1
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: [email protected]
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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.