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Abstract
PbS quantum dots (QDs) with mean radii of 3.7 nm to 9.0 nm are precipitated in silicate glasses. Upon above-band-gap excitation, photoluminescence from QDs is strongly dependent on their size and excitation wavelength, exhibiting photo-darkening (PD) or photo-brightening (PB). Photoluminescence of PbS QDs exhibits strong darkening by short excitation wavelength but the darkening gradually mitigated as the excitation wavelength increased and even turns to be photo-brightening at room temperature. But PD and PB show a much more complicated variation tendency under the same excitation condition when the size of QDs increased. The dependence of PD and PB on the QDs’ size and excitation wavelength indicates that electron/hole trap states of PbS QDs, defect states in surrounding glass matrix as well as on the interface between the glass matrix and PbS QDs all have strong effects on the photoluminescence properties of PbS QDs. These findings are important to modulate the photoluminescence and promote the potential applications of PbS QDs embedded glasses towards various optoelectronic devices.
© 2019 Optical Society of America under the terms of the OSA Open Access Publishing Agreement
1. Introduction
Quantum dots (QDs) have attracted much attention due to the unique optical and electronic properties induced by the quantum confinement effect [1,2]. Since bulk lead chalcogenides have narrow band gap energies and large exciton Bohr radii (0.41 eV/18 nm for PbS and 0.28 eV/46 nm for PbSe, respectively) [3,4], strong quantum confinement effect can be easily achieved with a moderate quantum dots sizes, and as a result, effective bandgap energies of lead chalcogenide QDs can be adjusted in a wide range, covering the most important telecommunication window and atmospheric transmission window [5,6]. Therefore, lead chalcogenide QDs have great potentials for applications in near-infrared and mid-infrared lasers, optical fiber amplifiers, bio-labeling and environmental monitoring [7–9].
Glasses have been utilized as host materials for various luminescence centers, such as rare-earth ions and transition metal ions [10–12], due to the good thermal, chemical and mechanical stabilities. Embedding QDs into glasses can keep the QDs from the rigorous conditions and maintain the chemical and optical properties, leading to the successful application of QDs embedded glass in color filters and saturable absorbers for generation of laser pulses [13,14]. However, due to the presence of defects on the surface or interface between quantum dot and glass matrix, photoluminescence (PL) efficiency of QDs embedded in glasses is generally low. For II-VI QDs, this problem has been partially solved by forming the core/graded shell structured CdSe/Cd1-xZnxSe QDs [15] and sandwich structured CdS/Cd1-xZnxS QDs [16] in glasses. However, for lead chalcogenide QDs in glasses, such passivation has not been reported, and PL efficiency of these QDs in glasses still remains low. Therefore, it is necessary to understand the optical response of lead chalcogenide QDs in glasses in order to develop highly photoluminescent QDs embedded glasses.
Due to the presence of surface defects or defects at the interface between QDs and amorphous matrix, PL properties of QDs in glasses are strongly dependent on the excitation conditions. PL intensity of QDs in glass either increases (photo-brightening, PB) or decreases (photo-darkening, PD) with time upon continuous-wave light excitation, and such phenomena have been observed in II-VI QDs and IV-VI QDs embedded glasses [17–19]. In our previous work, photo-darkening and photo-brightening of PbS QDs in glasses are found to be strongly dependent on the excitation intensity and experimental temperature. At low temperature and with low excitation intensity, photo-brightening prevails, and photo-darkening starts to dominate the PL with increases in temperature and excitation intensity [20]. Such temperature- and excitation-dependence make it possible to modulate the PL behavior of PbS QDs in glasses, and control the PL intensity using one additional light beam [21]. Several models have been proposed to explain the fundamental mechanism of PD and PB of QDs. Mitsunaga et al. proposed a three-level model to explain the PD effect [22]. Later, Jin et al. put forward a model where the PD effect was explained by the reduction of the number of trap levels and creation of “new” trap levels that have a shorter lifetime of nonradiative recombination [23]. Microscopically, these PD and PB phenomena are closely related to the fluorescence intermittency behavior of QDs which is known as blinking and was first reported in CdSe nanocrystals in 1996 [24,25]. PL blinking of QDs depends on the QDs itself and their local environment. That is, it’s affected by the particles’ structure (material, shape and multiple layers) [26,27], its passivation [28], surrounding matrix, its polarity [29,30], etc. The trapping model was proposed to account for most features of blinking according to which the blinking of nanocrystals was assigned to the electron tunneling towards a uniform spatial distribution of traps [31]. What’s more, both excitation intensity and excitation wavelength have strong effects on the blinking statistics. It has been reported that off-statistics is not influenced by the excitation wavelength or excitation intensity, but the on-statistics are sensitive to the excitation wavelength [32].
In this work, a series of PbS QDs with different diameters were precipitated in silicate glasses. For further understanding of the PD and PB of QDs in glasses, effects of excitation wavelength, excitation intensity, and temperature on the PL properties of PbS QDs with different sizes are investigated. The mechanism of photo-darkening and photo-brightening of PbS in glasses was discussed.
2. Experimental
Glass with nominal composition of 50SiO2-24Na2O-5Al2O3-10BaO-8ZnO-2ZnS-1PbO (mol%) was prepared. ZnS and PbO were used as the sources of sulfur and lead. Chemical powder with purity of >99.9% was weighted and mixed thoroughly before melting. A 50 g batch was melted in an alumina crucible at 1350 °C for 30 min under the ambient atmosphere, and the melt was poured onto a brass mold and pressed with another one for quenching. Glass thus obtained was transferred swiftly into a muffle furnace and annealed at 350 °C for 3 h to release the thermal stress. The annealed glass was cut into small pieces with a dimension of ~1 cm × 1 cm × 0.2 cm for thermal treatment to precipitate the PbS QDs at temperatures determined from the simultaneous differential scanning calorimetry (STA449c/3/G, NETZSCH, Germany).
Optical absorption spectra of glass samples at wavelength range from 300 nm to 2500 nm were recorded by using a UV/Vis/NIR spectrophotometer (UV3600, Shimadzu, Japan). X-ray diffraction (D8 Advance, Bruker, Germany) patterns were recorded to illustrate the structural changes before and after heat-treatment. Cu-Kα radiation with a scanning rate of 2°/min was used for the measurement with a resolution of 0.02°. High-resolution scanning transmission electron microscope (HR-TEM, TITAN THEMIS 200, FEI, USA) was used to characterize the size and shape as well as the nanostructure of the QDs. Lasers with wavelength of 532 nm, 800 nm, 900 nm, and 980 nm laser from a Ti: Sapphire laser (3900S, Spectra-Physics, USA), and diode lasers with wavelength of 655 nm, 1319 nm, and 1532 nm were used as excitation light. Excitation laser beam was focused into specimens using a silica lens with focal length of 5 cm. PL was collected along the direction perpendicular to the excitation beam. Before dispersing into a 0.25 m monochromator (77200, Oriel, USA), the collected PL was modulated by one mechanical chopper at a frequency of 25 Hz. InGaAs and InSb detectors were used to record the PL intensity. Low temperature experiments were conducted in a close-cycle cryostat system cooled by compressed helium gas (Optistat AC-V12, Oxford, UK).
3. Results and discussion
Formation of PbS QDs in the heat-treated specimens was confirmed using the XRD and TEM (Fig. 1), absorption and PL spectra (Fig. 2). With the increase of heat-treatment temperature, color of specimens changed from light yellow to black as a result of the formation of PbS QDs. XRD patterns (Fig. 1(a)) of the AP and heat-treated specimens were consistent with their color change (Fig. 2(a)). For AP specimen with light yellow color, only broad halo was observed, indicating the amorphous nature of the as-prepared glass and absence of detectable nanocrystalline phases in the matrix. When the specimens were heat-treated at 480 °C (for 10 h and 20 h) and 500 °C (for 10 h), brown and black colors were developed (inset in Fig. 2(a)), however, no obvious diffraction peak was observed, mainly due to the small size and low concentration of nanocrystalline phase in the glass matrix. When the specimens were heat-treated at 500 °C (for 20 h) or above, diffraction peaks corresponding to the cubic structure phase (JCPDS No. 78-1057) appeared and gradually became prominent without any shift in the diffraction angle, confirming the precipitation of PbS nanocrystals in the glass matrix. It has to be pointed out that no other diffraction peaks were observed in the XRD patterns, indicating the PbS nanocrystals were the single phase precipitated in the glasses. Considering the exciton Bohr radius of PbS crystal [3], PbS nanocrystals formed in these specimens were in the quantum confinement range, i.e., PbS QDs were formed in the glasses. Precipitation of these PbS QDs was further confirmed by TEM. These PbS QDs were homogeneously distributed in the glass matrix homogeneously (Fig. 1(b)), and the lattice constant measured from one single nanocrystal was found to be 2.04 Å, consistent with (220) crystal plane distance (d220 = 2.0965 Å, JCPDS No.: 78-1057). Average diameter of PbS QDs in glasses formed in the glass heat-treated 530 °C/10h was calculated to 8.9 ± 1.1 nm.
Fig. 1 (a) X-ray diffraction patterns of as-prepared glass and heat-treated glasses. Bottom gray line is the diffraction patterns of PbS crystal (JCPDS#: 78-1057). All XRD patterns were shifted vertically for clarity. (b) TEM image, (c) high-resolution TEM image of one nanocrystal, and (d) size distribution of nanocrystals formed in the glass heat-treated at 530 °C for 10 h.
Fig. 2 (a) Absorption spectra and (b) normalized photoluminescence spectra of as-prepared glass and heat-treated glasses. Inset in (a) shows the photograph of the as-prepared and heat-treated glasses.
Precipitation of PbS QDs in the glasses made it possible to tune the absorption and PL properties of these specimens induced by the quantum confinement effect [1,2]. Compared with
AP specimen, absorption shoulder or peaks were developed upon heat-treatment (Fig. 2(a)), consistent with gradual growth of PbS QDs in glass as confirmed by the XRD patterns and color changes (Fig. 1(a)). As the heat-treatment temperature increased from 480 °C to 530 °C, absorption peaks shifted from ~700 to ~2000 nm (Table 1). Based on the hyperbolic band model calculation [33,34], average diameter of PbS QDs increased from 3.7 ± 0.6 to 9.0 ± 1.2 nm (Table 1). At the meantime, peak wavelengths of these PbS QDs doped specimens shifted from 1000 to 1980 nm as the heat-treatment and duration increased (Fig. 2(b) and Table 1), demonstrating the quantum confinement effect. Full widths at half maximum (FWHM) of PL bands (135 to 273 nm, or 59 to 177 meV) were comparable to those observed from colloidal PbS QDs [35] and PbS QDs precipitated in other glasses [36,37]. Compared with the peak wavelengths of the absorption bands, Stokes shift of PL from PbS QDs embedded in these specimens decreased from ~260 nm (~440 meV) to ~10 nm (43 meV) (Table 1), and were also consistent with those from colloidal PbS QDs [38,39] and those reported for PbS QDs precipitated in glasses [40,41]. These PbS QDs with different sizes and wide tunable absorption/PL wavelengths made them very suitable to investigate the size-dependent and excitation-dependent photo-darkening and photo-brightening properties.
Table 1. Average diameter, peak wavelengths of first excitonic absorption bands, PL bands, Stokes shift, and FWHM of PL bands from PbS QDs doped in glass-ceramics heat-treated at various conditions.
PD and PB of PL from PbS QDs embedded in glasses were strongly dependent on excitation intensity and temperature, and these PD and PB phenomena were completely reversible due to the inert nature of the inorganic glass matrix [20,42]. These reversible PD and PB made it possible to control the PL of PbS QDs in glasses for practical applications [21]. Since the effective band gap energy of QDs was determined by their sizes, PD and PB of PL from PbS QDs in glasses were expected to be dependent on the size of QDs and the excitation wavelength (excitation energy), besides the excitation intensity and temperature. Upon 800 nm laser excitation (inset in Fig. 3(a)), PL intensity of 6.0 nm PbS QDs in glasses gradually decreased from (the initial intensity at , when the laser was turned on) to a stable value (indicated by the red dashed line in the inset), typical for PD of PL from QDs [19,43,44]. For simplicity, was used to describe the PD () and PB ()From the normalized time traces of PL intensity (), PbS QDs with different diameters showed the similar PD phenomena, but degree of PD was dependent on the diameters of PbS QDs (Figs. 3(a) and 3(b)). Under low intensity excitation (excitation photon flux: 40 × 1016 photons/s), values increased from ~68% (for 3.7 nm PbS QDs) to ~93% (for 7.7 nm PbS QDs) (Fig. 3(a)), showing that smaller PbS QDs exhibited more serious PD than large PbS QDs. Increase in excitation intensity enhanced the PD of PL from PbS QDs (Fig. 3(b)), consistent with our previous report [20,21,45]. However, at high excitation level (excitation photon flux: 200 × 1016 photons/second), degree of PD for large and small PbS QDs was reversed.values of large PbS QDs (5.0-9.0 nm) varied between 4 and 9.4%, and those of small PbS QDs (3.7 nm and 4.3 nm) varied between 13 and 16% (Fig. 3(b)). It can be expected that PL from PbS QDs in glasses can be completely quenched at excitation level high enough. Excitation intensity dependence of all sized PbS QDs was summarized in Fig. 3(c), and continuous and monotonic decrease invalues with the increase in excitation intensity was clearly illustrated, with the crossover of values for small and large PbS QDs. Even though the values of PbS QDs decreased monotonically, the initial PL intensity showed more interesting behaviors (Fig. 3(d)). For 3.7-nm and 4.3-nm sized PbS QDs, increased with the increase in excitation intensity until the excitation photon flux reached 280 × 1016 photons/s. For larger PbS QDs, increased initially and then decreased with the increase in excitation intensity. In addition, at certain excitation level, increased firstly and then decreased with the increase in diameter of PbS QDs with the maximum recorded from 5.0-nm sized PbS QDs. These phenomena clearly showed that size of PbS QDs had strong effects on the PD behaviors.
Fig. 3 Normalized PL traces of various sized PbS QDs recorded under 800 nm laser excitation with an excitation flux of (a) 40 × 1016 photons/second and (b) 200 × 1016 photons/second. (c)and (d) initial as a function of excitation photon flux for different sized PbS QDs.
Several factors may contribute to the size-dependent PL of PbS QDs in glasses. Firstly, excitation energy (1.55 eV for 800 nm laser) was larger than the band gap energies of PbS QDs in glasses used in this work. As the diameters of PbS QDs increased, the “normalized confinement energy” [] (whereis the exciton energy upon the immediate absorption of the excitation photon,is bulk bandgap energy of PbS crystal, and is the lowest exciton energy) [46,47], increased from 0.89 (3.7 nm) to 1.20 (4.3 nm), 1.60 (5.0 nm), 2.27 (6.0 nm), 3.05 (6.9 nm), 3.77 (7.7 nm), and 5.24 (9.0 nm) under 800 nm laser excitation. These large “normalized confinement energy” showed that electrons in PbS QDs were excited to high energy levels, and as a result, excitons generated in PbS QDs were hot excitons, especially in the large PbS QDs. Therefore, these excitons had large probability to escape from the PbS QDs and be trapped by the defects in glass matrix, resulting in the formation of ionized QDs and decreased PL quantum efficiency (QE) of PbS QDs in glasses. The excitation spectra (Fig. 4) showed that excitation efficiency greatly decreased when the photon energy was larger than ~1.2 eV for 7.7 nm-sized PbS QDs in glasses, confirming that hot carriers generated by high energy photons were more energetic to escape from the QDs. From this sense, the larger PbS QDs should exhibit much more serious PD than smaller PbS QDs, contradict with the results shown in Fig. 3(a). Secondly, surface states of PbS QDs in glasses could make significant contribution to the PD observed in Fig. 3. It is well known that there exist a lot of defects on the surface of QDs embedded in glasses, since these surface defects remain unpassivated [17,48]. These surface defects trap the photo-generated charge carriers, and decrease the PL QE of QDs. For smaller QDs, surface to volume ratio is very large, and this ratio decreases with the increase in size of QDs. Therefore, PD of the smaller PbS QDs should be more obvious than the larger ones, consistent with the PL time traces of PbS QDs recorded under low excitation intensity (Fig. 3(a)), where PD is alleviated with the increase in diameter of PbS QD. While the surface to volume ratio cannot account for the crossover of PD at higher excitation intensities recorded from large and small PbS QDs (Figs. 3(b) and 3(c)). In addition, the initial PL intensityrecorded from small and large PbS QDs showed different behaviors (Fig. 3(d)). For small PbS QDs (3.7 nm and 4.3 nm), the initial intensities monotonically increased with the excitation intensity, while for large PbS QDs, the initial intensitiesreached the maximal values when the excitation intensities were ~100 × 1016 photons/s. The crossover of PD (Fig. 3(c)) between small and large PbS QDs and the changes in the initial PL intensitywith excitation intensity (Fig. 3(d)) indicated that additional processes, besides the photon-induced ionization of PbS QDs and surface trapping of charge carriers, were evoked.
Fig. 4 Absorption spectrum, photoluminescence spectrum by 800 nm and excitation spectrum from 7.7 nm-PbS QDs doped glass sample.
These additional processes were considered to be closely related to the complex surface states of PbS QDs and interface between PbS QDs and glass matrix. For PbS QDs precipitated in glass matrix, surface defects such as dangling bonds formed the surface trap states (STS) below the state [49–51]. Due to the discontinuity in amorphous structure, there existed defects located at the interface between QDs and glass matrix, forming the defect states (DS). Relative energy positions of these STS and DS were strongly dependent on the size of QDs. STS located at higher energy than DS for small PbS QDs, and vice versa for large PbS QDs [49,52,53]. As a result, normal Stokes shift of PL from small PbS QDs in glasses changed into anti-Stokes shift for large PbS QDs, and spectral position of emission from DS gradually shifted from the red-side to the blue-side of the band edge emission of PbS QDs as the diameter of PbS QDs increased [52]. Presence of these DS has been confirmed by the near-infrared anti-Stokes PL [54], and intensities of the anti-Stokes PL induced by the DS showed similar dependence on the excitation intensity shown in Fig. 3(d). Distribution of electrons (or holes) in these DS was found to be dependent on local temperature, which in turn influenced the emission from these states [46,54]. In fact, when the excitation intensity increased, PL spectra of PbS QDs in glasses recorded at the stable state (when reached the stable state) showed blue-shift (Fig. 5). According to our previous works [46,54], this blue-shift was induced by the redistribution of charge carriers trapped in DS [54] and change in the effective band energy of PbS QDs in glasses with the increase in temperature [55]. Therefore, these size-dependent PL time traces should also rely on the excitation photon energy and experimental temperature.
Fig. 5 Normalized photoluminescence spectra of PbS QDs recorded at stable state under different illumination intensity from PbS QDs of (a) 5.0 nm excited by 800 nm, (b) 7.7 nm excited by 655 nm.
As expected, changes in the excitation photon energy (Fig. 6) and experimental temperature (Fig. 7) did affect the PD and PB of PL from PbS QDs in glasses. For 7.6 nm-sized PbS QDs,values increased with the increase in excitation wavelength (decrease in excitation photon energy), and they increased from 0.48 to 0.75, 0.88, 0.93, 0.98 and 1.03 when the excitation wavelength increased from 532 nm to 655 nm, 900 nm, 980 nm, 1319 nm, and 1532 nm, respectively (Fig. 6(a)). When the excitation wavelength further increased to 1532 nm, PB of PL from PbS QDs was observed with avalue of 1.03 (Fig. 6(a)). Changes in values for 6.0 nm- and 7.7 nm-sized PbS QDs with excitation photon flux were shown in Figs. 6(b) and 6(c), respectively. These values decreased with the increase in excitation photon flux for the excitation wavelength employed, same as that observed in Fig. 3(c). Compared to that shown in Fig. 3(c) (which was recorded using 800 nm excitation light), curves of these values recorded shift towards large photon flux as the excitation wavelength increased. This shift showed the PD of PL from PbS QDs was reduced when the excitation photon energy decreased, which led to less thermal energy generated and smaller local temperature rise around PbS QDs upon absorption of each photon. At low temperatures, PB of PL from PbS QDs was observed (Fig. 7(a)), similar to our previous work on PbS QDs [20,21]. With the increase in experimental temperature, PL intensities of PbS QDs grew faster and reached the stable state in a shorter time, indicating that this process was promoted by thermal energy. At room temperature, PD similar to that observed in Fig. 3. Figure 7(b) summarized the changes in values with temperature and excitation intensity. Different from that observed in Fig. 3(d), the values showed initial increase to the maximal values at ~100 × 1016 photons/second, and then started to decrease with further increase in the excitation photon flux. At higher temperature (293 K), the values only showed one shoulder when the excitation photon flux increased. For large PbS QDs shown in Fig. 3(c) and Fig. 6(c), similar shoulders were observed at low excitation level. These observations confirmed that the size-dependent STS and DS exerted large effects on the PD and PB of PL from PbS QDs in glasses.
Fig. 6 (a) Normalized PL traces of 7.7 nm-sized PbS QDs recorded using different excitation wavelength with photon flux of 60 × 1016 photons/second. values of (b) 6.0 nm- and (c) 7.7 nm-sized PbS QDs recorded using different excitation wavelength.
Fig. 7 (a) Low temperature PL traces and (b) values of 7.7 nm-sized PbS QDs excited by 1532 nm laser.
Based on the above results and discussion, the PD and PB of PL from PbS QDs in glasses can be described using the following schematic diagram (Fig. 8) [43,52,54]. For small PbS QDs, DS locates below STS. The photo-generated charge carriers nonradiatively relax to the STS and DS, giving the band edge emission (P1, dashed lines) and defect emission (P2, dotted lines) [52]. With the increase in the excitation photon flux, more thermal energy will be generated in the local environment of PbS QDs and lead to the temperature rise. Increase in the local temperature of PbS QDs in glasses will lead to redistribution of trapped charge carriers in DS, resulting in the broadening and blue-shift of the PL band (Fig. 5) and near-infrared excitation induced anti-Stokes PL from PbS QDs [54]. For small PbS QDs, energy gap between STS and DS is relatively large, the trapped charge carriers in DS can be ejected into the defect states in the surrounding glass matrix, instead of being thermalized into the STS. This thermal ejection of trapped charge carriers can be accelerated by increasing the excitation photon flux or decreasing the excitation wavelength, leading to the enhanced PD of PL from PbS QDs. While, for large PbS QDs, the DS is located on top of STS, leading to the defect emission with higher energy (P2 dotted line) and band edge emission with lower energy (P0 and P1, solid and dashed lines) [52]. These photo-generated charge carriers can be either trapped by DS or relax to STS andlevel of PbS QDs, giving rise to the defect emission and band edge emission. Upon low intensity excitation, these photo-generated charge carriers can be easily trapped by the DS, resulting in very low initialvalue. Under continuous excitation, these DS are gradually populated and more charge carriers can reach the STS andlevel, leading to the gradual enhancement in PL intensity, i.e., photo-brightening. However, under intense light excitation, the light-induced local heating can easily eject these trapped charge carriers into defect states in glasses, leading to the PD of PL of PbS QDs. These size-, excitation-photon-energy- and temperature-dependent photo-darkening and photo-brightening of PL from PbS QDs in glasses can be described using the above proposed model.
Fig. 8 Schematic diagram for PD and PB of PbS QDs in glasses. STS is the surface trap states (red dashed lines), and DS is the defect states located at the interface between QDs and glass matrix (thin gray lines). Dashed down arrows and solid down arrow represent the band edge emission; dotted down arrows represent defect emission. The red up arrow represents the excitation. Waved arrows are the nonradiative relaxation, and curve-dashed arrows are the thermal ejection of charge carriers.
4. Summary
In conclusion, a series of PbS QDs with average diameters from ~3.7 nm to ~9.0 nm were prepared in glass by thermal treatment. Photoluminescence from those PbS QDs were strongly dependent on their size and excitation condition. Dependence of PL on QDs’ sizes essentially resulted from the different relative location between surface trap states on surface of QDs and defect states at interface from different sizes QDs. Photodarkening and photobrightening phenomena were observed at room temperature and photoluminescence tends to be brightening from darkening with excitation wavelength increasing because thermal ejection of trapped charge carriers can be accelerated by decreasing the excitation wavelength, leading to the enhanced PD of PL from PbS QDs. Charge carriers trapped in DS and STS even ejected to the defect states in glass matrix worked together and led the final photoluminescence of PbS QDs. Optimum fluorescence properties of PbS QDs can be achieved via the proper selection of excitation condition.
Funding
Natural Science Foundation of Hubei Province (2018CFA005); National Natural Science Foundation of China (51602235).
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