# Interface effects on titanium growth on graphene

Georg Zagler1,2,∗, Alberto Trentino1,2, Kimmo Mustonen1, Clemens Mangler1,
and Jani Kotakoski1,∗
1University of Vienna, Faculty of Physics, Boltzmanngasse 5,
1090 Vienna, Austria
2University of Vienna, Vienna Doctoral School in Physics, Boltzmanngasse 5,
1090 Vienna, Austria
∗Email<EMAIL_ADDRESS>(GZ<EMAIL_ADDRESS>(JK)

###### Abstract

Poor quality interfaces between metal and graphene cause non-linearity and
impairs the carrier mobility in graphene devices. Here, we use aberration
corrected scanning transmission electron microscopy to observe hexagonally
close-packed Ti nano-islands grown on atomically clean graphene, and establish
a 30∘ epitaxial relationship between the lattices. Due to the strong binding
of Ti on graphene, at the limit of a monolayer, the Ti lattice constant is
mediated by the graphene epitaxy, and compared to bulk Ti, is strained by ca.
3.7% to a value of $0.306(3)$ nm. The resulting interfacial strain is slightly
greater than what has been predicted by density functional theory
calculations. Our early growth stage investigations also reveal that, in
contrast to widespread assumptions, Ti does not fully wet graphene but grows
initially in clusters with a thickness of 1-2 layers. Raman spectroscopy
implies charge transfer between the Ti islands and graphene substrate.

## Introduction

The outstanding properties of two-dimensional (2D) solids such as the linear
dispersion relation and high carrier mobility render them attractive for
device applications. However, such applications necessitate contacting the
device via bulk electrodes, and resulting in metal-graphene interfaces that
are known to have an influence on graphene’s intrinsic properties [1, 2, 3],
which may reduce its usability by a significant factor [4, 5]. In addition to
electronics, the interfaces of metal and graphene play a role e.g. in
applications such as ultrathin catalysts [6, 7] and hydrogen storage [8, 9].
Further on, the growth of metal nanoclusters on graphene provide a relatively
straightforward way to engineer the intrinsic electronic properties of 2D
materials [10], and graphene-metal composites have been put forward in order
to create materials with enhanced mechanical strength [11, 12].

In addition to physical and chemical vapour deposition (PVD/CVD), atomic layer
deposition and wet chemistry approaches are used in the synthesis of
heterostructures based on 2D materials and nanoclusters. Independent of the
used growth method, however, in order obtain surface contamination free
structures, conditions that are void of any spurious chemical species are
required, which is best achieved by PVD based on ultra-high vacuum (UHV)
systems. The ratio of homogenous and defect mediated nucleation in the PVD
process are substrate mediated and hence, the choice of substrate has a
substantial effect on the growth dynamics [13]. The inert surface of 2D
materials and chiefly that of graphene, provides a system for the study of
nucleation and growth with weakly interacting substrates [14]. Indeed, as
confirmed by experiments carrier out with Au and CuAu [15, 16, 17], the inert
surface suppresses the nucleation of nanoclusters on graphene and necessitates
the use of anomalously high supersaturation to nucleate the crystal growth.
These results are in a good agreement with earlier simulations that, with the
exceptions of Ti [3, 18], imply a weak interaction between graphene and
metals.

Some recent reports imply that Ti may form covalent bonds with graphene and
produce TiC [19], while others have claimed that Ti cannot do so on defect-
free graphene [20, 21]. It has also been reported that Ti wets graphene in
Frank–van der Merwe (layered) growth mode, and that it grows in a (roughly)
epitaxial relationship with the graphene substrate [21].

Transport [22] and photocurrent [1] measurements imply that significant
$n$-type doping of graphene may be observable at the Ti-graphene interfaces,
while Raman studies have been much less conclusive [18, 21]. In one of these
studies a low-yield Ti deposition onto graphene was also investigated at the
atomic-scale [23], but due to contaminated graphene, the authors reported the
Ti nucleation exclusively on the most reactive, contaminated parts of their
samples. Nevertheless, although the structure and properties of graphene-
titanium interface are still not fully resolved [24], due to the strong
binding of Ti on graphene and its suitable work function, Ti is being used in
device experiments as an intermediate material to fabricate electrical
contacts with graphene and external circuits [5],

Here, we report the growth of Ti on clean graphene via physical vapour
deposition. We demonstrate that Ti grows in Volmer–Weber (island) mode,
resulting in clusters with hexagonal close packed (hcp) structure and a clear
epitaxial relationship to the graphene substrate. At early growth stage, Ti
forms atomically thin islands that are epitaxially aligned to the graphene
substrate and that have a lattice constant of $0.306(3)$ nm, corresponding to
a strain of ca. 3.7% compared to the bulk structure. This is somewhat higher
than what has been previously predicted via density functional theory
calculations [3]. Continuation of the growth leads to laterally larger and
thicker islands, which at $5-7$ layers reach a lattice constant of $0.296(3)$
nm, close to the bulk value. Additional Raman spectroscopy measurements show
signs of significant charge transfer. Overall, our results bring light to the
previously unexplored early stages of Ti growth on graphene, and highlight the
importance of clean surfaces and atomic-scale characterization for the study
of interfaces.

## Experimental

Monolayer free-standing graphene substrates were prepared from commercially
available (Graphenea, ”Easy Transfer”) chemical vapour deposition-grown
graphene. Graphene, covered with a sacrificial layer (100 $\mu$m polymer
film), was placed into a beaker filled with deionized water, and fished out
using a perforated SiN TEM-grid (Ted Pella, 2.5 $\mu$m diameter holes). The
sacrificial layer was removed with thermal etching in an oven over 4 h at
400∘C in underpressure with a Ar-H gas flow. Additional experiments were
conducted using commercially available graphene-covered holey carbon membranes
on gold grids (Graphenea). The specimen were then baked overnight in vacuum at
nominal 150∘C before being inserted into the interconnected ultra-high vacuum
system (base pressure 10-8 mbar), where all subsequent experimental steps
(including sample preparation and microscopy) were undertaken [16]. Before
physical vapor deposition of Ti, the samples were further cleaned of
ubiquitous contamination using a high-powered laser (ca. 450 $\mu$m spot size
with powers between $10-25$ mW and exposure times between $500-5000$ $\mu$s)
[25].

Ti was deposited onto the specimen using two different methods. For early
experiments a titanium sublimation pump was first cycled multiple times to
ensure removal of possible contaminants from the evaporation material, after
which Ti was deposited by placing the specimen under the pump. For a more
controlled deposition, Ti was in later experiments deposited using electron-
beam physical vapour deposition from a Ti rod. With a well-defined and better
reproducible deposition rate the growth of Ti islands was observed deploying
multiple deposition steps with intermediate imaging. The graphene was not
heated during deposition and the deposition was performed at a base-pressure
at or below $10^{-10}$ mbar. After deposition, the specimen were transported
in the interconnected vacuum system to the Nion UltraSTEM 100 dedicated
scanning transmission electron microscopy (STEM) instrument operated here at
60 kV. The convergence semiangle was 30 mrad, and images were acquired with
medium- and high-angle annular dark-field detectors (MAADF and HAADF, annular
ranges of 60–200 mrad and 80–280 mrad, respectively). Electron energy-loss
spectroscopy (EELS) was carried out in the same microscope using a Gatan PEELS
666 spectrometer fitted with an Andor iXon 897 electron-multiplying charge-
coupled device camera [26]. STEM image simulations were performed using AbTEM
[27].

The specimen pre-characterized with Raman spectroscopy were cleaned with the
laser as described above, with each free-standing graphene patch (covering a
hole of the TEM-grid) having between 10 and 60% of the surface free of
contamination, as established via STEM imaging. After pre-characterization, Ti
was deposited using the sublimation pump. All Raman spectroscopy was conducted
outside the vacuum system, transferring the specimen in air, but minimizing
the air exposure to a maximum of 3 h. For collecting the spectra, a 532 nm
laser with 5 mW of power, 0.5 s exposure time, 30 accumulations and a 600
gr/mm grating were used to measure the Raman response. D, G and 2D peaks were
fitted with Lorentzian peaks.

## Results and Discussion

STEM MAADF images in Fig. 1 show the growth of Ti on free-standing monolayer
graphene. Before deposition (Fig. 1a), the specimen is cleaned with the laser,
which creates large areas (blue, some 100 nm in diameter) of atomically clean
graphene. After Ti deposition (Fig. 1b), changes in both the contamination-
free and contaminated graphene can be seen. On contamination (not shown), Ti
grows in small connected clusters that display no uniform orientation or
lattice parameter, most likely due to the strong interaction between Ti and C
atoms in the contamination that in contrast to C atoms in graphene do not have
ideal bonding configuration. In contrast, previously clean graphene areas are
now covered with Ti structures with a uniform size distribution, whereas
contaminated areas show a more uniform increase in contrast. The smaller field
of view STEM HAADF image in Fig. 1c reveals the actual morphology of the
deposited Ti structures on clean graphene. They show a hexagonal crystal
structure, as is apparent also in the Fourier transform shown in the inset,
and intensities that take only discrete values with a constant step height. In
the STEM HAADF images, the contrast can be directly interpreted as a
thickness, which means that the Ti structures have an island structure with
clearly defined atomic layers (with a typical thickness of ca. 5–6 layers in
the case of this image, as can be seen from the adjacent histogram that shows
the intensity values for each pixel from the image after a Gaussian blur with
a $\sigma$ of seven pixels). The Ti island are epitaxial with the graphene
substrate with a 30∘ mis-orientation, and have a lattice constant of 0.296(3)
nm, close to the bulk value of 0.295 nm [28]. We note that bulk Ti has a
hexagonal close packed structure, where one layer consists of two atomic
layers (AB stacking), which is also consistent with our results for the small
Ti islands on graphene.

Fig. 1: Ti clusters on Graphene and Contamination. a) STEM HAADF pre-
characterization image of suspended graphene after laser cleaning. In brighter
areas surface contaminants cover the otherwise clean graphene. b) STEM HAADF
image of titanium deposited onto graphene, where it grows large, mostly
uniform islands on the clean sample areas. c) STEM HAADF image of Ti islands
showing layered growth. Fourier transform (inset) reveals the hexagonal
crystal structure of the Ti islands and the epitaxial relationship with the
graphene substrate at 30∘ misalignment. The corresponding intensity histogram
(after a Gaussian blur of the image in d)) shows peaks corresponding to a
thickness of $4-7$ layers.

To better understand the growth of Ti on graphene, low quantities of Ti were
next deposited in three steps. Depositing low amounts of Ti, small Ti clusters
grow on clean defect-free graphene (Fig. 2). The clusters have an average size
of 9.2 nm2 (logarithmic distribution) and display monolayer (ML) areas with a
hexagonal structure (see Fourier transform in the inset, sharp peaks
correspond to graphene, broader peaks to Ti), as well as some Ti atoms on the
ML region. These atoms can form a second or a third layer (Fig. 2b), but are
not always perfectly aligned (Fig. 2a). After the first deposition, the
coverage is 0.7% with 700 islands/$\mu$m2 of clean graphene. This confirms the
island-type Volmer–Weber growth mode. It should be noted that all clusters
also align with the graphene lattice, and no exclusively monolayer Ti island
was found.

Fig. 2: Ti nanocluster growth with step-wise deposition. a-b) STEM HAADF
example images of Ti islands after the first deposition step. They have a
hexagonal crystal structure with some monolayer areas (single-atom contrast)
and some thicker areas. Insets are the Fourier transforms of each image,
confirming the hexagonal structure (hexagonal pattern with broadened peaks)
and the graphene substrate (sharp hexagonal pattern). c) Island size
distribution after the first deposition step. d-e) STEM HAADF example images
after the second deposition step. Some clusters still have some monolayer
areas, while many are covered with another layer, which typically does not
display the perfect hexagonal structure. f) Size distribution after the second
deposition step. g) STEM HAADF example image after the third deposition step.
h) Size distribution after the third deposition step.

The same specimen is then deposited with more Ti (Fig. 2 d-f). Contamination
visible in the images (red in color, covering large parts of the graphene) is
due to mobile contamination being pinned down by the electron beam during
imaging [15]. The island growth is not influenced by this, as the
contamination only appears after deposition and only during imaging, and after
each deposition step, different sample area was imaged. The average cluster
size is 9.4 nm2, and many clusters appear to have more layers. At the same
time the coverage increases to 1.8% and the cluster density increases to 1900
islands/$\mu$m2. The distribution of the clusters is similar to the first
deposition (Fig. 2c), with a small increase in the number of both larger and
smaller clusters. The Fourier transform still displays the same hexagonal
pattern with epitaxial alignment between the islands and the graphene
substrate.

Another Ti deposition on the same specimen (Fig. 2 g-h) yields similarly sized
Ti clusters (8.4 nm2 on average) with a distribution (Fig. 2h) similar to
after the first deposition (Fig. 2c). The coverage (1.5%) as well as the
density (1800 islands/$\mu$m2) are comparable to the second deposition. It
should be noted that some pinned-down hydrocarbon contamination was already
present and appears to have not altered the growth significantly (if it did,
one would expect to find many smaller disordered clusters similar to what was
described above for contaminated sample areas). This is in line with previous
experiments have shown that thin contamination coverage does not necessarily
act as preferred nucleation site [15].

A feature common to many multilayer Ti nanoislands is a lattice mismatch in
the hexagonal structure, which is seen as an apparent lattice mismatch within
each Ti island, which is typically most visible at its edges (see for example
the island marked with numbers in Fig. 1c). We assume that this is due to the
strong interaction between graphene and the first Ti layer which imposes
strain on the structure, but decreases layer by layer as the distance from
graphene increases. This is in line with simulations that predict strain at
the Ti-graphene interface [3, 21]. Specifically, Hsu and co-workers [21] used
experimental input to create a Ti-graphene superstructure and calculated a 2%
interface strain that was assumed to be in graphene. In the same study, the
authors also reported a substantial shift in the Raman G-peak of graphene,
which can be attributed to strain in graphene caused by a Ti layer.

To investigate this further, we evaluated the Ti lattice constant from the
Fourier transform, by comparing it to the lattice constant of graphene
estimated from the corresponding sharp peaks. Interestingly, contrasting the
lattice constant of thin regions (Fig. 2a) to mostly $5-6$-layered islands
(Fig. 1c), reveals that there is a significant strain in the thin Ti islands.
Indeed, the lattice constant in the thinnest structures is about 0.306(3) nm,
whereas in the thicker islands it is 0.296(3) nm, that is close to bulk Ti in
$\alpha$-phase at 0.295 nm [28]. This corresponds to a strain of ca. 3.7%. The
data is shown in Fig. 3a. To confirm our hypothesis that the appearance of the
thicker structures could be explained by gradually changing strain within the
Ti islands, we created an artificial Ti structure with five AB layers (ten
atomic layers), where the lattice constant is varied by 3% linearly between
the first and the last layer (Fig. 3b). In Fig. 3c an image simulation of the
part of this structure highlighted with a rectangle in Fig. 3b is overlaid on
top of an experimental STEM HAADF image of an island with a comparable
thickness. As can be immediately recognized, the simulated image is in
excellent agreement with the experimental one confirming our hypothesis.

Fig. 3: Strain in Ti nanoclusters. a) Lattice constant of Ti islands decreases
by 3.4% from 0.306(3) nm ($1-2$ layers) to 0.296(3) nm ($5-6$ layers) with
increasing island thickness. The red points are means calculated from all
experimental images (gray points), and the error bars correspond to the
standard deviation of the data and maximum uncertainty in determining the
number of layers from the STEM HAADF image contrast. b) Ti slab consisting of
five Ti layers (each consisting of two atomic layers in AB-stacking) with 3%
linear change in the lattice constant from the lowest to the highest layer. c)
STEM HAADF image of Ti island overlaid with a simulated image from the
structure shown in panel b from the area marked with a rectangle.

Raman spectroscopy was conducted before loading the specimen into the vacuum
system, after laser cleaning (10-60% of the graphene surface cleaned), and
after Ti deposition. Most suspended areas covered with Ti islands were 20-40%
free of surface contamination (with direct contact between Ti and graphene).
The pristine free-standing graphene has a Raman response with a G peak at
1584.1 cm-1 and a 2D peak at 2670.1 cm-1 (Fig. 4a). The peaks are sharp and no
D peak is present. After laser cleaning and Ti deposition both G and 2D peaks
are shifted up (to 1592.0 cm-1 and 2675.9 cm-1, respectively). The peaks are
also broadened, and the D peak appears.

Fig. 4: Raman spectra of graphene before and after Ti deposition. a) Example
Raman spectra before cleaning and Ti deposition, and after them. b) Raman G
and 2D peak positions for two samples (804 and 842). Before Ti deposition both
samples show a linear dependency between the peaks with an approximate slope
of 2.2 (black dashed line). After Ti deposition the peaks show another linear
distribution with slopes of 1.0 and 0.87, respectively.

The G and 2D peaks show a linear dependency (Fig. 4b) for the as-prepared
samples with an approximate slope of 2.2. While the linearity remains after
laser cleaning and Ti deposition, the slope decreases significantly down to
values close to 1 and below (depending on the sample). Slopes close to 2.2 are
well-established to arise from uniaxial strain in graphene [29, 30, 21, 31],
whereas values close to 0.75 have been attributed to $p$-type doping [31]. For
$n$-type doping, a non-linear global behavior is expected, although it would
appear nearly linear for the range of values investigated here [32, 31]. We
point out that for our samples, the situation is complicated by the laser
cleaning step, which can lead to either negative or positive strain in the
sample due to thinning and possibly partial cracking of the sample support.
Other influence arising from heating, such as hole doping or substrate
interactions [33, 31] should however be absent for suspended monolayer
graphene samples with a holey SiN support, laser heated in high vacuum.
Overall, the observed change in the G/2D peak behavior is a clear indication
of charge transfer in our Ti deposited samples. Significant widening of the G
peak or sharp decreases in the 2D peak intensities that have been previously
reported [21, 18] were not observed in our samples.

## Conclusion

Ti deposited onto demonstrably clean, monolayer graphene nucleates on clean
sample areas already at low densities and grows in the Volmer–Weber (island)
fashion. The resulting Ti islands display the expected hcp crystal structure
with different layers and step edges clearly visible. For the thinnest
structures directly in contact with graphene, the formation of an epitaxial
interface with a 30∘ mis-orientation leads to a significant strain of up to
over 3%, and a corresponding lattice constant of 0.306(3) nm. By island
thickness of ca. 5–6 AB layers (10–12 atomic layers) or more, the lattice
constant decreases to 0.296(3) nm ($<0.5$% strain compared to bulk value).
Based on Raman spectroscopy, Ti deposited onto graphene leads to clear charge
transfer independent of surface contamination. In contrast to island growth on
clean graphene, Ti deposited onto contaminated graphene grows as non-
crystalline clusters. These results highlight how important characterization
of graphene and removal of the ubiquitous surface contamination is for the
complete understanding of the interface effects between graphene and other
materials.

## Acknowledgments

We acknowledge the Austrian Science Fund (FWF) for funding through project
P31605.

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Competing interests: The authors declare no competing interests.

Data availability: All data needed to evaluate the conclusions in the paper
are present in the paper or the supplementary materials.