Figure 1 shows low-temperature photoluminescence (PL) spectra of monolayer MoS2 on hexagonal boron nitride (hBN). In addition to exciton (0XA) and trion (-XA),11 as-exfoliated MoS2 exhibits a prominent L-band from approximately 1.5 eV to 1.9 eV (Fig. 1a). Mild annealing (Tannealing = 500 K) in vacuum results in a striking reduction of the L-band (Fig. 1b), presumably due to desorption of adsorbates. Further annealing at high temperature (Tannealing = 800 K) introduces a narrow peak XL at 1.75 eV (Fig. 1c). A similar, yet even sharper, spectral signature is observed in fully encapsulated MoS2 after He-ion irradiation (Fig. 1d). The improved inhomogeneous broadening agrees with previous studies of fully hBN encapsulated heterostructures23. In the following, we show that adsorbates introduce a continuum of defect states, which is responsible for the L-band emission, whereas pristine sulfur vacancies introduce a deep center, which is very likely also the origin of recently discovered single photon emission in He-ion treated MoS2.3,24
Figure 2 shows low-temperature PL of MoS2 after stepwise in-vacuo annealing up to 800 K. In each cycle, the samples were rapidly annealed in a customized cryostat for 30 min and then cooled back to cryogenic temperature (Tsample ~ 20 K) for PL characterization maintaining a high vacuum (p < 10-4 mbar) at all times (Supplementary Information S1). Figure 2a illustrates spectra of as-exfoliated MoS2 on hBN and after several mild annealing steps to 420 K, 450 K, and 510 K. Again, the as-exfoliated flake exhibits a prominent L-band (cf. Fig. 1a). The intensity of the L-band decreases relative to the intensity of the free exciton emission by one order of magnitude after annealing to Tannealing = 420 K, and it gradually disappears for higher annealing temperatures (Fig. 2b) . Furthermore, the trion emission decreases initially compared to the free exciton indicating a reduced doping, as observed previously in TMDs during desorption of physisorbed gas25 and chemical dopants26. Hence, we attribute the L-band to adsorbates, which are progressively removed during the mild annealing steps. The desorption does not follow a simple Arrhenius law, since it depends on the total number of adsorbates, which is unknown. Therefore, we can estimate only an upper bound of ~100 meV for the desorption barrier, which agrees with ab-initio studies for molecular adsorbates on MoS227 and temperature programmed desorption on bulk MoS228.
At Tannealing = 510 K, a spectrally narrow emission line XL appears around 1.75 eV. In contrast to the L-band, the intensity of XL increases with higher annealing temperatures, until the whole PL signal disappears at Tannealing > 700 K (see Supplementary Information S2). To extract the thermal activation barrier of XL, we use an hBN/MoS2/hBN heterostack, where the L-band is already suppressed in the as-prepared structures. As seen in Fig. 2c, XL brightens in the encapsulated monolayer with increasing annealing temperature, and further narrows after annealing to 800 K indicating a complete removal of residual adsorbates. At even higher temperatures (Tannealing = 900 K), the intensity of XL decreases drastically, followed by the complete disappearance of the overall PL (Supplementary Information S2). Consistent with saturating a finite density of localized defect levels, XL exhibits a saturating behavior as function of excitation power (see Supplementary Information S3)24. In a simple rate equation model, the saturated defect emission is then proportional to the total number of emission centers. In this case, we can readily determine the energy barriers for defect generation since it is proportional to the difference in the integrated spectral weight of XL compared to the previous annealing step. For example, the number of defects generated during annealing at 600 K is proportional to the difference in integrated PL intensity, which we label ΔInt(XL), after annealing to 600 K and 510 K. In the limit of low density, defect generation is independent of the number of existing defects, and we expect a simple Arrhenius law. From Fig. 2d, we find an activation energy of (0.71 ± 0.13) eV for the XL-peak consistent with theoretical predictions for the formation energy of mono-sulfur vacancies (approximately 1 eV) 29,30. Interstitial sulfur defects should not form under high-vacuum, i.e. sulfur-poor, conditions29. The formation energies for transition metal vacancies are much larger (3 eV – 8 eV) 29,30. Consequently, the formation of sulfur vacancies during annealing is thermodynamically the most favorable and, therefore, most likely process.
We continue to corroborate the dominant generation of sulfur vacancies during annealing by atomic-scale characterization. Here, we perform high-resolution low-temperature scanning tunneling microscopy (STM) and atomic force microscopy (AFM) of single-layer MoS2 before and after high-temperature annealing in vacuum (Tannealing > 500 K) as well as before and after He-ion irradiation. These experiments are conducted on graphene/SiC heterostructures. The graphene substrate is essential as conductive support, but it quenches. the optical emission from defects (Supplementary Information S7)31. For STM, all samples were prepared in-vacuo by a mild annealing step in vacuum (Tannealing < 500 K) to remove adsorbates. In agreement with our optical studies, the surface of MoS2 is virtually free from adsorbates after mild annealing. By far the most dominant defects in pristine material were oxygen atoms substituting sulfur, but no sulfur vacancies were observed (Supplementary Information S5)9,19,20.
After additional high-temperature annealing, we observed only two additional types of defects within our experimental statistics (Fig. 3). Their high-resolution STM topography exhibits a trigonal symmetry (Figure 3a and 3b), consistent with a vacancy either in the top or bottom sulfur lattice. For both vacancies, the charge densities calculated by DFT (Supporting Information S7) at a distance of 4.5 Å above the MoS2 layer (similar to experimental conditions in STM) are shown in Figs. 3c and 3d. The calculations are in good agreement with the STM topography. The top vacancy appears trigonally ring-shaped, whereas the bottom vacancy appears triangular-shaped with bright maxima at the corners. These same defects, i.e. top and bottom sulfur vacancies, were also confirmed in the He-ion irradiated samples (Fig. 3e and Fig. 3f). For He-ion treated samples, sulfur vacancies are the dominant defect type, among other defects that are generated with lower yield9. Furthermore, we performed CO-tip AFM on the sulfur vacancies (Figure 3g and Fig. 3h). For the top vacancy, we observe an apparent depression at the sulfur site, whereas for the bottom vacancy structural relaxation results in a slight protrusion, in agreement with previous studies on WS220. Most importantly, AFM conclusively assigns the defect onto the sulfur sublattice, which is difficult from STM alone20. Overall, the scanning probe experiments confirm the composition and surface condition of MoS2 derived from optical characterization (cf. Fig. 1). From our different annealing experiments, we conclude that defect luminescence XL at 1.75 eV arises from pristine, i.e. undecorated, sulfur vacancies. We note that, there are at least two pathways for formation of sulfur vacancies, which are desorption of a sulfur atom or desorption of an oxygen atom from a sulfur site. The latter substitute for sulfur in as-prepared TMDs 19,20. Based on the similarity of the optical spectra and the abundance of sulfur vacancies in thermally annealed as well as in He-ion irradiated MoS2 (cf. Figure 1), we propose that also the origin of quantum emission from individual He-ion induced defects is due to the (non-passivated) sulfur vacancy 24.
From ab-initio calculations of defective MoS2 (5×5 supercell corresponding to 2 % of vacancies), three types of excitonic transitions are qualitatively distinguished, as shown in Fig. 4a. Transitions at the optical gap (XA) arise mainly from unbound electronic states between valence and conduction band at K and K’ valleys. The next series of transitions (D2) occur predominantly between the resonant defect state overlapping with the valence band and the localized in-gap states. In the dilute limit, the localized defect states are k-independent, such that defect-defect transitions cannot exhibit valley-selectivity12. The lowest series of transitions (D1) couples the localized in-gap state to dispersive states in the valence band near K and K’, and they are predicted to show valley selectivity. Overall, the strong electron-hole interaction leads to a vast manifold of excitonic transitions for the defective crystal with varying eigen energies and mixed degrees of pristine and defect-like character, and in general many different excitonic states contribute to the absorbance spectrum12,13. Figure 4b compares the absorbance spectrum calculated within the GW plus Bethe-Salpeter equation (GW-BSE) approach32,33 and experimental photoluminescence of a MoS2 monolayer with sulfur vacancies. The calculated sub-gap resonances located about 0.45 eV and 0.3 eV below the free exciton transition (0XA) can be associated with transitions of type D1 and D2, respectively. We note that the absorbance includes a phenomenological broadening of 0.1 eV, which accounts for the numerical uncertainty of the used approach12. Furthermore, the localization at the atomic scale defect results in a large spread of the exciton transitions in k-space, which further broadens the absorption resonances12.
For the experimental curves, vacancies were introduced in fully encapsulated MoS2 both by in-vacuo thermal annealing and by ex-situ helium ion modification. All spectra are referenced to the neutral exciton transition, where we found 1.89 eV for the calculated spectrum and 1.96 eV (1.94 eV) for the measured annealed (ion-modified) spectrum. Notably, while the dominant defect emission (XL) occurs about 0.2 eV below 0XA, we consistently observe weak emission features (red arrows) at even lower energies3. The multiple sub-gap emission peaks are qualitatively in agreement with the ab-initio calculations of excitonic defect states (top panel of Fig. 4b). Although absorbance cannot strictly be used to infer emission properties, we identify possible transitions based on the correspondence between calculated absorbance and experimental emission spectrum. Within the computational uncertainty, the defect emission at 1.75 eV agrees with the energy range where calculations predict dominant contributions from defect to defect transitions (D2). Figure 4c further corroborates this assignment. Here, we plot the emission spectrum of a single defect, which was generated by He-ion bombardment9,24, for co- and cross-circularly polarized excitation and detection. We do not detect a valley polarization, as expected for a transition that occurs between localized defect levels. Therefore, we conclude that the XL peak observed in our thermally annealed as well as He-ion treated MoS2 monolayers arises due to a localized excitonic transition between the defect orbitals of the pristine sulfur vacancy.
Typically, the dominant emission process should involve the lowest energy state of the system, i.e. transitions of type D1. However, in our experiments, the defect emission is governed by transitions of intermediate energy, i.e. transitions of type D2, although the calculated oscillator strength varies only weakly in the relevant regime. A naïve scenario to explain our observations involves a relaxation cascade after the absorption process: optical excitation creates a free exciton, which gets localized, and then both hole and electron decay into a defect state (type D2). If further relaxation of the captured exciton into an excitonic state of type D1 is slow or prohibited, the emission will occur dominantly from the fully localized electron and hole state. From a theory point of view, the above picture demands to include further interactions, such as exciton-exciton or exciton-phonon coupling34.
In summary, by combining far-field optical spectroscopy, atomic-resolution scanning probe microscopy, and ab-initio theory, our study provides compelling evidence of optical defect emission from pristine sulfur vacancies in single layer MoS2. In contrast to previous studies, these pristine sulfur vacancies are generated in-vacuo or capped by hBN, and therefore, neither passivated by oxygen nor decorated with adsorbates. Similar to previous reports, we observe a broad L-band luminescence due to adsorbates in as-prepared MoS2 monolayers, which can be suppressed by a combination of h-BN encapsulation and mild annealing. In as-prepared layers and after mild annealing, pristine sulfur vacancies are absent, and oxygen passivated vacancies are the dominant defect. We suggest that oxygen-passivated vacancies form active sites for adsorption of molecular species, since many previous studies established a positive correlation between sulfur deficiency and defect emission4,5,15. Pristine vacancies are created in h-BN/MoS2/h-BN heterostructures either via in-vacuo thermal annealing or ex-situ helium ion bombardment, whereby the latter allows generating single photon emitters on demand24 with a position accuracy below 10 nm9. Guided by ab-initio calculations, we identify transitions between a localized in-gap defect state and a localized resonant defect state as the most likely candidate.