Introduction

Transition metal dichalcogenides (TMDs) have attracted tremendous attention due to their extraordinary optoelectrical properties, setting the stage for new breakthroughs in materials science, especially when the thickness of these materials approaches atomic-level thicknesses1,2,3,4,5,6,7. Molybdenum disulfide (MoS2), as a typical TMD, has experimentally demonstrated high mobility and high on/off ratio when used as an active material in electronics devices, while the circularly polarized photoluminescence (PL) for single layer MoS2 makes it promising for valleytronic devices7,8,9,10. Though single layer MoS2 is a good candidate for its application in electronic devices, the size of samples fabricated by ME method limits its practical use. Therefore, CVD method has been used to synthesize large-area samples although the electrical properties of CVD grown samples are poorer than that of ME samples11,12,13,14,15. However, CVD grown monolayer MoS2 displays abnormally optical quality with a stronger PL intensity compared to that of ME samples, which has been observed in previous reports and attributed to the proof of high crystal quality16,17, since optical performance at room temperature is usually used to evaluate the crystal quality of two dimensional (2D) semiconductors, based on the premise that intense PL emission is observed for low-defective compound semiconductors18. However, this conclusion cannot work here, because the crystal quality of ME samples is usually better than that of CVD grown samples and is expected with a better optical quality. Till now, it is still a controversial issue that the precious explanation for the stronger PL in CVD grown samples than that of ME sample is due to its higher crystal quality16,17,19,20,21,22,23. Thus, the mechanism for the strong PL in CVD grown samples is still lacking and the clarification for this doubt in this work will be desirable and helpful to recognize its nature of stronger PL in CVD grown samples. In this work, by investigating the optical spectra of monolayer MoS2 samples with different amounts of defects, as well as their electrical performance, we uncover the emission mechanism of abnormally strong PL observed in CVD grown samples and attribute it to the p doping effect from the larger amounts of adsorbates on the defect sites of MoS2, which rebalance the radiative emission intensities of excitons.

Results and Discussion

Figure 1a shows the PL spectra of typical CVD grown and ME monolayer MoS2. Their corresponding PL mappings show the uniform intensities. Clearly, the PL intensity of the former is stronger than that of the latter, which was previously attributed to the high crystal quality of CVD grown crystal16,17. The X-ray photoelectron spectroscopies for these two groups of samples are carried out to show the atomic ratio (S/Mo), as shown in Fig. S1. For CVD grown sample, the peaks at 162.3 and 163.4 eV are assigned to the S2− 2p3/2 and 2p1/2, respectively. The peaks at 229.5 and 232.7 eV are attributed to the Mo4+ 3d5/2 and 3d3/2. While for ME sample, the peaks at 162.6 and 163.6 eV are assigned to the 2p3/2 and 2p1/2, of divalent sulphide ions (S2−) and the peaks at 229.67 and 232.8 eV are from the 3d5/2 and 3d3/2 of the core levels of Mo4+. The S/Mo ratio for CVD sample is 1.66 and that for ME sample is 1.86, which is consistent to the results of EDS and more defects are introduced in CVD grown samples. Figure 1b,c show typical high resolution atomic structure generated by high angle annular dark field (HAADF) scanning transmission electron microscopy (STEM), and they clearly display the defects in CVD grown monolayer MoS2, as indicated in the lines. The lattice spaces of hexagonal lattice structure for MoS2 is 2.7 Å ((100) plane) and 1.6 Å ((110) plane). The corresponding intensity line profiles shows in Fig. S2 display most of defects are single S vacancy. Fig. S3a gives a directly visual sense of the effect of defects on the PL intensities of ME MoS2. The marked circle area with pre-exposed by Argon (Ar) plasma shows stronger PL intensity. Fig. S3b shows the PL spectra of area A (exposed by Ar plasma) and B (without Ar plasma treatment), respectively. Therefore, it seems that the more defects contribute to the stronger PL emission in MoS2. Nevertheless, whether the PL spectra observed in MoS2 is its intrinsic emission or defect-induced emission remains unclear. The previous reports have revealed the energy of defect-induced PL is around 1.78 eV, at the low energy side of the reported A exciton22. However, as shown in Fig. 1, there is no extra emission peak except a broadened shoulder, which has been identified as the trion emission, marked as A− exciton emission. Moreover, the temperature dependent PL spectra (shown in Fig. S4) indicates the PL peak observed is an intrinsic emission, and the match among the electron spin, the layer pseudospin and the valley pseudospin reported in the literature display the features of intrinsic emission6,7,8,9,24. Thus the observed PL peak in Fig. 1 is not the defect-induced emission in MoS2. Another potential factor that may lead to the difference of optical properties between CVD grown and ME samples is the strain effect from the substrate since the contact behaviours between the substrate and the sample are different for these two cases. The CVD grown samples are inevitably affected by the strain effect because of the different thermal expansion coefficients between deposited layer and substrate during the cooling process25,26,27. Figure S5 compares the Raman spectra of both of CVD grown and ME sample. The \({{\rm{E}}}_{2{\rm{g}}}^{1}\) and A1g modes of CVD grown samples show a red-shift and blue-shift, respectively, when compared with those of ME 1 L MoS2 sample. This phenomenon differs from the variation trend when strain is applied28. Also, if the strain effect is the main factor in the PL spectra of MoS2, the energies of PL would shift clearly towards higher energy (compressive strain) and lower energy (tensile strain), respectively29,30, but there is no obvious shift of peak position shown in Fig. 1. Thus, the strain effect could be excluded in this study. So far, defects are most likely to be responsible for the enhancement of photon emission in CVD grown monolayer MoS2.

Figure 1
figure 1

(a) PL spectra of single layer ME and CVD grown MoS2. Left column: optical image of a typical monolayer sample fabricated by ME method and its PL mapping. Right column: optical image of a typical CVD-grown monolayer sample and its PL mapping. Both of two PL mappings are normalized to the intensity of A exciton. (b) and (c) are the HAADF image of ME and CVD samples, respectively. The lines indicate the defects observed in this image.

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Since defects in atomically layered materials usually could act as adsorption centres due to their higher active adsorption energy, the PL spectra of CVD and ME MoS2 sample were measured in vacuum to exclude the role of adsorbates. Surprisingly, as shown in Fig. 2a, the PL intensity of CVD grown MoS2 decreased dramatically in vacuum. At the same time, the emission intensities are reversible when the CVD grown samples exposed in air again. However, for the PL spectra of ME sample (Fig. 2b), its intensity does not show obvious variations with the ambient varies. The reduction of PL intensity in vacuum for CVD grown sample implies that the adsorbates adsorbed on the surface of MoS2 greatly affect its photon emission. According to the vacuum-level dependent PL intensities, it seems that more adsorbates contribute to stronger PL intensity. Moreover, it is found that in Fig. 2c, the ratio of PL intensities between CVD grown and ME MoS2 samples is even less than 1 when the adsorbates level is at the lowest vacuum level (1 × 10−5 mbar). This result shows that the PL intensity of CVD grown sample is comparable or even weaker than that of ME sample, which verifies that the room-temperature optical performance in vacuum could reflect the crystal quality. Figure 2d,e show the PL images of a typical CVD grown sample with grain boundary in air and vacuum, respectively. The sites at the grain boundary are expected with more defects and its decreased PL intensity measured in vacuum provides clear evidence that PL intensities at defect sites could be decreased when measured in vacuum.

Figure 2
figure 2

Vacuum level dependent PL spectra of (a) CVD grown and (b) ME MoS2 samples. (c) The ratios of PL intensities between CVD grown and ME MoS2 samples at different vacuum levels shown in Figure (a,b). The PL mappings of a typical CVD-grown monolayer MoS2 measured in air (d) and vacuum (e), respectively. Both of PL mappings are normalized to the intensity of A exciton.

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Meanwhile, the in-situ Raman spectra have been obtained for these two groups of MoS2 samples, as shown in Fig. S6. There is no obvious change for the E1 2g mode (almost fixed at ~385.7 cm−1 for ME sample and ~384.4 for CVD grown sample), but a clear blue-shift of A1g mode (from 404.94 cm−1 to 403.35 cm−1) for the CVD sample in vacuum compared with that in air, shown in Fig. S6a, which is larger than that of ME sample with a smaller shift from 404.21 cm−1 to 403.71 cm−1 (Fig. S6b). This result is different from the reported strain effect on the Raman spectra for monolayer MoS2 sample28, which could be further exclude the strain effect. The similar variation trends for A1g and E1 2g mode have been reported by the gated Raman spectra from a 1 L MoS2 31 . The results show the softening of the A1g mode with electron doping, while the frequency of E1 2g mode remains unchanged, which are due to the stronger electron-phonon coupling of the A1g mode than that of the E1 2g mode32,33. So whether the shift of A1g modeobserved in this work is due to the doping effect will be discussed later.

The vacuum-level dependent electrical performances of CVD grown and ME MoS2 are investigated. Figure 3a shows the schematic of a monolayer MoS2 based transistor fabricated in this work. The fabrication process is described in the experimental section. The optical images of FETs based on ME and CVD grown samples are shown in Fig. S7. Figure 3b,c show the electrical performance of single layer ME and CVD grown MoS2 based transistors, respectively. Both of these two groups transistors display n type behaviour, which mean excess electrons introduced in MoS2. Moreover, with the vacuum level decreasing, the threshold voltage shifts towards the positive direction, which means that the adsorbates in air play a p doping effect on MoS2 itself. Additionally, comparing the current at Vg = 0 V, the current is larger under vacuum when compared with that in air, which represents a stronger n doped effect in vacuum. While in air, due to the p doping effect from absorbates, the current is decreased. Both of CVD grown and ME samples show this variation trend. The mobility for CVD grown samples in vacuum is 0.69 cm2/VS while it is 44.2 cm2/VS for ME samples in vacuum, which means a poor crystal quality for CVD grown samples without doping effect from adsorbates. While this p-doping effect due to the adsorbates in air reduced the amounts of the excess electron in MoS2 and thus increased the radiative decay of exciton emission, which could be described by a three-level model including the excitation, exciton emission and trion emission processes4,34. As shown in Fig. S8a, G represents the generation rate of optical excitons. The radiative decay rates of the exciton and trion are marked as \({\Gamma }_{ex}\) and \({\Gamma }_{tr}\) respectively. When the mass action law is considered with trions together to evaluate the doped electron density in MoS2 samples, the integrated intensity ratio of trion emission (I tr ) to electron emission (I ex ) can be expressed as34,35,36,37:

$$\frac{{I}_{tr}}{{I}_{ex}}=\propto \frac{{n}_{e}}{{k}_{b}T}$$
(1)

I tr and I ex represent the trion and exciton emission intensities, respectively. n e is the carrier concentration, k b is the Boltzmann constant, and T represents temperature.

Figure 3
figure 3

(a) The schematic of transistors used in this work. Id -Vg curves of (b) CVD grown and (c) ME MoS2 transistors for Vbg = −20 V to 20 V measured in air, vacuum, and back to air, respectively. The fixed source-drain voltage is 1 V. The black, blue and red colour lines represent the curves measured in vacuum, air and back to air, respectively.

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From equation (1), it is confirmed that the ratio of emission intensities from I tr and I ex indeed highly depends on the carrier concentration in MoS2, which corresponds to the experimental results shown in Fig. S8b. The fitting results for CVD grown and ME samples in air and vacuum are shown in Fig. S9, respectively, which directly reflect the vacuum level dependent intensity ratio of exciton and trion. Moreover, it is concluded that the binding energy of trion is ~24 meV. So their emission spectra overlapped with that of exciton emission since such negatively charged excitons usually have finite binding energies, ~20 meV4.

Moreover, such a p doing effect also explain well the Raman frequency shift of E1 2g and A1g mode observed in Fig. S6. According to the symmetrical group theory, the A1g mode has symmetrical lattice variation and would have a nonzero expectation value for the matrix element of the electron-phonon coupling, which leads to a larger electron-phonon coupling31. While for the E1 2g mode, the matric element of electron phonon coupling vanishes and its coupling with electrons is weaker on doping in ambient compared with A1g modes31,32,33. Moreover, we have calculated the phonon energies of E1 2g mode and A1g mode with and without considering the adsorbates, as shown in Fig. S10, the results are consistent to our experimental results and explanation above.

So far, the physical picture for the reason why the CVD grown sample shows a strong PL intensity is clear. The greater the amounts of adsorbates, the stronger the intensity of the PL emission in MoS2. Supposing there is no adsorption on the surface of MoS2, the PL intensity from CVD grown MoS2 should be weaker than that of ME sample due to its poorer crystal quality, which is corresponding to the experimental results on the intensity ratio (0.74) at higher vacuum shown in Fig. 2c.

Furthermore, the CVD grown monolayer MoS2 under different growth conditions marked as S1, S2, S3, S4, and S5, respectively, are investigated to confirm the vacuum level dependent PL spectra. Figure 4a shows the optical images of selected CVD samples. Their PL intensities are different in air (Fig. 4b), and their emission intensities are reduced measured in vacuum (1 × 10−5 mbar) by different extent, as shown in Fig. 4c. Figure 4d presents the ratios of PL intensities measured in air and vacuum for these samples. The different ratios when the vacuum level varies during measurement are due to their different amounts of adsorbates on the samples. The ratios of exciton and trion emissions for these five samples are also shown in Fig. 4d and the variation tread is almost the same as their vacuum-level dependent PL intensities. Additionally, compared with that in air, the ratios of exciton and trion emissions are decreased in vacuum which further confirms that the amounts of adsorbates are highly related to the PL intensity and the ratios of exciton and trion emission intensities in MoS2 and the optical performance of MoS2 in vacuum are positively related to its crystal quality.

Figure 4
figure 4

(a) Optical photos of CVD grown monolayer samples selected for PL measurement in (b), marked as S1, S2, S3, S4, S5. The scale bars shown in figures represent 15μm. (b) The corresponding PL spectra of CVD samples shown in (a). (c) The PL intensities of samples, shown in Figure (b), in air and in vacuum, respectively. The inset is the enlarged figure for the PL intensities in vacuum. (d) The ratios of PL intensities in air and vacuum for these five samples, and the intensity ratios of exciton and trion emissions for these five samples in air. The insert figure shows the intensity ratios of exciton and trion emission in vacuum.

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To clarify the most likely components of adsorbates in air and how the adsorbates contribute to the p doping effect on MoS2, we have evaluated the charge transfer between adsorbates (O2, N2, OH− and H2O molecules) and MoS2 by performing first-principle calculations as these three kinds of molecules are bountiful in air. We first consider the defects in MoS2. As shown in Fig. S11, we consider four typical types of defects by using a 4 × 4 supercell of MoS2 as the prototype: (1) a single S vacancy; (2) a single Mo vacancy; and (3) and (4) for two S vacancies at the same side and different sides. Our results show the single S vacancy is the most common defect with the lowest formation energies. Then we turn to focus on the charge transfer process between adsorbates and perfect (defective) MoS2. The binding energy (Eb) of adsorbates A (O2, N2, OH− and H2O) on pristine or defective monolayer MoS2 is given by

$${{\rm{E}}}_{{\rm{b}}}={\rm{E}}({{\rm{MoS}}}_{2}+{\rm{A}})-[{\rm{E}}({{\rm{MoS}}}_{{\rm{2}}})+{\rm{E}}({\rm{A}})]\,({\rm{A}}={{\rm{O}}}_{2},{{\rm{N}}}_{2},{{\rm{H}}}_{2}{\rm{O}})$$
(2)

All energies are calculated within the same supercell for comparison. For the pristine MoS2, three adsorption sites for O2, N2, OH− and H2O molecules are schematically shown in Fig. 5a,b, and their binding energies with MoS2 are shown in Fig. 5c. The most favourable adsorption configurations for O2, N2 and OH− are at B site, while for H2O it is at A cite, which are in good accordance with former calculations38. For the defective MoS2, we compare the adsorption energies of O2, N2, OH− and H2O molecules on its surface. Figure 5d shows O2 molecule possesses the lowest adsorption energy (−3.25 eV), indicating O2 molecules are the most likely component to be adsorbed on defective MoS2. The O-H bond has a strong adsorption energy because it can form covalent bond on the Mo-edge39, but the sample used in this work is S terminated40. Therefore, we only discuss the adsorption of O2 molecules on the defect-sites of MoS2. Figure 5e,f show the amounts of charge transfer between perfect (defective) MoS2 and O2, which are 0.002e and 1.025e, respectively. Obviously, the charge transfers can be greatly enhanced with defects introduced in MoS2 and charges are extracted from defective MoS2 by O2 molecules due to its strong electronegativity. Thus, O2 molecules play a major role in the p-doping effect for MoS2. Moreover, to further confirm the doping effect of O2 molecules, we anneal the samples at 400 °C in one hour with different oxygen atmosphere (20, 30, 40, 50 Torr) and find the introduction of O2 indeed enhance the PL intensity of CVD grown monolayer MoS2 samples, as shown in Fig. 6. The simulated charge transfer process is highly consistent with the experimental results. The more defects in CVD samples which act as the localization centre could form the stronger localized exciton and increase the radiative recombination of exciton of MoS2.

Figure 5
figure 5

(a) Top and (b) side views of optimized structure of monolayer MoS2. A, B, C represent the different adsorption sites of adsorbates. (c) The adsorption energies of O2, N2, H2O and OH− on A, B, C sites of pristine MoS2. (d) The adsorption energies of O2, N2, H2O and OH− on defective MoS2 with a S vacancy. (e) and (f) Show the charge density difference of O2 molecule absorbed on pristine and defective 1 L MoS2 with a single S vacancy, respectively. The positive and negative charges are shown in yellow and blue colors, respectively. Isosurface values are 7.5 × 10−2 e/Å3 for (e) and 5 × 10−3 e/Å3 (f).

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Figure 6
figure 6

The PL spectra of annealed MoS2 under different oxygen atmosphere. The arrow represents the increase of oxygen amounts. The purple line represents the PL spectra of pristine sample.

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Conclusion

In summary, we clarify the origin of the abnormally strong PL in CVD grown samples and attribute it to the p doping effects of adsorbates at the defect-sites. HRTEM and EDS are used to confirm more defects in CVD grown samples than that in ME samples, and the in-situ vacuum level dependent PL and Raman spectra, as well as the electrical performance confirms our proposal. The CVD samples under different grown conditions are expected with different crystal qualities, which show similar variation trends by different extent which further confirm the observed phenomenon. First principle calculations were carried out to investigate the charge transfer process between MoS2 and adsorbates and clarify the p-doping effect of adsorbates due to the strong electronegativity of adsorbates. Such a p-doping effect from adsorbates reduces the concentration of excess electron in MoS2 and contributes to the radiative recombination process of exciton. This work proposes a new vacuum technique to evaluate the crystal quality and can guide the future works on the engineering of PL emission intensity on monolayer two-dimensional layered semiconductor materials.

Methods

Samples Preparation and Optical Characterization

Two groups of MoS2 samples were used in this experiment: (1) monolayer MoS2 exfoliated from nature MoS2 crystals (from SPI), and (2) CVD grown monolayer MoS2 (sulphur and MoO3 powders are placed in furnace, with a SiO2/Si substrate located face down above the MoO3 powder). The growth temperature is at 650 °C and this temperature is kept with 15 min. Raman and PL spectra were conducted on Witec CRM 300 confocal Raman microscopy. The excitation laser was 532 nm laser with a spot size about 500 nm. The output power measured from the objective lens was controlled below 0.5 mW to avoid damaging and heating on samples. The accumulation time for PL and Raman spectra was 10 s. The Raman spectra were collected by 1800 g/mm grating while PL spectra were collected by 300 g/mm grating.

The size of the probe X-ray beam is 200 μm. For sample viewing, firstly, an optical image is taken by an external camera. Based on that optical image, we zoom in and position to roughly the area that we want. Then, similar to a SEM/EDX, a finely focused x-ray beam is used to create a secondary electron image for sample viewing, and actual positioning for interested analysis area.

Fabrication of Back-Gated MoS2 Transistors

Both of these two groups of MoS2 samples were deposited on 270 nm SiO2/Si substrates. The devices were fabricated by standard electron beam lithography, followed by the thermal evaporation of Cr/Au electrode, and lift-off process. The electric measurement system was carried out using Agilent B1500A semiconductor analyzer. The vacuum level is set at 1 × 10−5 Torr.

First Principle Calculations

The first-principles calculations were realized based on the density functional theory (DFT), as implemented in the Vienna Ab initio Simulation Package (VASP)41. The exchange-correlation potential is chosen as generalized gradient approximation (GGA), formulated by Perdew-Burke-Ernzerhof (PBE) functional42,43. The cut off energy of 400 eV is used for the plane-wave expansion of valence electron-wave functions. To avoid artificial interaction between layers, a vacuum spacing of >15 Å is built. In the calculations, DFT-D2 method is employed to describe the long-range van der Waals interactions44. For both defects identifications and gas molecules adsorptions on the monolayer MoS2, the calculations were performed on a 4 × 4 × 1 size of MoS2 supercell, with Monkhorst-Pack k-point meshes of 7 × 7 × 145. The convergence criteria for energy and force are set to be 10−5 eV and 0.01 eV Å−1, respectively. The amount of charge transfers between gas molecules and MoS2 monolayer is estimated by using the Bader charge method.