# Kinetically-controlled epitaxial growth of Fe3GeTe2 van der Waals
ferromagnetic films

Wenyi Zhou  The Ohio State University, Department of Physics, Columbus, Ohio
43210, USA  Alexander J. Bishop  The Ohio State University, Department of
Physics, Columbus, Ohio 43210, USA  Menglin Zhu  The Ohio State University,
Department of Materials Science and Engineering, Columbus, Ohio 43210, USA
Igor Lyalin  The Ohio State University, Department of Physics, Columbus, Ohio
43210, USA  Robert C. Walko  The Ohio State University, Department of Physics,
Columbus, Ohio 43210, USA  Jay A. Gupta  The Ohio State University, Department
of Physics, Columbus, Ohio 43210, USA  Jinwoo Hwang  The Ohio State
University, Department of Materials Science and Engineering, Columbus, Ohio
43210, USA  The Ohio State University, Department of Physics, Columbus, Ohio
43210, USA  Roland K. Kawakami  The Ohio State University, Department of
Physics, Columbus, Ohio 43210, USA<EMAIL_ADDRESS>
###### Abstract

We demonstrate that kinetics play an important role in the epitaxial growth of
Fe3GeTe2 (FGT) van der Waals (vdW) ferromagnetic films by molecular beam
epitaxy. By varying the deposition rate, we control the formation or
suppression of an initial tellurium-deficient non-van der Waals phase (Fe3Ge2)
prior to realizing epitaxial growth of the vdW FGT phase. Using cross-
sectional scanning transmission electron microscopy and scanning tunneling
microscopy, we optimize the FGT films to have atomically smooth surfaces and
abrupt interfaces with the Ge(111) substrate. The magnetic properties of our
high quality material are confirmed through magneto-optic, magnetotransport,
and spin-polarized STM studies. Importantly, this demonstrates how the
interplay of energetics and kinetics can help tune the re-evaporation rate of
chalcogen atoms and interdiffusion from the underlayer, which paves the way
for future studies of van der Waals epitaxy.

###### keywords:

van der Waals magnet, ferromagnetism, molecular beam epitaxy, growth kinetics

## Introduction

There is tremendous interest in ultrathin van der Waals magnets due to the
unprecedented tunability of magnetic properties via strain1, 2, 3,
electrostatic gating4, 5, and incorporation into heterostructures6, 7, 8, 9.
Fe3GeTe2 (FGT) (Figure 1) is one of the most well-studied vdW magnets due to
its strong perpendicular magnetic anisotropy (PMA), relatively high Curie
temperature ($T_{C}$) ranging from $\sim$200 K to above room temperature10,
11, 12, 13, 14, 15, and observation of novel phenomena such as spin-orbit
torque switching16, 14, 15, 17 and skyrmion spin textures18, 19, 20.

While many efforts have focused on small flakes exfoliated from bulk crystals,
epitaxial growth of thin films and heterostructures by methods such as
molecular beam epitaxy (MBE) provide unique opportunities. This includes
stabilizing materials far from equilibrium that cannot be realized in bulk
crystal form (e.g. vdW MnBi2Se421), performing atomic-scale manipulation of
interface structure and material properties, and enabling scaling up to large
areas for electronic and magnetic applications.

To date, there have been several reports of epitaxial growth of FGT and
related materials by MBE (e.g. FGT films on Al2O3(0001)22, GaAs(111)22,
Ge(111)23, and Bi2Te3(0001)13, 24, 17; Fe5GeTe2 on Al2O3(0001)25). However,
critical issues of the growth process such as the relative importance of
thermodynamic equilibrium (energetics) and growth/reaction rates (kinetics)
remain largely unexplored. Understanding these issues are important for
improving the material quality, as it relates to important properties of thin
films and heterostructures including interdiffusion and defect formation.

In this work, we demonstrate that kinetics play an important role in the
epitaxial growth of FGT films by MBE. For FGT grown on Ge(111), we find that
growth rate is a critical parameter that determines whether FGT grows directly
on Ge(111) with an atomically abrupt interface or whether a Te-deficient non-
vdW Fe3Ge2 alloy forms at the interface. Contrary to expectations for reactive
MBE growth26, 27, our studies show that the Fe3Ge2 alloy forms at lower growth
rates while the abrupt FGT/Ge(111) interface is obtained for higher growth
rates. This behavior suggests that energetic considerations favor the
formation of non-vdW Fe-Ge alloys, while the growth kinetics can favor the
stabilization of pure FGT. The high quality of the pure FGT films is
corroborated by their excellent magnetic properties, where $\sim$20 nm FGT
films on Ge(111) exhibit typical properties of bulk FGT including
perpendicular magnetic anisotropy with square hysteresis loops and a $T_{C}$
of $\sim$250 K. These results demonstrate that the interplay of energetics and
kinetics can help tune the re-evaporation rate of chalcogen atoms and
interdiffusion from the underlayer to facilitate the development of high
quality vdW epitaxial films.

## Experimental Section

The FGT films were deposited in a Veeco 930 MBE system with base pressure of
$2\times 10^{-10}$ Torr. The Ge(111) substrates (MTI Corporation) were
prepared by sonication in acetone and isopropyl alcohol for 5 min each before
being inserted into the MBE chamber. After degassing at 800 °C for 12 min, the
Ge substrates were cooled down to 325 °C for growth of the FGT films. For
synthesis, the films were grown by co-depositing Fe (99.99%, Alfa Aesar), Ge
(99.9999%, Alfa Aesar), and Te (99.9999%, United Mineral Corp.) with an atomic
flux ratio of 3:1:20, as measured by a beam flux monitor (BFM) and calibrated
by x-ray reflectivity (XRR) of elemental films. Prior to growth, the substrate
was exposed to the Te flux for 5 minutes.

We employed a wide range of characterizations including in situ by reflection
high energy electron diffraction (RHEED) and ex situ by x-ray diffraction to
probe the surface and bulk crystal structure, respectively. The atomic-scale
structure was investigated by scanning tunneling microscopy (STM) and cross-
sectional scanning transmission electron microscopy (STEM) with elemental
analysis (i.e. energy dispersive x-ray spectroscopy (EDX)). Magnetic
properties were investigated by anomalous Hall effect (AHE), optical magnetic
circular dichroism (MCD), and spin-polarized STM.

For cross-sectional STEM imaging, specimens were prepared using conventional
focused ion beam lift-out technique with Helios NanoLab 600. Argon ion milling
under 900 and 500 eV was used subsequently to clean the surface amorphous
layer and minimize subsurface damage. High angle annular dark field (HAADF)
STEM images and EDX maps were collected using a probe-aberration corrected
Thermo Fisher Scientific Themis Z S/TEM operated at 300 kV, 20 mrad
convergence semi-angle.

For STM measurements, samples mounted on flag-style sample holders were
transferred from the MBE chamber to the STM using a UHV suitcase to keep the
sample under vacuum during the entire sample transfer. The STM measurements
were performed in a Createc LT-STM/AFM system with the sample held at a
temperature of 5 K. Spin-polarized STM measurements utilized a Cr tip and an
out-of-plane magnetic field.

We performed AHE and MCD measurements in an Oxford Spectromag magneto-optical
cryostat. To obtain out-of-plane hystersis loops, we applied a magnetic field
perpendicular to the surface of the sample. The AHE measurements were
performed using lock-in detection at 987 Hz and an excitation current of 100
$\mu$A (rms). For MCD measurements, we utilized a 100 $\mu$W, 532 nm laser
beam focused to a spot size of $\sim$150 $\mu$m. The helicity of the incident
beam was modulated at 50 kHz by a photoelastic modulator and the MCD of the
reflectivity was measured using an amplified Si photodiode and lock-in
detection.

(a) (a) (b) (c) (d) (e)

Figure 1: MBE-grown Fe3GeTe2 (FGT) films on Ge(111) at different growth rates.
(a) Atomic lattice of FGT, where the crystal unit cell is outlined. (b) RHEED
pattern of Ge(111) substrate. (c) RHEED pattern for 40 nm FGT grown at 0.06
Å/s (“Slow Growth FGT”). (d) RHEED pattern for 20 nm FGT grown at 0.12 Å/s
(“Fast Growth FGT”). (e) $\theta-2\theta$ x-ray diffraction scan of Slow
Growth FGT (40 nm). The red arrow indicates the presence of an unexpected
additional peak. (f) $\theta-2\theta$ x-ray diffraction scan of Fast Growth
FGT (20 nm), which does not exhibit the additional peak.

## Results and Discussion

### Growth and Structural Characterization

In studying the epitaxial growth of FGT on Ge(111), we observe a fascinating
dependence of the materials quality on the growth rate. To illustrate, we
employ two representative growth rates, namely a “Slow Growth” rate of $\sim
0.06$ Å/s and a “Fast Growth” rate of $\sim 0.12$ Å/s. Starting from streaky
RHEED patterns of the Ge(111) substrate (Figure 1), in both cases, we observe
streaky RHEED throughout the growth as shown in Figures 1 and 1. In addition,
the streak spacing matches the FGT in-plane lattice constant in both cases.
However, the X-ray diffraction (XRD) shows different features for Slow Growth
and Fast Growth, even though their RHEED patterns are very similar. For Slow
Growth FGT in Figure 1, it can be seen that the $\theta-2\theta$ XRD scan has
an unexpected extra peak of unknown origin (red arrow) beside the FGT(006)
peak. But for Fast Growth FGT in Figure 1, the extra peak in the XRD scan no
longer exists. This indicates a pure phase FGT film for Fast Growth and a
mixed phase FGT film for Slow Growth.

Figure 2: Scanning transmission electron microscopy (STEM) of Slow Growth FGT
on Ge(111). (a) Cross-sectional STEM image of Slow Growth FGT ($\sim 0.06$
Å/s). (b) Energy dispersive x-ray spectroscopy (EDX) of Slow Growth FGT as a
function of depth, showing Te deficiency during the initial growth.

To better understand this phenomenon, we perform cross-sectional STEM for
atomic-scale characterization of the structure and chemical composition. For
the Slow Growth FGT (Figure 2a), we observe a non-vdW phase at the initial
growth on the Ge substrate. But good vdW layers subsequently grow on top of
this initial non-vdW phase. The vdW layers are clearly visible in the STEM
image by the bright horizontal lines. The chemical composition shown in the
EDX scan (Figure 2b) provides some key insights. In the top part of the film,
the stoichiometry of Fe, Ge and Te is 3.13:1.00:2.03 (averaged in between
position 20 nm to 26 nm) indicating the formation of the vdW magnet Fe3GeTe2.
In the bottom part of the film, the Te atomic fraction goes to zero and the Fe
and Ge atoms have a ratio of $\sim$1.56 (averaged in between position 36 nm to
39 nm), which suggests the alloy Fe3Ge2. This conclusion is further supported
by the extra peak in the XRD corresponding to a lattice spacing of $\sim$5.04
Å, which is close to the lattice spacing of 5.010 Å 28, 29 for Fe3Ge2 along
the [001] direction (hexagonal c-plane). In addition, Fe3Ge2 has an in-plane
lattice constant of $a=b=3.998$ Å 28, 29, which is very similar to that of FGT
($a=b=3.991$ Å 23, 30, 31). This explains why it is hard to tell the
difference between FGT and Fe3Ge2 in RHEED during growth. Thus, the EDX, XRD,
and RHEED provide strong evidence that the non-vdW interfacial phase is
(001)-oriented Fe3Ge2.

The Fast Growth FGT on Ge(111) exhibits quite different behavior. As shown in
Figure 3, cross-sectional STEM imaging reveals that vdW FGT layers form
immediately at the start of growth. Again, the vdW layers are clearly visible
in the TEM by the bright horizontal lines. Moreover, a close look at the
interface shows an atomically sharp transition between the Ge substrate and
the FGT film, which indicates that the vdW FGT layers form immediately at the
interface. This is a great improvement over the Slow Growth FGT (Figure 2)
which has a transition layer of Fe3Ge2 in between the Ge substrate and the FGT
film. However, there is still room for improvement, as some defects are
visible.

(a) (a) (b)

Figure 3: Atomic scale structural characterization of Fast Growth FGT on
Ge(111). (a) Cross-sectional STEM image of Fast Growth FGT showing atomically
abrupt interfaces between the van der Waals FGT and the Ge substrate. (b)
Surface morphology of Fast Growth FGT measured by STM. Inset: Height profile
line cut across a step-edge indicated by the solid red line (tunneling
conditions: sample bias V = -154 mV, I = 0.22 nA). (c) Atomically-resolved
imaging of Te terminated FGT measured by STM. Inset: line cut profile
indicated by the solid red line with a periodicity of 3.99 Å (tunneling
conditions: sample bias V = 100 mV, I = 0.15 nA).

STM measurements on a Fast Growth FGT sample, shown in Figures 3 and 3, also
support the high quality of films grown by this method. Figure 3 shows large
atomically flat terraces with a line cut across a step edge. The height
profile of the line cut, displayed in the inset, indicates that the two
terraces differ in height by 8.2 Å, which is in good agreement with the height
of a single vdW layer of FGT. Atomic resolution imaging (Figure 3) reveals a
well ordered hexagonal lattice with an in-plane lattice constant of 3.99 Å,
consistent with the topmost Te layer of an FGT vdW layer. These results
indicate that the growth method established here facilitates good layer-by-
layer growth with consistent Te termination.

(c) (a) (b) (c)

Figure 4:  Out-of-plane magnetic hysteresis loops of Fast Growth FGT. (a)
Anomalous Hall effect measurement at 2 K. (b) Magnetic circular dichroism
(MCD) measurement at 2 K. (c) MCD for a series of temperatures, showing a TC
of $\sim$250 K. Curves have been vertically offset. (d) Spin-polarized
scanning tunneling microscopy measurement at 5 K. The data has been filtered
for noise reduction. Curves have been offset for clarity.

### Magnetic Properties

For the Fast Growth FGT/Ge(111) samples, we investigate the out-of-plane
hysteresis loops by measuring the anomalous Hall effect (AHE), reflection
magnetic circular dichroism (MCD), and spin-polarized STM. Figures 4 and 4
show out-of-plane hysteresis loops of a 20 nm Fast Growth FGT film measured by
AHE and MCD, respectively, at 2 K. The AHE measurement exhibits a square
hysteresis loop and a coercivity of $\sim$0.379 T, similar to values reported
in the literature. Similarly, the hysteresis loop measured by MCD exhibits
virtually identical characteristics as the AHE hysteresis loop measured in the
same cryostat. The temperature dependence of the MCD hysteresis loops is shown
in Figure 4. From 10 K to 130 K, the hysteresis loops remain square and the
coercivity reduces from 0.355 T to 0.199 T. From 160 K to 250 K, the
magnetization reversal becomes more gradual and the coercivity continues to
decrease. The magnitude of the MCD signal reduces to nearly zero at 250 K,
leading to an estimated TC of $\sim$250 K. These magnetic properties reflect
the high quality of the epitaxial films.

Spin-polarized STM measurements32 are performed using a Cr tip and sweeping
the out-of-plane magnetic field. The magnetic constrast is observed in the
dI/dV signal at a sample bias of 100 mV and a tip current of 0.15 nA. As shown
in Figure 4, the dI/dV signal shows sharp switchings indicated by black arrows
at 0.42 for the positive field ramp (blue) and at -0.48 T for the negative
field ramp (red). Unlike the MCD measurement which averages the magnetization
within the $\sim$150 $\mu$m laser spot, the STM measurement is local at the
atomic scale. This explains why MCD hystersis loops are more uniform while the
STM hysteresis loop has switchings that occur at different field magnitudes.
The latter reflects the stochastic nature of magnetization reversal. Future
studies are planned to capture images of the magnetic domain structure during
a magnetization reversal.

### Discussion

We now consider why growth rate is such an important factor in determining the
material quality during the initial growth. The fact that the synthesis
depends strongly on growth rate implies the importance of kinetics, or the
rate of material processes (e.g. interdiffusion, chemical reaction, atomic
diffusion on surfaces, re-evaporation, etc.). The complementary factor that is
significant for material synthesis is the thermodynamic equilibrium phase
which minimizes the free energy. The balance between energetics and kinetics
can be controlled by the growth rate, where slower growth rates favor the
equilibrium phase and faster growth rates can produce materials further away
from equilibrium based on kinetic considerations.

In applying these concepts to our case, we first point out that the growth
conditions with a 10$\times$ overpressure of Te (i.e. Fe:Ge:Te beam flux ratio
of 3:1:20 for Fe3GeTe2 with a stoichiometry of 3:1:2) is a typical condition
for reactive MBE, where the volatile species (Te) reacts with the other
elements to form the desired material with proper stoichiometry due to re-
evaporation of excess Te. In reactive MBE, when the film has a deficiency of
the volatile species, typical solutions are to increase its beam flux or
reduce the growth rate. This allows more time for the reaction to occur and
can increase the incorporation of the volatile species into the material. But
our FGT growth appears to behave in an opposite way, so we would not consider
this process as traditional reactive MBE. Since slower growth favors the
equilibrium phase while faster growth is further from equilibrium, this
suggests that a Fe3Ge2 surface is more thermodynamically stable than the FGT
surface within a partial pressure of $10^{-8}$ to $10^{-7}$ Torr of Te and at
325 $\degree$C during the intial stages of growth. In other words, energetics
favor formation of Fe3Ge2. Thus, in order to stabilize the FGT phase that we
want, we must consider the rate kinetics associated with the growth. In
comparing the chemical composition of Fe3Ge2 with Fe3GeTe2, the Fe3Ge2 has a
higher relative content of Ge (Ge:Fe = 2:3 for Fe3Ge2 vs. 1:3 for FGT) and an
absence of Te. The absence of Te in Fe3Ge2 is achieved by the out-diffusion of
Te from the film into vacuum. For the increased Ge:Fe ratio (2:3) relative the
beam flux ratio (1:3), either the sticking coefficient of the Fe must become
lower than that of Ge or the excess Ge is supplied by the substrate. Since the
cell temperature for Fe of $\sim$1300 $\degree$C is much higher than the
substrate temperature of 325 $\degree$C, this argues against a reduced
sticking coefficient of Fe and strongly supports the scenario that the excess
Ge is supplied via interdiffusion from the substrate. This is also supported
by the fact that the growth eventually converts to a vdW FGT phase. Thus, to
stabilize the FGT phase we want (and avoid Fe3Ge2), the growth rate must be
sufficiently high to outpace the out-diffusion of Te from the film surface and
the interdiffusion of Ge into the film. One would expect the required minimum
growth rate to be temperature dependent. Lower temperatures have slower
diffusion so it could tolerate slower growth rates, while higher temperatures
would require faster growth rates. Considering that the optimized growth
temperatures are usually a balance between maximizing smoothness and reducing
defects (favoring higher T) and minimizing interdiffusion (favoring lower T),
our current results indicate that growth rate should also be strongly
considered in the materials optimization.

Identifying this kinetically-controlled growth regime for FGT is important for
the future optimization of the material and its incorporation into epitaxial
heterostructures. While the Fe3Ge2 phase might be specific to FGT on a Ge
substrate, the issues of interdiffusion and chemical reaction with a substrate
or underlayer are universally important for film growth. Our results suggest
that the FGT growth rate may be used to control interdiffusion and chemical
reaction with underlayers more generally, which will be useful for
incorporating FGT into epitaxially-grown vdW heterostructures.

## Conclusion

In summary, we discovered that an important factor for optimizing FGT
synthesis is the growth rate, which has an unexpected influence on thin film
quality. The phenomenon is explained by energetics and kinetics of MBE growth
where the stabilization of FGT is assisted by a higher deposition rate that
inhibits the non-vdW Fe3Ge2 phase from forming. Atomic-scale imaging by STM
and STEM confirm the structural quality of the surface and interface. The
excellent magnetic properties with square out-of-plane hysteresis loops is
confirmed by AHE, MCD, and spin-polarized STM measurements. Our results reveal
an important way to think about and optimize MBE growth, leading to
potentially better 2D materials.

We thank Shuyu Cheng, Katherine Robinson, and Ryan Bailey-Crandell for
technical assistance. This work was supported by AFOSR MURI 2D MAGIC Grant No.
FA9550-19-1-0390 (MBE, magneto-optics), U.S. DOE Office of Science, Basic
Energy Sciences Grant No. DE-SC0016379 (STM), and the Center for Emergent
Materials, an NSF MRSEC, under Grant No. DMR-2011876 (TEM, magnetotransport).
Electron microscopy was performed at the Center for Electron Microscopy and
Analysis (CEMAS) at The Ohio State University.

## Author Contributions Statement

W.Z., A.J.B., and R.K.K. conceived the experiments. W.Z. and A.J.B. performed
the MBE growth. M.Z. and J.H. performed the TEM measurements. A.J.B., R.C.W.,
and J.A.G. performed the STM measurements. I.L., W.Z., and R.K.K. performed
the magnetotransport and magnetooptic measurements. All authors participated
in data analysis and preparation of the manuscript.

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