Physico-chemical and optical properties of Er³⁺-doped and Er³⁺/Yb³⁺-co-doped Ge₂₅Ga_9.5Sb_0.5S₆₅ chalcogenide glass

Introduction

Rareearthdoped chalcogenide glasses exhibit many interesting properties therefore, they are promising materials in variety of applications. Chalcogenide glasses (ChGs) possess low phonon energy and high refractive index [1], [2], [3] which suppresses the multiphonon relaxation and promotes radiative recombination, respectively [4]. Moreover, ChGs are transparent from visible to mid-infrared spectral region and show large intra-4f cross sections when they are doped with rare-earth (RE) ions [5]. RE³⁺-doped ChGs can be applied such as optical amplifiers [6], waveguides [7], displays [8], sensors and detectors [9], [10], [11], lasers [12], [13].

The Er³⁺: ⁴I_13/2→⁴I_15/2 (λ≈1.5 μm) near-infrared emission has attracted the attention in telecommunication C-band [14], [15]. To improve the efficiency of the ≈1.5 μm emission, the Er³⁺ ions can be co-doped with the Yb³⁺ ions, where Yb³⁺ ions play role of the sensitizer at pumping wavelength of ≈980 nm and allow the energy transfer to neighboring Er³⁺ ions. It was demonstrated that the Er³⁺/Yb³⁺ co-doping has become an effective method for production of short amplifiers and efficient lasers in the long haul telecommunications [16], [17]. The efficient Yb³⁺→Er³⁺ energy transfer was first investigated by Snitzer and Woodcock in Na-K-Ba silicate glasses [17]. Recently, Er³⁺/Yb³⁺ -co-doped fiber lasers [17], [18] and planar waveguide amplifiers [14], [16], [19] have been demonstrated.

In this paper, we have chosen the Ga-Ge-Sb-S chalcogenide glasses as the promising host matrix for the RE³⁺ ions [20], [21]. The addition of Sb into Ge-Ga-S glasses improves the resistance against crystallization and resistance to moisture with still large solubility of the RE³⁺ ions due to the presence of Ga [20], [22]. Reference [21] reports the effect of compositional changes in Ge₂₅Ga_10−x Sb_xS₆₅: 0.5 at.% Er³⁺ on the upconversion photoluminescence (UCPL). In the present study, we aim on the UCPL and ≈1.5 μm photoluminescence (PL) emission in (Ge_0.25Ga_0.095Sb_0.005S_0.65)_99.9−x Er_0.1Yb_x, where x=0, 0.1, 0.5 and 1.0 at.% and (Ge_0.25Ga_0.095Sb_0.005S_0.65)_99.85Er_0.05Yb_0.1 (further denoted as Ga-Ge-Sb-S: Er³⁺/Yb³⁺) chalcogenide glasses with respect to their optical properties and structure.

Experimental

The (Ge_0.25Ga_0.095Sb_0.005S_0.65)_99.9−x Er_0.1Yb_x (x=0, 0.1, 0.5, 1) and (Ge_0.25Ga_0.095Sb_0.005S_0.65)_99.85Er_0.05Yb_0.1 ChGs were synthesized from high-purity elements of Ge (5N), Ga (5N), Sb (5N), S (4.5N), Er (3N) and Yb (3N). Elements were weighted into silica ampoules, which were sealed at residual pressure of ~10⁻³ Pa. The total weight of batch was 10 g. The sealed ampoules were heated at 1270 K for 24 h in a rocking furnace. The melt of the glass was quenched into water and the glass was annealed at 20 K below the glass transition temperature T_g for 2 h to relax the mechanical strain. Finally, the prepared samples were cut into discs with diameter of ≈10 mm and thickness ≈2–4 mm and polished into optical quality.

The non-crystalline nature of prepared glasses has been confirmed by X-ray diffraction (XRD) analysis using the Bruker D8 advance powder XRD diffractometer. The XRD investigation with Cu Kα radiation was carried out on the as-prepared and annealed glasses in the 2θ range of 5–65° with a step of 0.02°. The compositions of prepared glasses were characterized by the energy dispersive X-ray (EDX) microanalyzer Aztec X-Max 20, Oxford Instruments at accelerating voltage of 20 kV. Differential scanning calorimeter DSC Diamond, Perkin-Elmer was employed to investigate thermal properties of the studied glasses in the range of 570–870 K and at heating rate of 10 K min⁻¹. A piece of glass about ≈10 mg was sealed into an aluminum pan for the measurement. The glass transition temperature (T_g) was determined as half-height between extrapolated onset and endset, the crystallization temperature (T_c) as onset of crystallization peak.

The Raman spectra were measured under the Nd:YAG laser (λ=1064 nm) excitation at room temperature using the FT-IR spectrophotometer IFS 55 with the Raman FRA-106 accessory (Bruker) for back scattering geometry. Raman spectra were reduced by the Gammon-Shuker approximation [23] and decomposed into individual bands by the pseudo-Voigt functions.

Archimedes method was used for determination of the density of prepared glasses. Refractive index of glasses was determined by the variable angle spectroscopic ellipsometry (VASE, J.A. Woollam Co., Inc.) measured in the spectral region of 500–2300 nm with a spectral step of 20 nm and at angles of light incidence 65°, 70°, 75°. The ellipsometric data were parameterized by the Sellmeier model [24] in transparent spectral region of studied materials with obtained fit accuracy given by the mean square errors <1.5. Er³⁺ and Yb³⁺ absorption cross sections spectra were determined using the double-beam UV-Vis-NIR spectrophotometer (JASCO V-570) in the spectral region of 300–2500 nm with a spectral step of 2 nm. Intensities of the Er³⁺ intra-4f electronic transitions in the Ge₂₅Ga_9.5Sb_0.5S₆₅ host matrix were calculated by the Judd-Ofelt theory [25], [26], [27], [28] using the Er³⁺ absorption bands centered at 1538, 989, 815, 661, 545, and 522 nm and JOF program v. 2.3 [25], [26], [29]. The areas of absorption bands were fitted by Gausssians in Fityk program v. 0.9.8 [30]. The Er³⁺ 1.5 μm PL emission spectra of the Ga-Ge-Sb-S: Er³⁺/Yb³⁺ ChGs were measured in the spectral region of 1440–1650 nm with a 0.5 nm step and excited by the 980 nm diode laser (power density ≈80 W cm⁻²). PL signal was processed with a 1/2 m double grating monochromator and amplified by preamplifier and lock-in amplifier at chopping frequency of 30 Hz. The PL signal was detected under 90° by the two-step-cooled Ge detector. PL and UCPL emission spectra measured in the spectral region of 500–900 nm with a 0.5 nm step were acquired at pumping wavelengths of 490, 532, 810 and 980 nm using the Ti:sapphire laser (power density ≈3800 W cm⁻²) pumped with Nd:YVO₄. The PL signal was processed through 1/8 m monochromator and detected by the GaAs photomultiplier tube. All photoluminescence measurements were carried out at room temperature.

Results and discussion

All measured XRD patterns presented in Fig. 1 confirmed the amorphous state of the prepared Ga-Ge-Sb-S: Er³⁺/Yb³⁺ ChGs.

Fig. 1:

XRD patterns for Ga-Ge-Sb-S: Er³⁺/Yb³⁺ chalcogenide glasses. (color online).

The chemical composition of (Ge_0.25Ga_0.095Sb_0.005S_0.65)_99.9−x Er_0.1Yb_x (x=0, 0.1, 0.5, 1) and (Ge_0.25Ga_0.095Sb_0.005 S_0.65)_99.85Er_0.05Yb_0.1 ChGs determined by the EDX spectroscopy is shown in Table 1. The observed results show a small deviation between experimental and theoretical chemical composition.

The temperature difference ∆T between the crystallization temperature T_c≈837 K and the glass transition temperature T_g≈713 K is >100 K suggesting the good thermal stability of synthesized glass. By the thermal stability is meant a broader temperature range at which the glass can be processed such as for fibers drawing.

Densities of prepared ChGs determined by the Archimedes method are shown in the Table 2. Densities were further used to calculate the concentration of Er³⁺ and Yb³⁺ ions per unit volume in Ge₂₅Ga_9.5Sb_0.5S₆₅ glass (see Table 2). The density of Ga-Ge-Sb-S ChGs increases with increasing Yb and Er content [31].

Moreover, the refractive index determined by the VASE slightly increases with increasing concentration of Yb³⁺ ions (Table 2).

The structure of the Ge₂₅Ga_9.5Sb_0.5S₆₅ glass doped with 0.1 at.% Er was investigated by Raman spectroscopy, as shown in Fig. 2. The spectrum was decomposed into several bands. The main and the most intense Raman band near 340 cm⁻¹ can be assigned to a ν₁ vibrational mode of the corner-sharing GeS_4/2 and GaS_4/2 tetrahedra. The small band at 265 cm⁻¹ can be associated with vibrations of the metal–metal bonds in S₃Ge(Ga)-Ge(Ga)S₃ structural units. The band with maximum near of 375 cm⁻¹ can be assigned to the ν₁ mode of two edge shared tetrahedra Ge₂S₄S_2/2 and Ga₂S₄S_2/2. The last two bands located at 400 and 435 cm⁻¹ correspond to the ν₃ modes of the corner-sharing and edge- sharing GeS_4/2 and GaS_4/2 tetrahedra. Broad band in the region of 80–230 cm⁻¹ does not show any fine structure and thus, it is difficult to decompose it into individual bands. However, the ν₂ and the ν₄ vibrational modes of the corner-sharing GeS_4/2 and GaS_4/2 originate at this spectral region. We did not find any band which can be associated with vibrational modes of Sb-based structural units, probably due to the low concentration of Sb (<1 at.%) in the studied glasses.

Fig. 2:

Reduced Raman spectrum of the (Ge_0.25Ga_0.095Sb_0.005S_0.65)_99.9Er_0.1 glass. (color online).

Absorption coefficients of the Er³⁺-doped and the Er³⁺/Yb³⁺-co-doped Ge₂₅Ga_9.5Sb_0.5S₆₅ glasses are shown in Fig. 3. We observed the ground state absorption (GSA) bands of Er³⁺ centered approximately at 1538, 979, 815, 661, 545, and 522 nm. One strong GSA band at 992 nm originates from Yb³⁺: ²F_7/2→²F_5/2 transitions and this band overlaps with the Er³⁺: ⁴I_15/2→⁴I_11/2 GSA band at 979 nm. In the Fig. 3 is evident a red shift of the absorption edge of the Ge₂₅Ga_9.5Sb_0.5S₆₅ host matrix with increasing Yb³⁺ content.

Fig. 3:

The absorption coefficient of the Ge₂₅Ga_9.5Sb_0.5S₆₅: Er³⁺/Yb³⁺ ChGs with labeled Er³⁺ and Yb³⁺ GSA transitions from respective levels ⁴I_15/2 and ²F_7/2 to the upper manifolds. (color online).

The Judd-Ofelt phenomenological parameters for the Er³⁺ ions embedded in the Ge₂₅Ga_9.5Sb_0.5S₆₅ glassy host were found to be Ω₂=(12.58±1.10)×10⁻²⁰ cm², Ω₄=(3.25±1.28)×10⁻²⁰ cm² and Ω₆=(2.05±0.45)×10⁻²⁰ cm² which are comparable to similar chalcogenide glasses [21]. Root-mean-square (RMS) deviation between the theoretical and the experimental line strengths was 0.63×10⁻²⁰ cm². The large value of the Ω₂ parameter can be related to a high degree of the covalent bonding around the Er³⁺ ions and in some extent to the asymmetry. The spectroscopic quality factor Ω₄/Ω₆=1.59 can be used as a prediction for the stimulated emission ability for such intra-4f transitions, where the first terms of doubly reduced matrix elements are negligible or zero. The calculated radiative transition probabilities, branching ratios and radiative lifetimes of selected Er³⁺ transitions are presented in Table 3.

The proposed UCPL mechanism in the Er³⁺/Yb³⁺-co-doped samples under 980 nm excitation wavelength is schematically drawn in the Fig. 4. The 980 nm wavelength excites the Er³⁺ and Yb³⁺ by the ground state absorption (GSA) into Er³⁺: ⁴I_11/2 and Yb³⁺: ²F_5/2 levels, respectively. Subsequently, the Er³⁺: ⁴F_7/2 manifold can be populated: (1) by the excited state absorption (ESA) ⁴I_11/2→⁴F_7/2 within Er³⁺ ion, or (2) by the energy transfer (ETU1) between Er³⁺ and Yb³⁺ ions as Er³⁺: ⁴I_11/2, Yb³⁺: ²F_5/2→Er³⁺: ⁴F_7/2, Yb³⁺: ²F_7/2. The thermalized levels ²H_11/2 and ⁴S_3/2 responsible for the green UCPL will be populated from ⁴F_7/2 in particular by the multiphonon relaxation due to their small energy difference. On the other hand, the energy transfer (ETU2) Er³⁺: ⁴I_13/2, Yb³⁺: ²F_5/2→Er³⁺: ⁴F_9/2, Yb³⁺: ²F_7/2 can be assumed as well promoting the Er³⁺ red UCPL ⁴F_9/2→⁴I_15/2. It should be noted that the ETU1 and ETU2 processes can originate between neighboring Er³⁺ ions (without Yb³⁺) due to similar energy positions of the Er³⁺: ⁴I_11/2 and Yb³⁺: ²F_5/2. Revealing the UCPL dynamics is a difficult task and among the ChGs was recently studied in Ge-Ga-S: Er³⁺ ChGs [32], [33]. Moreover, experimentally observed Er³⁺: ⁴S_3/2→⁴I_13/2 (≈850 nm) UCPL emission populating the ⁴I_13/2 level is promising for the Er³⁺: ⁴I_13/2→⁴I_15/2 (≈1.5 μm) optical amplification.

Fig. 4:

Schematic energy level diagram for the Er³⁺/Yb³⁺ -co-doped Ge₂₅Ga_9.5Sb_0.5S₆₅ glass with observed emissions and proposed mechanisms standing behind the UCPL emissions. (color online).

The effect of Yb³⁺ addition to sensitize the UCPL emission was investigated for samples (Ge_0.25Ga_0.095 Sb_0.005S_0.65)_99.9−x Er_0.1Yb_x (x=0, 0.1, 0.5, 1) and (Ge_0.25Ga_0.095Sb_{0. 005}S_0.65)_99.85Er_0.05Yb_0.1 at pumping wavelength of 980 nm. The UCPL emission spectra are presented in the Fig. 5 from the visible to the near-infrared spectral region. The addition of Yb³⁺ ions lowers the total UCPL emission intensity and promotes the red-to-green UCPL intensity ratio. This is probably because of a red shift of the absorption edge of glassy host matrix with Yb addition (Fig. 3) which promotes the absorption of the excitation light (due to the ESA, ETU processes) by host matrix. Thus, the co-doping of the Ge₂₅Ga_9.5Sb_0.5S₆₅: 0.1 at.% Er³⁺ ChGs by the Yb³⁺ ions is inefficient to improve the UCPL emission intensity.

Fig. 5:

UCPL emission spectra of the (Ge_0.25Ga_0.095Sb_0.005S_0.65)_99.9Er_0.1 and the (Ge_0.25Ga_0.095Sb_0.005S_0.65)_99.8Er_0.1Yb_0.1 glasses at pumping wavelength of 980 nm laser. (color online).

Stokes and anti-Stokes (UCPL) PL spectra of the (Ge_0.25Ga_0.095Sb_0.005S_0.65)_99.9Er_0.1 glass at pumping wavelengths of 490, 532, 810 and 980 nm are presented in Fig. 6. The observed emission bands at 529, 551, 660, 808, and 850 nm can be assigned to the Er³⁺: ²H_11/2→⁴I_15/2, ⁴S_3/2→⁴I_15/2, ⁴F_9/2→⁴I_15/2, ⁴I_9/2→⁴I_15/2 and ⁴S_3/2→⁴I_13/2 electronic transitions, respectively. The PL emission intensities at 550 nm originating from ⁴S_3/2→⁴I_15/2 transitions are similar for excitation wavelength of 490 and 532 nm which directly excite the upper ⁴F_7/2 and ²H_11/2 levels. From these manifolds the ⁴S_3/2 level is populated by the multiphonon relaxation and thus, the emission characteristics are similar. On the other hand, the emissions at excitation wavelengths of 810 and 980 nm originate from the nonlinear UCPL processes and thus, they have lower intensity (approximately ≈5 times). However, the UCPL intensity is still sufficiently high, which makes these materials promising for the UCPL applications.

Fig. 6:

Emission spectra of the (Ge_0.25Ga_0.095Sb_0.005S_0.65)_99.9Er_0.1 glass measured at various excitation wavelengths λ_ex. (color online).

In addition, we investigated the ≈1.5 μm PL emission originating from Er³⁺: ⁴I_13/2→⁴I_15/2 transitions at 980 nm pumping wavelength against Yb³⁺ content which is depicted in the Fig. 7. The PL emission intensity decreases with increasing Yb³⁺ content. This is probably due to the concentration quenching and/or due to UCPL processes depopulating the ⁴I_13/2 energy level. Thus, the present study demonstrates that the Yb³⁺ addition into Ge₂₅Ga_9.5Sb_0.5S₆₅: Er³⁺ ChGs does not lead to appreciable sensitization of the UCPL as well as 1.5 μm PL emissions. However, we believe that this drawback can be overcome by using the host material with larger optical band gap energy. Such promising materials can be chalcohalide or oxychalcogenide glasses where the trade-off between intra-4f electronic intensities, solubility of Er³⁺/Yb³⁺ ions and optical band gap energy might be found.

Fig. 7:

(a) The PL Er³⁺: ⁴I_13/2→⁴I_15/2 emission spectra and (b) the same normalized spectra of the Ge₂₅Ga_9.5Sb_0.5S₆₅: 0.1 or 0.05 at.% Er³⁺ ChGs co-doped with various concentrations of Yb³⁺ ions under 980 nm excitation. (color online).

Conclusions

Er³⁺-doped and Er³⁺/Yb³⁺-co-doped Ge₂₅Ga_9.5Sb_0.5S₆₅ thermally stable chalcogenide glasses were synthesized by the melt-quenching technique. The Judd-Ofelt theory was used for calculation of the Er³⁺ intra-4f electronic transitions intensities. The Ge₂₅Ga_9.5Sb_0.5S₆₅ glass doped with 0.1 at.% Er³⁺ shows intense upconversion photoluminescence (UCPL) from green to near-infrared spectral region. The green UCPL emission is approximately 5 times lower than the green Stokes photoluminescence (PL) Stokes emission intensity, which makes this material attractive for the UCPL applications. However, it is presented that the Yb³⁺ ions sensitization is ineffective resulting in the decrease of the total UCPL emission intensity under 980 nm laser excitation. This is attributed to a red shift of the absorption edge of host matrix with Yb addition, which is merged with the upper Er³⁺ energy level manifolds ⁴F_7/2, ²H_11/2 and ⁴S_3/2. Moreover, the ≈1.5 μm PL emission intensity decreases with the increasing Yb content probably due to the concentration quenching. Thus, the addition of Yb³⁺ ions was shown to be inefficient to improve the UCPL and ≈1.5 μm PL, which could be overcome by the using of chalcohalide or oxychalcogenide glasses as host matrices. Thereafter, the observed Er³⁺: ⁴S_3/2→⁴I_13/2 (≈850 nm) emission could be utilized for the optical amplification at ≈1.5 μm (⁴I_13/2→⁴I_15/2) via upconversion processes.