancies.²⁷² Furthermore, Wang et al. reported that the PersL color of Sr₃SiO₅:Eu²⁺,Nb³⁺ can be finely tuned by Ba substitution. It can be tuned from 580–600 nm as x varies from 0 to 0.75 in Sr_3−xBa_xSiO₅:Eu²⁺,Nb³⁺ system, realizing warm-color PersL with x = 0.12. Surprisingly, the PersL duration is significantly prolonged to 22 h in Sr_3−xBa_xSiO₅:Eu²⁺,Nb³⁺ (x = 0.15) (Fig. 25h).²⁷⁴ The high PersL brightness, superlong PersL duration, and warm color PersL emission make Sr_3−xBa_xSiO₅:Eu²⁺,Nb³⁺ a potential candidate for applications in energy, environmental, and optical information storage fields.

Challenge and outlook

Although substantial progress has been achieved in LED and persistent phosphors, challenges remain with respect to fundamental questions within the science and technology of luminescent materials including establishing the coupling model of multiple structural effect, understanding the mechanism of quantum efficiency, and unveiling the mystery of trap centers. It is also tricky to be able to exert the progress in fundamental scientific understanding on realizing and optimizing the practical applications. Therefore, collaborative efforts from multidisciplinary researchers are required in order to understand the physical properties and enhance the performance in practical devices, so that Ce³⁺/Eu²⁺-activated phosphors progress to higher levels. For this to happen, we envision that the following pillars need to be strengthened in the future:

(1) Developing blue light excitable cyan and deep red-near infrared (NIR) emitting phosphors

To fulfill the requirements of practical applications in security monitoring, food safety, and solar-like LED health lighting, developing blue light excitable cyan and deep red-near infrared (NIR) emitting phosphors is currently under the spotlight. In order to accelerate the phosphor development process, significant progress will be required in understanding the coupling mechanism of multiple structural effects. Based on luminescence theory, by virtue of first principles calculations, establishing the multiple effects of crystal field splitting, centroid shift and Stokes shift is going to be helpful to effectively tune and even predict the luminescence properties, paving way to the aimed phosphors and practical applications.

(2) Improving thermal luminescence stability and quantum efficiency

Non-radiative process suppresses the luminescence efficiency of both LED phosphor and persistent phosphor. Host with rigid and ordered framework can efficiently restrain the possibility of luminescence quenching due to lattice-phonon coupling, contributing to higher luminescence efficiency, better thermal luminescence stability, and narrow band emission. Exploration of rigid structures will open the door for development of phosphors with advanced performance.

(3) Exploring sunlight excitable high performance persistent phosphors

Although several persistent phosphors have been successfully commercialized for practical applications, the number of persistent phosphors with sufficient duration and intensity is still limited, especially the sunlight excitable ones. To make the most use of solar energy and realize the ultimate “dream” of night lighting, exploring sunlight excitable high performance persistent phosphors is an emergent task. Since PersL properties including duration and intensity are mainly determined by the characteristic of the trap center, such as trap depth and density, exploration and functionalization of dopants and structural defects as efficient trap center and investigation on functional mechanism of multiple trap centers will be helpful to accelerate the development of sunlight excitable high performance persistent phosphors.

(4) Developing high performance red-emitting persistent phosphors

Relative to blue and green emitting persistent phosphors, efficient red-emitting persistent phosphors are still deficient. Based on electronic structure engineering, designing high performance red-emitting persistent phosphors is of vital importance for developing white PersL and broadening the application fields of persistent phosphors.

Summary

Over the past few decades, great progress has been achieved in the discovery and optimization of LED and persistent phosphors. Nowadays, Ce³⁺/Eu²⁺-activated phosphors have been developed as an indispensable component for various applications including daily lighting, LED backlights, emergency signage, and information storage. In this review, we summarize the state of art in Ce³⁺/Eu²⁺-activated LEDs and persistent phosphors from the perspectives of both theory and application, highlighting the design principles for new high-performance phosphors. We first introduced Dorenbos theory on f–d transitions in detail, and elaborated some of the latest developments on the effect of the local structure on f–d transitions, to give an overall vision on the composition–structure–property correlations in Ce³⁺/Eu²⁺-activated LED phosphors. We also demonstrated the electronic structure diagram as a powerful design tool towards promising properties. Then the luminescence mechanism for persistent phosphor is illustrated, addressing on the critical role of the band gap structure and trap characteristics on PersL properties. As templates, the development and optimization routines of some typical phosphors were expounded, with emphasis on design principles and how the chemical characteristic affects the luminescence properties, aiming at providing inspirations for tailoring and optimizing the properties of Ce³⁺/Eu²⁺-activated phosphors toward specific applications. Taking advantage of electronic structure engineering, employing structure manipulation of the host via doping control and cationic/anionic substitution, the luminescence properties of LED phosphors can be effectively tuned towards desired emission color, improved luminescence efficiency, and better thermal luminescence stability. As for persistent phosphors, led by band gap engineering and trap depth engineering, the intensity and PersL duration can be significantly enhanced by controlling defect concentration, intentionally introducing aliovalent or isovalent co-dopants as trap centers. Despite the prosperity of this well-established research filed, there are clear challenges ahead. We thus provide a look toward promising research directions, and offer suggestions for future development.

We anticipate that the luminescence theory elaborated and the strategy exemplified in this review will provide a huge boost in the advanced performance of Ce³⁺/Eu²⁺-activated phosphors for practical applications and, also, provide inspirations to drive the development of other functional materials.

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

This work was supported by the National Key Research and Development Program of China (No. 2022YFB3503800) and the National Natural Science Foundation of China (No. 51832005).

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