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Towards large area and continuous MoS2 atomic layers via vapor-phase growth: Thermal vapor sulfurization Hongfei Liu,1,* K. K. Ansah Antwi,1,2 Chengguo Li,2 Jifeng Ying,3 Soojin Chua,1,2 and Dongzhi Chi1

1 Institute of Materials Research and Engineering (IMRE), A*STAR (Agency for Science, Technology and Research), 3 Research Link, Singapore 117602, Singapore

2 Department of Electrical and Computer Engineering, National University of Singapore, 4 Engineering Drive 3, Singapore 117576, Singapore

3 Data Storage Institute (DSI), A*STAR (Agency for Science, Technology and Research), 5 Engineering Drive I, Singapore 117608, Singapore

Abstract

We report on the effects of the substrate, starting material, and temperature on the growth

of MoS2 atomic layers by thermal vapor sulfurization in a tube-furnace system. With Mo

as the staring material, atomic layers of MoS2 are obtained on sapphire substrates while

MoS2 nanoparticles embedded in SiO2 amorphous are obtained on native-SiO2/Si

substrates under the same sulfurizing conditions. An anomalous thickness-dependent

Raman shift (A1g) of the MoS2 atomic layers is observed in Mo-sulfurizations on sapphire

substrates, which can be attributed to the competition between the effect of thickness and

that of surface/interface. Both effects vary with the sulfurizing temperatures for a certain

initial Mo thickness. This anomalous frequency trance of A1g is missing when using

MoO3, instead of Mo, as the starting material. In this case, the lateral growth of MoS2 on

sapphire substrates is largely improved; meanwhile, the area density of the resultant

MoS2 atomic layers is significantly increased by increasing the deposition temperature of

the starting MoO3. The thickness of MoS2 is generally controlled by the thickness of the

starting material; however, the structural and morphological properties of the MoS2

crystallites, towards large area and continuous atomic layers, are strongly dependent on

the temperature of the initial material deposition as well as that of the post-deposition

sulfurization.

Keywords: MoS2, two-dimensional materials, vapor-phase growth, Raman scattering *Corresponding author; E-mail: liuhf@imre.a-star.edu.sg

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1 Introduction

The development of non-graphene two-dimensional (2D) materials, e.g.,

atomically thin transition-metal dichalcogenide semiconductor MoS2, makes it possible to

prepare electronic [1 , 2 , 3], photonic [4 , 5 ], and optoelectronic devices with novel

functions and/or improved performances [6, 7, 8, 9, 10]. Although many such devices

have recently been demonstrated, problems are met in attempts to make large area and

uniform device-quality MoS2 atomically thin layers. A typical example is exfoliation [11],

which has the potential of providing high quality 2D materials but lacks a systematic

control of the thickness, size, and uniformity of the product [ 12 , 13 , 14 ]. Similar

limitations exist with other top-down methods, e.g., laser thinning [15] and chemical

etching [16]. In contrast, the thickness and uniformity are feasibly controlled and the

growth is readily scaled up in bottom-up methods [17, 18, 19]. However, there is a great

challenge to control the crystal quality, especially the grain size, in the growth of 2D

MoS2 by bottom-up methods [19, 20].

Chemical vapor deposition (CVD), a typical bottom-up growth method, has long

been used for growing semiconductor thin films and nanostructures, and has recently

shown a great success in the growth of graphene [21]. For growing 2D MoS2, however,

this method now has a limit to control the lateral-to-vertical ratio of growth rates and thus

the lateral grain sizes with atomic thickness of the resultant crystal [19, 20]. Physically,

the lateral-to-vertical ratio of growth rates is dominated by the growth conditions,

including the source materials, temperatures, pressures, etc. [17-20]. Most importantly,

the effects of the substrate and/or the template on the growth of 2D MoS2, particularly the

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lateral-to-vertical ratio of growth rates which plays an important role in concurrently

controlling the lateral grain size and the layer thickness, are largely increased due to the

thickness reduction.

In this paper we have systematically studied the effects of the substrates, starting

materials, and growth temperatures on the structural and lattice dynamic properties of

MoS2 atomic layers grown by thermal vapor sulfurization in a tube-furnace system. This

method is also commonly named vapor-phase growth or CVD in the literature [18, 22].

Various techniques have been employed to analyze and characterize the resultant MoS2

atomic layers. By carefully comparing the sapphire and native-SiO2/Si substrates, the Mo

and MoO3 starting materials, and the low- and high- temperatures depositions of the

starting material, we observed a remarkable improvement in lateral grain size as well as

in area density of the MoS2 atomic layers on sapphire substrate with MoO3 deposited at

high temperatures as the starting material. These observations could have an important

consequence in fabricating large area and continuous none-graphene 2D materials.

2 Experimental

2.1 Preparation of starting material

Thin Mo and MoO3 films were used as the starting materials in this work. They

were deposited on c-plane sapphire and SiO2/Si (001) substrates in a magnetron-

sputtering chamber using pure argon (99.999%) and oxygen-argon-mixture as the

working gases, respectively. The SiO2 is unintentionally introduced but the native oxide

on the surface of the Si substrate. The oxygen-to-argon ratio was controlled by the gas

flow rates. For the Mo deposition the flow rate of argon was set at 10 SCCM and the

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chamber pressure was 5 mTorr during sputtering. On the other hand, for the MoO3

deposition, the chamber pressure was also 5 mTorr, however, the flow rate of argon was

reduced to 5 SCCM while oxygen was introduced at a flow rate of 5 SCCM. In the latter

case, low- and high-temperature depositions were carried out at room temperature and

700 °C, respectively. The film thickness, generally smaller than 4 nm, was controlled by

the deposition time with the deposition rates calibrated at certain conditions.

2.2 Sulfurization of Mo and MoO3 thin layers into MoS2

The post-deposition thermal vapor sulfurization of the starting material was

carried out in a tube-furnace system using nitrogen as the carrier gas. The general

configuration of the tube-furnace system and the sulfurization procedures have been

reported elsewhere [18]. It is worth mentioning that two separated crucibles containing

sulfur powders (99.999%) were set in the upstream at different distances from the

samples. So that the one near the samples starts to supply sulfur species via sublimation

at lower temperatures while the second source starts to supply sulfur via evaporation

before the entire consumption of the first source at higher temperatures. The sulfurization

temperature was varied from 650 to 950 °C while the sulfurization time was varied from

20 to 40 min.

2.3 Instruments and characterization methods

The MoS2 samples were characterized by employing various techniques,

including atomic force microscopy (AFM), Raman spectroscopy, optical absorption

spectroscopy, x-ray photoelectron spectroscopy (XPS), and transmission-electron

microscopy (TEM). The AFM images were recorded using a tapping mode in a Veeco

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Dimension-Icon AFM system. The Raman spectra were collected at room temperature in

a backscattering configuration using a micro-Raman system that was equipped with a

488-nm argon-ion laser. The optical absorbance was measured from the sulfurized

samples on sapphire substrates using a UV-VIS-NIR scanning photospectrometer. The

XPS experiments were carried out in a Quantera SXM XPS chamber employing Al-Kα

(hν =1486.6 eV) as the x-ray beam source. In its high-resolution spectroscopy, the energy

resolution is smaller than 0.5 eV.

3 Results and discussion

3.1 Mo sulfurization: Effect of substrate

To study the effect of substrate on vapor-phase growth of 2D MoS2, ultrathin Mo

layers (~2 nm) were deposited on c-plane sapphire and native-SiO2/Si substrates followed

by thermal sulfurization at 900 °C for 20 min. Figures 1(a) and 1(b) show the AFM

images while Figs. 1(c) and 1(d) show the Raman spectra collected from the MoS2 layers

grown on the sapphire and SiO2/Si substrates, respectively. Remarkable differences in the

surface morphologies are seen in Figs. 1(a) and 1(b), typically the larger surface coverage

of MoS2 on SiO2/Si than that on sapphire. An AFM image of the MoS2/Sapphire sample

with a larger magnification is shown in the inset of Fig. 1(a), which clearly shows the

incomplete surface coverage with MoS2 grains in the scale of 20-50 nm. A sectional

analysis across the grain edges revealed that the MoS2 grains are about 2.0 nm, i.e., 3-

monolayer (3L) in thickness. However, such MoS2 grains are not seen at all in Fig. 1(b).

Instead, Fig. 1(b) exhibits a continuous surface, which consisted of dots with different

sizes. The inset in Fig. 1(b) is an enlarged AFM image in a tilted-view configuration,

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where larger dots decorated by smaller ones seeding into the surface are observed. Such

an anomalous morphology of MoS2 grown by CVD on SiO2/Si has not been reported in

the literature. Nevertheless, for both samples, the characteristic in-plane ( 12gE ) and out-

plane (A1g) optical vibrational modes of MoS2 are clearly seen at 386.0 ± 0.5 cm-1 and

409.8 ± 0.2 cm-1, respectively. The frequency differences (Δ = 23.7-24.1 cm-1) between

the in- and out-plane phonon modes are larger than those of 3L MoS2 fabricated by

exfoliation but consistent with those of MoS2 in the same thickness grown by CVD [19,

23].

To work out the origin of the morphology differences between the MoS2 atomic

layers grown on the sapphire and the SiO2/Si substrates, a set of high-resolution XPS

spectra were collected from the ‘surface’ of both samples before and after ion-beam

sputtering for different durations. Figure 2 presents the XPS core level spectra, labelled ts

= 0 - 240 sec, collected by increasing the sputtering time with intervals of 30 sec. The

XPS spectra were aligned to the C1s core level (285 eV) detected on the as-

grown/sulfurized surfaces. The spectra in the left and right panels of Fig. 2 were collected

from the MoS2 layers grown on the sapphire and the SiO2/Si substrates, respectively. It is

seen that the core level binding energies of Mo3d and S2p measured from the as-grown

samples, i.e., those spectra labelled ts = 0 in Figs. 2(a)-(d), are almost identical to those of

atomically thin MoS2 reported in the literature [19, 20, 24]. Likewise, the core level

binding energies in Figs. 2(e)-2(h) are consistent with those of sapphire (i.e., Al2O3) and

SiO2 [25, 26, 27]. It is worth noting that the XPS signals of neutral Si (i.e., Si2p at a

binding energy of ~99.2 eV) were not detectable at all in this study due to the ‘thick’

native SiO2 layer.

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The XPS spectral evolutions upon sputtering, i.e., the peak shifts and intensity

variations in Fig. 2, reveal a remarkable difference between the MoS2 samples grown on

the sapphire and the SiO2/Si substrates. Apparent peak shifts are seen in the left panels

rather than in the right ones. The Mo3d and S2p peaks shift to lower binding energies at

the initial sputtering stage and then turn back to higher ones after sputtering for ts = 120

sec. On the contrary, the Al2p and O1s peaks first shift to higher binding energies and

then turn back to lower ones after sputtering for ts = 60 sec. The core levels of the anions

and cations of the MoS2 layer shift to the same direction with nearly the same values

upon sputtering [see Figs. 2(a) and 2(c)], so do those of the sapphire substrate [see Figs.

2(e) and 2(g)]. These observations indicate that an onset of charge transfer between the

MoS2 layer and the sapphire substrate occurred during the ion-sputtering. Physically, the

surface charge of the ultrathin MoS2 covered sapphire develops with the ion-sputtering

due to the emission of electrons, so does the space charges in the subsurface due to the

penetration of x-ray and the trapping of the electrons in the insulating substrate [28]. Both

developments lead to the charge transfer as a function of sputtering.

It is known that MoS2 molecule has two stacking structures [29]. One is trigonal

prism (2H) and the other is octahedral coordination (1T). From the electronic structure

point of view, the former is semiconductor while the latter is metal. Thermodynamically,

the former structure is more stable than the latter one. Phase transition from 1T to 2H of

MoS2 atomic layers during thermal annealing has been reported by Eda et al. [14]. They

observed that both Mo3d and S2p core levels shift to larger binding energies along with

the 1T-to-2H phase transition induced by thermal annealing. In this light, the parallel

shift of Mo3d and S2p peaks to lower binding energies in Figs. 2(a) and 2(c) could be

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attributed to 2H-to-1T phase transitions induced by ion-sputtering. This attribution is also

supported by the consistence between the sputtering-induced S2p1/2-S2p3/2 peak merging

observed in Fig. 2(c) and the anneal-induced S2p1/2-S2p3/2 peak splitting reported in Ref.

14. Certainly, the 2H-to-1T transition tends to convert MoS2 from semiconductor to metal

and thus promotes the change transfer between the MoS2 overlayer and the sapphire

substrate. In contrast, the absence of peak shift, together with the absence of S2p1/2-S2p3/2

peak merging, in Figs. 2(b) and 2(d) provide evidence that no phase transition occurred

during the ion-sputtering in the MoS2 structure grown on the SiO2/Si substrate.

It has been discussed above that the AFM studies revealed a dots-seeding

morphology for the MoS2-on-SiO2/Si structure, different from that of MoS2-on-sapphire.

To study the surface structure of MoS2-on-SiO2/Si, here, we make a comparison on the

evolutions of XPS-peak intensity upon sputtering. The spectra in the left panels of Fig. 2

show that the Mo3d and S2p peaks monotonically decrease, while the Al2p and O1s

peaks monotonically increase, with the increase in sputtering time. This is a normal

phenomenon, which is commonly observed in depth-profiling XPS measurements of

heterostructures with ultrathin layers covered on extrinsic substrates. However, in the

right panels the Mo3d and S2p peaks monotonically increase in intensity, while the Si2p

and O1s peaks do not exhibit any apparent intensity variations, when the sputtering time

is increased. These observations imply that the MoS2 structure is partially or completely

embedded in SiO2 on the Si substrate. Combined with the AFM studies, we may conclude

that nanoscale MoS2 dots/particles were seeded into SiO2 amorphous on the Si substrate,

which is hitherto unknown for CVD growth of 2D MoS2 [22]. The mechanism that the

structure of MoS2 nano dots embedded in SiO2 can prevent the 2H-to-1T transition

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during sputtering is not clear at this stage. However, it may relate to the lack of a

collective orientation, i.e., the MoS2 dots are randomly oriented in the SiO2 amorphous,

as well as the semiconducting Si substrate. Nevertheless, these comparisons clearly show

that sapphire ranks priority to native-SiO2/Si as the substrate for CVD growth of MoS2

2D layers with increased lateral grain size under certain conditions.

A careful comparison between the intensity evolutions in Figs. 2(a) and 2(c)

reveals that the loss of S from the surface upon sputtering is much faster than that of Mo.

As we have mentioned above that the transition point of the Mo3d peak shift is at ts = 120

sec and after that S is not detectable any more (supporting materials, S1). However, the

transition point of the parallel Al2p and O1s shifts is at ts = 60 sec. This result indicates

that there is another mechanism beside the 2H-to-1T phase transition that affects the

charge transfer between the MoS2 overlayer and the sapphire substrate. A reasonable

explanation is that the interface between the MoS2 layer and the sapphire substrate is

occupied by Mo. The fast loss of S during sputtering leaves Mo remaining on the surface,

which may react with the sapphire substrate upon sputtering and modify the charger

transfer. On the other hand, the fast decrease in the S2p-to-Mo3d intensity ratio with

increasing the sputtering time is also an indication of the layered structure of MoS2 grown

on the sapphire substrate. In contrast, the S2p-to-Mo3d intensity ratio of MoS2 grown on

the SiO2/Si substrate does not show any apparent variations upon sputtering because of

the orientation randomly MoS2 dots of the dots-seeding structure.

3.2 Mo sulfurization: Effect of thickness

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To control the thickness of MoS2 atomic layers sulfurized on sapphire substrates,

a sulfurization of Mo layers with varied thicknesses was carried out at a higher

temperature (i.e., 950 °C). The MoS2 samples hereafter labelled S120, S60, S30, and S10

are referred to those with the initial Mo deposited for tMo = 120, 60, 30, and 10 sec,

respectively. Their absorption spectra, collected using a bare sapphire wafer as the

reference, are presented in Fig. 3. Exciton resonance absorption features are clearly seen

and the features at about 660, 610, and 440 nm, assigned to exciton A, B, and C, are

consistent with those of 2D MoS2 semiconductor nanosheet fabricated by chemical

exfoliation [14]. The exciton energies of A and B derived from the absorption spectra are

plotted in the inset of Fig. 3 as a function of tMo. When tMo is reduced, the absorbance

decreases in the entire measurement wavelength range while the exciton absorption

resonances (typically exciton C) shift to shorter wavelengths (i.e., blue shift). Since the

sulfurizations were carried out simultaneously under the same conditions, the blue shifts

in the exciton absorption resonances thus provide clear evidence that the thickness of the

resultant MoS2 atomic layers is strongly dependent on the thickness of the starting Mo

layers.

Figure 4(a) presents the Raman spectra collected from the set of tMo-varied MoS2

samples grown on sapphire substrates. The characteristic in- and out-plane phonon

vibration modes are clearly seen along with the Raman features of sapphire (indicated by

the asterisk). Both modes of MoS2 decrease in intensity with decreasing tMo. The nearly

linear decrease in mode intensity, as well as the monotonic decrease in 12gE -to-A1g

intensity ratios, with decreasing tMo (supporting materials, S2) is consistent with those

observed in exfoliated few-layer MoS2 [23]. Figure 4(b) plots the Raman shifts and the

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frequency differences, obtained by spectral fittings, as a function of tMo. It is seen that the

frequency differences are smaller than 22 cm-1, indicating that the MoS2 layers are not

thicker than 3L [23, 30]. Also seen is that the in-plane mode 12gE exhibits a monotonic

blue shift with decreasing the nominal layer thickness, consistent with those reported in

the literature [19, 23, 31]. However, the out-plane mode A1g exhibits a blue shift as tMo is

decreased from 120 to 30 sec followed by a red shift from tMo = 30 to 10 sec. The

frequency difference thus increases, followed by decreases, when the MoS2 thickness is

reduced in atomic layers [see Fig. 4(b)]. The frequency trend of A1g is anomalous to those

reported in the literature, where the A1g mode monotonically decreases in frequency with

decreasing the thickness of MoS2 from bulk to monolayer [19, 23, 31].

In classic theory, assuming that layer stacking does not affect intralayer bonding,

both 12gE and A1g modes of MoS2 are expected to red shift with decreasing the layer

numbers. In general, most of the reported shifts of A1g agree with, while the shifts of

12gE are opposite to, the classic theory except for the results reported by Ramakrishna et

al. [32], where they observed a red shift for both A1g and 12gE modes as the thickness of

MoS2 is decreased. The anomalous shift of 12gE was attributed by Molina-Sáanchez and

Wirtz to the competition between the long-range and the short-range Coulomb

interactions [33]. As the layer thickness is increased, the long-range part is decreased due

to an enhanced dielectric screening, which overcompensates for the increase in the short-

range interaction due to the weak interlayer interaction, leading to the observed red shift

of 12gE . Without the enhancement of dielectric screening, the in-plane mode 1

2gE would

blue shift, i.e., in the same trend as that of the out-plane mode A1g, with increasing the

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layer numbers of MoS2. However, Luo et al. recently attributed the anomalous 12gE shift

to the surface effect of MoS2 atomically layers [34], where the surface effect refers to the

larger Mo-S force constants at the surface of MoS2 due to a loss of neighbors in adjacent

MoS2 layers. Basically, these two explanations are consistent in terms of the atomic force

constant variations. In fact, a detailed study by Luo revealed that if the modified surface

constants are not taken into account both A1g and 12gE increase in frequency with the layer

numbers and the frequency increase of A1g is much faster than that of 12gE . However,

when the modified surface constants are taken into account, an additional blue shift is

introduced to both A1g and 12gE , and the additions are larger for smaller layer numbers,

leading to the reduced blue shifts of A1g while the red shifts of 12gE with decreasing the

layer numbers of MoS2. Unfortunately, neither the dielectric screening model nor the

surface effect mode can explain the anomalous frequency trend of A1g observed in Fig.

3(b). One may suspect that variations in crystal strain may be of concerns. However, both

experiments and theoretical calculations indicate that the A1g mode has the same shift

direction as that of 12gE when the remaining crystal strain is changed, moreover, the

former mode is less sensitive than the latter one to the strain variations [35].

Buscema et al. have recently reported that the substrate has an important effect on

the vibration frequency of A1g rather than 12gE [36], the similar effect of high-k dielectrics

covered on top of monolayer MoS2 has been observed earlier [37]. These observations

are well explained by charge transfer, e.g., via doping due to charged impurities at the

interface [36, 37]. By electrically charging a monolayer MoS2 in a top-gate field effect

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transistor (FET) structure, Chakrabory et al. directly observed a larger red shift of A1g

while a negligible shift of 12gE [38]. It has been discussed above that the interface between

the MoS2 overlayer and the sapphire substrate is occupied by Mo. The lacking of S at the

bottom of this interfacial Mo layer could negatively charge MoS2 via a direct charge

transfer. Physically, the presence of Mo at the MoS2/sapphire interface increases with

increasing the initial Mo thickness because of the same sulfurization conditions. As a

result, the doping of MoS2 increases with increasing the initial Mo thickness, leading to a

larger red shift in A1g, which is indeed what we observed in Fig. 4(b).

3.3 Mo sulfurization: Effect of temperature

Figures 5(a) and 5(b) show the Raman spectra collected from the MoS2 samples

sulfurized on sapphire substrates with the initial Mo layers deposited for tMo = 20 and 10

sec, respectively. The sulfurizations were carried out at T = 650, 750, 850, and 950 °C for

30 min. The increased sulfurization time chosen here is to minimize, if possible, the

density of interfacial Mo. Figures 5(c) and 5(d) plot the phonon linewidths and the

integrated intensities as a function of sulfurization temperatures. The linewidths and

integrated intensities were derived by spectral fittings from Figs. 5(a) and 5(b),

respectively. A typical spectral fitting is shown in the inset of Fig. 5(b), which was

carried out for the S10 sample sulfurized at 650 °C. It is seen that the characteristic

phonon intensity of both samples increases with increasing T. Also seen is that the in-

plane mode 12gE of both samples exhibits an apparent linewidth reduction with increasing

T. The full-width at half-maximum (FWHM) of 12gE decreases from 7.5 cm-1 (S20) and

4.4 cm-1 (S10) at T = 650 °C to 3.0 cm-1 (S20) and 2.5 cm-1 (S10) at T = 950 °C.

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However, the linewidth of the out-plane mode A1g is insensitive to the sulfurization

temperatures, which varies in 5.5-5.0 cm-1 and 5.0-4.5 cm-1 for S20 and S10, respectively.

Both the linewidths of 12gE and A1g are comparable to those of high quality exfoliated

monolayer MoS2 flacks [23]. In general, the monotonic increase in intensity, together

with the monotonic decrease in the linewidths of 12gE , indicates that the crystal quality of

MoS2 is improved by the increase in sulfurization temperatures. In Fig. 5(a), one sees that

the frequency difference, Δ = 21.2 cm-1, between the in- and out-plane modes is

independent of the sulfurization temperature. This observation indicates that the thickness

of MoS2 in S20 is dominated by the initial Mo thickness regardless of the sulfurization

temperatures in the range of 650-950 °C. The increased intensity of the Raman features

accompanied by the peak narrowing of 12gE as well as the unchanged layer thickness of

MoS2 provide evidence that the lateral grain sizes of the 2D materials are increased by

increasing the sulfurization temperatures. The remarkable difference between S20 and

S10 [see Figs. 5(a) and 5(b)] is that the former exhibits no frequency variations at all for

both the in- and out-plane modes while the latter shows opposite shifts of 12gE and A1g

with increasing T. The MoS2 thicknesses as determined from the frequency difference,

i.e., Δ = 21.2 cm-1 for S20 and Δ = 19.8-21.3 cm-1 for S10 [see Figs. 5(a) and 5(b)], are

not more than 2 monolayers (2L) for both samples.

A direct evidence of the increase in lateral grain sizes of MoS2 is further provided

by AFM analysis, which is shown in Fig. 6. The left and right panels are from S20 and

S10, respectively. The AFM images in Figs. 6(a)-6(b) and 6(c)-6(d) are recorded from the

samples sulfurized at 750 °C and 950 °C, respectively. Figures 6(c’) and 6(d’) are

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sectional analysis of the grain height as indicated by the lines in Figs. 6(c) and 6(d),

respectively. It is clearly seen that the lateral grain sizes are apparently increased for both

samples by increasing T from 750 to 950 °C. However, the surface coverage is

dramatically reduced, particularly for S10 [Figs. 6(b) and 6(d)], by increasing the

sulfurization temperature. The sectional analysis in Figs. 6(c’) and 6(d’) revealed that the

MoS2 grains are generally 0.6-0.7 nm in thickness (i.e., monolayer) for both samples.

Moreover, the surface coverage comparisons between T = 650 and 950 °C indicate that

the onset of Mo evaporations from the sapphire surface occurred at elevated

temperatures, especially when the initial Mo thickness is smaller. On the other hand, a

comparison between the left (S20) and the right (S10) panels in Fig. 6, typically the

larger lateral grain sizes in the right panels than those in the left ones, revealed that the

thickness of the initial Mo layer plays an important role in surface mobile of the reactive

species at elevated temperatures. These comparisons provide evidence that both the

sulfurization temperature and the initial Mo thickness are key parameters in controlling

the lateral grain size of MoS2 atomic layers upon sulfurization. At high sulfurization

temperatures, a reduced thickness of the initial Mo not only increases the surface mobility

but also enhances the surface evaporation, leading to the growth of MoS2 atomic layers

with increased lateral grain size but reduced surface coverage.

3.4 MoO3 sulfurization: Effect of initial thickness

The competition between the lateral grain size and the surface coverage makes

Mo unsuitable as the initial material for growing high quality, larger area, and continuous

2D MoS2 layers via thermal vapor sulfurization at elevated temperatures. MoO3, another

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potential initial material for growing 2D MoS2 via sulfurization may somehow solve the

problem [24]. For this purpose, we have deposited MoO3 layers with different thicknesses

at 700 °C on c-plane sapphire substrate. The deposition times were varied from 300 sec

to 120, 100, and 60 sec, and the sulfurized samples are named SO300, SO120, SO100,

and SO60, respectively. The sulfurizations were carried at 950 °C for 20 min. Figure 7(a)

presents the absorbance spectra collected from these sulfurized MoS2 samples. The

characteristic exciton absorption resonances are clearly distinguished and assigned to A,

B, and C of MoS2. Their energies were analyzed and plotted in Fig. 7(b) as a function of

the initial MoO3 deposition time. It is seen that the overall absorbance decreases [see Fig.

7(a)] and the energy of the exciton absorbance resonances increases with decreasing the

initial MoO3 thickness. These observations are consistent with those of MoS2 atomic

layers grown by Mo sulfurizations discussed above [see Fig. 3], indicating that the

thickness of the resultant MoS2 is also strongly dependent on the thickness of the initial

MoO3 layer.

Figure 8(a) presents the Raman spectra collected from the MoS2 samples grown

by sulfurizing the MoO3 layers. Both the characteristic in- and out-plane phonon modes

of MoS2 are clearly seen and the phonon energies derived by detailed spectral analysis

are shown in Fig. 8(b) as a function of the deposition time of the initial MoO3. As

expected, one can clearly see that the frequency difference between the A1g and 12gE

monotonically decreases with the decrease in the thickness of the initial MoO3. This

correlation indicates that the resultant MoS2 layers are in a few monolayers thick and

their thicknesses are strongly dependent on the initial MoO3 thicknesses [23]. A careful

comparison between the Raman spectra in Fig. 8(a) revealed that although the A1g mode

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monotonically shifts to lower frequencies (i.e., from 410.5 cm-1 down to 408.2 cm-1) as

the MoS2 thickness is decreased by reducing the deposition time of the initial MoO3, the

frequency variations of the 12gE mode (< 0.4 cm-1) are apparently smaller than those

observed in Mo-sulfurizations (see Fig. 4). The frequency differences, in terms of the

thickness of MoS2 (e.g., SO60, see later discussions), are larger than those observed in

exfoliated MoS2 flakes. A possible reason could be associated with the residual lattice

strain. It is well known that biaxial tensile strain remaining in a crystal can induce a red

shift to its optical phonons [39 ]. In the case of 2D MoS2, both experimental and

theoretical studies show that the tensile strain induced red shift of 12gE (-2.1 cm-1 per %

strain) is much larger than that of A1g (-0.4 cm-1 per % strain) [35]. In this regard, the

small blue shift of 12gE with decreasing the thickness of MoS2 observed in Fig. 8 could be

the result of compensation by a red shift induced by tensile strain. It is quite reasonable

that the residual tensile strain in the MoS2 atomic layers increases with the decrease in the

layer thicknesses because of the increased effect of the substrate.

AFM images taken from samples SO300, SO120, SO100, and SO60 (i.e., those

with the initial MoO3 layers deposited for 300 sec, 120 sec, 100 sec, and 60 sec,

respectively) are presented in Figs. 9(a)-9(d), respectively. The sectional analyses, shown

below the individual AFM images, revealed that the surface roughness of the resultant

MoS2 monotonically decreases with decreasing the initial MoO3 thickness. However, one

can see that the significant surface coverage reduction that observed in the case of Mo

sulfurizations at 950 °C, induced by a reduced thickness of the initial Mo [see Figs. 6(c)

and 6(d)], is missing here. The reduced surface roughness, together with the improved

18

Raman features of MoS2, i.e., the increased intensity accompanied by the narrowed

linewidth [typically those of 12gE , see Fig. 8(a)], provides evidence that the crystal quality

of the resultant MoS2 atomic layers is improved by the decrease in the thickness of the

initial MoO3.

3.5 MoO3 sulfurization: Effect of MoO3 deposition temperature

A more detailed study of the SO60 sample discussed in section 3.4, which had the

initial MoO3 deposited at 700 °C, is further carried out with a comparison to the

sulfurization of MoO3 deposited at room temperature. Figures 10(a) and 10(b) show the

AFM images of the MoS2 atomic layers grown by sulfurizing MoO3 that were deposited

at high- and low-temperatures (indicated by HT and LT), respectively. Although the

initial MoO3 thickness and the sulfurization conditions are the same for both samples, one

sees that the surface coverage is significantly improved by increasing the deposition

temperature of the initial material. Also seen in Fig. 10(c) is that both 12gE and A1g shifted

to lower frequencies by increasing the deposition temperature of the initial MoO3.

Moreover, the shift of 12gE is obviously larger than that of A1g, suggesting that tensile

strain was introduced into the MoS2 layers with increasing the deposition temperature of

the initial MoO3. These results further support the observations in Fig. 8 that the small

blue shifts of 12gE were caused by the compensation of strain induced red shift.

Furthermore, the Raman spectra in Fig. 10(c) shows that the frequency differences

between the A1g and 12gE modes are almost the same for both MoS2 samples. As a

consequence, we can conclude that the thickness of both samples should be the same,

19

which is thus readily obtained (1.4 nm, i.e., 2L) by a sectional analysis across the grain

edges in Fig. 10(b). A careful look into the AFM image in Fig. 10(a) revealed that most

of the MoS2 “grains” are more or less connected with one another. A typical chain of

grain connections is indicated by the dots in Fig. 10(a). Such grain connection chains may

promote the obtained MoS2 atomic layers for practical applications. The increased

surface coverage, the increased crystal quality, and the long grain connection chains thus

provide a clue to further improve the growth of high quality 2D MoS2 towards continuous

monolayer in large scale.

4 Conclusions

In conclusion, we have systematically studied the vapor-phase growth of MoS2

atomic layers via thermal vapor sulfurization. On a native-SiO2/Si substrate the Mo-

sulfurization (i.e., using Mo as the initial material) resulted in a structure of MoS2

particles seeded into SiO2 amorphous. A possible reason is that the native SiO2 is not

dense enough, the existence of surface pits and/or voids could result in atomic diffusions

of Mo and S into SiO2 and thus the formation of MoS2 nanoparticles therein. The same

Mo-sulfurization process on c-plane sapphire substrates led to the desired MoS2 atomic

layers, but there exists a challenge in increasing the lateral grain sizes. It is found that the

initial Mo thickness and the sulfurization temperature play a key role in controlling the

lateral grain sizes. An increase in sulfurization temperature can improve the crystal

quality as well as the lateral grain sizes. However, the surface coverage is significantly

reduced with reducing the initial Mo thickness. On the other hand, an increased Mo

thickness tends to suppress the surface mobility and thus results in smaller lateral grain

20

sizes of MoS2. By using MoO3 deposited at room temperature as the starting material, 2

monolayer thick MoS2 flakes of ~100 nm in lateral sizes with well-developed grain edges

(i.e., triangular or hexagonal) are obtained. However, the surface coverage so obtained is

quite small. An increase in the deposition temperature of the initial MoO3 (with the same

thickness as that deposited at room temperature), i.e., to 700 °C, can significantly

increase the surface coverage without increasing the layer thickness of MoS2 at all. The

coalescence of MoS2 grains in one direction or another throughout the wafer can promote

the MoS2 atomic layers for practical applications.

Acknowledgement

The authors would like to acknowledge Junguang Tao and Jisheng Pan (both are

currently working at IMRE, A*STAR) for fruitful discussions.

21

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Figure Captions:

Figure 1: AFM images and Raman spectra collected from the MoS2 samples grown on c-

plane sapphire (a)-(c) and native-SiO2/Si substrates (b)-(d) by thermal vapor

sulfurization of Mo.

Figure 2: XPS spectra of core levels collected from the MoS2 samples grown on c-plane

sapphire (a)-(c)-(e)-(g) and native-SiO2/Si substrates (b)-(d)-(f)-(h) by thermal

vapor sulfurization of Mo.

Figure 3: Absorbance spectra measured from the MoS2 atomic layers grown on c-plane

sapphire substrates via sulfurizing (at 950 °C) Mo layers that were deposited

by sputtering for variable times. The inset plots the exciton absorption

resonance energies as a function of the initial Mo deposition times.

Figure 4: Raman scattering spectra (a) and the frequency trends (b) measured from the

MoS2 atomic layers grown on c-plane sapphire substrates via sulfurizing (at

950 °C) Mo layers that were deposited by sputtering for variable times. The

asterisk indicates the Raman feature from the sapphire substrate.

Figure 5: Raman scattering spectra and the feature evolutions of MoS2 atomic layers

grown on c-plane sapphire substrates by Mo-sulfurizations at different

temperatures. The Mo layers were deposited for 20 sec (a)-(c) and 10 sec (b)-

(d). The asterisks indicate the Raman features from the sapphire substrate.

27

Figure 6: AFM images taken from the MoS2 samples grown by Mo-sulfurizations at 750

°C (a)-(b) and 950 °C (c)-(d) with the initial Mo deposited for 20 sec (a)-(c)

and 10 sec (b)-(d). The profiles in (c’) and (d’) are section analyses from the

parts indicated by the straight lines in (c) and (d), respectively.

Figure 7: Absorbance spectra (a) and the exciton absorption resonance energies (b)

measured from MoS2 atomic layers grown on c-plane sapphire substrates by

sulfurizing (at 950 °C) MoO3 layers that were deposited by sputtering at 700

°C.

Figure 8: Raman scattering spectra (a) and the frequency trends (b) measured from MoS2

atomic layers grown on c-plane sapphire substrates by sulfurizing (at 950 °C)

MoO3 layers that were deposited by sputtering at 700 °C.

Figure 9: AFM images taken from the MoS2 samples grown on c-plane sapphire

substrates by sulfurizing (at 950 °C) MoO3 layers that were deposited by

sputtering at 700 °C for 300 sec (a), 120 sec (b), 100 sec (c), and 60 sec (d).

The panels below the individual images are the section analyses, showing the

improved surface smoothness with reducing the MoO3 deposition time.

Figure 10: Morphological comparisons between the MoS2 atomic layers grown on c-

plane sapphire substrates by sulfurizing (at 950 °C) MoO3 that were deposited

at 700 °C (a) and room temperature (b). Plotted in (c) are the Raman spectral

comparisons, the asterisk indicates the Raman feature from the sapphire

substrate.

28

Figure 1

29

Figure 2

30

Figure 3

31

Figure 4

32

Figure 5

33

Figure 6

34

Figure 7

35

Figure 8

36

Figure 9

37

Figure 10