Luminescent ion-doped transparent glass ceramics for mid-infrared light sources [invited]

Jing Ren; Jing Ren; Xiaosong Lu; Changgui Lin; Changgui Lin; R. K. Jain; R. K. Jain

doi:10.1364/OE.395402

1. Overview of transparent glass ceramics

1.1 Background on glass ceramics (GCs) for mid-infrared (MIR) light sources

1.1.1 Basics of GCs and MIR luminescence issues

Glass ceramics (GCs) were first developed in the middle ages as “colored glasses” or “stained glasses” for decorative purposes, notably for the use as stained glass windows in cathedrals all over Europe [1], although it is clear that the initial fabricators of such “stained glasses” were not aware of the microscopic structure of these materials, as first proposed and expounded by Stookey [2,3]. Subsequent optical applications include their use as nonlinear optical materials [4], in which the presence of semiconductor nanocrystallites (in such “glasses”) were discovered via four-wave mixing experiments. As understood in modern times, GCs are essentially multi-composite or hybrid materials with semiconductor, metal, oxide, or other crystallites uniformly dispersed or “embedded” in a “super-cooled liquid” or an essentially “amorphous” solid matrix [1–7], as shown schematically in Fig. 1, which depicts a specific method (namely “internal crystal nucleation and growth”, described in Sect. 1.2.1) of creating such a GC by adding appropriate constituents into the ingredients of the original glass melt, followed by appropriate heat treatment [6–8]. Although it is not clear from the micrograph of Fig. 1(d), the crystallites in such GCs can be made to vary in sizes from a few nm to several hundred µm – and up to single-component materials [7] – with appropriate choices of glass constituents and thermal processing methods.

Fig. 1. Schematic of “internal” transformation of (a) an amorphous glass at a temperature just below its melting point showing a possible arrangement of atomic constituents to (b) a glass ceramic upon thermal processing by method such as “annealing” at a temperature between the glass transition point and melting point; (c) Dark field transmission electron microscopic (TEM) images of (c) an “untreated” glass; and of (d) a glass ceramic, depicting the presence of nanometer-scale crystallites [8].

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GCs can easily be tailored in a way that their key optical properties are determined primarily by the embedded nano/microcrystals while their macroscopic or “bulk” properties, including physical, chemical, and mechanical properties, are determined primarily by the glass matrix. In particular, transparent glass ceramics (TGCs) are created by using high purity constituents and controlling the fabrication processes and parameters to minimize spurious absorption and optical scattering losses [6–11]. Scattering losses in GCs are usually caused by refractive index differences between the glass and crystal phases and by morphological effects [11]. GCs containing isotropic cubic phase crystals tend to exhibit reduced scattering losses because the cubic phase crystals are free of birefringence and thus have the same refractive index in all crystal directions [11], which can be more easily matched to the glass matrix. Scattering losses can also be minimized by reducing pores and grain boundaries, and by making sure that the nanocrystals (NCs) are very small compared to the wavelengths of interest, i.e. typically less than 100 nm for MIR TGCs, or by matching the refractive index of the embedded crystal closely to that of the host glass matrix [12]; the latter case enables achievement of TGCs even with the use of large “micron-sized” microcrystallites (MCs).

Many advanced functional ceramics, mainly those based on silicates were developed around the turn of the century, both to improve the mechanical properties – such as hardness and fracture toughness – of these “glass-like” materials, as well as to develop new optically, electrically, and magnetically active materials based on the nano-/micro-crystallite constituents of these GCs [13]. In contrast with silicate GCs, the glassy matrix of MIR (2.0 - 5.0 µm) GCs are based on heavy atomic element constituents to make low phonon energy (LPE) materials such as tellurite and chalcogenide (ChG) glasses [14], and alkali halides [15]. Functionalization of such LPE “soft glasses” for efficient MIR luminescent materials is then achieved by developing methods for incorporating appropriate trivalent rare earth (RE³⁺, see Fig. 2) or divalent transition metal ions “dopants” (TM²⁺, see Fig. 3) with appropriate MIR emitting transitions that are optically excited by appropriate pump wavelengths [16–20], as depicted in Figs. 2 and 3. For TM²⁺, it is also critical that the ions are located at optimal crystal sites in MIR-transparent LPE crystallites of the desired composition (usually II-VI wide bandgap semiconductors) [16–18]. The challenge in the last step lies in the extreme difficulty of developing the best materials processing techniques that enable formation of stable high quality TGCs such that the desired dopant ions with the correct valence are located at appropriate sites in the NCs or MCs (such as + 2 valence of Cr²⁺ in tetrahedral sites of II-VI ChG NCs or MCs) with relatively high dopant densities (> 10¹⁸/cm³); as such, methods for achieving these goals are a key subject of this review paper.

Fig. 2. (a) Energy level diagrams of common luminescent triply-ionized rare earth ions (RE³⁺). The pump and emission wavelengths are indicated by blue and red arrows, respectively. (b) Typical emission bands for different RE³⁺ [19].

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Fig. 3. (a) Energy level diagrams of several common doubly-ionized transition metal ions (TM²⁺). The pump and emission wavelengths are indicated by blue and red arrows, respectively. (b) Typical emission bands for TM²⁺ in ZnSe crystal [20].

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As stated above, achievement of the desired luminescent LPE crystallite and host matrix-based GCs is a difficult task, with no a priori method of predicting achievement of desired structures and of precise doping densities in the NCs/MCs in the TGCs, and fabrication processes are currently developed empirically and laboriously by multiple cycles of trial and error. As such, there is a strong need for predicting and controlling the design and fabrication processes a priori via theoretical understanding of the crystal nucleation and growth processes [6–7], based largely on molecular dynamics computations [21] and detailed nanoscale heterogeneous structure models of the glasses and the resultant GCs [22,23]. Despite considerable effort expended on these molecular dynamics-based models [21], and the use of advanced machine learning techniques [24,25] to accelerate these extremely complex computations, the theoretical basis for predictive development of processing recipes based on theoretical methods is still very much in its infancy. Significant advances in the development of the best MIR luminescent glasses – including those suitable for MIR lasers based on materials such as Fe²⁺: ZnSe NCs in appropriate glass matrices, which appear to be extremely promising for next-generation MIR fiber lasers [26] – will most likely be developed initially by developing novel glass compositions [18] empirically, as elaborated below.

Figure 4 depicts a schematic illustration of the nature of MIR Luminescent Ion-Doped Nanocrystallite-Embedded Glasses (LIDNEGs), which are the primary focus of this review paper. The blue circles characterize different categories of ion-doped luminescent materials, while the red circles characterize different categories of TGCs, with a general decrease in crystallinity from left to right (Note that we have ignored luminescence from the band edge of semiconductors, such as the interband luminescence from narrow band gap (NBG) semiconductors in this paper to allow focus on the more pervasive research on MIR emission from ion-doped NCs and MCs embedded in TGCs).

Fig. 4. Schematic Venn diagram of MIR Luminescent Ion-Doped Nanocrystallite-Embedded Glasses (LIDNEGs) that are the primary focus of this review article. LPE and HDD stand for low phonon energy and high doping density, respectively. NBG semiconductors and ceramics containing MCs are ignored in this figure for simplicity. Note that the areas designated by the green (A) and orange (B) regions in the Venn diagram represent the optimal target areas for many of the most practical applications.

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As depicted in Fig. 4, the materials of greatest relevance lie in the intersection of the smallest red and blue circles. As such, TGCs in Regions A and B of this Venn diagram – corresponding to the use of materials with a high doping density (HDD) and a high degree of crystallinity in hosts with LPEs – are the most promising materials for MIR LIDNEGs, and likewise for LIDMEGs (based on index-matched MC-embedded TGCs), and are thus the primary materials of focus in this review. Nevertheless, TGC materials outside these regions are also of great interest in the overall discussion, because critical information can be gleaned from detailed understanding of such materials systems, since the overall development of optimized MIR light-emitting TGCs (namely LIDNEGs and LIDMEGs) is still very much in a relative state of infancy.

1.1.2 Key properties of some MIR-relevant TGCs

Tables 1 and 2 summarize the phonon energies and degree of crystallinity for several oxyfluoride, tellurite, and chalcogenide (ChG) material systems (constituent glasses and crystals in Table 1 and TGCs in Table 2) corresponding to Regions A and B, or in close proximity to these regions in the Venn diagram of Fig. 4. Note that the LPE is important not only to minimize multiphonon absorption [27] in these glasses and improve MIR transparency, but also to reduce nonradiative relaxation rates from the emission levels (the upper laser levels for laser transitions) that are inevitably spaced relatively close to the next lower level (usually the lower laser level for laser transitions) [27]. The latter feature is critical for efficient radiative emission, and is absolutely essential to achieve population inversion in MIR laser sources. While the term GCs was originally supposed to describe only materials that contain significant volume fractions of crystallites in a hosting glass matrix (> 50%), GCs of significantly lower crystal volume fractions have also been developed over the last 60 years [28]. In fact, as shown in Table 2, the volume fraction of crystallites (or “crystallinities”) may vary from a few percent to almost 100%. Note that, although crystallinity is an important parameter, many authors do not explicitly measure or state the crystallinities of their TGCs. The quantitative value of crystallinity can be estimated from X-ray diffraction (XRD) measurements, and is given by the ratio of the total areas under the indexed diffraction peaks to the area under the entire XRD plot [29]; alternatively, it can be measured from Transmission Electron Microscopy (TEM) measurements [8]. One drawback of XRD lies in the difficulty of measuring low crystallinities (< 2%).

Table 1. Summary of key properties of several glasses and crystals relevant to MIR applications. n^† at 1.55 µm; n* at 3 µm; n^‡ at 0.593 µm.

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Table 2. Summary of the reported crystallinities of TGCs including the phonon energies of the embedded nanocrystals (NCs). The crystallinity is defined as the volume fraction (vol.%) of the embedded crystals in TGCs [28]. Internal crystal growth (ICG)

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As stated earlier, there has been a large amount of experimental effort and success in the development of MIR LIDNEGs in recent years. High quality TGCs in numerous glass material systems, including oxyfluoride, tellurite and ChG glasses have been successfully demonstrated, and relatively efficient broad MIR emissions in the spectral range of 2–5.5 µm have also been demonstrated from these GCs. The next section (Section 1.2) provides a brief introduction of fabrication methods used for optimized MIR LIDNEGs. Subsequent sections describe the development of specific TGCs used for LIDNEGs as well as key results and their significance, followed by some thoughts on key limitations of results to date, key needs, and anticipated future developments in this field. For the convenience of the reader, a list of key acronyms is given in Table 3.

Table 3. List of key acronyms and their meanings

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1.2 General methods for fabrication of TGCs for MIR LIDNEGs

The various methods of fabrication of GCs relevant to MIR LIDNEGs can be grouped into two broad categories, namely (a) Internal Crystal Growth (ICG) and (b) External Crystallite Embedding (ECE) methods. Only the basics of these methods will be described here, with an emphasis on issues relevant to the fabrication of oxyfluoride, tellurite, and ChG GCs.

1.2.1 Internal crystal growth methods for fabricating LIDNEGs

From a thermodynamic point of view, crystallization (nucleation and crystal growth) of glass is energetically favored under the right processing conditions, resulting in the formation of GCs. For example, when a glass is heated to temperatures between the glass transition temperatures (T_g) and the liquidus points (T_l) of the glass, ICG (also called “natural or thermal crystallite nucleation and growth”) can occur (see Fig. 1 and Fig. 5(a)), resulting in an agglomeration-free homogenous dispersal of NCs with high optical uniformity. The nucleation process is dictated by the fact that the free energy in creating a certain volume of crystal is lower than the free energy of the interface between the liquid and the crystalline phase caused by spontaneous creation of ordering of the atomic constituents of the liquid. For our application, volume crystallization [7] – based on bulk nucleation – is usually preferable because of favorable optical homogeneity. Also, as will be discussed in Section 2 below in which specific experiments on ICG of oxyfluoride and ChG GCs are discussed in detail, Type II GCs [7] – in which network formers do not participate in crystallization – are preferred, because they readily enable the precipitation of LPE crystallites such as CaF₂, LaF₃, NaYF₄, ZnS, and ZnSe, as needed to reduce multi-phonon relaxation (see Fig. 4 above) from the MIR transitions in RE³⁺ or TM²⁺ ions that are naturally embedded in these crystallites during the GC formation process. Although theoretical efforts are being pursued actively, with a focus on accurate prediction of the precise experimental processing steps for ICG of the precise crystallites and dopant concentrations needed for LIDNEGs of specific interest, especially those based on complex multicomponent glasses, state-of-the-art research for the fabrication of such LIDNEGs largely relies on empirical trial-and-error approaches. Direct ICG of metal or semiconductor NCs can also be achieved in glasses in a spatially-selective manner with localized heating with lasers, as has been used [6,7,41] to create different crystalline structures inside such glass ceramics with a high three-dimensional spatial resolution that is not achievable by means of conventional thermal heating-induced crystallization.

Fig. 5. Schematics of (a) internal crystal growth (ICG) and (b) external crystallite embedding (ECE) approaches of making LIDNEGs [5]. (a) The basic ICG method consists of preparing a precursor glass with appropriate compositions by melt-quenching methods (casting) followed by a controlled crystallization process via thermal treatment; (b) One ECE method – as depicted here – is as follows: the precursor glass is pulverized into fine glass powder, which is mixed and homogenized with pre-fabricated NCs; the mixture is heated at temperatures above the softening point of the glass matrix but below the decomposition temperature of the NCs; after a brief heating period, the molten (viscous) mixture is cast into a mold forming the GCs. [NP: nanoparticle; OANP: optically active NP].

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1.2.2 External crystallite embedding methods for fabricating LIDNEGs

As depicted schematically in Fig. 5(b), GCs can also be produced by external crystallite embedding (ECE) methods, also called NC direct doping (NDD) [42], ex situ [5], glass powder doping (GPD) [12] and co-melting [7] methods, in which NCs are synthesized or fabricated separately before incorporating them into the glass matrix; in principle, any type of NC – with an arbitrary shape, morphology, size, and nanostructure – can be incorporated into a baseline glass matrix to make a “synthetic” GC via this method. As such, unlike the ICG method, the various ECE methods allow an unlimited choice of NC/MC-host glass matrix combinations, limited only by: (1) a requirement of avoidance of a strong chemical interaction (or dissolution) of the NC/MCs in the matrix glasses (which dictates a preference for low melting temperatures and high viscosities during the dispersal process), and (2) index matching between the crystallites and the glass matrix to minimize light scattering and improve optical transparency. Note that the ECE method depicted schematically in Fig. 5(b) involves adding the optically active nanoparticles (OANPs) to the precursor glass powder or the powder constituents of the glass melt; other ECE methods, including addition of OANPs directly to the vitrified glass, either in the molten or liquidus phase, or at lower temperatures closer to T_g corresponding to higher viscosities and lower levels of dissolution, may also be used. However, because of the challenges of dissolution and chemical reaction of the pre-fabricated NCs added to the powder mix (or to the molten/semi-molten glass matrix at elevated temperatures), ECE methods have so far been used in fluorophosphates (FPs) and tellurite glass hosts for MIR LIDNEGs with limited success [5,12]. Controlling the viscosity of the glass to minimize dissolution of the NCs is apparently a viable – yet unproven – strategy, but achieving a good trade-off between NC survival and homogeneous dispersal in the glass matrix is a big challenge. As such, much of this review focuses on intrinsic nucleation and crystallite growth, i.e., on ICG methods for making MIR LIDNEGs, since these appear to hold significant promise as the most likely candidate materials for the first demonstrations of LIDNEG-based solid state coherent and incoherent MIR light sources.

2. Mid-IR luminescent oxide TGCs

2.1 Oxyfluoride TGCs

Oxyfluoride TGCs – which belong to the family of Type II GCs [7] – are based on mixtures of fluoride and oxide compounds, and can be referred to as fluoro-silicate/borate/phosphate/germanate/tellurite GCs, depending on the specific glass forming oxides used. As stated earlier, in these TGCs the fluoride crystallites precipitate readily from the glass network-forming oxides, enabling entrapment of the luminescent ions in a LPE crystallite environment. A large variety of oxyfluoride TGCs have been developed which show superior luminescence properties for numerous photonic applications [15]. Our review of oxyfluoride TGCs is distinguished by the type of NCs embedded, namely (a) lead fluoride NCs, (b) alkali earth fluoride NCs, and (c) mixed metal fluoride NCs.

2.1.1. Oxyfluoride TGCs containing PbF₂ NCs

Since the first demonstration of upconversion emission from RE³⁺-doped PbF₂ NC-containing oxyfluoride TGCs based on SiO₂-Al₂O₃-CdF₂-PbF₂-YF₃ system by Wang [43] in 1993, much attention has been paid to investigating the upconversion luminescence properties of PbF₂ NC-containing oxyfluoride TGCs. Incorporation of RE³⁺ in PbF₂ NCs – by substitution of RE³⁺ in the Pb²⁺ sites – has been suggested as a mechanism for enhancing MIR emissions. The first observation of intense 2.7 µm MIR emission of Er³⁺ was reported in 2006 [44] in a fluorosilicate TGC based on a modified version of Wang’s system, by replacing YF₃ with ZnF₂, i.e., with the use of the SiO₂-Al₂O₃-CdF₂-PbF₂-ZnF₂ glass system. It was theoretically proven that laser action at 2.7 µm was feasible in such Er³⁺-doped fluorosilicate TGCs, and as such, increasing attention has been paid to studying MIR luminescent properties of RE³⁺-doped oxyfluoride TGCs, and bright emission originating from the Dy³⁺: ⁶H_13/2 - ⁶H_15/2 transition at 2.87 µm has been demonstrated by Chung et al [45].

2.1.2. Oxyfluoride TGCs containing MF₂ (M = Ca, Sr, Ba) NCs

Increasing concern for environmental and health issues has stimulated research for alternative oxyfluoride TGCs that are free of toxic compounds such as PbF₂ and CdF₂. Bright 2.7 µm emission (from Er³⁺) was observed by Wu et al [46] in fluorosilicate TGCs containing cubic phase CaF₂ of particle sizes less than 20 nm. The formation of CaF₂ NCs in fluorosilicate glasses occurs via a self-limiting crystallization process which controls the overgrowth of NCs and helps improve the optical transparency [47]. ErF₃ was also observed to play a role as a nucleating agent in the SiO₂-Al₂O₃-Na₂O-CaF₂ system, and enable doping of Er³⁺ into the CaF₂ NCs. Cubic phase BaF₂ NCs have also been used to develop oxyfluoride LIDNEGs, and exhibit self-limiting crystallization behavior [48–49] similar to that observed in CaF₂.

2.1.3 Oxyfluoride TGCs containing ALnF₄ (A = Na, K; Ln = Y, La, Lu) NCs

Much attention has also been paid in recent years to RE³⁺ doped fluorosilicate TGCs containing ALnF₄ NCs, because of the very efficient incorporation (more than 90%) of RE³⁺ in ALnF₄ NCs. The RE³⁺ dopants have the same ionic charge and similar ionic radii to those of Ln³⁺, requiring minimal energy expenditure to get these ions incorporated into the ALnF₄ crystal lattice [39]. This is in sharp contrast to RE³⁺ doping into MF₂ NCs which results in significant lattice distortion and cation vacancy formation, e.g., 2RE³⁺ → 3M²⁺ + M vacancy [50]. Several groups have reported that Er³⁺-doped fluorosilicate TGCs containing cubic phase NaYF₄ yield bright 2.7 µm Er³⁺ emissions (Fig. 6) [51–53]. The emission cross section was estimated to be as high as 7.05 ×10⁻²¹ cm², which is slightly higher than that observed in Er³⁺-doped ZBLAN (6.57 × 10⁻²¹ cm²) glass, a “gold standard” for MIR luminescent glasses. Table 4 shows a comparison of the emission cross-sections in various oxyfluoride TGCs. Positive gain is expected when the population inversion value is larger than 0.5 (see Fig. 6) [51]. The presence of RE³⁺ in NaYF₄ NCs has been directly confirmed [54] with advanced TEM analysis. While maintaining good transparency, the NaYF₄ TGCs (Fig. 6) showed a higher degree of crystallinity than the PbF₂ and CaF₂ TGCs, which was qualitatively inferred from a comparison of the XRD intensities of the TGCs [55]. In NaYF₄ – embedded TGCs, because of the metastable nature of NaYF₄ NCs, a crystal phase transition [from cubic (α) to hexagonal (β)] was observed to occur when the glass was heated to high temperatures (650 °C). The α → β phase transition resulted in a dramatic (an order of magnitude) enhancement of the upconversion luminescence of Er³⁺ [55], which is related to different coordination environment of Er³⁺ in these two phases (Fig. 7). Since the electronic transitions of RE³⁺ belong to forced oscillator transitions, the transition probability of RE³⁺ increases when located in crystalline sites that are more asymmetrical [56].

Fig. 6. MIR emission spectra of NaYF₄ TGCs doped with different concentrations of Er³⁺ pumped at 980 nm. The left inset illustrates the transparency and the right inset depicts gain spectra as a function of pumping of such Er³⁺ doped TGCs [51].

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Fig. 7. Schematic presentation of Er³⁺ ions doped in hexagonal (a) and cubic (b) phase NaYF₄ crystals. In (a) hexagonal NaYF₄, an ordered array of F⁻ offers two types of cation sites: one occupied by Na⁺ ions, and the other occupied by Y³⁺ and Na⁺ ions, the Y³⁺ is located in 9-coordinated sites with low symmetry, whereas in (b) cubic phase NaYF₄, Y³⁺ occupies the 8-coordinated cubic symmetry site. Adapted from [60].

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Table 4. Representative mid-IR transparent oxyfluoride TGCs. σ_em: emission cross-section (×10⁻²⁰ cm²); The NC size is in nm unless otherwise specified; Internal crystal growth (ICG); External crystallite embedding (ECE).

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2.1.4. Oxyfluoride TGCs obtained by the external crystallite embedding (ECE) method

By carefully matching the refractive indices between the host glass and the externally embedded crystallites, it is possible to obtain TGCs containing fluoride crystallites that are several micrometers in size (Fig. 8(a)) [12,59]. Er³⁺-doped CaF₂ (with a refractive index 1.434) MCs were embedded in fluorophosphates GCs (with a refractive index 1.439) by Fan et al [12], in one of the first implementations of the ECE method. Specifically, the GCs were obtained by ball-milling Er³⁺: CaF₂ NCs (to ∼ tens of nm), and mixing these NCs into the FP glass powder, and then heating the mixture to a temperature of 820 °C, at which temperature the FP glass was molten and the NCs were homogeneously dispersed in the FP glass melt. Subsequent thermal treatment led to an ICG type of growth of the doped NCs into spherical MCs of ∼10 µm size. Strong 2.7 µm emission was observed in the Er³⁺: CaF₂ GCs, but not in the Er³⁺ doped FP glasses (see Fig. 8(b)), indicating the importance of doping erbium ions into the CaF₂ crystallites, and providing evidence of the importance of the TGCs and the related phonon energy considerations. However, at higher temperatures (1100 °C), the CaF₂ MCs dissolve into the FP glass melt, as evidenced by XRD measurements (Fig. 9) [61].

Fig. 8. (a) Transmittance spectrum of the Er³⁺: CaF₂-FP glass ceramic (2 mm thick) obtained by the external crystallite embedding method. Inset: digital photo of the sample. (b) MIR emission spectra of the Er³⁺ doped FP glass (dashed line) and glass ceramic (solid line) samples excited by an 808 nm laser diode [12].

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Fig. 9. XRD measurements of the Er³⁺: CaF₂-FP glass ceramics obtained by the ECE method. The figure clearly shows the absence of CaF₂ crystallites (due to dissolution) in the glass melt at increased co-melting temperatures (1100 °C) [61].

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2.1.5. MIR luminescent oxyfluoride TGC fibers

It is usually much more challenging to produce TGC fibers by the conventional “rod-in-tube” method, because uncontrolled crystallization can occur during the fiber drawing at the glass softening temperatures (which are higher than the crystallization temperatures of the core glass). Because of the high degree of crystallization and increased light scattering from such crystallites, these fibers exhibit optical attenuation losses that are so high that the resulting opacity prevents quantification of the transmission loss [62–63]. Recently, a so-called “melt-in-tube” method (see Fig. 10), which involves inserting a core “rod” into a tube (Fig. 10(a)) with a higher melting point, and drawing fibers (Fig. 10(b)) at conditions where the core glass was molten while the cladding glass was relatively soft – followed by heat treatment (to enable formation of NCs in the core glass) – was proposed and used to produce TGC fibers containing cubic NaYF₄ NCs [54,64]. As seen in Fig. 10(c), although the GC fibers had higher transmission loss at 1310 nm than that of the precursor glass fibers (11.81 dB/m vs. 7.44 dB/m), they showed improved transparency than fibers made by the rod-in-tube method, and also exhibited strong 2.7 µm emission spectra characteristic of Er³⁺ (Fig. 10(d)) which was absent in the non-heat-treated glass fiber. A theoretical simulation predicted a lasing threshold of 1.12 W and an optical conversion efficiency of 0.27 under 980 nm pumping [54].

Fig. 10. Images of (a) the Er³⁺-doped glass preform and (b) precursor glass fiber fabricated using melt-in-tube technique; (c) 1310 nm loss measurement of the precursor glass and glass ceramic fibers; (d) 2.7 µm emission spectra (980 nm LD pump) as a function of the heat treatment temperature [54].

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2.1.6. Summary of progress in oxyfluoride TGCs

Oxyfluoride TGCs combine the LPE of fluoride crystals and the mechanical robustness and chemical durability of oxide glasses very favorably. GCs with good optical transparency have been obtained by ICG and ECE methods. However, the latter method is limited to FP glasses with low melting temperatures. New methods have been proposed to produce TGC fibers with acceptably low optical loss. So far, MIR emissions in oxyfluoride TGCs have been limited to wavelengths shorter than 3 µm, since the high attenuation of borate-, phosphate- and silicate-based GCs in the MIR wavelength region poses an obvious obstacle in their use for longer wavelength MIR emission. In comparison, fluorogermanate and tellurite TGCs, having lower phonon energies and extended MIR transparency windows (Table 1), appear to be more promising materials for LIDNEGs at longer wavelengths.

2.2 Tellurite (tellurium dioxide) TGCs

2.2.1. Overview

Among oxide glasses, tellurite glasses are characterized by relatively large refractive indices (1.93-2.3), high Raman gain coefficients (∼ 0.86 × 10⁻¹² m/W at 2.8 µm) [65], and LPEs (700-780 cm^-1) [19,66]. Other favorable attributes of tellurite glasses, as compared with non-oxide glasses (e.g., fluoride and ChG glasses), include the ease of large volume production, robustness and good chemical durability. Since the first demonstration of tellurite TGCs in the TeO₂-Nb₂O₅-K₂O system in 1995 [67], tellurite TGCs have found use in several photonic applications, including second harmonic generation [68], solid state lasers [69] and magnetic field sensing [70]. We briefly describe preparation of tellurite TGCs via the ICG and ECE techniques here since there are several significant distinctions between tellurite and other oxide TGCs.

2.2.2. Tellurite TGCs obtained by conventional internal crystal growth

Bright MIR luminescence of RE³⁺ has been observed in tellurite TGCs containing cubic phase PbTe₃O₇ [71] or Li₂TeO₃ [72] NCs with particle sizes in the 8-30 nm range. TEM measurements show that Er³⁺ ions preferentially accumulate in these NCs. As a result: (1) 2.7 µm emissions with a higher emission cross-section (0.80×10⁻²⁰ cm²) than that from Er³⁺-doped ZBLAN, and (2) prolonged emission lifetimes – from 0.2 ms to 1.1 ms – have been observed. The increase in the emission lifetimes are presumably because the RE³⁺ ions are incorporated in a well-ordered crystalline environment and are located far from OH^- impurities (which are known to quench MIR luminescence) [73]. A comparison of the emission cross-sections in various tellurite GCs is given in Table 5. The emission spectra of RE³⁺ located in crystalline environments often exhibit marked spectral splittings and structuring. For example, Stark splitting has been observed in the emission spectra of Ho³⁺-doped TeO₂-Nb₂O₅-K₂O tellurite TGCs, in which an order of magnitude enhancement of the 3 µm emission of Ho³⁺ was also observed [74].

Table 5. Representative mid-IR luminescent tellurite TGCs. σ_em: emission cross-section (×10⁻²⁰ cm²); Size of nanocrystals is in nm. Internal crystal growth (ICG); External crystallite embedding (ECE).

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Very recently, congruent crystallization has been realized in the TeO₂-Nb₂O₅-Bi₂O₃ system [40,67,77,78]. Nearly full crystallization was achieved in 12.5Bi₂O₃-12.5Nb₂O₅-75TeO₂ (in mol.%) with the crystallinity approaching 100% [78]. In such glasses, because of their high polarizability, the heavy cations can easily self-assemble into an ordered arrangement, while the anions are located in a disordered sublattice, resulting in a metastable “anti-glass” phase. The anionic network experiences rearrangement after subsequent thermal treatment, leading to a stable fully crystalline phase [40,78]. The tellurite TGCs thus obtained display excellent optical transparency (Fig. 11(a)) [40], which is attributed to the close match of the refractive index of the NCs to that of the host tellurite glass [67], and contain Bi_0.8Nb_0.8Te_2.4O₈ NCs in a previously unreported cubic phase, with a typical particle size of ∼ 20 nm (Fig. 11(c)). Optical fibers (Fig. 11(b)) were drawn from such tellurite TGCs with preforms made by the rod-in-tube method, and weak but clear broadband MIR emission in the 4.3–4.95 µm spectral range was observed (Fig. 11(d)) from this fiber (the origin of this MIR emission is shown in Fig. 2(a) for Er³⁺); this result represents the first time that MIR emission at wavelengths longer than 4 µm has been reported in oxide-based TGCs.

Fig. 11. (a) Transmission spectra of the TeO₂-Nb₂O₅-Bi₂O₃ glass and glass ceramics. (b) Photo of the glass ceramic fiber transmitting 532 nm light. Scanning electron microscopy images of the glass ceramic (c) and fiber (Inset in (b)). (d) Emission spectrum of Er³⁺-doped GCs pumped at 808 nm. (a) adapted from [77]; (b)-(d) adapted from [40].

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In separate experiments in tellurite TGCs doped with Nd³⁺, 1064 nm continuous-wave (CW) and gain-switched laser operation has been achieved [69] at record slope efficiencies (54%), which imply the possibility of ultralow light scattering losses from such tellurite TGCs, and their utility as MIR LIDNEGs and laser gain media.

2.2.3. Tellurite TGCs obtained by external crystallite embedding method

The ECE method has been employed to incorporate Er³⁺ doped Y₃Al₅O₁₂ (YAG) or NaYF₄ NCs in tellurite glasses because of its favorable viscosity at moderate temperatures [76,79]. The high melting temperature (1940 °C) of the YAG crystal is presumed to be helpful in minimizing dissolution during the melting (at 570 °C) of the glass powder (containing Er: YAG NCs in the mix of melt constituents). 2.7 µm emission (from Er³⁺) has been observed from such Er: YAG tellurite GCs [76]. However, the GCs suffered from significantly reduced transparency (Fig. 12, the thickness of the samples is ∼ 1 mm) because of the presence of air bubbles, and further work is necessary to reduce these air bubbles for future laser development.

Fig. 12. Transmission spectra of tellurite GCs containing different amounts of NCs, which are obtained by ECE method [76].

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2.2.4. Summary of tellurite TGC-based LIDNEGs

Enhanced MIR emissions can be obtained in transparent tellurite TGCs and fibers embedded with a number of cubic-phase NCs. Due to the low melt temperatures of tellurite glasses, tellurite TGCs containing NCs with desired compositions can be made using the ECE method. The very low phonon energies of tellurite glasses also enable MIR emissions at longer wavelengths (> 4 µm). Tellurite TGCs with a crystallinity approaching 100% can be achieved from the TeO₂-Nb₂O₅-Bi₂O₃ system via congruent crystallization [78]. Because of the possibility of incorporation of RE³⁺ ions [69] in the embedded LPE NCs, development of tellurite TGCs with high crystallinities are promising for MIR LIDNEGs. In addition, tellurite TGCs containing metal halide NCs such as MH₂ (M = Pb²⁺, Ba²⁺; H = Cl^-, Br^-, I^-), LnH₃ (Ln = Y³⁺, La³⁺) – and more recently, CsPbH₃ perovskites NCs [80] (with particle sizes of 6-10 nm obtained by the ICG method) – appear to be very promising candidates for MIR LIDNEGs. The low preparation temperatures of tellurite glass make them particularly attractive for making glasses containing highly volatile components such as metal halides [81].

3. Mid-IR luminescent chalcogenide (ChG) TGCs

3.1 Rare-earth doped ChG TGCs

Due to the heavy masses of the constituent atoms, the vibration frequencies of metal-chalcogen (S, Se) bonds are very low, i.e., ChG glasses possess extremely LPEs (typically in the range of 150-450 cm^-1). As stated in Section 1, LPE hosts are critical for MIR optical transitions from dopant ions (RE³⁺, TM²⁺, or other), particularly at transitions corresponding to wavelengths longer than 3.5 µm, where the corresponding luminescence is strongly quenched in oxide and fluoride hosts due to fast multiphonon relaxation [27]. Thus, most observations of MIR luminescence at wavelengths longer than 3.5 µm have been reported primarily in ChG crystals or glasses [82]. As listed in Table 6, the MIR optical transitions of the following RE³⁺: Dy³⁺, Er³⁺, Pr³⁺, Tm³⁺, Tb³⁺, and Ho³⁺ can cover the entire spectral range between 3 and 5 µm (although laser inversion is much more difficult at the longer wavelengths).

Table 6. Representative mid-IR luminescence of rare-earth ions in chalcogenide glasses. σ_em: emission cross-section (×10⁻²⁰ cm²)

View Table | View all tables in this article

ChG TGCs are Type II GCs, in which the II-VI crystallites segregate from the network formers via precipitation of appropriate crystalline phases that contain the luminescent RE³⁺. ChG TGCs have been widely studied because of their promising properties of increased mechanical strength, permanent IR frequency-doubling, and enhanced luminescence – with appropriate RE³⁺ – in the MIR region [14]. Enhanced luminescence of RE³⁺ in ChG TGCs was firstly observed in Nd³⁺ doped 70GeS₂·8Ga₂S₃·12Sb₂S₃·10CsCl nanocrystallized samples [88]. In addition, with the precipitation of Ga₂S₃ NCs, a large enhancement (∼20×) of the upconversion luminescence was achieved in Er³⁺ doped 70GeS₂·20Ga₂S₃·10CsCl TGCs [89], resulting in several studies of the enhanced luminescence of RE³⁺ in ChG TGCs during the past decade [90–93].

In 2011, Dai et al [94] reported the first observation of enhanced 3.8 µm luminescence from the ³H₅→³F₄ transition of Tm³⁺ ions in 65GeS₂·25Ga₂S₃·10CsI ChG TGCs doped with 0.6 wt% Tm³⁺ ions; the presence of ∼50 nm Ga₂S₃ NCs was shown to cause a two-fold increase in MIR emission; this enhancement was further increased to more than 5-fold in well-crystallized 80GeS₂·20Ga₂S₃ TGCs [95–96]. In the GC sample with a similar composition of Ge_28.125Ga_6.25S_62.625, emissions at 2.9 µm and 3.5 µm originating from appropriate optical transitions of Dy³⁺ ions (see Fig. 2), were also enhanced by ∼12 times because of the precipitation of Ga₂S₃ NCs [97].

As seen in the above–described reports, controlled crystallization of ChG TGCs not only enhances the MIR emission, but also the upconversion and NIR emissions from these RE³⁺ because the LPEs in these crystal hosts inhibit non-radiative relaxation from the excited levels that are the source of the MIR emission. Balda et al reported narrowing and structuring of upconversion emission of Er³⁺ ions as a clear indicator of the fact that the RE³⁺ ions were incorporated into the Ga₂S₃ NCs [89], and did not remain in the surrounding glass matrix. However, in other publications, broad “amorphous glass-like” spectral emission profiles of the enhanced emission bands indicated that the RE³⁺ was largely located in the residual glass matrix [94–97]. Because of the contradictory nature of these different observations, the fundamental cause of enhancement of emissions from RE³⁺ in ChG GCs is still unclear. Different physical mechanisms may be responsible for the observed emission enhancements, including optical “multi-reflection” effects between the precipitated NCs, compositional variations in the residual glass matrices during crystallization, a variation of the local chemical environment around the RE³⁺ ions, and bonding of the RE³⁺ ions to NC surfaces [90–97].

In recent years, Ge-Ga-S TGCs have been investigated quite thoroughly to study and clarify the fundamental mechanisms leading to enhanced MIR emissions [95–97]. It is worth mentioning that the Ga-based ChG glasses were selected specifically because Ga not only favors increasing the doping concentration of RE³⁺ (the limited doping concentration of RE³⁺ in ChG glasses has always been a headache) [98], but also provides essential structural units related to control of the crystallization of Ga₂S₃ NCs [14]. In particular, in Ge-Ga-S glasses, the Ga-related units (i.e., [GaS₄] tetrahedral and ethane-like [S₃Ga-GaS₃]) play a vital role – due to their structural similarity – for the formation of isochemical Ga₂S₃ nuclei, which grow subsequently under further thermal treatments. Microstructural studies [95–97] have revealed that nano-polycrystalline structure – composed of interconnected crystallites sharing grain boundaries or a narrow amorphous coalescence neck – formed in glass systems in which the glass network former (Ga-related structural units) participates in the crystallization process, similar to the crystallization mechanism described in 80GeS₂·20In₂S₃ TGCs [99].

The nano-polycrystalline structure makes the analysis of the distribution of RE³⁺ ions very complicated, leading to uncertainty in the origins of the observed emission enhancements. In addition, the low doping concentrations (< 0.5 mol%) used in previous studies of Dy³⁺/Tm³⁺ - doped ChG GCs embedded with Ga₂S₃ NCs made it difficult to get a clear picture of the distributions of the RE³⁺ [95,97]. Thus, in order to reveal how Tm³⁺ ions disperse in Ga₂S₃ TGCs, a heavily Tm³⁺ doped (2.0 mol%) sample was prepared for conclusive elemental mapping by scanning TEM (STEM) equipped with an energy dispersive spectrometer (EDS), as shown in Fig. 13 [96]. It is seen that the NCs isolated by the Ge-rich glassy matrix (Fig. 13(a)) belong to the Ga₂S₃ crystalline phase, which can be clearly discerned in the Ga-rich areas in Fig. 13(b). Tm³⁺ ions are dispersed homogeneously among the glassy matrix (Fig. 13(c)), and seem to be complementary to the distribution of Ga. The line scanning measurement (Fig. 13(d)) further confirmed that Tm³⁺ ions disperse exclusively from the Ga₂S₃ NCs, in line with the Ge distribution. This is different from Wang’s demonstration that RE³⁺ (Dy³⁺) were strongly bonded to the surface of Ga₂S₃ NCs [97]. Based on the studies to date, it can be concluded that Tm³⁺ ions are expelled out of the precipitated Ga₂S₃ NCs and reside primarily in the glassy matrix during the onset of crystallization, resulting in their being dispersed in a smaller net volume (that of the residual glass matrix), and thus to a decrease in the ion-to-ion spacing. The enhanced cross-relaxation between the Tm³⁺ ions (due to the closer interionic spacing), along with increased multiple scattering effects, is believed to be the reason for the observed enhanced MIR emissions in such ChG TGCs.

Fig. 13. High-Angle-Annular-Dark-Field (HAADF) scanning TEM (STEM) image (a) of 2.0 mol% Tm³⁺ doped 80GeS₂·20Ga₂S₃ glass-ceramics with the precipitation of Ga₂S₃ nanocrystals, and corresponding STEM-EDS elemental mappings of (b) Ga and (c) Tm. (d) Line scanning analysis for Ge, Ga, S, and Tm elements [96].

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3.2 Transition metal ion-doped ChG TGCs

Divalent transition metal ions (TM²⁺, such as Cr²⁺, Co²⁺, Fe²⁺) doped II-VI ChG (such as ZnS, ZnSe) semiconductors have emerged as excellent MIR gain media, due to their broad MIR emission bands (∼ 0.5λ₀, where λ₀ is the emission peak wavelength), large stimulated emission cross-sections (∼ 10⁻¹⁸ cm²), high quantum efficiencies (> 70%) and the ready “pumpability” of the upper laser levels by commercially available cost-effective laser sources (including diode lasers, and diode-pumped Er³⁺ and Tm³⁺ doped high power fiber lasers [100]. As such, TM²⁺: II-VI crystals now appear to rival the Ti³⁺-doped sapphire crystals (Ti³⁺: Al₂O₃) used for NIR solid state lasers as the dominant broadband MIR solid state gain materials. However, adverse thermal lensing effects due to their relatively large thermo-optic coefficient (six times that of Ti³⁺: Al₂O₃) and the uneven distribution of TM²⁺ ions (due to surface diffusion doping methods commonly used) hinder the generation of laser beams of excellent quality and beam-pointing stability. A viable solution is to make glass-based waveguides (either in planar or fiber form) embedded with TM²⁺: II-VI crystals, which have a large surface area-to-volume ratio, and thus good heat-dissipation capabilities [101].

As₂S₃ GC fibers embedded with Cr²⁺: ZnSe and Cr²⁺: ZnS crystallites [102] were first demonstrated in 2010. The fiber preforms were based on GCs fabricated by the ECE (glass powder doping) method, and contained crystallite of about 1 µm in size [103]. These fibers emitted broadband luminescence in the 1.8-3 µm (centered at 1.9 µm) range, but as seen in Fig. 14, were still far from satisfactory due to significant scattering losses (2- 4 dB/m). Hot pressing based ICG methods have also been used [104] to prepare II-VI crystal/ChG glass composites, and emission in the 1.8 - 2.8 µm spectral range has been observed in Cr²⁺:ZnSe/As₄₀S₅₇Se₃ glass composites. However, due to the presence of strong scattering, these materials are still very lossy, with optical transmissions of less than 40% (Fig. 15(a)) in the 2-10 µm wavelength region in a sample that is only 2 mm thick. Nevertheless, strong interparticle scattering caused by MCs has been used favorably to demonstrate room temperature (RT) random lasing at 2.4 µm in Cr²⁺: ZnSe/As₂S₃: As₂Se₃ micro-composites [105]. Emission in the 1.9 - 3.1 µm emission range has also been reported in Cr²⁺: ZnSe/95%As₂S₃:5%As₂Se₃ composites prepared by a uniaxial hot pressing method [106]. Reduced scattering loss TGCs, obtained by matching the refractive indices between the embedded II-VI crystals and the ChG glasses, are expected to result in much more efficient ChG based MIR LIDNEGs, as needed for fiber lasers based on this material system.

Fig. 14. Optical losses in pure glass As₂S₃ (dotted line) and glass ceramic fibers containing 0.1 wt% of Cr²⁺:ZnS MCs [102].

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Fig. 15. (a) Transmission spectra of ZnSe: Cr²⁺/As₄₀S₅₇Se₃ ChG glass composites prepared by the hot-pressing method at different synthesis temperatures [104]; (b) Transmission spectra of TM²⁺ (Cr²⁺, Co²⁺, Co²⁺/Fe²⁺) doped ChG GCs containing ZnS or ZnSe NCs obtained by in-situ crystallization (sample thickness is 2 mm) [16–18].

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Lu et al reported the first synthesis of ChG TGCs containing ZnS NCs by the ICG method in the As₂S₃-ZnSe glass system [36]. The GCs showed higher than 60% transmission in the MIR wavelength region (Fig. 15(b)). However, limited by the poor solubility of the ZnSe compound in As₂S₃, the crystallinity of the GCs was only about 1%. In a subsequent study whose results are depicted in Fig. 16, Lu et al showed that ChG glasses based on Ge-As-S system with a three-dimensional glass structure could incorporate much greater amounts of ZnSe (up to 15 mol.%) [16], and as such, the crystallinity of the GCs was increased to about 4%; the larger crystallinity enabled accommodation of higher doping densities of TM²⁺ ions, which – along with higher dopant densities per crystallite – leading to higher MIR luminescence efficiencies. An intense broadband MIR (1.8 - 2.8 µm) emission of Cr²⁺ in such ChG GCs (Fig. 16(e)) was demonstrated by these researchers [16]. Elemental mapping (Fig. 16(c)) suggested that most of Cr²⁺ ions (> 90%) were doped in ZnS NCs, most likely by substitution for the tetrahedrally-coordinated Zn²⁺ sites (Fig. 16(f)).

Fig. 16. (a) Dark field transmission electron microscope (TEM), and (b) high angle annular dark field scanning TEM (HAADF-STEM) images of Cr²⁺ doped ChG glass ceramic and its corresponding elemental mappings of (c) Zn, (d) Cr. (e) MIR emission spectra of the Cr³⁺ doped ChG glass ceramic under the 1570 nm excitation and some selected RE³⁺. (f) Crystal structure of Cr²⁺-doped ZnS crystal, and electronic states responsible for the MIR emission [16].

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In contrast to RE³⁺ ions, the electronic transitions of TM²⁺ ions are extremely sensitive to local crystal-field environments [107]. Recently, Lu et al [17] reported tunable broadband MIR (2.5 - 4.5 µm) emissions of Co²⁺ in ChG GCs containing a variety of II-VI NCs (such as ZnS, ZnSe, ZnSSe and ZnCdS). Cost-effective and commercially available Er³⁺-doped fiber lasers can be used as the excitation source for Co²⁺. By crystal-field engineering of the embedded NCs through cation- (Zn²⁺ ↔ Cd²⁺) or anion-substitution (S^2- ↔ Se^2-) (Fig. 17(a)), the emission properties of Co²⁺ including its emission peak wavelength and bandwidth can be varied over a broad spectral range (∼ 700 nm) (Fig. 17(b)) [17].

Fig. 17. (a) Scheme of crystal field engineering of TM²⁺ (e.g., Co²⁺) emission via cation or anion substitution. CB: conduction band; VB: valence band. The indicated energy (in eV) correlates with the emission peak wavelength of Co²⁺ in the corresponding crystals. (b) Emission spectra of Co²⁺-doped ChG GCs containing different II-VI NCs [17].

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For Fe²⁺-doped crystals, the desired pump wavelengths (around 3 µm) are often difficult to obtain, and necessitate the use of Er³⁺ or Cr²⁺ doped crystalline (e.g., YSGG: Er³⁺, Cr²⁺: ZnS) solid state, or Er³⁺-doped ZBLAN fiber pump lasers [108]. The large spectral overlap of 2.5 - 4.5 µm emission band of Co²⁺ and the 2 - 5 µm absorption band of Fe²⁺, suggests the use of Co²⁺ as a sensitizer for Fe²⁺. Based on the efficient Co²⁺ → Fe² energy transfer (Fig. 18(a)), Lu et al [18] demonstrated broadband MIR (2.5 - 5.5 µm) emission in Co²⁺-Fe²⁺ codoped ChG GCs pumped by a commercially available erbium doped fiber amplifier emitting at 1.57 µm (Fig. 18(b)). Lu et al [18] also described the first observation of a unique “anomalous” increase in the MIR luminescence intensity as a function of temperature (Fig. 18(d)). In addition, these researchers also demonstrated [18] using these ChG LIDNEGs for gas sensing of multiple target analytes (butane and carbon dioxide).

Fig. 18. (a) Schematic of energy transfer between Co²⁺ and Fe²⁺. (b) Room temperature MIR emission spectra of Co²⁺-singly doped (multiply by 1/5), Fe²⁺-singly doped, and Co²⁺/Fe²⁺ codoped ChG GCs under the 1570 nm excitation. Temperature dependence of MIR emission spectra of the Co²⁺-singly doped (c), and (d) the Co²⁺/Fe²⁺ codoped samples [18].

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In summary, ChG TGCs and fibers embedded with TM²⁺ doped II-VI crystals can be obtained by ICG and ECE methods. Hot pressing has also been suggested to make ChG GCs, which has the potential of mitigating glass dissolution of NCs (also referred to as “melt corrosion” by some researchers). Because of the large range of compositions accessible to practical ChG glasses, it is possible to match the refractive indices between the II-VI crystals and the ChG glasses. As such, GCs and fibers with low scattering losses can be expected even for particle sizes that are of micron scales [102]. Because of their sensitivity to the crystal field environment, MIR emissions of TM²⁺ can be tuned over a very broad spectral range by cation- (Zn²⁺ ↔ Cd²⁺) or anion-substitution (S^2- ↔ Se^2-) [17]. Broad MIR (1.9 - 5.5 µm) emissions of TM²⁺ have been observed in TM²⁺: II-VI/ChG TGCs and fibers at room temperature, offering a platform of gas sensing for multiple target analytes. However, only random lasing has been realized in ChG LIDNEGs so far [105]. Future studies should focus on exploring new ChG TGCs with a much larger crystallinity – and ultrahigh dopant densities – for enhanced MIR emissions of TM²⁺, as needed for efficient lasing in such materials.

4. Summary and anticipated future directions

TGCs containing functional NCs/MCs can meet requirements that are impossible to fulfill by their glass and crystal counterparts. This is particularly true for MIR luminescence, which is generally significantly quenched in high phonon energy glasses. The crystalline nature of the sites is critical for TM²⁺ ions, because of the sensitivity of these ions to the local field effects and the need for sites of specific coordination. For photonic applications, GCs with low optical attenuations – i.e., high optical transparencies – are critical. This need for high optical transparency requires that: 1) the particle size of the embedded crystallites be kept as small as possible, ideally of the order of 200 nm or smaller, i.e., much smaller than the MIR wavelength of interest, and significantly smaller than the pump or excitation wavelength; or 2) the refractive index of the embedded crystal be very well matched (ideally with a difference of less than 0.001) to that of the surrounding glass matrix. However, the difficulty in molecular dynamics computational methods and the lack of fundamental knowledge to predict which types of crystals can be grown in GCs by internal nucleation and crystallization precludes the possibility of designing GCs materials based on first principles at present. On the other hand, because the refractive index is generally well known for the NC or MC that is to be synthesized before its incorporation into the host glass, it is possible to match the refractive indices by tuning the glass composition using ECE methods.

Although extensive studies have clearly demonstrated that enhanced MIR emissions of RE³⁺ and TM²⁺ can be achieved in various TGCs, and theoretical simulations point to their feasibility for coherent MIR emissions, MIR lasing in ion-doped TGCs has been rather limited so far. This is due in part to the following issues: 1) poor material quality and significant inhomogeneity in multi-component TGCs, and high scattering losses in current TGC-based fibers, 2) the presence of substantial amount of impurities such as [OH^-], [S-H^-] or [Se-H^-] groups [82], 3) surface defects in NCs due to their large surface-to-volume ratio, which may act as electron trap centers [109], and 4) limitations in doping concentrations of ICG-doped NCs [18]. With regard to the latter, note that doping densities of the order of 10²⁰ ions/cm³ are highly desirable in TGCs in Regions A and B of Fig. 4, especially if the crystallinities are of the order of 10 percent or less, as may be critical if these TGCs are to be drawn into advanced MIR laser fibers.

Nevertheless, both incoherent and coherent, discrete and broadband MIR light sources based on TGCs are anticipated to have significant impact on future photonic applications. Promising target areas for research include:

(1) Efficient MIR emissions at a wavelength longer than 3 µm. These are anticipated to occur with improved tellurite TGCs, ChG TGCs, and in fluorogermanate TGCs (which are much more mechanically robust than tellurites and ChG TGCs). In order to realize longer emission wavelengths, it will be advantageous to develop new TGCs (such as those based on new tellurites) that are embedded with (a) RE³⁺-doped ultralow phonon energy crystallites, such as metal halide NCs (e.g., CsPbCl₃/Br₃/I₃, or LaCl₃) whose phonon energies (< 200 cm^-1) are much smaller than those of most oxyfluoride NCs, or with (b) heavily TM²⁺-doped ZnSe or other LPE II-VI crystallites.
(2) TGCs with a high degree of crystallinity containing crystallites with high ionic doping densities, which will be beneficial to enhancement of MIR emission for bulk MIR emitters or bulk material based solid state lasers. For instance, ChG GCs containing TM²⁺-doped II-VI crystals are hindered by very low crystal volume density (less than 4%), and an increase in crystallinity – and ionic doping densities – will be critical for enhanced MIR emissions. Research on finding new tellurite TGCs – with a potential crystallinity approaching 100% – by congruent crystallization, also appears very promising, as long as these are accompanied by high ionic doping densities; however for applications involving shaping of the glasses by casting, growing thin films, and drawing into fibers, a suitable value of between 10 to 25% may be preferred for reduced scattering losses.
(3) Improved theoretical understanding of underlying molecular dynamics leading to improved materials design, including an in-depth understanding of the glass topological structures, which can be achieved with the aid of theoretical simulation using first principles molecular dynamics or machine learning [21–25,110].
(4) Development of advanced waveguide and microstructured geometries. Most of the studies and advances to date in LIDNEGs have involved bulk materials. In the future, more advances are expected in making TGC waveguides, either in fiber or planar form, for example, by using femtosecond laser writing in bulk TGCs [111–112], which should lead to promising applications in integrated optics, remote real-time gas sensing and non-invasive medical diagnosis. Fabricating microresonators, such as microspherical gain media based on TGCs [75] appears to have large potential for many applications, such as unique comb generators and extremely sensitive sensors because of their potentially ultrasmall footprints and accessibility to on-chip integration. Laser action has been demonstrated recently at NIR wavelengths in singlemode TGC fibers [113] and microspheres [75,114]. Extending microspherical lasing capabilities to the MIR spectral region at new wavelengths beyond the 2.7 µm lasers previously demonstrated in RE³⁺doped fluorozirconate glass microspheres [115–117] – based on new RE³⁺- or TM²⁺- doped TGCs – should also unleash the potential of TGCs for ultracompact photonics device applications, particularly for ultrasensitive molecular sensing.

Funding

National Natural Science Foundation of China (51872055, 61775110); Natural Science Foundation of Heilongjiang Province (F2017006); Fundamental Research Funds for the Central Universities; Guangdong Provincial Key Laboratory of Fiber Laser Materials and Applied Techniques; Harbin Engineering University (111 project (B13015)).

Acknowledgments

We also thank Dr. Zhigang Gao for assistance with several aspects of preparation of this manuscript.

Disclosures

The authors declare no conflicts of interest.

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Materials	Phonon energy (cm^-1)	Refractive index (n)	Transparency range (µm)	Ref.
Fluorosilicate Glass	∼1100	1.75^†	0.2-2.5	[30]
Fluorogermanate Glass	∼900	1.73-1.85^†	0.38-5.0	[31]
TeO₂-based Glass	∼700	1.9-2.3	0.4-5.0	[19]
GeSe₂-based Glass	∼250	2.4-2.8	0.7-20	[32]
CaF₂ Crystal	328	1.42*	0.2-7.0	[33]
LaF₃ Crystal	300	1.48*	0.2-10.5	[34]
NaYF₄ Crystal	360	1.47^‡	0.2-8.0	[35]
ZnS Crystal	350	2.26*	0.4-14	[20]
ZnSe Crystal	250	2.44*	0.5-20	[20]

Dopants	Glassy host	Wavelength (µm)	σ_em	Nanocrystals (NC)	Size (nm)	method	Ref.
Er³⁺	SiO₂-Al₂O₃-CdF₂-PbF₂-ZnF₂	∼2.7	∼0.4	Cubic PbF₂	∼20	ICG	[44]
Dy³⁺	SiO₂-Al₂O₃-CdF₂-PbF₂-ZnF₂	2.87	-	Cubic PbF₂	∼30	ICG	[45]
Er³⁺	SiO₂-Al₂O₃-Na₂O-CaF₂	∼2.7	-	Cubic CaF₂	<20	ICG	[46]
Er³⁺	GeO₂-Ga₂O₃-NaF-BaF₂-LaF₃	∼2.7	-	Cubic BaF₂	∼7.5	ICG	[57]
Er³⁺	SiO₂-Al₂O₃-Na₂O-YF₃-NaF	∼2.7	0.71	Cubic NaYF₄	∼13.2	ICG	[51]
Er³⁺	SiO₂-Al₂O₃-Na₂O-LuF₃-NaF	∼2.7	0.52	Orthorhombic LuOF	16-28	ICG	[52]
Er³⁺	SiO₂-B₂O₃-Na₂O-LaF₃	∼2.7	0.72	Hexagonal LaOF	25-42	ICG	[58]
Er³⁺	MgF₂-CaF₂-SrF₂-BaF₂-NaF-AlF₃-NaPO₃-Al(PO₃)₃	∼2.7	-	Cubic CaF₂	10 µm	ECE	[12]
Er³⁺	MgF₂-CaF₂-SrF₂-BaF₂-NaF-AlF₃-NaPO₃-Al(PO₃)₃	∼2.7	1.2	Cubic SrF₂	15 µm	ECE	[59]

Active ions	Glassy host	Wavelength (µm)	Transition	σ_em	Ref.
Er³⁺	Ge-Ga-Sb-S	4.6	⁴I_9/2→⁴I_11/2	0.285	[83]
Dy³⁺	Ga-La-S	4.3	⁶H_11/2→⁶H_13/2	1.17	[84]
Pr³⁺	Ba-In-Ga-Se	3.4, 4.5	¹G₄→³F₄, ³H₆→³H₅	-	[85]
Tm³⁺	Ga-La-S	3.88	³H₅→³F₄	0.27	[86]
Tb³⁺	Ga-La-S	3.0, 4.8	⁷F₄→⁷F₆, ⁷F₅→⁷F₆	0.89	[86]
Ho³⁺	Ge-Ga-S-CsI	3.86	⁵I₅→⁵I₆	-	[87]

Materials	Phonon energy (cm^-1)	Refractive index (n)	Transparency range (µm)	Ref.
Fluorosilicate Glass	∼1100	1.75^†	0.2-2.5	[30]
Fluorogermanate Glass	∼900	1.73-1.85^†	0.38-5.0	[31]
TeO₂-based Glass	∼700	1.9-2.3	0.4-5.0	[19]
GeSe₂-based Glass	∼250	2.4-2.8	0.7-20	[32]
CaF₂ Crystal	328	1.42*	0.2-7.0	[33]
LaF₃ Crystal	300	1.48*	0.2-10.5	[34]
NaYF₄ Crystal	360	1.47^‡	0.2-8.0	[35]
ZnS Crystal	350	2.26*	0.4-14	[20]
ZnSe Crystal	250	2.44*	0.5-20	[20]

Dopants	Glassy host	Wavelength (µm)	σ_em	Nanocrystals (NC)	Size (nm)	method	Ref.
Er³⁺	SiO₂-Al₂O₃-CdF₂-PbF₂-ZnF₂	∼2.7	∼0.4	Cubic PbF₂	∼20	ICG	[44]
Dy³⁺	SiO₂-Al₂O₃-CdF₂-PbF₂-ZnF₂	2.87	-	Cubic PbF₂	∼30	ICG	[45]
Er³⁺	SiO₂-Al₂O₃-Na₂O-CaF₂	∼2.7	-	Cubic CaF₂	<20	ICG	[46]
Er³⁺	GeO₂-Ga₂O₃-NaF-BaF₂-LaF₃	∼2.7	-	Cubic BaF₂	∼7.5	ICG	[57]
Er³⁺	SiO₂-Al₂O₃-Na₂O-YF₃-NaF	∼2.7	0.71	Cubic NaYF₄	∼13.2	ICG	[51]
Er³⁺	SiO₂-Al₂O₃-Na₂O-LuF₃-NaF	∼2.7	0.52	Orthorhombic LuOF	16-28	ICG	[52]
Er³⁺	SiO₂-B₂O₃-Na₂O-LaF₃	∼2.7	0.72	Hexagonal LaOF	25-42	ICG	[58]
Er³⁺	MgF₂-CaF₂-SrF₂-BaF₂-NaF-AlF₃-NaPO₃-Al(PO₃)₃	∼2.7	-	Cubic CaF₂	10 µm	ECE	[12]
Er³⁺	MgF₂-CaF₂-SrF₂-BaF₂-NaF-AlF₃-NaPO₃-Al(PO₃)₃	∼2.7	1.2	Cubic SrF₂	15 µm	ECE	[59]

Luminescent ion-doped transparent glass ceramics for mid-infrared light sources [invited]

Abstract

1. Overview of transparent glass ceramics

1.1 Background on glass ceramics (GCs) for mid-infrared (MIR) light sources

1.1.1 Basics of GCs and MIR luminescence issues

1.1.2 Key properties of some MIR-relevant TGCs

1.2 General methods for fabrication of TGCs for MIR LIDNEGs

1.2.1 Internal crystal growth methods for fabricating LIDNEGs

1.2.2 External crystallite embedding methods for fabricating LIDNEGs

2. Mid-IR luminescent oxide TGCs

2.1 Oxyfluoride TGCs

2.1.1. Oxyfluoride TGCs containing PbF₂ NCs

2.1.2. Oxyfluoride TGCs containing MF₂ (M = Ca, Sr, Ba) NCs

2.1.3 Oxyfluoride TGCs containing ALnF₄ (A = Na, K; Ln = Y, La, Lu) NCs

2.1.4. Oxyfluoride TGCs obtained by the external crystallite embedding (ECE) method

2.1.5. MIR luminescent oxyfluoride TGC fibers

2.1.6. Summary of progress in oxyfluoride TGCs

2.2 Tellurite (tellurium dioxide) TGCs

2.2.1. Overview

2.2.2. Tellurite TGCs obtained by conventional internal crystal growth

2.2.3. Tellurite TGCs obtained by external crystallite embedding method

2.2.4. Summary of tellurite TGC-based LIDNEGs

3. Mid-IR luminescent chalcogenide (ChG) TGCs

3.1 Rare-earth doped ChG TGCs

3.2 Transition metal ion-doped ChG TGCs

4. Summary and anticipated future directions

Funding

Acknowledgments

Disclosures

References

Cited By

Figures (18)

Tables (6)

Optics Express

Compositions (in mol.%)	NC	Phonon energy (cm^-1)	Size (nm)	Crystallinity (%)	Method	Ref.
95As₂S₃-5ZnSe	ZnS	350	40-80	1-2	ICG	[36]
85Ge_1.5As₂S_6.5-15ZnSe	ZnS	350	200-1000	4	ICG	[16]
64SiO₂-23Ga₂O₃-13Li₂O	Ga₂O₃	770	20	15	ICG	[8]
30SiO₂-15Al₂O₃-29CdF₂-22PbF₂-4YF₃	PbF₂	250	9-18	20-30	ICG	[37]
44SiO₂-28Al₂O₃-17NaF-11YF₃	YF₃	400	15-25	32	ICG	[38]
50SiO₂-25YF₃-25KF	KY₃F₁₀	370	60	35	ICG	[39]
75TeO₂-12.5Bi₂O₃-12.5Nb₂O₅	Bi_0.8Nb_0.8Te_2.4O₈	880	20	97	ICG	[40]

Acronym	Meaning
ChG	Chalcogenide
ECE	External Crystallite Embedding
EDS	Energy Dispersive Spectrometer
FP	Fluorophosphate
GC	Glass Ceramic
HAADF	High Angle Annular Dark Field
HDD	High Doping Density
ICG	Internal Crystal Growth
LIDMEG	Luminescent Ion-Doped Microcrystallite Embedded Glass
LIDNEG	Luminescent Ion-Doped Nanocrystallite-Embedded Glass
LPE	Low Phonon Energy
MC	Microcrystallite
MIR	Mid-Infrared
NBG	Narrow Bandgap
NC	Nanocrystallite
NIR	Near-Infrared
NP	Nanoparticle
OANP	Optically-Active Nanoparticle
RE³⁺	Trivalent Rare Earth Ion
STEM	Scanning Transmission Electron Microscope
TEM	Transmission Electron Microscope
TGC	Transparent Glass-Ceramic
TM²⁺	Divalent Transition Metal Ion
XRD	X-Ray Diffraction

Dopants	Glassy host	Wavelength (µm)	σ_em	Nanocrystals (NC)	Size (nm)	method	Ref.
Er³⁺	TeO₂-Nb₂O₅-Bi₂O₃	4.3-4.95	-	Cubic Bi_0.8Nb_0.8Te_2.4O₈	∼20	ICG	[40]
Ho³⁺	TeO₂-Nb₂O₅-K₂O	2.858	1.76	Unidentified Cubic phase	< 100	ICG	[74]
Er³⁺	TeO₂-TiO₂-BaO-Li₂O	∼ 2.7	0.85	Li₂TeO₃	8-20	ICG	[72]
Er³⁺	TeO₂-PbO	∼ 2.7	0.8	Cubic PbTe₃O₇	8-30	ICG	[71]
Tm³⁺	TeO₂-Bi₂O₃-ZnO	∼2.0		Cubic Bi₂Te₄O₁₁	16-27	ICG	[75]
Er³⁺	TeO₂-ZnO-Na₂O	∼2.7	-	Y₃Al₅O₁₂	30	ECE	[76]

Abstract

1. Overview of transparent glass ceramics

1.1 Background on glass ceramics (GCs) for mid-infrared (MIR) light sources

1.1.1 Basics of GCs and MIR luminescence issues

1.1.2 Key properties of some MIR-relevant TGCs

1.2 General methods for fabrication of TGCs for MIR LIDNEGs

1.2.1 Internal crystal growth methods for fabricating LIDNEGs

1.2.2 External crystallite embedding methods for fabricating LIDNEGs

2. Mid-IR luminescent oxide TGCs

2.1 Oxyfluoride TGCs

2.1.1. Oxyfluoride TGCs containing PbF2 NCs

2.1.2. Oxyfluoride TGCs containing MF2 (M = Ca, Sr, Ba) NCs

2.1.3 Oxyfluoride TGCs containing ALnF4 (A = Na, K; Ln = Y, La, Lu) NCs

2.1.4. Oxyfluoride TGCs obtained by the external crystallite embedding (ECE) method

2.1.5. MIR luminescent oxyfluoride TGC fibers

2.1.6. Summary of progress in oxyfluoride TGCs

2.2 Tellurite (tellurium dioxide) TGCs

2.2.1. Overview

2.2.2. Tellurite TGCs obtained by conventional internal crystal growth

2.2.3. Tellurite TGCs obtained by external crystallite embedding method

2.2.4. Summary of tellurite TGC-based LIDNEGs

3. Mid-IR luminescent chalcogenide (ChG) TGCs

3.1 Rare-earth doped ChG TGCs

3.2 Transition metal ion-doped ChG TGCs

4. Summary and anticipated future directions

Funding

Acknowledgments

Disclosures

References

Cited By

Figures (18)

Tables (6)

Optics Express

2.1.1. Oxyfluoride TGCs containing PbF₂ NCs

2.1.2. Oxyfluoride TGCs containing MF₂ (M = Ca, Sr, Ba) NCs

2.1.3 Oxyfluoride TGCs containing ALnF₄ (A = Na, K; Ln = Y, La, Lu) NCs