Efficacy of boron nitride encapsulation against plasma-processing of 2D semiconductor layers

Abstract

Two-dimensional (2D) transition metal dichalcogenides (TMDCs) are the subject of intense investigation for applications in optics, electronics, catalysis, and energy storage. Their optical and electronic properties can be significantly enhanced when encapsulated in an environment that is free of charge disorder. Because hexagonal boron nitride (h-BN) is atomically thin, highly crystalline, and is a strong insulator, it is one of the most commonly used 2D materials to encapsulate and passivate TMDCs. In this report, we examine how ultrathin h-BN shields an underlying MoS 2 TMDC layer from the energetic argon plasmas that are routinely used during semiconductor device fabrication and postprocessing. Aberration-corrected scanning transmission electron microscopy is used to analyze defect formation in both the h-BN and MoS 2 layers, and these observations are correlated with Raman and photoluminescence spectroscopy. Our results highlight that h-BN is an effective barrier for short plasma exposures (<30 s) but is ineffective for longer exposures, which result in extensive knock-on damage and amorphization in the underlying MoS 2.

Full text

J. Vac. Sci. Technol. A 39, 032201 (2021); https://doi.org/10.1116/6.0000874 39, 032201
© 2021 Author(s).
Efficacy of boron nitride encapsulation
against plasma-processing of 2D
semiconductor layers
Cite as: J. Vac. Sci. Technol. A 39, 032201 (2021); https://doi.org/10.1116/6.0000874
Submitted: 17 December 2020 . Accepted: 22 February 2021 . Published Online: 16 March 2021
Pawan Kumar, Kelotchi S. Figueroa, Alexandre C. Foucher, Kiyoung Jo, Natalia Acero, Eric A. Stach, and
Deep Jariwala
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Efficacy of boron nitride encapsulation against
plasma-processing of 2D semiconductor layers
Cite as: J. Vac. Sci. Technol. A 39, 032201 (2021); doi: 10.1116/6.0000874
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Submitted: 17 December 2020 · Accepted: 22 February 2021 ·
Published Online: 16 March 2021
Pawan Kumar,
1,2
Kelotchi S. Figueroa,
3
Alexandre C. Foucher,
2
Kiyoung Jo,
1
Natalia Acero,
4
Eric A. Stach,
2,5,a)
and Deep Jariwala
1,b)
AFFILIATIONS
1
Department of Electrical and Systems Engineering, University of Pennsylvania, Philadelphia, Pennsylvania 19104
2
Department of Materials Science and Engineering, University of Pennsylvania, Philadelphia, Pennsylvania 19104
3
Department of Physics and Electronics, University of Puerto Rico, Humacao, Puerto Rico 00791
4
Vagelos Integrated Program for Energy Research, University of Pennsylvania, Philadelphia, Pennsylvania 19104
5
Laboratory for Research on the Structure of Matter, University of Pennsylvania, Philadelphia, Pennsylvania 19104
a)
Electronic mail: stach@seas.upenn.edu
b)
Electronic mail: dmj@seas.upenn.edu
ABSTRACT
Two-dimensional (2D) transition metal dichalcogenides (TMDCs) are the subject of intense investigation for applications in optics,
electronics, catalysis, and energy storage. Their optical and electronic properties can be significantly enhanced when encapsulated in an
environment that is free of charge disorder. Because hexagonal boron nitride (h-BN) is atomically thin, highly crystalline, and is a strong
insulator, it is one of the most commonly used 2D materials to encapsulate and passivate TMDCs. In this report, we examine how ultrathin
h-BN shields an underlying MoS
2
TMDC layer from the energetic argon plasmas that are routinely used during semiconductor device
fabrication and postprocessing. Aberration-corrected scanning transmission electron microscopy is used to analyze defect formation in both
the h-BN and MoS
2
layers, and these observations are correlated with Raman and photoluminescence spectroscopy. Our results highlight
that h-BN is an effective barrier for short plasma exposures (<30 s) but is ineffective for longer exposures, which result in extensive
knock-on damage and amorphization in the underlying MoS
2
.
Published under license by AVS. https://doi.org/10.1116/6.0000874
I. INTRODUCTION
The isolation of layered van der Waals crystals into atomically
thin, two-dimensional (2D) structures has led to significant new
insights into condensed matter physics,
1,2
which have in turn led to
fundamentally new electronic device designs.
3,4
Significant effort
has been spent on the study of synthetic routes, fundamental physi-
cal phenomena, and device properties in these systems. However,
there has been less focus on device processing. Ultimately, the
applications of all 2D materials will require precise control over
crystalline quality, thickness (layer number), and, thus, over device
processing conditions. With this background, we address an old
but relevant problem associated with a ubiquitous process in
microelectronics: plasma processing.
Plasma processing is widely used to clean, functionalize, and
passivate surfaces, as well as to etch materials.
5–11
It has been
applied to 2D materials since the early days of graphene device
research.
12–15
However, it was soon realized that energetic plasmas
could affect the structural and chemical stability of 2D materials
and thereby degrade lateral transport in electronic devices.
16–18
Therefore, high-quality, stable, and scalable encapsulants are
needed to protect 2D device channels, and they continue to be an
object of research. However, it is unclear if plasma processing leads
to charge incorporation in encapsulating layers that degrades
carrier transport in a structurally intact 2D material or if it directly
damages the material. Often charge transport measurements are
used to infer the role of defect formation in these cases, but trans-
port measurements only provide indirect evidence of defect intro-
duction. These facts motivate us to explore how plasma etching
conditions affect both the encapsulation layer and the active
channel using a direct approach.
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In all high-performance 2D semiconductor devices, it is essen-
tial to isolate or encapsulate the active semiconductor layers to
limit charge inhomogeneities and exposure to processing chemicals.
Encapsulation is needed to protect electronic devices such as tran-
sistors, where the channel is buried under a dielectric insulator,
and optoelectronic devices, where the junction is buried under con-
tacts and barrier layers. Charge inhomogeneity results from
trapped charges, dangling bonds, and dipoles of ionic bonds, all of
which impede electronic transport and inhibit radiative recombina-
tion in 2D layers. Organic layers (polymers) and flat, highly crystal-
line, and nearly covalent materials such as hexagonal boron nitride
(h-BN)
19
have proven to be effective substrates and encapsulants
for 2D channels and active layers.
20–22
However, polymers and
small organic molecules are susceptible to thermal damage as well
as swelling and dissolution upon solvent exposure. This means that
they are unsuitable as permanent encapsulants during semiconduc-
tor device fabrication and processing. However, h-BN possesses
high chemical and thermal stability, making it a potentially supe-
rior encapsulant. Several schemes for direct growth and transfer of
h-BN encapsulated graphene and 2D semiconductor devices have
been developed.
23–26
In these studies, a postlithography etching
step is essential for defining channels and contacts. Although
several studies assume that h-BN is adequately protecting the
underlying active 2D layer,
27–30
there have been no systematic
studies that provide mechanistic insight into its effectiveness. In
this study, we perform a systematic investigation of the efficacy of
the h-BN layer as an encapsulant. We correlate optical spectroscopy
and atomic-resolution imaging analysis to understand how plasma
dose variation, sequential plasma exposure, and encapsulant and
underlayer thickness affect the rates of damage accumulation.
II. EXPERIMENTAL SECTION
A. Materials and methods
Mechanically exfoliated 2D MoS
2
layers were prepared using
the conventional scotch tape method, as described elsewhere.
31
The thickness of the exfoliated layers was intentionally chosen such
that each sample could be reproduced for a number of repetitive
analyses. Exfoliated h-BN and MoS
2
layers were transferred to the
SiO
2
/Si substrate by dry transfer technique utilizing a poly-
dimethyl siloxane (PDMS) stamp. The dry transfer technique uses
a motorized micromanipulator stage (X-Y-Z axes), attached with
an optical microscope along-with with a home-made heating stage
(based on pyroelectric material). After transferring MoS
2
layers to
SiO
2
/Si substrates, samples were annealed in a quartz tube furnace
in a closed gas (Ar + H
2
) environment to remove all PDMS con-
taminants. Annealing was performed at 300 °C for 4 h to clean con-
taminants as well as release the strain developed during the
pressure-based dry transfer method. It is worth noting that the 2H
phase MoS
2
is extremely stable at 300 °C in a reducing atmosphere.
This has been confirmed by several prior studies and no evidence
of structural phase changes or defect formation has been observed
at our chosen annealing temperatures.
32,33
Similarly, h-BN and
MoS
2
layers were transferred to a dedicated SiNx TEM grid
(Norcada Inc., 3 × 3 array of 100 μm diameter holes) and annealed
in the same manner to remove PDMS contaminations prior to irra-
diation and subsequent scanning transmission electron microscopy
(STEM) characterization. Two different plasma irradiation systems
are used here for the different samples at different exposure times.
Ultra-pure Ar gas (99.995%) was used in all the plasma exposure
analysis. A dedicated plasma cleaning system was available for
holding TEM holders such that we can treat samples for multiple
exposure times while the TEM grid installed into TEM holder.
Plasma exposure time for TEM analysis (sample on TEM grid) vs
Raman analysis (sample on SiO
2
/Si substrate) is different as sample
positioning distance from the plasma ignition source is different in
both the systems based on the available configuration.
B. Characterization
Raman spectroscopy as well as photoluminescence measure-
ment is carried out using the LabRAM HR Evolution HORIBA
system. A 633 nm laser with a spot size ∼0.5 μm, 1% laser power
was used for diagnosed defect related analysis in h-BN as well as
MoS
2
layers. For the case, PL analysis, 405 nm laser, was used with
0.1% laser power (0.25 μW) with 1 s acquisition time. Optical
microscope from Olympus, USA was used to capture and analyze
all the samples before and after plasma treatment. The plasma irra-
diation system (Tergeo Plasma Cleaner, PIE Scientific) was used
with 65 sccm of Ar (99.995% purity) flow and 50-W transmitted
RF power under 0.35 Torr base vacuum environment. The same
parameters were utilized in the case of a dedicated TEM plasma
cleaner system (Gatan, Solarus 950). High-Angle Annular Dark
Field (HAADF)-STEM has been used to directly visualize all the
samples for defects evolution. An aberration-corrected JEOL
NEOARM STEM, operating at an accelerating voltage of 200 kV
and convergence angle of 25–29 mrad, is used for all samples. For
the JEOL NEOARM STEM, the condenser lens aperture was 40 μm
with a camera length of 4 cm for imaging and the probe current
was 120 pA. All of the captured STEM images were collected using
GATAN GMS software and associated Gatan bright-field and high-
angle annular dark field detectors. Experimentally acquired STEM
images are smoothed using the adaptive gaussian blur function
(with a radius of 1–2 pixels) available in IMAGEJ.
III. RESULTS AND DISCUSSION
During sample preparation, device fabrication, and postpro-
cessing, a 2D material or heterostructure is likely to encounter
several forms of energetic radiation sources. Radiation sources that
are used include electron, ion (Ar, He, and Xe), and photon (laser,
UV) fluxes. However, ion beams and UV light are the most rou-
tinely adopted in etching and lithography processes, respectively.
Here, we investigate the etching behavior of h-BN, MoS
2
, and their
heterostructures when subject to Ar
+
ion plasma exposure. Argon
ion plasmas are the most commonly used as argon does not form
chemical bonds with the sample due to its inherent inertness.
Argon is also a heavy enough ion that can provide sufficient kinetic
energy to etch samples at reasonable accelerating voltages.
Figure 1 shows a schematic of the physically stacked 2D heter-
ostructure samples used in this study. Thin flakes of h-BN and
MoS
2
were mechanically exfoliated from the bulk using Scotch
tape. They were subsequently transferred onto oxidized silicon
wafers or dry PDMS stamps. Both heterostructures and samples
transferred onto STEM heating platforms were created using direct
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transfer via viscoelastic PDMS stamps. Further details of the
sample preparation and transfer methods are provided in Sec. II A.
Raman spectroscopy is used to provide a global measure of
defect formation as a function of plasma exposure for both h-BN
shield and unshielded MoS
2
(Fig. 2). Figure 2(a) presents an optical
microscopy (OM) image of a stacked heterostructure of h-BN
(circled). There are also regions of pure h-BN and pure MoS
2
adja-
cent (arrows). Height images and corresponding h-BN thickness
profiles are presented for the same heterostructure (h-BN/MoS
2
)in
Fig. 2(b). Raman spectra as a function of Ar
+
exposure time
from the marked unshielded and shielded MoS
2
regions shown in
Fig. 2(c). The distinct peak at 226 cm
−1
[labeled LA (M)] has been
associated in prior work with the appearance of defects arising
from the scattering of phonons at the Brillion zone edge.
34–36
Similarly, E
2g
and A
1g
mode intensities decrease with increase of
plasma exposure time, confirming that the lattice experiences
increasing damage with exposure time. Control data from the h-BN
layer is presented as Figs. S1 and S2 in the supplementary mate-
rial,
49
and changes in optical contrast in the OM image are pre-
sented in Fig. S3 in the supplementary material.
49
Pristine MoS
2
has a very small LA (M) Raman mode as a monolayer and is negli-
gible for few-layer MoS
2
: this LA mode most likely originates from
defects created during the exfoliation and transfer process. Two
and half minutes of plasma exposure leads to a significantly
enhanced LA (M) signal, indicating the formation of a significant
quantity of defects. After 5 min of plasma exposure, the intensity of
the LA (M) Raman mode is of the same magnitude as the lattice
A
1g
(M) –LA (M) phonon mode, indicating significant damage
accumulation.
We have also examined the Ar
+
plasma exposure effects for
shorter intervals of time (within a 1-m duration, using 20 s step
sizes) across different layer thicknesses of unshielded MoS
2
and
report those results in Fig. S4 in the supplementary material.
49
We
find that defect formation occurs at levels measurable via Raman
Spectroscopy upon Ar
+
plasma exposure for time scales as short as
30 s for an unprotected monolayer MoS
2
sample. In contrast, Ar
+
plasma exposure times of 2.5 and 5 min showed a less significant
rise in the LA (M) Raman signal for the region of the MoS
2
flake
that was shielded under a few layers of h-BN, as shown in Fig. S5 in
the supplementary material.
49
While Raman spectroscopy is a
reliable way to ascertain lattice damage, Raman signals from 2D
materials are inherently inefficient. Photoluminescence is a far more
sensitive measure of crystal quality and defect density for a high-
quality, direct band gap semiconductor. The smallest deviation from
perfect crystalline order can induce nonradiative recombination that
reduces the primary PL efficiency.
37
Deviations from crystalline
order can also introduce trap states that lead to photoluminescence
that is red-shifted from the original PL peak location.
37,38
We have recorded the PL response after plasma exposure for
shielded and unshielded monolayer MoS
2
.Figure 2(a) presents an
FIG. 1. (a) Schematic representing MoS
2
/h-BN partially overlapping van der Waals heterostructures under investigation in this study before and after Ar
+
plasma exposure.
(b) Cross-sectional schematic of the heterostructure. (c) Corresponding STEM images show atomic-scale structural changes across shielded and unshielded regions after
the 15 s exposure of plasma.
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OM image of the heterostructure sample from which we obtained
PL spectra. Region 1 has a monolayer of MoS
2
that is shielded
under a thin h-BN flake (∼11.2 nm thin), and region 2 is
unshielded monolayer MoS
2
. The PL data are presented as
Figs. 2(d) and 2(e). The primary peak at ∼652 nm (∼1.9 eV) is
caused by X
0
-exciton emission,
39
and the peak at ∼740 nm
(∼1.67 eV) is known to be a defect bound exciton peak.
40,41
Similar
PL spectra for the low incident light intensity have observed previ-
ously in the literature for MoS
2
.
42–44
The PL intensity of the
unshielded MoS
2
region before irradiation (“pristine”) is shown in
red in Fig. 2(d). The PL intensity of the unshielded MoS
2
flake is
comparatively weaker than the shielded region for pristine samples.
This is an optical effect related to the formation of the heterostruc-
ture: the h-BN has a higher index of refraction, while the SiO
2
wafer has a lower index of refraction. The higher index medium
thus enables increased light extraction. With increasing plasma
exposure, we see two effects, which we summarize in Fig. 2(e).
First, with increasing Ar
+
plasma exposure, the peak associated
with defect bound exciton emission (X
B
) increases, clearly indicat-
ing the increase in defect content in the layer as the incident ener-
getic Ar
+
ions penetrate the h-BN shield. Second, there is a
concurrent, correlated decrease in the primary neutral exciton (X
0
)
peak intensity. Interestingly, the h-BN/MoS
2
heterostructure shows
a sudden increase in the MoS
2
-X
B
intensity with 3.5 min and
greater plasma exposure time. These data suggest that the h-BN
layer is an effective shield in the heterostructure up to a certain
threshold of plasma exposure only.
Optical and vibrational spectroscopies are an effective way to
track general trends in damage accumulation. However, they do not
provide atomic-scale information about the accumulation of indi-
vidual defects. To complement these spectroscopies, we have per-
formed extensive characterization using aberration-corrected
high-angle annular dark field (HAADF) scanning transmission
electron microscopy (STEM) imaging. Flake samples were prepared
using the same methodology described in Fig. 1 and transferred
onto a SiNx TEM grid that contains separated 100 μm diameter
holes. This allowed the heterostructures to be irradiated with Ar
+
plasma without damaging the underlying SiNx membrane. Details
of the plasma irradiation system used for analyzing all STEM
samples are described in Sec. II.
A basic MoS
2
structural model and low magnification as well
as atomically resolved HAADF-STEM imaging for pristine and
FIG. 2. Raman and photoluminescence characteristics of shielded vs unshielded monolayer MoS
2
for different Ar
+
plasma exposure times. (a) Optical microscope image
of few-layer MoS
2
shielded by few-layer h-BN, and the corresponding (b) AFM height image of heterostructure region along-with height profile of the h-BN layer
(∼11.2 nm) in its pristine form. (c) Raman spectra at different plasma exposures for unshielded and shielded monolayer MoS
2
. (d) PL response of monolayer MoS
2
ana-
lyzed for the same h-BN shielded monolayer MoS
2
region shown in the optical micrograph and corresponding (e) PL peak of X
0
and X
B
intensity variation with the increas-
ing plasma exposure.
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different plasma exposed MoS
2
regions are shown in Figs. S6 and S7
in the supplementary material.
49
HAADF-STEM imaging was
performed after successive plasma exposures with varying time inter-
vals (5, 10, and 15 s). This leads to four sample conditions: “pristine”
(0 s exposure) and 5, 15, and 30 s of cumulative exposure, respec-
tively, as shown in Fig. 3. We compare the defect evolution analysis
for two different regions of the same sample, an h-BN shielded
MoS
2
heterostructure versus an unshielded (bare) MoS
2
region. It is
worth noting that even though a vertical van der Waals heterostruc-
ture of h-BN over MoS
2
is being imaged in transmission in
Figs. 3(a)–3(d), nearly all of the image signals are from the MoS
2
layer. There are two reasons for this. First, aberration-corrected
STEM images have a very small depth of field, and we maintained
the focus on the MoS
2
layer.
45
Second, the intensity in HAADF
images scales with Z
1.65
,
45
and both Mo and S are significantly
heavier than B and N. Therefore, nearly all the signal that comprises
images in Fig. 3 arises from the MoS
2
layer. Nonetheless, a
fast Fourier transformation (FFT) diffractogram can detect the
periodicity in the image from both the MoS
2
and h-BN lattices, and
it indicates that there is a 12.2° twist angle between them.
Figures 3(a)–3(c) show that the h-BN shielded MoS
2
heterostructure
experiences negligible defect formation until a cumulative 15 s of
plasma exposure. However, after 30 cumulative seconds of plasma
exposure, the h-BN/MoS
2
heterostructure shows visible lattice
damage (circled).
In contrast, bare MoS
2
shows signatures of visible lattice
damage after merely 5 s of exposure [Fig. 3( f )]. Contrast in the dif-
ferent HAADF-STEM images (from pristine to varying plasma
exposure time) is not quantitatively comparable since each individ-
ual image is self-normalized by the image acquisition software
(GATAN GMS) during image acquisition to its maximum intensity.
Gray-scale histograms for the images in Fig. 3 are presented in
Fig. S8 in the supplementary material.
49
Each histogram has been
normalized to the respective maximum image intensity and scaled
to a common distribution based on the mean intensity and stand-
ard deviation. All the image processing and normalization carried
out using “Sci-kit”as well as “Fiji”to read and process the
images.
46,47
Lattice damage is induced by the plasma in both the
bilayer and the few-layer portion of the MoS
2
flake following just
5 s of exposure. Additional plasma exposure [Figs. 3( f )–3(h)] leads
to progressively more severe lattice damage, resulting in near
amorphization of the bilayer region after 30 s of plasma exposure
[Fig. 3(h)]. These images also indicate that the damage grows in
spatially localized regions with increasing exposure. In other words,
damage accumulates at defects introduced at earlier times not
through continued renucleation. We summarize that the defective
FIG. 3. Aberration-corrected HAADF-STEM images from pristine and Ar
+
plasma exposed samples. (a)–(d) Images from the h-BN/few-layer MoS
2
heterostructure region
showing minimal defect/damage creation up to 15 s (a)–(c) followed by significant damage at 30 s (d). In contrast, (e)–(h) STEM imaging from unshielded MoS
2
shows sig-
nificant lattice damage starting from 5 s plasma exposure (f ) with near amorphization of the bilayer region by 30 s exposure to Ar
+
plasma (h). Insets [(a) and (e)] are FFT
diffractograms indicating the presence of the h-BN/MoS
2
heterostructure vs bare unshielded MoS
2
, respectively; all images are at the same scale with a representative
scale bar shown in (a).
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regions created by that plasma are passivated by ambient atmo-
spheric species that accumulate during transfer from the plasma
chamber to the TEM column. This chemical passivation stabilizes
these regions, and subsequent plasma exposure leads to increased
damage at these undercoordinated sits. Additional low magnifica-
tion STEM images for both the heterostructure and bare sample
regions are shown in Fig. S6 in the supplementary material,
49
and
another set of atomic-scale STEM image showing clustered patches
of defected areas after 30 s of plasma exposure are presented in
Fig. S9 in the supplementary material.
49
We have performed elec-
tron energy loss spectroscopy (EELS) measurements after 15 s of
plasma exposure time (Fig. S10 in the supplementary material).
49
The atomically resolved EELS data showed the presence of
adsorbed oxygen, which are found in defective 2D MoS
2
basal
planes.
48
Mo and S signals are weaker since the thickness of the
MoS
2
layer (∼4–5 nm) is very thin in comparison with the total
sample thickness which includes a 20 nm SiNx membrane
underneath as well as the h-BN layer (∼10 nm) covering MoS
2
on
the top. We have also studied EELS for a heterostructure (h-BN/
MoS
2
) which undergoes continuous plasma exposure up to 20 s.
We have first acquired the core loss EELS spectra for oxygen in the
pristine state and then after 20 s plasma exposure from the same
region for unshielded as well as the shielded region of MoS
2
flake,
as presented in Fig. S11 in the supplementary material.
49
A slight
increased amount of oxygen has only observed. In the process of
transferring the samples from plasma processing system into the
TEM vacuum column, our samples get exposed to ambient air for
a maximum of 10 min. This observation supports our hypothesis
that ambient exposure stabilizes defects during transfer in and out
of the microscope.
To understand the effect of air exposure during sequential
plasma bombardment, we subjected a different heterostructure
stack to a continuous plasma exposure for 15 s. We observed wide-
spread damage and etching of the basal plane in the unshielded
FIG. 4. Characterization of a sample region following 15 s of Ar
+
plasma exposure. (a) Lower magnification HAADF-STEM imaging of a region containing just the h-BN
layer. Darker patches of damage are circled. (b) Atomic-scale HAADF-STEM image from a sample region containing both h-BN shielded and unshielded MoS
2
. The region
marked in yellow corresponds to the h-BN shielded MoS
2
region (left), while the right part of the image is the unshielded region (c). Atomic-scale image of the h-BN
shielded MoS
2
region shows no strong damage. (d) Atomically resolved STEM image of an unshielded MoS
2
region (from the pink dashed rectangular box shown in b)
shows substantial damage with an attached (inset) FFT profile.
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MoS
2
, as shown in Fig. 4. This is in stark contrast with the sequen-
tial plasma exposure combined ambient air exposure [vide supra,
Fig. 3(g)], where the lattice damage is more uniform. We see that
increasing the total plasma exposure creates a large number of
defects in the h-BN as well as the unshielded MoS
2
layer, as shown
in Figs. 4(a) and 4(b), respectively. Furthermore, the defects created
here are much larger and are not individual point defects, instead,
holes of 3–5 nm diameter form. This suggests that without the sta-
bilization provided by oxygen during the transfer process, there is
an accelerated accumulation of defects during irradiation.
Following the nucleation of an individual point defect, the under-
coordinated atoms can readily be knocked off their lattice sites,
which, combined with extended migration during the continuous
plasma exposure, could lead to the growth of substantially large
voids.
The h-BN/MoS
2
heterostructure remains remarkably intact, as
we can see in Figs. 4(b) and 4(c). This is consistent with the ability
of few-layer h-BN to be an effective shield, as shown in the PL data
of Fig. 2. We have also performed STEM imaging for a continuous
plasma exposure of 20 s. Figure S12 in the supplementary
material
49
presents AFM height images and corresponding thick-
ness profiles from the different section of h-BN/MoS
2
heterostruc-
tures. Furthermore, the same flake was analyzed using Raman
spectroscopy and the defect mode is clearly seen in the case of
unshielded MoS
2
(Fig. S13 in the supplementary material).
49
Afterward, STEM imaging from the same region of two different
sample configurations—unshielded and shielded MoS
2
—is pre-
sented after 20 s continuous plasma exposure in Fig. S14 in the
supplementary material.
49
We have also separately examined the
defect evolution, and corresponding lattice damage within just
the h-BN layers that are being used as a shield in the heterostruc-
ture the contrast on bare h-BN in Fig. 4(a) is not as clear. However,
the voids are visible and marked. We have observed similar void
formation and damage patches in h-BN as well upon continuous
plasma exposure, see Figs. S15–S17 in the supplementary material
49
for further details. We hypothesize that these localized voids act as
a channel for further penetration of plasma ions through layer by
layer etching, ultimately reaching the underlying layer, which then
causes localized damage in the MoS
2
(Fig. S9 in the supplementary
material).
49
Again, this is in stark contrast with sequential exposure
to plasma in multiple small steps with ambient air contact between
steps. The creation of defective patches and corresponding voids
can be easily seen in the AFM height image of h-BN layers
(Fig. S18 in the supplementary material).
49
IV. SUMMARY AND CONCLUSIONS
We have studied the dynamics of Ar
+
plasma-induced defect
generation and etching in atomically thin van der Waals hetero-
structures of h-BN and MoS
2
. We observe that h-BN effectively
shields underlying layers from plasma damage. An atomic-scale
imaging suggests that plasma-induced lattice damage is instanta-
neous for unshielded MoS
2
, whereas shielded MoS
2
is protected by
the h-BN until a certain extent of exposure, and that the extent of
damage as a function of the time of exposure depends upon the
h-BN thickness. Finally, we conclude that continuous plasma expo-
sure is more damaging as opposed to via sequential exposure.
These results indicate that h-BN encapsulation does provide
limited protection to underlying MoS
2
layers during plasma pro-
cessing but that the level of protection varies on a range of
parameters.
ACKNOWLEDGMENTS
This work was carried out in part at the Singh Center for
Nanotechnology at the University of Pennsylvania, which is
supported by the National Science Foundation (NSF) National
Nanotechnology Coordinated Infrastructure Program Grant No.
NNCI-1542153. D.J., E.A.S., and P.K. acknowledge primary support
for this work from via the NSF DMR Electronic Photonic and
Magnetic Materials (EPM) core program (Grant No. DMR-1905853)
as well as the University of Pennsylvania Laboratory for Research
on the Structure of Matter, a Materials Research Science and
Engineering Center (MRSEC) supported by the National Science
Foundation (No. DMR-1720530). K.S.F. was supported by the LRSM
MRSEC REU and the Penn-UPRH Partnership for Research and
Education in Materials (PREM), Program No. NSF-DMR-1523463.
N.A. and D.J. acknowledge support from Vagelos Integrated
Program for Energy Research at the University of Pennsylvania as
well as the Center for Undergraduate Research and Fellowships at
Penn. D.J. also acknowledges support for this work by the US Army
Research Office under Contract No. W911NF1910109. A.C.F. and
E.A.S. would like to acknowledge the Vagelos Institute for Energy
Science and Technology at the University of Pennsylvania for a grad-
uate fellowship to A.C.F. The authors thank James Horwath for
help/assistance in the STEM image histogram analyses. The authors
also thank Douglas Yates and Jamie Ford in the Singh Center for
Nanotechnology for help with the TEM/STEM measurements. There
are no conflicts of interest to declare.
DATA AVAILABILITY
Most of the data sets generated or analyzed during this study
are included in this article and its supplementary material.
Additional findings of this study are available from the correspond-
ing author upon reasonable request.
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See supplementary material at https://doi.org/10.1116/6.0000874 for detailed and
extra characterization of the used 2D layers utilizing optical microscopic images,
scanning transmission electron microscope, Raman, as well as PL spectroscopy.
ARTICLE avs.scitation.org/journal/jva
J. Vac. Sci. Technol. A 39(3) May/Jun 2021; doi: 10.1116/6.0000874 39, 032201-8
Published under license by AVS.