Continuous Heteroepitaxy of TwoDimensional Heterostructures Based on
Layered Chalcogenides
Yu Kobayashi,† Shoji Yoshida,‡ Mina Maruyama,‡ Hiroyuki Mogi,‡ Kota Murase,‡ Yutaka Maniwa,†
Osamu Takeuchi,‡ Susumu Okada,‡ Hidemi Shigekawa,‡ and Yasumitsu Miyata*,†
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†

Department of Physics, Tokyo Metropolitan University, Hachioji, Tokyo 192-0397, Japan
Faculty of Pure and Applied Sciences, University of Tsukuba, Tsukuba 305-8573, Japan‡S Supporting Information*

ABSTRACT: The in-plane connection and layer-by-layer stacking of atomicallythin layered materials are expected to allow the fabrication of two-dimensional(2D) heterostructures with exotic physical properties and future engineering
applications. However, it is currently necessary to develop a continuous growth
process that allows the assembly of a wide variety of atomic layers without
interface degradation, contamination, and/or alloying. Herein, we report thecontinuous heteroepitaxial growth of 2D multiheterostructures and nanoribbons based on layered transition metal dichalcogenide (TMDC) monolayers,
employing metal organic liquid precursors with high supply controllability. This versatile process can avoid air exposure
during growth process and enables the formation of in-plane heterostructures with ultraclean atomically sharp and zigzagedge straight junctions without defects or alloy formation around the interface. For the samples grown directly ongraphite, we have investigated the local electronic density of states of atomically sharp heterointerface by scanningtunneling microscopy and spectroscopy, together with ﬁrst-principles calculations. These results demonstrate an approachto realizing diverse nanostructures such as atomic layer-based quantum wires and superlattices and suggest advanced
applications in the ﬁelds of electronics and optoelectronics.
KEYWORDS: transition metal dichalcogenides, two-dimensional materials, semiconductor, heterostructures,chemical vapor depositionH

composition-dependent tunable electronic properties andintriguing spin-valley physics.24−28 In previous studies,
TMDC growth has been primarily obtained using vaporphase deposition with solid precursors such as metal oxides
and chalcogens.14−17 These solid precursors typically have low
vapor pressures at room temperature and thus require hightemperatures (often above 600 °C) for vaporization. In theearly studies, this requirement frequently makes it diﬃcult torepeatedly supply diﬀerent precursors while maintaining acontrolled growth rate during epitaxial growth. To solve thisissue, several approaches have been recently reported tocontrol precursor supply rates and to realize multiheterojunctions.20,21,23 For example, Zhang et al. have achieved thesequential vapor deposition of solid precursors of various
multiheterostructures of MoS2, MoSe2, WS2, and WSe2 using areverse ﬂow.20 In this case, the regrowth is conducted with aseparate vapor deposition system and the continuous growth
has never been demonstrated. Similar multiheterostructures

eteroepitaxial growth is widely used to fabricate thinﬁlms of conventional solid materials on substrates.
During this process, the crystal orientation of the ﬁlm
can be controlled based on the orientation of the substrate.
This technique also enables the formation of highly crystalline
heterointerfaces,1 various semiconductor quantum well structures,2−4 and superlattice structures.5,6 All these are importantin modern condensed matter physics and have applications inelectronics and optoelectronics.7−10 Recently, similar epitaxial
growth processes and the resulting heterostructures have beenrealized with atomically thin layers of so-called two-dimensional (2D) materials, including graphene, hexagonal boron
nitride (hBN), and transition metal dichalcogenides(TMDCs).11−23 In such cases, epitaxial growth can be initiatedat the edges of the 2D material, resulting in the formation of anin-plane 2D heterostructure with a one-dimensional interface.
These 2D heterostructures have attracted much attention aspromising systems that allow the investigation of lowdimensional quantum phenomena and may also lead to thedevelopment of high-performance devices.
TMDC atomic layers are ideal materials for the fabricationof functional 2D heterostructures because they exhibit© 2019 American Chemical Society

Received: October 19, 2018
Accepted: May 31, 2019
Published: May 31, 2019
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Figure 1. (a) Schematic illustration of the present MOCVD system. (b) Optical microscopy images, (c) PL spectra, and (d) Raman spectra ofmonolayers composed of MoS2, MoSe2, WS2, and WSe2 grains. (e) High-resolution STM image of a MoS2 monolayer grown on graphite,acquired using the settings Vs = +2.0 V and It = 0.1 nA, where Vs and It are the sample voltage and tunneling current, respectively. (f) PL
spectra of WS2 monolayers grown on graphite using (i) liquid precursors at 640 °C (red) and (ii) WO3 and S at 1100 °C (blue).

to elucidate the band diagram of the type I straddling gap at
the heterointerface.

based on MoS2/WS2 and MoSe2/WSe2 have been continuously grown by water-assisted supply control of solid metal
precursors.21 Instead of such solid precursors, organic metal/
chalcogen precursors with high vapor pressure have been also
used to fabricate the multiheterostructures.23 For these organic
precursors, the heterostructures are prepared with the precise
ﬂow control of all metal and chalcogen precursors without air
exposure. Even though there has been such outstanding
progress, it still remains necessary to develop a sophisticatedtechnique for forming 2D heterostructures with an atomicallystraight interface and to characterize their interface atomic/
electronic structures.
In the present work, we demonstrated a versatile, repeatable
method for the heteroepitaxial growth of 2D multiheterostructures and semiconductor nanoribbon structures, based onusing liquid precursors. TMDC growth was performed with alab-made metal organic chemical vapor deposition (MOCVD)
system in conjunction with liquid organic metal and chalcogenprecursors, as shown in Figure 1a. Using this system, we wereable to grow four diﬀerent types of atomic TMDC layers,
consisting of WS2, WSe2, MoS2, and MoSe2 monolayers. The
heteroepitaxial growth of these monolayers allowed thefabrication of multiheterostructures, as well as nanoribbon
structures 20 nm wide and with six diﬀerent TMDC-based
semiconductor heterostructures. We have investigated the local
electronic density of states of atomically sharp heterointerface
by scanning tunneling microscopy (STM) and spectroscopy(STS), together with ﬁrst-principles calculations. Fine tuningof the growth rates enabled the formation of atomically straightinterfaces with perfect metal zigzag edges in MoS2/WS2
heterostructures. These high-quality interfaces were employed

RESULTS AND DISCUSSION
The liquid precursors used in this work were found to permit
the rapid growth of high-quality TMDC monolayer single
crystals having wide areas at low temperature underatmospheric pressure. As shown in the optical microscopy
images provided in Figure 1b, MoS2, MoSe2, WS2, and WSe2
monolayers were grown on SiO2/Si substrates under optimized
conditions. These monolayer grains generally exhibited typical
3-fold symmetry with straight edges, indicating the preferential
formation of metal zigzag edges.29 These grains also showed
the expected photoluminescence (PL) and Raman peaks(Figure 1c,d). We note that the use of alkali metal halides as
growth promotors was critical to increasing the crystal size andquality (Figure S1). Similar results have been observed in priorwork using diﬀerent precursors.22,23,30−32 This could bederived from the formation of intermediate products of alkali
metals and transition metals, and the details will be reported
elsewhere. Both grain size and quality were also aﬀected by therate at which the metal and chalcogen precursors were supplied(Figure S2). The STM image provided evidence for highly
crystalline lattices in the TMDC monolayers (Figure 1e). The
quality of each sample could also be conﬁrmed based on the
narrow full-width at half-maximum (fwhm) values of the PL
peaks. It has been previously reported that the PL line widths
become much narrower when the samples are grown ongraphite substrates.33 As an example, the PL peak generated by
the WS2 monolayer on graphite had an fwhm value of 20 meVand could be ﬁtted by a single Lorentzian function (Figure 1f).
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Figure 2. PL intensity maps and PL spectra of (a) MoS2/WS2, (b) WSe2/WS2, (c) MoS2/MoSe2, (d) WSe2/MoSe2, (e) WS2/MoSe2, and (f)MoS2/WSe2 single heterostructures. In the PL intensity maps, red, yellow, cyan, and green correspond to the intensities of the MoS2, MoSe2,
WS2, and WSe2 peaks, respectively. Sharp peaks indicated by asterisks correspond to the Raman modes of the graphite substrate.

A similar narrow line width has been observed for the highquality monolayer WS2 grown on graphite with solid
precursors at 1100 °C.33 Overall, these results indicate thatthe present method can produce high-quality samples with
reduced quantities of structural defects and charged impurities
even at a low temperature of 640 °C and under atmosphericpressure.
The heteroepitaxial growth of these monolayers produced
six diﬀerent in-plane heterostructures. Figures 2 and S3 display
the PL intensity maps, PL spectra, optical images, and Raman
spectra of MoS2/WS2, WSe2/WS2, MoS2/MoSe2, WSe2/
MoSe2, WS2/MoSe2, and MoS2/WSe2 monolayers grown ongraphite used as a substrate to allow for STM observations. In
each spectrum, PL peaks attributable to the inner and outer
TMDCs can be clearly seen. This result indicates that theinitially grown grains were able to maintain their structure andcomposition during the second growth step, as a result of thehigh degree of crystallinity of the initial grains. In almost every
sample, the second-growth TMDCs were preferentially formed
around the edges of the ﬁrst-growth grains rather than on thesurface. It is therefore evident that the edges were the mostactive sites in terms of incorporating diﬀusing atoms under thepresent conditions.
The edge-initiated heteroepitaxial growth observed in thiswork would be advantageous when fabricating multiheterostructures and nanoribbon structures. As a demonstration, we
prepared a multiheterostructure having a WS2/MoS2/WS2
structure by varying the precursors, as shown in Figure 3a−f.
The optical microscopy image and PL maps/spectra clearly
demonstrated that an MoS2 monolayer was sandwiched
between inner and outer WS2 monolayers within a single
triangular grain. Furthermore, the growth rate could be tuned
by adjusting the ﬂow rate of the gas used to transport the liquid
precursors, thus enabling control of the size of the epitaxially
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Figure 3. (a) Schematic illustrations of the heteroepitaxial growth of a WS2/MoS2/WS2 in-plane multiheterostructure. (b−f) Optical
microscopy image, PL intensity maps, and PL spectra of a WS2/MoS2/WS2 in-plane multiheterostructure. In the PL intensity maps, red andcyan correspond to the intensities of the MoS2 and WS2 PL peaks, respectively. Sharp peaks indicated by asterisks correspond to the Raman
modes of the graphite substrate. (g) Schematic illustration and (h) STM image (Vs = +1.0 V, It = 0.05 nA) of a MoS2 monolayer nanoribbon
grown from the edge of a WS2 monolayer. The white arrows indicate the MoS2/WS2 heterointerface.

reﬂected the crystal lattice of the underlying WS2, suggesting
the van der Waals heteroepitaxial growth of the WSe2 on theWS2. In addition, the grains exhibited a distinct moiré patternin the triangular lattice, as expected for a lattice-mismatchedvertical heterostructure such as that in Figure 4c. The period ofeach moiré pattern was approximately 8.2 nm, which isconsistent with the diﬀerence between the lattice constants ofWSe2 (3.28 Å) and WS2 (3.15 Å). These results are similar to
the moiré pattern previously observed in a WSe2/MoS2 vertical
heterostructure.35,36
To obtain further insight into the moiré pattern, STS was
performed. Figure 4d is the color scale map of dI/dV at Vs =−1.8 V reﬂecting change of the valence band LDOS
corresponding to the moiré pattern. We note that the LDOSof the conduction band hardly changed. Figure 4e exhibits thedI/dV curves of three diﬀerent points in Figure 4d. In thesecurves, the threshold voltages at positive and negative voltages
correspond to conduction band minimum (CBM) and valenceband maximum (VBM). Depending on the position, only the
VBM shifts by approximately 0.2 eV. The periodic change ofthe VBM can be clearly seen from color scale map based on
dI/dV curves in Figure 4f. Because the VBM of multilayer
TMDCs is located at the Γ point,25 these results indicate thatthe interlayer interaction aﬀects the Γ point of the valencegrown TMDC grains. This same process could also be used to
produce nanoribbon structures consisting of TMDC monolayers. In the present study, WS2 grains were initiallysynthesized, after which MoS2 monolayers were grown aroundthese earlier grains, applying a slow growth rate, as illustrated
in Figure 3g. This process resulted in the formation of
nanoribbon structures made of MoS2 monolayers, having
widths of 20 nm, as shown in the STM image in Figure 3h.
Here, the white arrows indicate the MoS2/WS2 heterointerface.
In this image, the diﬀerence in contrast between the MoS2 andthe WS2 regions is derived from their local density of state(LDOS) values.15,34 This nanoribbon could have a zigzag edge,
as discussed further on.
In addition to in-plane heterostructures, vertically stacked
heterostructures based on TMDCs could be produced using
the same system. As shown in Figure 4a, a second growth
TMDC layer was applied to the surface of an initial monolayer.
During these trials, we found that the nucleation density at the
TMDC surface could be enhanced by cooling the substrate toroom temperature before the second growth. In the case of thegrowth of WSe2 on a ﬁrst-growth WS2 monolayer on graphite,
STM observations revealed that small triangular WSe2 grains
were formed on the WS2 surface and that these grains had two
primary orientations (Figure 4b). This result clearly demonstrates that the crystal orientation of the WSe2 monolayer
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Figure 4. (a) Structural model and (b) wide-area (Vs = +1.5 V, It = 0.05 nA) and (c) magniﬁed (Vs = +1.5 V, It = 0.1 nA) STM images of a
WSe2/WS2 vertical heterostructure. (d) Color scale map of dI/dV at Vs = −1.80 V. (e) dI/dV curves of three diﬀerent points (A, B, C) in (d).
(f) Color scale map based on dI/dV curves calculated from the spatially resolved scanning tunneling spectroscopy (STS) spectra acquiredalong the white dotted line in (d).

Figure 5. (a and b) STM images (Vs = +1.0 V, It = 0.08 nA) and (c) a spatial map of the calculated total LDOS (top) and structural model(bottom) of the MoS2/WS2 heterointerface. The total LDOS was obtained by the integration of the DOS within the energy range from EF to
EF + 3 eV. (d) Averaged dI/dV spectra obtained from the MoS2, WS2 and heterointerface regions. (e) Color scale map of the dI/dV spectra
measured in the vicinity of the heterointerface. In (e), the red dotted line indicates the position of MoS2/WS2 heterointerface.

(W) atoms can be seen as bright (dark) dots under the present
measurement conditions, due to the diﬀerence in LDOS
between the two.15,34 The STM images are in good agreementwith a spatial map of the total LDOS obtained by ﬁrstprinciples calculations (Figure 5c). In this sample, the WS2 wasgrown from the edge of the MoS2, and it is noteworthy that the
Mo and W atoms were perfectly separated along theband and formation of an electronic superlattice corresponding
to the moiré pattern.35−37
To investigate the interface structures, atomically resolved
images of the interfaces of in-plane MoS2/WS2 monolayer
heterostructures were obtained by high-resolution STM, as
shown in Figures 5a,b and S4. In these images, the left andright sides correspond to MoS2 and WS2, respectively, and Mo
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nm zone around the interface, as shown in the color scale mapof the dI/dV spectra (Figure S6b). The CBM and VBM of the
MoS2 monolayer were located below the CBM and VBM of
the WS2 monolayer, respectively, indicating the formation of a
type II staggered gap at the heterointerface. A similar type II
staggered gap is observed for the MoS2/MoSe2 heterojunction(Figure S5b,c). Interestingly, at a diﬀerent grain, the CBM and
VBM of the MoS2 were located within the bandgap of the WS2,
indicating that the interface has a type I straddling gap asshown in Figure 5d,e, and the bandgaps are also changed to be
2.1 eV for the MoS2 and 2.5 eV for WS2. The LDOS of MoS2/
WS2 was uniform at least within the scanning area (around 6
nm × 6 nm) along the interface (Figure S7). This modulationof band structure is probably caused by grain-dependent latticestrain. The grain-dependent strain is actually found in the PL
spectra. As shown in Figure S8, the PL peak positions of MoS2
and WS2 at the WS2/MoS2 interfaces are shifted in diﬀerent
triangle-shaped grains and even within the identical grain. The
peak positions of MoS2 varies more than these of WS2. The
variations suggest that the lattice strain depends on the size of
MoS2 and WS2 in individual grains. In other words, theseresults suggest that the band alignment of the MoS2/WS2 inplane heterostructure can be tuned by applying lattice strain.
The strain-induced conversion from type II to type I bandalignment was also observed for the WSe2/MoS2 heterostructure by STM/STS measurements.36 In the case of WSe2/
MoS2, the bandgap of MoS2 is reduced by approximately 0.5
eV due to large strain induced by the lattice mismatch between
MoS2 and WSe2. Our results provide additional evidence that,even for the WS2/MoS2 in-plane heterostructure, the latticestrain aﬀects the band alignment at the interface.
These experimental results were subsequently comparedwith the calculated LDOC and band diagram for the MoS2/
WS2 interface (Figure S6c,d). The calculated bandgaps of MoS2
and WS2 were 1.8 eV . The associated band edges weresmoothly shifted within a region of 1 nm around the interface,
and the MoS2 was found to have lower CBM and VBM thanthe WS2. This theoretical result suggests that the MoS2/WS2
heterointerface intrinsically has the type II band alignment, andthe formation of type I band alignment (Figure 5d,e) could bederived from the strain eﬀect. We note that the type II
staggered gap coincides with the band alignment in MoS2/WS2
vertical heterostructures.39
The presence of strain eﬀect is also supported by a small dip
in the experimental VBM data in the vicinity of the interface(Figure 5d,e). This phenomenon can possibly be attributed tothe presence of a positive ﬁxed charge, likely resulting from thepiezoelectric eﬀect induced by lattice strain at the interface, as
discussed in our previous report.15 This potential modulationof the interface could have applications to the switching ofelectronic transport across the interface as well as carrier
accumulation along the interface.

heterointerface. Furthermore, the STM images demonstratethat the heterointerface junction consisted of atomicallystraight zigzag edges of Mo and W atoms and that the MoS2
and WS2 honeycomb lattices were smoothly connected with
one another. From these results, it is apparent that the present
process enabled heteroepitaxial growth initiated at the edges of
the ﬁrst-growth grains. We note that the formation of
atomically sharp interface and the isotropic growth ofindividual triangle-shaped grains during heteroepitaxy (Figure2) indicate that individual metals and chalcogens aredecomposed from the metal−organic precursors and then are
adsorbed on the edge while diﬀusing. In other words, thevapor−solid−solid (VSS) mechanism plays a major role in thepresent epitaxial growth rather than the vapor−liquid−solidmechanism as observed in recent salt-assisted CVD.38
In previous studies, the interfaces tended to have disorderedstructures, likely due to the low degree of crystallinity as well asdegradation during the second growth process, partly as aresult of exposure to air.16,18,19,37 In the case of the technique
reported herein, a high-quality interface was possible because
the process was continuous, proceeded without exposure to air,and exposed the liquid precursors to a relatively low growthtemperature. It should be noted that optimization of precursor
supply rates and growth temperature for each TMDC was also
necessary to obtain atomically straight edges. The formation of
atomically straight interface can be explained by the following
reasons. First, the suppression of kink formation could beachieved by smaller thermal ﬂuctuation under low temperature
growth. Low temperature growth is enabled by the present
liquid precursors, which are easily decomposed at lower
temperature compared with solid precursors. Additionally, thecontinuous growth without air exposure of samples can avoid
the adsorption of impurities and oxidation at reactive grainedges. The controlled supply rates of liquid precursors would
be also important to improve uniformity of edge structures. We
note that a similar advantage is expected for the metal−organic
precursors used in the previous works.22,23,31 A majordiﬀerence between the present liquid precursors and thewell-used organic precursors such as Mo(CO)6 and W(CO)6 is
the presence of oxygen atoms, which may inﬂuence the quality
and grain size of samples. For the heterostructures with
diﬀerent chalcogen atoms such as MoS2/MoSe2, both MoS2
and MoSe2 tend to alloy around the heterointerfaces probably
due to the chalcogen substitution and migration during
growth, as observed in our STM result (Figure S5). To
fabricate an atomically straight interface of the heterostructureswith diﬀerent chalcogen atoms, it would be necessary to
conduct the growth at a lower temperature.
The spatial variations in the LDOS of MoS2/WS2 andMoS2/MoSe2 monolayer heterointerfaces were investigated by
STS analyses. Figure S6a presents the spectra obtained fromaveraged dI/dV values (where I and V are the tunnelingcurrent and bias voltage, respectively) for the MoS2, WS2, andheterointerface regions. Based on the threshold voltages for the
dI/dV values, the MoS2 and WS2 bandgaps were estimated tobe 2.2 and 2.4 eV, respectively, both of which are consistent
with previous reports (2.2 eV for MoS2 and 2.4 eV for WS2 atroom temperature,39 2.4 eV for MoS2 at 77 K40). The
diﬀerence of bandgaps could result from the presence of latticestrain in the in-plane heterostructures at 83 K, which can be
introduced by a mismatch in the lattice constants and/or
thermal expansion coeﬃcients.41−43 The valence and conduction band edges were smoothly connected within the 2−3

CONCLUSIONSIn summary, we have demonstrated the continuous heteroepitaxial growth of various 2D heterostructures with atomicallystraight heterointerfaces by switching between liquid precursors. The present liquid precursors enable a continuously
and controllable supply and the growth of high-quality
specimens even at a relatively low temperature of 640 °C, as
recently demonstrated by other metal organic precursors.23
Our study provides another choice for the growth of TMDCbased heterostructures. This technique also allowed us to7532

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strips with widths of 10 zigzag chains of MoS2 and WS2 molecules inan alternating arrangement. The unit cell used in these calculations isshown in Figure S9. In this model, the TMDC heterosheet was
separated from its periodic images by a 0.89 nm vacuum space. The
geometric structure of the heterosheet was optimized based on theexperimental lattice constant of MoS2 until the remaining force acting
on the atoms was less than 0.005 Ryd/Å.

demonstrate the formation of MoS2 nanoribbons with a widthof 20 nm via epitaxial growth from the edges of TMDC
monolayers, as well as vertical heterostructures. The present
process has several advantages as follows. We have achieved a
direct growth of samples on graphite, which facilitates the
STM/STS study without the need for wet or dry transfer. We
found that the heterointerfaces exhibit perfect zigzag edges
without alloying and revealed the LDOS of both the lateral andthe vertical heterostructures. The obtained band diagram of
MoS2/WS2 in-plane heterostructure was compared with theresults of ﬁrst-principles calculations. This work suggests ameans of fabricating a wide variety of in-plane and verticalsuperlattices, nanoribbons, and nanowires.

ASSOCIATED CONTENT
S Supporting Information*

The Supporting Information is available free of charge on the
ACS Publications website at DOI: 10.1021/acsnano.8b07991.
Table of growth parameters for WS2, WSe2, MoS2, andMoSe2 monolayers in the present study (Table S1);
optical microscopy images of WS2 grown on SiO2/Si via
MOCVD with and without NaCl (Figure S1); optical
microscopy images of MoSe2 grown with diﬀerent
supply rates for the metal and chalcogen precursors(Figure S2); optical image, PL intensity maps, andRaman spectra of MoS2/WS2, WSe2/WS2, MoS2/MoSe2,
WSe2/MoSe2, WS2/MoSe2, and MoS2/WSe2 heterostructure (Figure S3); STM images of the MoS2/WS2
heterointerface from diﬀerent regions (Figure S4); STM
image and color scale map of the dI/dV spectra acquiredin the vicinity of the heterointerface and averaged dI/dV
spectra obtained from the MoS2/MoSe2 heterostructure(Figure S5); averaged dI/dV spectra, color scale map of
the dI/dV spectra acquired in the vicinity of theheterointerface, calculated LDOS, and associated color
scale map for the MoS2/WS2 heterostructure (FigureS6); STM image and dI/dV map of the area around the
MoS2/WS2 heterointerface (Figure S7); PL spectra andpeak position plot of MoS2 and WS2 at the WS2/MoS2
interfaces in diﬀerent triangle-shaped grains and within
the same grain (Figure S8); and unit cell of a superlatticeconsisting of WS2 and MoS2 strips used for theoretical
calculations (Figure S9) (PDF)

EXPERIMENTAL METHODSIn this work, four liquid precursors were employed during MOCVD
growth: (t-BuN) 2 W(NMe 2 ) 2 , (t-BuN) 2 Mo(NMe 2 ) 2 , (tC4H9)2S2, and (C2H5)2Se2 (Gas-Phase Growth, Ltd.). TMDC crystals
were grown on either SiO2/Si or Kish graphite ﬂakes (Type B,
Covalent Materials Co.), as reported previously.33 The substrates
were placed in a quartz tube (3 cm diameter, 70 cm long) with NaCl
(100 mg, Wako Pure Chemical Industries, Ltd.), after which the tube
was ﬁlled with a N2/H2 gas mixture at a ﬂow rate of 594/6 cm3/min.
NaCl powder was palced at 1−3 cm upstream of substrates, and its
temperature is around 540−590 °C. The temperature of the substrate
was gradually increased to the growth temperature (590 to 640 °C)
over 20 min using an electric furnace. When the substrate temperature
was at the set point, the precursors were supplied by bubbling N2
through precursor reservoirs to the substrate for 60 min underatmospheric pressure. During this step, the liquid precursors were
kept at room temperature (24 to 25 °C). The growth conditions used
to obtain MoS2, WS2, WSe2, and MoSe2 monolayers are summarizedin Table S1. Following this growth process, the quartz tube wasimmediately cooled using an electric fan.
In-plane heterostructures and nanoribbon structures were obtainedby conducting a second MOCVD immediately after the ﬁrst growth,without cooling the furnace. The quartz tube was ﬁlled with N2/H2
gas at a ﬂow rate of 594/6 cm3/min for 2 min at the desired growthtemperature (when making MoS2/WS2, WSe2/WS2, MoS2/MoSe2,
and WSe2/MoSe2) or for 30 min at room temperature (when making
WS2/MoSe2 and MoS2/WSe2) to remove residual precursor from theﬁrst growth. To keep the room temperature, the furnace was moved
during removal of residual precursor. The second TMDC growth wasthen conducted in the same manner as the ﬁrst over a span of 30 min.
The multiheterostructures were prepared by an initial WS2 growth
step over 60 min, a second MoS2 growth over 30 min, and a third WS2
growth over 30 min.
For the growth of vertical heterostructures, the quartz tube was
ﬁlled with N2/H2 gas at a ﬂow rate of 594/6 cm3/min for 30 min atroom temperature after the growth of a WS2 monolayer on graphite.
WSe2 growth was then conducted in the same manner as the ﬁrst
growth over 30 min.
Samples were observed using optical microscopy (Nikon, EclipseLV100D). Raman and PL spectra and PL intensity maps wereacquired using a micro-Raman spectroscope (Renishaw, inVia) with a
532 nm excitation laser. STM/STS assessments were carried out using
a low-temperature STM system (Omicron, LT-STM) operating at 83
K.
All calculations were based on the density functional theory(DFT)44,45 and implemented using the STATE software package.46
We employed the generalized gradient approximation47 and ultrasoft
pseudopotentials48 to describe the exchange-correlation potential andthe electron−ion interaction, respectively. The valence wave functions
and deﬁcit charge density were expanded in conjunction with a planewave basis set, applying cutoﬀ energies of 25 and 225 Ryd,
respectively. The Brillouin zone integration was performed with a 1
× 16 × 1 uniform k-point grid. To simulate the border between WS2
and MoS2, we constructed a superlattice consisting of WS2 and MoS2

AUTHOR INFORMATION
Corresponding Author*(Y. Miyata) E-mail: ymiyata@tmu.ac.jp.
ORCID

Yu Kobayashi: 0000-0001-8305-0991
Hidemi Shigekawa: 0000-0001-9550-5148
Yasumitsu Miyata: 0000-0002-9733-5119
Author Contributions

Y.K. prepared all experimental samples and performed theoptical analyses. S.Y., H.M., and K.M. carried out the STM/
STS observations, with H.M. and O.T. providing technicalassistance. M.M. and S.O. performed the theoretical
calculations. Y.Miyata developed the concept and supervised
the project. Y.K., S.Y., M.M., S.O., and Y.Miyata prepared theﬁgures and wrote the paper. All authors discussed the results
and commented on the manuscript.
Notes

The authors declare no competing ﬁnancial interest.

ACKNOWLEDGMENTS
This work was supported by a JST CREST grant (number
JPMJCR16F3) and by JSPS KAKENHI grants (numbers
JP18H01832, JP17H06088, JP15H05412, and JP16H00918).
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