Self-trapped exciton emission in inorganic copper(I) metal halides

Abstract

The broad emission and high photoluminescence quantum yield of self-trapped exciton (STE) radiative recombination emitters make them an ideal solution for single-substrate, white, solid-state lighting sources. Unlike impurities and defects in semiconductors, the formation of STEs requires a lattice distortion, along with strong electron–phonon coupling, in low electron-dimensional materials. The photoluminescence of inorganic copper(I) metal halides with low electron-dimensionality has been found to be the result of STEs. These materials were of significant interest because of their lead-free, all-inorganic structures, and high luminous efficiencies. In this paper, we summarize the luminescence characteristics of zero- and one-dimensional inorganic copper(I) metal halides with STEs to provide an overview of future research opportunities.

Graphical abstract

Keywords

self-trapped exciton (STE) / low electron-dimensional / inorganic copper(I) metal halides

1 Introduction

Halide perovskites have been recognized as a promising semiconductor material for next-generation optoelectronics, benefiting from high photoluminescence quantum yield (PLQY), tunable emission, and good charge mobility [1,2]. Photoelectric devices based on these materials, such as solar cells, light emitting devices, photodetectors, and lasers, have made unprecedented progress [2]. However, the long-term stability and toxicity of lead are major challenges to their commercialization [2]. To solve these problems, the strategy of replacing toxic lead with non-toxic ions (such as tin, germanium, antimony, and bismuth) has been extensively developed to maintain the advantageous properties of lead perovskites while improving the stability and reducing the toxicity [3–7]. However, some disadvantages still persist. Tin(II)-based perovskites have poor chemical stabilities, and Ge- or Bi-based metal halides have unsatisfactory photoelectric properties, which hamper their further optoelectronic applications [3,6,7].

Luo et al. reported a lead-free double perovskite, Cs₂Ag(Na)In(Bi)Cl₆, which is an important breakthrough [8]. This perovskite provides highly efficient and stable white light emission from self-trapped excitons (STEs). STEs are usually formed in semiconductors with low electron-dimensions and soft lattices [8]. Low electron-dimensionality is conducive to the carriers’ localization, and the soft lattice can elastically deform under photoexcitation, which leads to strong electron–phonon coupling [9]. STEs have a large exciton binding energy [9], which is one of their remarkable characteristics. Ordinary exciton luminescence has slightly lower energy than the bandgap, but as an unconventional exciton, STEs have much smaller emission energies than the bandgap [10], because the emission energy of STEs is equal to the difference between bandgap energy and the exciton binding energy, self-trapped energy, and lattice deformation energy, which leads to the large Stokes shift of the STEs emission energy [10].

STEs are found widely in halide crystals, condensed rare gases, and organic molecular crystals [11]. Since the discovery of STE emission in inorganic copper(I)-based metal halides (Cu(I)MHs), with Cu(I)-X (X= Cl, Br, I) polyhedra as the basic unit, further research of the STE emission from these materials has been rapidly developed [12–17]. After Jun et al. first synthesized single crystals of Cs₃Cu₂I₅ [12], the high PLQY (̴90%) and air stability of low electron-dimensional Cu(I)MHs immediately attracted great attention. Subsequently, reports regarding the regulation of the luminescence properties of such low electron-dimensional compounds by ion substitution have emerged. Through anion mixing, the emission of zero-dimensional (0D) Cs₃Cu₂X₅ (X=Cl, Br, I) can be modulated to green light (in Cs₃Cu₂Cl₅) from blue light (in Cs₃Cu₂I₅), whereas the emission of the one-dimensional (1D) CsCu₂X₃ (X= Cl, Br, I) analog can be tuned from green to yellow [15,18,19]. However, the emission and PLQY of Cu(I)MHs show non-monotonic changes. The emission of 1D A₂CuX₃ (A= Rb, K; X= Br, Cl) is mainly concentrated in the ultraviolet region with a small tuning range [13,16,20]. The compound Cs₅Cu₃Cl₆I₂, another 1D Cu(I)MH, produces pure blue light [21]. By changing the anions and cations in the material, the bandgap, exciton binding energy, self-trapped energy, and lattice deformation energy are changed, resulting in the modulation of the emission color, Stokes shift, full width at half maximum (FWHM), luminescence lifetime, and other properties. In addition to color modulation, unusual luminescence phenomena including pressure-induced luminescence enhancement, polarization-dependent PL emission, and multi-photon absorption have also been observed [22–24].

In this paper, we have summarized the current progress of inorganic Cu(I)MHs with STEs emission. The basic characteristics of STEs and their emission in these materials are described. The photophysical properties of STEs are explored from the perspective of the materials’ crystal structures, spectral properties, and electronic structures. We introduce the applications of these materials in optoelectronic devices. Consequently, we discuss the challenges faced in the future use of all-inorganic Cu(I)MHs.

2 STEs in 0D inorganic Cu(I)MH materials

Jun et al. reported the novel lead-free 0D luminescent material Cs₃Cu₂I₅, which has a PLQY of nearly 90% for its single crystal and 60% for its thin film [12]. They attributed the emission to an excited-state structural reorganization mechanism caused by the Jahn-Teller distortion of tetrahedral Cu [12]. The inorganic and lead-free nature and efficient emission characteristics of Cs₃Cu₂I₅ attracted great attention. As a typical 0D inorganic Cu(I)MH, Cs₃Cu₂X₅ (X= Cl, Br, I) crystallizes in the orthorhombic structure with a Pnma space group. In this structure, a trigonal planar CuX₃ shares an edge with a tetrahedral CuX₄ unit to constitute a [Cu₂X₅]³⁻ cluster. The [Cu₂X₅]³⁻ dimers are completely separated by Cs⁺ cations, as shown in Fig. 1(a).

Sebastia-Luna et al. suggested that Cu(I)MHs can form a continuous solid solution by exchanging the halogen from Cl⁻ to Br⁻ and Br⁻ to I⁻ [18]. Their XRD patterns are fitted using a Pnma space group. The unit cell parameters increase with increasing anion radius (Cl⁻<Br⁻<I⁻) (Fig. 1(b)) [18]. Correspondingly, the emission shifts to higher energy overall; however, the shift in maximum excitation and emission wavelengths are not monotonic when increasing the halide ionic radius (from Cl⁻ to I⁻) as shown in Fig. 1(c), which is in agreement with the PL characteristics of Cs₃Cu₂Br_5−nI_n (0≤n≤5) reported in Refs. [15,18]. In general, the semiconductor bandgap decreases and the inter-band emission shifts to lower energy with increasing halogen anion radius. The emission shift trend here is different from that of inter-band radiation recombination, but analogous to that of CuX (X= Cl, Br, I) where the emission results from strong exciton-phonon coupling [15]. On the contrary, Cs₃Cu₂X₅ shows very large Stokes shifts, as shown in Table 1. Typically, PL spectra with such large Stokes shifts are attributed to STEs as a result of strong exciton-phonon coupling [15]. To verify this inference, the excitation power dependences of the PL for Cs₃Cu₂X₅ (X= Cl, Br, I) were measured [15,26,27]. As shown in Figs. 1(e)–1(g), the linear relationship between the emission intensity and the excitation power indicates the absence of permanent defects; hence, the emission likely originates from STEs [26]. There are some other effective experimental techniques to detect STE, such as characterizing the coupling strength of excitons and phonons, photoconversion spectroscopy [30]. Additionally, the broad excited-state transient absorption plateau across the probe region is a direct evidence of the formation of STEs [27]. Sebastia-Luna et al. suggested that the emission shift with the change of components may be caused by either halide segregation or phase segregation [18]. The XRD pattern of Cs₃Cu₂Cl₅ is better fitted using the space group Cmcm. Compared to other Cs₃Cu₂X₅ materials, Cs₃Cu₂Cl₅ has a significantly red-shifted PL (516 nm) because of the different crystal structure. The dual excitation and emission properties of Cs₃Cu₂(Cl_0.75Br_0.25)₅ suggest the coexistence of the Cmcm phase (as Cs₃Cu₂Cl₅) and Pnma phase (as Cs₃Cu₂Br₅) [18].

To obtain further information about the emission mechanism, the temperature-dependent steady and transient state PLs of Cs₃Cu₂Cl₅ were measured over the range 300–20 K [26]. The results showed that the FWHM decreases gradually, the PL intensity increases, and the emission redshifts as the temperature decreases. The lack of free excitons emission with the change in temperature indicated that the self-trapped barrier is low and the photogenerated excitons are easily trapped [26]. The PL lifetimes of Cs₃Cu₂Cl₅ increased significantly from 112.4 µs at 300 K to 877.7 µs at 20 K, while the mono-exponential decay was maintained. The mono-exponentiality of the lifetime indicates that there is no change in the radiative and nonradiative channels [26]. This temperature-dependent PL characterization is essential for the analysis of the relevant mechanism. It is commonly used to calculate the exciton binding energy and the S factor associated with the electro-phonon coupling strength [10]. The formation of STEs usually requires materials with large exciton binding energies and S values. A large S value means that there is strong electro-phonon coupling [10].

Theoretical calculations found that the valence and conduction bands of Cs₃Cu₂X₅ (X= Cl, Br, I) are mainly consist of Cu 3d and Cu 4s states (Figs. 2(a)–2(c)), indicating that they have narrow bandgaps [15,26]. Roccanova et al. suggests that the structural distortion breaks the two Cu-halogen bonds and moves the two Cu atoms closer to each other in the [Cu₂X₅]³⁻ cluster, as shown in Fig. 2(d) [15]. The hole has Cu 3d character and is localized on the two Cu atoms in the [Cu₂X₅]³⁻ cluster, whereas the electron is localized mostly in the bond-center position between the two Cu atoms [15]. The spatial distribution of the electron and hole of the exciton in Cs₃Cu₂I₅ is similar to that in Cs₃Cu₂Br₅ [15]. Lian et al. suggested that the photoexcitation induces a local structural distortion of Cs₃Cu₂X₅ (X= Cl, Br, I), and the [Cu₂X₅]³⁻ cluster is converted into a spindle-like species (two tetrahedrons sharing a face), decreasing the Cu-Cu distance as illustrated in Fig. 2(e) [26]. The distortion enhanced the local symmetry of the spindle-like [Cu₂X₅]³⁻ and decreased the energy barrier for the formation of STEs, which explains the origin of the single emission from the STEs in Cs₃Cu₂X₅ (X= Cl, Br, I) [26]. In this situation, the valence band maximum (VBM) shows stronger charge localization. However, as the anion goes down the halogen group, the increase of anionic p-component mixing with the Cu 3d orbitals leads to weakening of the VBM. This result partially explains the abnormal blue-shift in the emission of Cs₃Cu₂X₅ (X= Cl, Br, I) when the anion changes from Cl⁻ to Br⁻ to I⁻ [26]. The calculations of Zhang et al. reveal that the electron and hole are separated spatially. This spatial separation impedes the recombination of the electron and hole and thus prolongs the PL lifetimes of Cs₃Cu₂X₅ (X= Cl, Br, I) [15,26]. Because of the highly localized electron and hole, the emission of Cs₃Cu₂Cl₅ is insensitive to defects, giving rise to the high PLQY and the mono-exponential PL decays.

3 STE in 1D inorganic Cu(I)MH materials

Three typical crystal structures of 1D Cu(I)MH are shown in Figs. 3(a)–3(f). Specifically, CsCu₂X₃ (X= Cl, Br, I) belongs to the orthorhombic Cmcm space group, where CuX₄ tetrahedra share common edges, forming 1D chains separated by Cs⁺ ions [18]. Furthermore, A₂CuX₃ (A= Rb, K; X= Br, Cl) crystallizes in an orthorhombic structure with a Pnma space group, where CuX₄ tetrahedra share a common corner, forming 1D chains separated by A⁺ ions along the b-axis [20]. The mixed halide compound Cs₅Cu₃Cl₆I₂ belongs to the orthorhombic Cmcm space group, where the alternately connected [CuCl₂I₂]₂ unit and a single [CuCl₂I₂] unit form 1D zigzag chains of [Cu₃Cl₆I₂]_n⁵ⁿ⁻, where only I⁻ ions bridge between these units as shown in Fig. 3(e). Similar to CsCu₂X₃, the 1D [Cu₃Cl₆I₂]_n⁵ⁿ⁻ chains were separated by Cs⁺ ions [21].

In contrast to the blue-shift of the Cs₃Cu₂X₅ (X= Cl, Br, I) emission with increasing anion radius, the emission of the single halide in CsCu₂X₃ (X= Cl, Br, I) shows a continuous red-shift with an increase in the Stokes shift [19]. However, the emission shifts of the mixed halides CsCu₂Cl_1.5Br_1.5 and CsCu₂Br_1.5I_1.5 do not follow a linear trend, and their maximum emission is red-shifted with lower PL efficiency compared to that of the single halide compounds (Table 2) [19]. Exceptionally, mixed Br–Cl compounds, which are rich in Cl⁻, demonstrate minor phase segregation between CsCu₂Br₃ and CsCu₂Cl₃ [19]. For mixed halides, the emission behaviors are similar to those observed for MAPb(IBr)₃ [19]. These facts could be explained by the presence of greater structural distortion in the mixed halide samples, which can affect the PL efficiency as well as the self-trapped depth within the bandgap of the material and result in a red-shift of the PL maximum [19]. Particularly, the PL lifetimes of CsCu₂X₃ (X= Cl, Br, I) increase from 13.8 ns for CsCu₂Cl₃ to 62 ns for CsCu₂I₃, following the trend opposite to the PL lifetimes of Cs₃Cu₂X₅ (X= Cl, Br, I). The PL lifetimes of CsCu₂X₃ (X= Cl, Br, I) are also one to two orders of magnitude shorter than those of Cs₃Cu₂X₅ (X= Cl, Br, I) [19,26]. The temperature-dependent PL shows that the emission of CsCu₂I₃ is continuously blue-shifted with increasing temperature [14]. A similar blue-shift was also observed in almost all CsCu₂X₃ (X= Cl, Br, I) compounds [19]. As mentioned above, such an abnormal blue-shift phenomenon is caused by the high lattice distortion and the electron–phonon renormalization.

To better understand STE emission in non-octahedral units, Li et al. investigated the pressure-induced PL enhancement in 1D Cu(I)MH CsCu₂I₃ [22]. The experimental results show that the slight structural distortion of the CuCl₄ tetrahedra leads to bandgap broadening and STE enhancement under mild compression from 1 atm to 4.1 GPa. Such changes result in the increase of emission energy and intensity for CsCu₂I₃ in phase I. Meanwhile, the slight structural distortion results in a lower detrapping barrier, weaker electron-phonon coupling strength, and lower activation energy [22]. Therefore, the STE emission efficiencies of phase I are relatively low under this mild compression. The emission intensity is significantly enhanced with further continuous compression to 8.0 GPa. However, it is quenched with further compression beyond this limit, because the strong structural distortions lead to a phase transition and greatly increase the electron-phonon coupling, inducing a higher detrapping barrier and larger activation energy [22]. Thus, less STEs can be detrapped from the ST state to the free exciton (FE) state in phase II (Fig. 3(i)). This relatively increases the concentration of STEs, thus enhancing the STE emission [32,33]. The emission quenching phenomenon at high compression is usually attributed to the synergistic effect of deviatoric stress and structural amorphization [34].

The electronic structures of 1D Cu(I)MH materials are similar to the 0D Cu(I)MH mentioned above. The valence band (VB) and conduction band (CB) are predominantly composed of Cu 3d orbital and mixed Cu 4s orbitals. The M⁺ (Cs⁺, Rb⁺, K⁺, etc.) ions are far away from the VB and CB, so their contribution to the energy band can be ignored [13,15,16,20,26]. In addition, the electronic structures of Cu(I)MHs significantly change upon photoexcitation, which is strongly correlated with their optical properties [26]. Du also found, using first principles calculations, that the emission energy of CsCu₂Cl₃, CsCu₂Br₃, and CsCu₂I₃ STEs decreased gradually as the anion goes down the halogen group [35]. The red-shift trend experimentally observed for the emission spectra of the three materials is not caused by the electronegativity of anions, but rather from the different STE properties [35]. Du’s calculation showed that each CsCu₂X₃ (X= Cl, Br, I) has three different STE states. Figure 4(a) shows the partial charge density contours of both the hole and electron wave functions for STE1 and STE2 in CsCu₂Cl₃ and STE3 in CsCu₂I₃. The calculated emission of CsCu₂Br₃ and CsCu₂I₃ originates from the irradiative transfer of the lowest energy level of the STE, whereas CsCu₂Cl₃ emits from the STE’s metastable energy level [35]. These results are in excellent agreement with the experimental results. Based on Du’s calculations, Kentsch et al. suggest that the exciton self-trapped process in CsCu₂I₃ films was directly observed by femtosecond transient absorption [36]. The edge absorption bimodal structure is derived from the 130 meV spin orbital splitting of the Cu d electrons. The formation of the lowest level of the STE state results in the disappearance of band edge absorption. The experiment showed that the time constant of the self-trapped process of the thermal relaxation free exciton is 12 ps and the energy barrier is not lower than 60 meV [36]. Each process is obtained by species-associated spectra synthesis. It is probably the most powerful experimental description of STE at present.

The emission from CsCu₂X₃ (X= Cl, Br, I) has a large FWHM and Stokes shift, similar to the above-mentioned characteristics of the STEs in Cs₃Cu₂X₅ (X= Cl, Br, I), but it exhibits the shorter PL lifetime. It is unusual that although a STE mechanism is observed, the FWHM and Stokes shift of the A₂CuX₃ (A= Rb, K; X= Cl, Br) emission are much smaller than those of CsCu₂X₃ (X= Cl, Br, I). In addition, Rb₂CuCl₃ showed an upconversion emission phenomenon [13]. In this process, photons with lower energy than the bandgap interact with the lattice, causing absorption of phonons in the lattice and subsequent emission of photons with higher energy than before. This mechanism leads to the extraction of heat energy from the lattice to emit the higher energy photon, which cools down the material. The advantages of high PLQY and low trapped state density in these Cu(I)MHs are important parameters for optical cooling [13]. Among all the reported inorganic 1D Cu(I)MHs, Cs₅Cu₃Cl₆I₂ was the only one to emit pure blue emission, with a near-unity quantum yield of ≈95%. The flat valence bands promote the formation of localized holes, leading to STE formation. The photo-deformable nature of the copper halogen polyhedra facilitates highly efficient STEs, leading to efficient radiative recombination [21].

4 Applications of inorganic Cu(I)MH STE materials

Based on the advantages mentioned above, inorganic Cu(I)MH STE materials are very suitable for applications in lighting and displays as a phosphor. Cs₃Cu₂I₅ and CsCu₂I₃ (Fig. 5) can coexist in the synthesis process; this gives the opportunity to obtain tunable white light emission using a simple mechanochemical method to quickly synthesize mixed-phase materials with different proportions [37]. Fang et al. used this method to obtain white luminescent samples by mixing CsCu₂I₃ and Cs₃Cu₂I₅ with different ratios. When the mixture had a mass ratio of 8:3, the sample emitted with a Commission Internationale de L’Eclairage (CIE) chromaticity coordinate of (0.36, 0.36) [37]. In other reports, pure CsCu₂I₃ and Cs₃Cu₂I₅ materials (powders and nanocrystals) were synthesized first and then mixed in suitable proportions to produce white light emitters [25,28]. Furthermore, CsCu₂I₃ single crystal (SC) has great potential in energy-saving white lighting because of its bimodal emission [14,41]. However, ensuring the stability of the light efficiency and color at different temperatures is a key problem to be solved in the lighting and display applications of such materials.

Inorganic Cu(I)MH STE materials can also be directly used as an electroluminescent layer in LED devices. Ma et al. fabricated CsCu₂I₃-based LEDs with an ITO/PEDOT:PSS/poly-TPD/CsCu₂I₃/TPBi/LiF/Al structure [38]. The best device had a turn-on voltage of about 5.0 V, a maximum luminance (L_max) value of 47.5 cd/m² at an applied bias voltage of 9.2 V, and maximum external quantum efficiency (EQE) of 0.17% [38]. Then they constructed a high color-rendering index (CRI= 91.6) and stable white light-emitting diode (WLED) with an ITO/PEDOT:PSS/poly-TPD/PVK/CsCu₂I₃@Cs₃Cu₂I₅/TPBi/LiF/Al structure [40]. Liu et al. also reported a WLED with an ITO/PEDOT:PSS/V-NPB/CsCu₂I₃@Cs₃Cu₂I₅/SPPO13/LiF/Al structure. This WLED had a CIE chromaticity coordinate of (0.327, 0.348) and a CRI of up to 94 [29]. Wang et al. reported deep-blue LEDs using Cs₃Cu₂I₅ nanocrystals as the electroluminescent layer [39]. These blue LEDs had a glass/ITO/p-NiO/Cs₃Cu₂I₅/TPBi/LiF/A structure, a turn-on voltage of about 4.5 V, a L_max value of 262.6 cd/m² at 7.5 V, and an EQE of ~1.12% at an applied bias voltage of 7.5 V [39]. These results indicated that copper(I) halides are promising as an environment-friendly and stable emitter for LEDs and are compatible with the practical applications of the devices.

Photodetection is another important application of these high-efficiency inorganic Cu(I)MH STE materials. Li et al. reported Cs₃Cu₂I₅-based ultraviolet photodetectors with high spectral selectivity [42]. The Cs₃Cu₂I₅/GaN heterojunction device had a narrow spectral response “window” of 300–370 nm, and the response range can be adjusted by changing the thickness of Cs₃Cu₂I₅ film and bias voltage [42]. At zero bias, the responsivity, detection rate, maximum current on/off ratio, and response speeds were 0.28 A/W, 1.4 × 10¹² Jones, 1.2 × 10⁵, and 95/130 µs, respectively [42]. These values are comparable to those of the numerous previously reported lead halide photodetectors [42]. In addition, Zhang et al. reported a Cs₃Cu₂I₅-based deep ultraviolet photodetector with a responsivity of 64.9 mA/W, a specific detectivity of 6.9 × 10¹¹ Jones, a turn-on ratio of 127, and response speeds of 26.2/49.9 ms for the rise/fall time [43]. Li et al. also constructed a polarization-sensitive and flexible ultraviolet photodetector using 1D CsCu₂I₃ nanowires, which had a high photocurrent anisotropy ratio up to 3.16 because of the electrical and optical anisotropy of the asymmetric structure and the external morphological anisotropy [23]. The device had an on/off ratio of 2.6 × 10³, a response speed of 6.94/214 µs, a photoresponsivity of 32.3 A/W, and a specific detectivity of 1.89×10¹² Jones [23]. Similarly, Yang et al. fabricated a deep ultraviolet photodetector based on a CsCu₂I₃ film [31]. The on/off ratio, responsivity, specific detectivity, and EQE of this device were 22, 22.1 mA/W, 1.2 × 10¹¹ Jones, and 10.3%, respectively, under 265 nm illumination with a light density of 0.305 mW/cm² [31]. Furthermore, Fang et al. observed the facet-dependent photoresponse of a CsCu₂I₃ single crystal, whose morphology consists of 010, 110, and 021 crystal planes (Fig. 6) [44]. The lower dark current of the 010 crystal plane results in the on-off ratio being higher than that of 110 crystal plane [44]. This work gives new insights into the 1D electronic structure associated with high anisotropy.

Other possible applications of inorganic Cu(I)MH STE materials have also been explored, such as X-ray scintillators, fluorescent inks, image sensors, and memristors and neuromorphic computing applications [20,43,45–49]. The excellent luminescence performance and lack of self-absorption in RbCu₂Br₃, RbCu₂Cl₃, K₂CuBr₃, and Cs₃Cu₂I₅ led to a good scintillation response to X-ray signals, with a high light yield of 91056, 16600, 23806, and 79279 photons per meV, respectively [20,45–47]. Zhang et al. reported a fluorescent ink developed from a Cs₃Cu₂I₅/polyvinylidene fluoride precursor solution. This work showed potential applications of Cs₃Cu₂I₅ for anti-counterfeiting and encryption fields (Fig. 7) [48]. Zhang et al. explored the image-sensing capability of Cs₃Cu₂I₅ crystalline films, where purple and light purple images can be readily formed under 265 and 365 nm illuminations, respectively [43]. Furthermore, Zeng et al. investigated the application of Cs₃Cu₂I₅ films in the field of memristors and neuromorphic computing [49]. Ag/PMMA/Cs₃Cu₂I₅/ITO structure memristors exhibited a bipolar resistive switching operating voltage lower than ±1 V, a large on/off ratio (up to 10²), stable endurance over 100 cycles, and a retention time longer than 10⁴ s [49].

5 Conclusions and perspective

We have summarized the recent progress made in the field of STE emission from inorganic Cu(I)MHs. Novel 0D and 1D inorganic Cu(I)MHs STE luminescent materials, such as Cs₃Cu₂X₅ (X= Cl, Br, I), CsCu₂X₃ (X= Cl, Br, I), A₂CuX₃ (A= Rb, K; X= Cl, Br), and Cs₅Cu₃Cl₆I₂, have been developed and researched. The basic features of STE luminescence have been summarized, methods for identifying STE luminescence have been proposed, and the application of these materials has been introduced. Evidently, the STE luminescence can be regulated by substituting cations and anions. In the process of ion modulation, the emission and light efficiency usually change non-monotonically, but the essential reason for this is currently unclear. Among the 1D materials, only CsCu₂X₃ has a PL lifetime of the order of nanoseconds, whereas other materials have a PL lifetimes of the order of microseconds. In addition, A₂CuX₃ (A= Rb, K; X= Cl, Br) have incomprehensible small emission FWHMs and Stokes shifts. These issues have not been thoroughly studied. Meanwhile, the formation process of STEs in this kind of material is rarely studied, and the direct experimental observation has not been realized. Therefore, the modulation mechanism has not been completely understood. A profound understanding of STE formation processes and modulation mechanisms is conducive to the accurate judgment and regulation of material properties and the development of similar materials with desired target properties. These studies will be essential for the performance optimization and extend the application of inorganic Cu(I)MH STEs and similar materials.

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