Introduction

Two-dimensional (2D) Transition Metal Dichalcogenides (TMDs) provide an appealing platform for emerging atomic scale research in nanophotonic and optoelectronic applications1,2,3,4. Monolayer molybdenum disulfide (MoS2), in particular, gains considerable attention due to its direct band gap and potential integration with other nanostructures to form nanoscale van der Waals heterojunctions with intriguing physical and optical properties5. Indeed, it has been shown that the optical properties of MoS2 monolayers, such as photoluminescence (PL), can be manipulated through its coupling with nanomaterials of various dimensionalities. In particular, zero dimensional (0D) quantum dots and nanoparticles6,7, one-dimensional (1D) nanowires and nanorods8,9,10, as well as other 2D materials5,11 had been combined with monolayer MoS2 to manipulate its emission intensity and/or quantum yield. Besides this, polymeric spacing12, defect engineering13, doping14, and chemical modification15 approaches were employed to manipulate the emission properties. However, monolayer MoS2 suffers from low intrinsic photoluminescence (PL) quantum yield (0.01–0.6%), dominated by the A-excitonic emission, due to its sub nanometer thickness and defect density mediated nonradiated recombination1. The low PL yield was overcome (more than 95%) with chemical treatment by an organic superacid16. In contrast to a monolayer MoS2, few layers of MoS2 have several orders of magnitude lower PL quantum yield1. On the other hand, few-layers MoS2, as an indirect semiconductor, have significantly larger optical density, which enhances its external quantum efficiency17. Owing to this advantage, research on the PL properties in few layers MoS2 has received significant attention. For example, metallic and other nanostructures5,7 were used to manipulate the A-excitonic emission in few layers of MoS218. However, this approach has only been limited to the enhancement of the A-excitonic emission.

On the other hand, transparent thermoplastic glasses (TTG) were extensively used for homogeneous incorporation of 2D layered materials. However, the relevant studies were limited to measure the nonlinear optical response of the embedded 2D nanoflakes19,20. On a rather different manner photonic crystal cavities21,22,23, as well as Mie-resonant metasurfaces6, have been employed to tailor the optical properties of MoS2. Similar to the case of nanostructures, the manipulation of PL emission has been only limited to A-exciton. Mikkelsen and co-workers24,25 were the first who carried out a systematic study to manipulate the B- excitonic emission of a single-layer of MoS2. However, the study of the emission properties was limited to the ground A-exciton state. Nevertheless, a detailed investigation of the B-exciton state in the ultrafast regime is crucial to shed light on the physical phenomena that take place.

In this study, we present the development of a nanohybrid heterojunction system composed of few layers of MoS2 embedded into a silver metaphosphate glass (AgPO3), as a means to enhance and control the MoS2 exciton emission. The selection of AgPO3 glass as a host matrix is prompted by several reasons: First, its transparency in most of the visible range (Fig. S1) enables the full exploitation of the AgPO3:MoS2 photoluminescence properties towards various nanophotonic applications26. Moreover, the presence of silver nanoparticles (NPs) within the glass matrix gives rise to interesting optical phenomena that can be exploited towards enhancing and manipulating the PL properties of the incorporated MoS2 layers. Finally, the AgPO3 glass exhibits a very low glass transition temperature of 192 °C, which is indicative of its soft nature. As a consequence, the MoS2 integration process is performed at low temperatures, suitable to avoid any oxidation. On top of that, an advanced 2D exciton–plasmon system composed of a few-layer TMD integrated with a semiconducting metal-phosphate glass is realized. It is shown that the layered TMDs create nanoscale van der Waals heterojunctions with the metallic nanostructures of the glass, which can be exploited to tailor light-matter interactions at the nanoscale.

Results

Fabrication and characterization of MoS2 and nanoheterojunctions

The MoS2 flakes were obtained by liquid exfoliation (see “Methods”)27. The lateral dimensions of the MoS2 nanoflakes, as determined by SEM imaging, were found to lie within the micrometer range (Fig. S2a), while, the average thickness measured by AFM was ~ 4 nm (Fig. S2b, c). A schematic representation of the composite glass, comprising numerous AgPO3:MoS2 nano-heterojunctions is illustrated in Fig. 1 (see “Methods”).

Figure 1
figure 1

Schematic representation of the composite transparent silver metaphosphate glass (AgPO3) matrix incorporating the MoS2 flakes. The right panel depicts a single AgPO3:MoS2 nanoheterojunction.

Full size image

Optical spectroscopy

Absorption spectroscopy was employed to confirm the formation of nano-heterojunctions between MoS2 and AgPO3. The pristine AgPO3 glass exhibits two characteristic peaks at 2.0 and 2.5 eV (Fig. S3a,b), which correspond to the Ag plasmonic bands and are attributed to a bimodal distribution of isolated nanoparticles or their clusters attained due to the phosphate matrix. Another reason of the emergence of the plasmonic band at 2.0 eV is the clustering/agglomeration of Ag NPs observed. Indeed, as the effective nanoparticles size increases, a corresponding red shift in the plasmon band occurs. This is also indicated by the broad plasmonic band centered at 2.0 eV. The absorption spectrum of bare MoS2 flakes exhibits the two characteristic excitonic peaks at 1.84 eV (A-exciton) and 2.03 eV (B-exciton) respectively (Fig. 2a, red line)1. Both peaks were also present in the absorption spectrum of the AgPO3:MoS2 heterojunctions of composite matrix (green line of Fig. 2a), i.e. in which the MoS2 is incorporated within the glass. At the same time, the absorption intensity of AgPO3:MoS2, is enhanced compared to the pristine AgPO3. The small blue shift of the A- and B-exciton peak positions by 13 meV and 17 meV, respectively, is due to the change of the dielectric environment rather than any oxidation process. The integration of MoS2 within AgPO3 glass took place at 170 °C, which is below the glass transition (TG), and thus hard to cause significant oxidation of the phosphate network. Instead, the modification of the dielectric constant, from that of the solvent to the higher dielectric constant of the surrounding glass matrix, could be the reason for the observed shift of the exciton states. It is noted that these findings are not observed when the MoS2 flake is positioned on the surface of AgPO3 glass, i.e. AgPO3/MoS2 spectrum in Fig. 2a (blue line).

Figure 2
figure 2

Spectroscopic characterizations of MoS2 and their heterojunctions. (a) Optical absorption spectra and (b) X-ray diffraction pattern, of MoS2 (MoS2 on glass substrate), AgPO3, AgPO3/MoS2 (MoS2 on AgPO3), and AgPO3:MoS2 (MoS2 embedded into AgPO3).

Full size image

It is widely acknowledged that the trigonal prismatic phase (2H-phase) integrity plays a crucial role in PL emission of exfoliated MoS2. Aiming to identify the phase integrity in MoS2 flakes dispersed into AgPO3, a series of structural studies have been carried out. In particular, the X-ray diffraction pattern of MoS2 (Fig. 2b) exhibits a strong peak located at 14.34°, corresponding to the (002) plane, which agrees well with the hexagonal MoS228,29. Besides this, the examination of the AgPO3/MoS2 and AgPO3:MoS2 matrix showed a primary peak at 2θ ~ 14.38°, indicating that the liquid exfoliation of few layers did not change the MoS2 structure. Since the peak position in XRD is not changing for both structures (Fig. S3), there is no significant strain induced when MoS2 is incorporated into the phosphate matrix. However, a reduction in the full width at half maximum (FWHM) has been observed (Fig. S3c and Table S1) and this could be due to the variation in the microstructure, the grain distortion, and dislocation density of the crystal30,31. The possibility that higher crystallinity may have reduce the FWHM of MoS2 (because of the heating process) was excluded by performing a controlled experiment in MoS2 treated under the same conditions used for the fabrication of AgPO3:MoS2. No changes in the peak position as well as in the FWHM were observed (Fig. S3b).

In addition, the Raman spectra (Fig. 3) of AgPO3:MoS2 composite glass were obtained and compared with that of a bare MoS2. The Raman spectra of bare MoS2 depict two characteristic peaks at 382.83 and 407.12 cm−1, corresponding to the in-plane (\({\text{E}}_{{2{\text{g}}}}^{1}\)) and out of plane (\({\text{A}}_{{1{\text{g}}}}\)) vibrational modes. The Raman frequency difference (Δω = ω(\({\text{A}}_{{1{\text{g}}}}\))—ω(\({\text{E}}_{2g}^{1}\))) between these two modes dependents on the number of layers, which is used to determine the MoS2 thickness32,33; this difference (Δω) is measured to be 24–25 cm−1, indicating that the MoS2 flakes have several layers32,33,34. This number is found to be similar for all the samples studied (Fig. S4). Besides this, the full width at half maximum (FWHM) of \({\text{E}}_{2g}^{1}\) is ~ 5.67 cm−1, suggests good crystallinity of the exfoliated MoS235,36. Furthermore, a small red shift of about 2 cm−1 in both Raman modes has been observed in AgPO3:MoS236. This shift is unlikely to be due to strain since it is the same for both modes and not only for the in-plane one (a signature of induced strain in the system). The inset of Fig. 3 also shows the obtained Raman spectrum of the pristine AgPO3 glass, i.e. prior to any MoS2 incorporation. The spectrum of AgPO3 glass exhibits a major band at around 1142 cm−1, whereas a broader band at ~ 675 cm−1 is also present. The first Raman signature is attributed to the symmetric stretching vibration of terminal PO2− groups, vs(PO2−), while the latter features originates from the symmetric stretching movement of P-O-P bridges within the phosphate backbone, vs(P-O-P)37,38.

Figure 3
figure 3

Raman spectra of (a) AgPO3 glass, (b) MoS2 flakes on Si (Si/MoS2), (c) MoS2 on AgPO3 (AgPO3/MoS2) and (d) MoS2 embedded into AgPO3 (AgPO3:MoS2). The inset presents the Raman spectrum of the pristine AgPO3 glass in a wide range.

Full size image

µ-photoluminescence (µ-PL) spectroscopy was employed to investigate the emission properties of MoS2 flakes embedded into the AgPO3 matrix. Figure 4a presents the steady state PL spectra of all the samples. The red, black, and blue curves correspond to the PL spectra of AgPO3, Si/MoS2, and AgPO3:MoS2, respectively. As a reference, we first measured the intrinsic PL spectra of MoS2 flakes deposited on Si substrates, using excitation energy of 2.28 eV (543 nm); considering that the direct excitonic transition is weakened with increasing the layer number in MoS2, a broad emission peak with weak intensity was observed (Fig. S4)18,39. It is notable that the MoS2 spectrum has been dramatically changed upon its incorporation in AgPO3. Indeed, the spectrum exhibited two well-defined emission peaks at 1.89 and 2.05 eV, corresponding to A- and B- excitonic transition of MoS2, respectively. At the same time the PL emission is significantly enhanced, corresponding to a 5- and 6-fold enhancement in the A- and B-exciton peak intensities (Fig. 4b). It should be noted that the PL spectrum of AgPO3 glass, presented in Fig. 4a shows no emission within the MoS2 excitonic emission range. Notably, besides the enhancement in the PL intensity, and contrary to the conventional PL properties of few-layered MoS2, a dominant B-excitonic peak is observed in the AgPO3:MoS2 emission spectra (Fig. 4b). The corresponding intensity ratio of B- and A-excitons (IB/IA) equals to 1.2.

Figure 4
figure 4

Room temperature spectroscopic characteristics. (a) PL emission spectrum of Si/MoS2, AgPO3, and AgPO3:MoS2. (b) PL intensity ratio of AgPO3:MoS2 and Si/MoS2 samples. (c) Absorption spectra of NaPO3, and NaPO3:MoS2 (Inset: Magnified spectra of MoS2 excitonic peaks in NaPO3:MoS2) and (d) PL emission spectrum of NaPO3:MoS2.

Full size image

To understand the effect of AgPO3 matrix on the PL properties, we have investigated the internal structure of AgPO3 by means of TEM microscopy. TEM studies reveal a bimodal size distribution of Ag nanoparticles with dominant average sizes of 8.4 nm and 14.5 nm while a broad size variation was also observed (Fig. S5c). The elemental composition of Ag was confirmed by EDX mapping (Fig. S5c). In this context, the large enhancement of MoS2 PL intensity observed in the AgPO3:MoS2 system can be attributed to the localized surface plasmon effect due to the presence of Ag NPs. In order to provide concrete evidence that the observed enhancement of AgPO3:MoS2 PL intensity is induced by the presence of surface plasmon of Ag particles, we prepare a similar NaPO3:MoS2 heterojunction, i.e. in which the silver is replaced by sodium, while the phosphate glass network remains unchanged.

Figure 4c shows that the steady state absorption spectra of NaPO3:MoS2 glass exhibits the two characteristic features at 1.84 and 2.03 eV, which are attributed to intrinsic A- and B-excitonic peaks of MoS2, respectively. Moreover, Raman spectroscopy reveals the presence of a few MoS2 layers within the fabricated NaPO3:MoS2 composite glass (Fig. S6). The NaPO3 glass exhibits its own characteristic vibrational modes at around 1155 cm−1 and ~ 681 cm−1, respectively. Contrary to the case of AgPO3:MoS2, no enhancement of the MoS2 PL is found for the NaPO3:MoS2 system. Indeed, the corresponding room temperature PL spectrum of NaPO3:MoS2 (Fig. 4d), displays only a very broad and extremely weak emission in the range of the direct A- and B-excitonic transitions. It is therefore clear that the remarkable enhancement on the emission properties of the AgPO3:MoS2 is induced by the presence of Ag NPs and their plasmon resonance.

The surface plasmon resonance (\(\omega_{LSPR}\)), which can be tuned by varying the size and shape of the nanostructures and surrounding dielectric medium, is known to strongly modify the excitonic emission25. In particular, the plasmon resonance was tuned by changing the nanostructure size, which enhanced the intrinsically weakly emitting B exciton of a MoS2 flake25. We investigated how the silver plasmon resonance affects the emission properties of the developed MoS2 glass heterojunctions upon changing silver content and particle size in the glass. To this aim, an additional glass-MoS2 heterojunction was fabricated upon employing the ternary silver-rich 0.3AgI–0.7AgPO3 glass instead of the binary AgPO3 glass, i.e. for the development of 0.3AgI–0.7AgPO3:MoS2 architecture. In one of our previous studies it was demonstrated that the incorporation of AgI in the AgPO3 glass results to the agglomeration of silver nanoparticles for the formation of larger silver phases26, while the phosphate network connectivity remains unaffected. Namely, it was reported that for the aforementioned nominal glass composition silver clusters (larger particles formed from the agglomeration of many nanoparticles) with an average size of 2.78 μm are formed while randomly positioned along the glass network (Fig. S7). The absorption spectrum of bare 0.3AgI–0.7AgPO3 exhibited the broad feature of absorption with a hump at ~ 2.47 eV (Fig. S7).

We now consider the effect of these large silver phases on the exciton emission properties of the so-formed 0.3AgI–0.7AgPO3:MoS2 heterojunctions. The measurement conditions were kept identical to these employed for the AgPO3:MoS2 measurements. However, the MoS2 emission spectrum has changed upon its incorporation in 0.3AgI–0.7AgPO3 when compared to AgPO3 (Fig. 5a). Specifically, the PL spectrum exhibits two well defined excitonic emission peaks at 1.75 eV and 1.94 eV, corresponding to A- and B-excitonic transitions of MoS2, respectively. The obtained shift in B-exciton peak position has been appeared to be 110 meV. These excitonic peaks are red shifted when compared to the considerably smaller size Ag NPs of the AgPO3:MoS2 architecture, and their intensities are almost identical. The corresponding intensity ratio of B- and A-excitons (IB/IA) is 0.99 in this case, compared to 1.2 of AgPO3:MoS2. Huang et al.25 have observed similar red-shifts in the peak position for the A and B excitons, due to the nanocavity resonance controlled by the size of plasmonic nanostructures. Thus, the enhancements in the peak intensities and positions of the excitonic transitions are strongly modified by the plasmon nanostructure size in metaphosphate glass.

Figure 5
figure 5

(a) Room temperature PL emission spectrum of 0.3AgI-0.7AgPO3:MoS2; Schematic (b) of the mechanism illustrating the exciton–plasmon coupling (c) in MoS2 and AgPO3 via dipole–dipole interaction.

Full size image

Ultrafast carrier dynamics in MoS2 nanoheterojunctions

Finally, in order to further shed light on the observed enhancement in A- and B-excitonic emissions we investigated the corresponding charge carrier relaxation dynamics by means of ultrafast pump probe time-resolved transient absorption spectroscopy (TAS) (Fig. S8)40,41. Figure 6a presents optical density (ΔOD) vs. wavelength plots at various time delays following photo-excitation of the AgPO3:MoS2 glass using a pump fluence of 2.8 mJ cm−2. Figure 6b shows the photo-bleaching recovery kinetics of A- and B- exciton states at 680 nm (1.82 eV) and 620 nm (2 eV), respectively. In agreement to previous findings42, it is observed that the formation of the A exciton is around 0.5 ps slower when compared to that of the B exciton. In particular, the maximum photo-bleaching of the latter is obtained instantly upon photo-excitation at almost 0 ps. This finding is attributed to the electron–hole cooling time from the upper valence state to the lower valence state within the valence band of the MoS242.

Figure 6
figure 6

Transient absorption study of AgPO3:MoS2 nanoheterojunctions. (a) Optical density (ΔOD) vs. wavelength plots at various time delays following photoexcitation of AgPO3-MoS2 glass at 1026 nm with a pump fluence of 2.8 mJ cm−2. (b) Normalized transient bleach kinetics (symbols) of the A- and B-excitons at 680 and 620 nm, respectively. The solid lines represent the corresponding decay exponential fits.

Full size image

Furthermore, upon following typical exponential fittings, we were able to distinguish the physical mechanisms of A- and B-exciton decay dynamics42,43. For the latter exciton state, the bi-exponential fitting procedure based on the equation y = yo + A1 exp(− x/τ1) + A2 exp(− x/τ2), clearly reveals the presence of two distinct times. Namely, an ultrafast component (τ1) of around 0.5 ps that corresponds to electron transfer from the MoS2 exciton to AgPO3, and a slightly slower time component (τ2) of around 2 ps that is attributed to carrier-carrier interactions42,43. Rather differently, in the case of A-exciton the ultrafast time component is apparently absent, a finding that implies no electron transfer from the lower energy excitation state towards the metallic particles of the hosting glass. The fast charge transfer present only in the B-exciton, explains why the PL enhancement for the B-exciton is only sixfold and comparable to the five-fold observed for the A-exciton (Fig. 4b). There are two effects taking place in the B-exciton during the photoexcitation process (i) a PL enhancement due to the efficient dipole coupling of exciton–plasmon and (ii) a fast charge transfer from the MoS2 to the AgPO3. These effects are antagonistic and lead to the observed enhancement.

Based on the aforementioned results, the plasmon coupling in silver based glasses and MoS2 heterojunction can be facilitated by either electromagnetic field enhancement due to localized surface plasmon (LSP) effect in Ag NPs, and/or via efficient charge injection between the Ag NPs and MoS2 flakes18,44,45. Screening and scattering effects due to the presence of metallic NPs could also slightly influence the PL intensities18. Moreover, heating and strain effects induced by the glass matrix could also contribute to the change in the PL spectrum observed39,46,47. However, such effects should have negligible influence on the PL enhancement in our case due to the inherent indirect band gap18, coupled with large thermal conductivity48,49 of the few-layered MoS2 flakes. It can thus be concluded that the exciton (in MoS2)-plasmon (in Ag) coupling (or LSP) is the most plausible explanation for the observed enhancement in PL intensity in Ag based heterojunctions.

To this date, the investigation of surface plasmon induced PL enhancement is only reported in the case of monolayer TMDs24,50. In particular, it is observed that the exciton–plasmon coupling is greatly influenced by the contact area between the plasmon nanostructure and 2D material9,47,51,52. In our case, it is obvious that the AgPO3 glass comprises plenty of nanoheterojunctions among MoS2 and Ag NPs, than in the 0.3AgI–0.7AgPO3 glass. The AgPO3 glass contained smaller diameter nanostructures than 0.3AgI–0.7AgPO3 (2.78 μm nanocluster), which can significantly enlarge the contact area and the spatial distribution of the localized electromagnetic field. Besides this, the exact modification of the A- and B- excitonic peaks should strongly depend on the nanoheterojunctions cavity resonance (dipole–dipole interaction), which is controlled by nanostructure size24. The steady state photoluminescence enhancement (η) is explained by the change in the quantum yield of the MoS2 in the presence of plasmonic nanostructures. The PL quantum yield Y is defined by53

$$ {\text{Y}} = \left( {\frac{{k_{r} }}{{k_{r} + k_{nr} }}} \right) $$

where, \(k_{r}\) and \(k_{nr}\) is the radiative and the nonradiative decay rates. The radiative decay rate is affected by the localized surface plasmonic fields whereas the nonradiative decay rate depends on plasmonic losses and exciton quenching.

The significant PL enhancement in both Ag—glass based nanoheterojunctions indicates an effective coupling between MoS2 excitons and LSP resonances in nanostructures with large increase in the radiative decay rate. The plasmonic absorption transition dipole moment which is the collective oscillations of the surface electrons in AgPO3 nanostructures that interacts with the transition dipole moments of MoS2 (excitonic states of A and B) leading to collective states (Fig. 5b). Such states are often called hybrid states and result in stronger optical PL than the isolated states of the TMD. Since the plasmonic absorption band of 2.0 eV in AgPO3 (there are two bands, one at 2.0 eV and the other at 2.5 eV) is in the vicinity of B- exciton transition of MoS2 (2 eV) there is a higher probability for B-exciton plasmon dipole–dipole interaction due to the local field enhancement. The physical mechanism behind this process is illustrated in Fig. 5b, c. Altogether, the appearance in discrete and enhanced excitonic emission is led by the exciton (MoS2) and surface plasmon (glass) coupling in the nanoheterojunctions (Glass:MoS2) system. Further work is currently in progress to optimize this coupling via tuning of the Ag NPs size and fraction26 into the AgPO3 matrix.

Conclusion

We have fabricated and demonstrated novel hybrid nanoscale heterojunctions of layered MoS2 and metaphosphate glasses. The MoS2 phase integrity and excitonic bands are preserved inside the glasses. The developed AgPO3:MoS2 composite heterojunctions exhibit a remarkably enhanced PL intensity with the presence of well-defined excitonic transitions. A strong modification of A- and B- exciton peak intensity by plasmonic nanostructure has been adopted. We have obtained a six-fold enhancement factor for the intrinsically weak B exciton peak. Such enhancement factor for the B excitonic emission is explained with the help of dipole–dipole interaction via exciton–plasmon coupling. The ultrafast electron transfer process and carrier-carrier interaction in the nanoheterojunction system support the enhancement in the B excitonic emission. No doubt, the efficient dipole coupling of exciton–plasmon and tunability of B- excitonic emission find application in emerging valleytronic devices working with B excitons. Moreover, the presented fabrication process might be promising for large scale production of inexpensive nanophotonic, valleytronics and optoelectronic devices with tunable B excitonic emissions.

Methods

Sample preparation

MoS2 nanoflakes

MoS2 flakes were prepared from bulk MoS2 powder (grain size < 2 µm, Sigma Aldrich) using liquid phase exfoliation (LPE) method, as reported elsewhere54,55. In detail, 40 mg of bulk MoS2 powder was dissolved in 10 ml IPA. The solution was ultrasonicated for 60 min in a Elma S 30 H bath sonicator (Elma Schmidbauer GmbH, Germany) under 80 W power and 37 kHz frequency. Room temperature (< 30 °C) was maintained throughout the exfoliation process (bath sonicator). After ultrasonication the dispersion was centrifuged to exclude the unexfoliated bulk MoS2. The supernatant of the resulting dispersion was collected and used for subsequent experiments.

AgPO3 glass and AgPO3:MoS2 nanoscale heterojunction formation

The development of AgPO3:MoS2 heterojunction glasses relies on the incorporation of MoS2 flakes within silver metaphosphate glass (AgPO3). First, a previously described procedure was followed for the preparation of the AgPO3 glass substrates37,38. Namely, equimolar amounts of high-purity AgNO3 (99.995%) and NH4H2PO4 (99.999%) dry-powders were melted in a platinum crucible. All weighing and mixing manipulations of the two powders were performed within a glove bag purged with dry nitrogen gas. After thorough mixing of the two powders, the melting batch was transferred to an electrical furnace initially held at 170 °C, while slowly heated up to 290 °C for the smooth removal of the volatile gas products. The furnace temperature was then raised to 450 °C and kept steady for 30 min, while performing frequent stirring in order to ensure melt homogeneity. AgPO3 glasses were obtained in the form of 1 mm thick disk specimens with a diameter of around 10 mm, upon splat-quenching the melt. This well-established procedure results in AgPO3 glasses with negligible water traces of less than 0.3 mol%, i.e. incapable of causing any optical or structural property modifications. Moreover, the glasses remain unaffected of room humidity (25–30%) for several months.

For the incorporation of MoS2, the AgPO3 glass substrate was positioned on a silicon wafer while a heating plate was employed in order to maintain a temperature around 80 °C. Ten drops of a previously prepared MoS2 solution (0.76 mg/ml) were drop-casted on the surface of the AgPO3 glass, while allowing 10 s intervals between each drop in order to ensure smooth solvent vaporization. After solvent removal the residual MoS2 flakes were randomly distributed on the AgPO3 surface. Then, the temperature was raised to 170 °C for 2 min, i.e. 22 °C below the glass transition temperature of the AgPO3 glass. At this temperature, the AgPO3 glass becomes viscous and allows readily the smooth incorporation of the MoS2 flakes within the glass matrix. Following MoS2 immersion, the AgPO3:MoS2 nano-hybrid glass was splat-quenched between two silicon wafers, while instantly removed from the heating plate and left to cool down to room temperature. The employment of silicon wafers allows the formation of smooth surfaces on both sides of the composite glass specimens and renders them suitable for optical characterization. The MoS2 incorporation process is presented in Schematic S1, while the samples used are presented in Table 1.

Table 1 Description of the samples used in this study.
Full size table

Optical measurements

The optical UV–Vis absorption spectra of the dispersion and solid films were carried out with a PerkinElmer, Lamda 950 UV/VIS/NIR spectrometer, USA. The Raman spectra were recorded under 473 nm laser excitation (Thermo Scientific) in the back-scattering geometry at ambient conditions at 300 K. The Si substrate peak at 520 cm−1 was used for calibration purposes.

For optical spectroscopy measurements, we used a Micro-Photoluminescence (μ-PL) setup and the spectra were collected in a backscattering geometry at 300 K. As excitation source was used a continuous wave (CW) He–Ne 543 nm (2.28 eV) laser. An iHR-320 spectrometer (Horiba Scientific/Jobin Yvon Technology) equipped with a Syncerity multichannel charge-coupled device (CCD) Camera was employed to collect the spectra.

For the XRD measurements an X-Ray Rigaku (D/max-2000) diffractometer was employed, while being operated with a continuous scan of Cu Ka1 radiation with λ equal to 1.54056 Å. The morphology of the Ag NPs was studied by transmission electron microscopy (TEM, LaB6 JEOL 2100), after depositing drops of glass-powder/toluene solution onto a carbon-coated TEM grid. Finally, a field emission scanning electron microscope (SEM, JEOL, JSM-7000F) was used for the examination of the lateral dimension of dispersed 2D MoS2 flakes, while atomic force microscopy (AFM) was employed to obtain the flakes’ thickness (Digital Instruments with controller Nanoscope IIIa).