Direct writing of gold nanostructures with an electron beam: On the way to pure nanostructures by combining optimized deposition with oxygen-plasma treatment

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This work presents a highly effective approach for the chemical purification of directly written 2D and 3D gold nanostructures suitable for plasmonics, biomolecule immobilisation, and nanoelectronics. Gold nano- and microstructures can be fabricated by one-step direct-write lithography process using focused electron beam induced deposition (FEBID). Typically, as-deposited gold nanostructures suffer from a low Au content and unacceptably high carbon contamination. We show that the undesirable carbon contamination can be diminished using a two-step process – a combination of optimized deposition followed by appropriate postdeposition cleaning. Starting from the common metal-organic precursor Me 2 -Au-tfac, it is demonstrated that the Au content in pristine FEBID nanostructures can be increased from 30 atom % to as much as 72 atom %, depending on the sustained electron beam dose. As a second step, oxygen-plasma treatment is established to further enhance the Au content in the structures, while preserving their morphology to a high degree. This two-step process represents a simple, feasible and high-throughput method for direct writing of purer gold nanostructures that can enable their future use for demanding applications.

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# Direct writing of gold nanostructures with an electron beam: On the way to pure nanostructures by combining optimized deposition with oxygen-plasma treatment

Domagoj Belić 1,2, Mostafa M. Shawrav 1,3, Emmerich Bertagnolli 1

and Heinz D. Wanzenboeck *1

Full Research Paper [Open Access](http://www.beilstein-journals.org/bjnano/about/openAccess.htm)

Address:
1Institute of Solid State Electronics, TU Wien, Floragasse 7, 1040
Vienna, Austria, 2 University of Liverpool, Department of Chemistry,
Crown Street, Liverpool L69 7ZD, United Kingdom and 3 Institute of

Sensors & Actuator System, TU Wien, Gusshausstrasse 27–29, 1040

Vienna, Austria

Email:
Heinz D. Wanzenboeck * [- Heinz.wanzenboeck@tuwien.ac.at](mailto:Heinz.wanzenboeck@tuwien.ac.at)

- Corresponding author

Keywords:

FEBID; gold nanostructures; oxygen plasma; postdeposition

purification

_Beilstein J. Nanotechnol._ 2017, _8,_ 2530–2543.

[doi:10.3762/bjnano.8.253](https://doi.org/10.3762%2Fbjnano.8.253)

Received: 30 June 2017

Accepted: 08 November 2017

Published: 29 November 2017

This article is part of the Thematic Series "Chemistry for electron-induced

nanofabrication".

Guest Editor: C. W. Hagen

License and terms: see end of document.

## Abstract

This work presents a highly effective approach for the chemical purification of directly written 2D and 3D gold nanostructures suit
able for plasmonics, biomolecule immobilisation, and nanoelectronics. Gold nano- and microstructures can be fabricated by one
step direct-write lithography process using focused electron beam induced deposition (FEBID). Typically, as-deposited gold nano
structures suffer from a low Au content and unacceptably high carbon contamination. We show that the undesirable carbon contam
ination can be diminished using a two-step process – a combination of optimized deposition followed by appropriate postdeposi
tion cleaning. Starting from the common metal-organic precursor Me2-Au-tfac, it is demonstrated that the Au content in pristine
FEBID nanostructures can be increased from 30 atom % to as much as 72 atom %, depending on the sustained electron beam dose.

As a second step, oxygen-plasma treatment is established to further enhance the Au content in the structures, while preserving their

morphology to a high degree. This two-step process represents a simple, feasible and high-throughput method for direct writing of

purer gold nanostructures that can enable their future use for demanding applications.

## Introduction

Focused electron beam induced deposition (FEBID) is an addi
tive direct-write method for making complex 2D and 3D nano
structures [1-5] with an ultimate resolution of less than 1 nm

[6,7]. It is frequently referred to as "3D nanoprinting" because

of its capabilities that are similar to the increasingly popular 3D

printers that operate on the micrometer scale. FEBID has been

only recently widely utilized [8-10] due to the growing number

of modern scanning electron microscopes (SEMs) which can be

2530

combined with focused ion beam (FIB) systems that are

equipped with gas injection systems by default. Rather than

being merely characterization instruments, such SEMs are now

becoming important fabrication tools for the production of

nanoscale structures and devices in magnetics [11-15], super
conductors [16], and nanoelectronics [17-19]. The use of

FEBID processing as a nanosurgery method for novel nanoma
terials, such as graphene, is particularly appealing [20-24].

Hence, any advancement made in the production [25-28] or in

situ purification [29] of FEBID materials attracts great atten
tion.

Among all FEBID materials, Au is of special interest because

its superior properties that make it highly suitable for use in

micro- and nanotechnology, as well as in nanomedicine [30].

Many applications require precise positioning of Au structures

on the nanoscale, e.g., for fabrication of interconnects [31,32]

and field emission tips [33] or in direct writing of plasmonic

and photonic nanostructures and devices [4,34-40]. We recently

utilized FEBID of Au to directly write metal-oxide-semicon
ductor capacitors [41] and to deposit Au nanocatalyst templates

for Si nanowire syntheses [42].

The main issue which is hindering the widespread use of

FEBID is the low Au content in the produced nanostructures:

when using the common, commercially available metal-organic

FEBID Au precursors, such as dimethylgold(III)-acetyl
acetonate (Me2-Au-acac), dimethylgold(III)-trifluoroacetylacetonate (Me2-Au-tfac) or dimethylgold(III)-hexafluoroacetylacetonate (Me2-Au-hfac), deposits typically contain less than
30 atom % of Au [33,37]. This value was only slightly in
creased under deposition in reactive environments [43,44].

Logically, it would seem that the best way to avoid carbon con
tamination is to use a carbon-free precursor. Indeed, FEBID ex
periments using the PF3AuCl precursors yielded almost pure Au
structures [32], but this precursor was eventually abandoned

due to its poor stability. Other gold complexes are under consid
eration [45,46], but have not been adopted in practice due to

significant experimental deficits. Therefore, the low Au content

in FEBID structures obtained from commonly used commercial

precursors remains an outstanding issue. For that reason, this

study focuses on the optimization of deposition mechanism

using a commercially available metal-organic gold precursor

and explores oxygen-plasma treatment as a potentially viable

large-scale postdeposition purification method for 2D and 3D

nanostructures.

Over the last few years several methods to purify FEBID

deposits have been presented. For instance, in the case of tung
sten, it was found that just by applying a smart writing strategy

one could produce single-crystal nanowires [47] whereas the

_Beilstein J. Nanotechnol._ 2017, _8,_ 2530–2543.

use of a supersonic carrier gas can also yield nanostructures of a

high purity [48]. For cobalt, a combination of heat, H2 exposure and electron irradiation improved the metal content in the

deposits [49]. Similarly, in the case of platinum, several purifi
cation approaches have been studied, including laser assisted

purification [50,51], a sequential EBID process and additional

postdeposition electron irradiation in the presence of H2O or O2

[52-56]. The mechanisms behind these processes point to

fundamental differences between molecular oxygen and water

as oxidative species [57].

Similar purification methods have also been investigated for

gold, but different surface mechanisms prevented the direct ap
plication of the methods used for Pt purification to Au purifica
tion. These differences between the platinum precursor

MeCpPtMe3 and the acetylacetonate-based gold precursors
Me2-Au-acac, Me2-Au-tfac, and Me2-Au-hfac become evident
by the fact that atomic layer deposition is feasible for platinum

[58,59], but has proved very challenging for gold [60,61].

While the encouraging results from Pt purification can provide

inspiration, it is still necessary to develop new procedures for

the purification of FEBID gold.

Roughly, all these procedures that strive for purer FEBID gold

can be categorized into two groups – ones that deal with (in

situ) optimization of deposition processes and the others that

deal with various postdeposition (ex situ) treatments to further

improve the purity of the deposits. Prime examples of the latter

group include heat treatment in a reactive atmosphere of air [62]

or oxygen [63], but these methods are compromised by an

(often significant) loss of structural fidelity. Similarly, room

temperature electron beam postirradiation in the presence of O2

[64] or H2O [4] has been used to purify FEBID gold in a second
step, but this method is limited to smaller structures and areas

that can be efficiently processed by a scanning electron beam.

As a possible solution to this issue, oxygen-plasma treatment,

has been explored, yielding promising results [65,66]. Never
theless, the ultimate goal is direct deposition of high-purity gold

nanostructures, removing (or minimizing) the need for any post
deposition purification. Recently, notable advances have been

made in this line of work. For example, in situ purification

during deposition using simultaneous flows of both O2 and precursor yielded up to a 1:10 C/Au ratio [67]. This one-step

process requires a special custom-designed dual nozzle-in
nozzle setup that is currently not readily available to the

majority of FEBID users. Our recent work on simultaneous

injection of gold precursor and water vapour during deposition

outlines an experimentally simpler approach that seems to be

very effective for realizing pristine nanostructures of high purity

[68]. However, injection of water directly into the microscope

chamber may not always be feasible or preferred, so conceiving

2531

an even simpler deposition strategy is desired. The simplest

possible method would be to optimize the deposition parame
ters, making use of conventional equipment and a standard pre
cursor. The ultrahigh vacuum/surface science approach to

studying the effect of electron beam irradiation on nanometer

thin films of a common Me2-Au-acac precursor from Wnuk et
al. [69,70] represents a valuable starting point for optimization

of FEBID Au nanoarchitectures. Similar to their findings, our

previous study on FEBID Au nanopillars (NPs) revealed that

unusually large Au crystals could be found at the bases of the

NPs, which was attributed to in situ focused electron beam in
duced curing (FEBIC) during deposition [65]. That investiga
tion explored a wide range of experimental FEBID parameters

in a systematic manner, in order to yield a recipe for deposition

of pristine 3D nanostructures with desired structural and

compositional properties. With a systematic experimental ap
proach to FEBID with metal carbonyl precursors it was possible

to obtain outstanding pure pristine Fe and Co nanostructures

[71]. Although the decomposition reaction of pyrolytic precur
sors is essentially different to the Au chelate complex used as

precursor in this study, we have used an analogous systematic

approach to improve the metal content of the deposit.

The carbon contamination, which mainly originates from the

imperfect chemical decomposition of the precursor, is a strong

_Beilstein J. Nanotechnol._ 2017, _8,_ 2530–2543.

limitation for possible usage of FEBID nanomaterials in optical,

magnetic, electronic or information storage devices. It is there
fore essential to improve their purity by optimizing the chemi
cal reaction pathway. This experimental study focuses on the

advancement of the material purity of FEBID Au by combining

optimized (in situ) deposition with (ex situ) oxidative plasma

treatment. We revisit the commonly used metal-organic Me2Au-tfac precursor and show that by smartly setting the deposi
tion parameters, one can achieve pristine FEBID nanostruc
tures of exceptionally high Au content without the use of addi
tional reactive species within the SEM. In a second step, we

employ chemical purification by oxidation. We demonstrate

that such postdeposition treatment of Au nanostructures can

further boost the Au content without compromising their struc
tural stability, indicating that the procedure could be suitable for

larger and more complex FEBID architectures.

## Results and Discussion

We deposited 2D and 3D Au nanostructures as a proof-of-prin
ciple because such structures are desirable for devices for bio
nanotechnology, nanoelectronics, plasmonics and photonics.

The nanostructures were prepared using FEBID, as schemati
cally depicted in Figure 1a. In brief, precursor molecules

(dimethylgold(III)-trifluoroacetylacetonate, chemical formula

C7H10AuF3O2) were introduced into the high-vacuum chamber

Figure 1: (a) Schematics of the FEBID experimental set-up. Note that the direction of electron beam movement, i.e., fast scanning axis (FSA) and
slow scanning axis (SSA), can be adjusted relative to the gas flux vector (GFV). (b) SEM image (45° tilt) of a Au nanopillar (NP) fabricated by FEBID.
(c) TEM image of a FEBID Au NP, with the corresponding electron diffraction pattern given in the inset, confirming the polycrystalline nature of the
NP. (d) SEM image (45° tilt) showing a group of various Au nanostructures produced by FEBID. (e) AFM topographic image of a relatively large
planar Au structure made by FEBID. Electron beam scanning strategy was adjusted so that FSA is perpendicular to GFV, whereas SSA was opposite
to GFV. Note the height variation of the deposit in direction of SSA. (f) HAADF-STEM image (top view) of a section of a planar FEBID Au structure,
deposited on a thin silicon nitride membrane.

2532

of a custom-modified SEM via home-built gas injection system

(GIS) [42], in such a way that a nozzle was brought into close

vicinity (≈200 μm) of the intended deposition spot on a pre
pared p-type Si substrate. When a beam of primary electrons

impacts on the surface, secondary electrons are generated by

inelastic interactions, having energies suitable for breaking the

bonds of the precursor molecules adsorbed to the substrate.

Consequently, a metal-containing material was deposited with a

considerable amount of incorporated carbonaceous species

mostly due to the partial break-up of precursor molecules. By

varying the experimental parameters one can alter the deposi
tion process and hence change the properties of deposited mate
rial, such as its shape, structure, and elemental composition. We

note that no etching by the F found in the precursor was ob
served.

Firstly, we investigated the FEBID of vertical Au NPs as model

structures to elucidate how the variation of the experimental pa
rameters influenced the metal content in the NPs. The electron

beam impinging on a single spot on the substrate allowed us to

produce free-standing NPs of a relatively large height-to-width

aspect ratio (up to 50), measuring several micrometers in height

while only having ≈100 nm in diameter (Figure 1b). Transmis
sion electron microscopy (TEM) revealed the polycrystalline

nature of the NPs having 3–8 nm crystallites embedded in

an amorphous matrix (Figure 1c), which seems to be a common

feature of Au NPs produced by FEBID in non-reactive environ
ments [38,43]. The corresponding electron diffraction

pattern given in the inset of Figure 1c is consistent with poly
crystalline Au, revealing interplanar distances of (2.4 ± 0.1) Å,

(2.1 ± 0.1) Å, (1.4 ± 0.1) Å, and (1.2 ± 0.1) Å that correspond

to the spacing between {1,1,1}, {2,0,0}, {2,2,0}, and

{3,1,1} planes in the face centered cubic (fcc) Au lattice, re
spectively.

Apart from the deposition of NPs, we used FEBID to directly

write horizontal Au nanowires (NWs), NW networks, and

planar structures (Figure 1d,e). In the latter case, the deposited

structures had an area measured in hundreds of square microm
eters, making them suitable for applications in electronics [41].

However, FEBID of such large structures is very challenging

and requires very stable experimental conditions over a

long time. For instance, the planar structure presented in

Figure 1e had been deposited for 8.5 h while the working pres
sure during the introduction of precursor gas was maintained at
1.2 × 10 −5 mbar. In a calibration experiment with nitrogen gas a
chamber pressure of 1.2 × 10 −5 mbar corresponded to a flux of

0.11 sccm. During such a long experiment the precursor supply

varied up to 15 rel %, as derived from the chamber pressure

readout. The precursor supply was manually readjusted via the

temperature setting of the precursor reservoir.

_Beilstein J. Nanotechnol._ 2017, _8,_ 2530–2543.

The depositions were performed with a single pass scan with

the e-beam writing line-by-line from top left to bottom right, as

is indicated in Figure 1e. In prolonged experiments a small

beam drift present during deposition resulted in rectangular

structures instead of the original square pattern. The uneven

(wrinkled) surface of planar FEBID structures can be the result

of an inhomogeneous beam drift or of precursor flux variation,

which we deem more likely.

To maximize the deposition rate, i.e., to reduce the total deposi
tion time, the electron beam parameters – acceleration voltage,

i.e., high tension (HT), current ( _I_ ), dwell time (DT), and pixel
to-pixel spacing (PPS) – had to be properly adjusted. It was ex
perimentally found that using HT = 3 kV and _I_ = 1.0 nA (i.e.,

60 μm condenser aperture) resulted in the fastest growth of both

2D and 3D nanostructures, for a given DT and/or PPS. We note

here that the maximum deposition rate usually does not lead to

the purest deposit (in fact, quite the opposite was observed).

Once the deposition rate was optimized and the samples were

produced, high-angle annular dark field scanning transmission

electron microscopy (HAADF-STEM) revealed that the crystal
lites in the planar deposits were considerably larger than those

in NPs (compare Figure 1f to Figure 1c). Although both sam
ples were produced using the same HT and _I_ values, the elec
tron dose that the material sustained during deposition was

approximately two times higher in the case of planar structure
(2.3 × 10 −7 C·μm −2 vs ≈1 × 10 -7 C·μm −2 ). It was the first clear

evidence that a change of electron beam parameters could sig
nificantly alter the inner structure of FEBID Au deposits,

deserving greater attention. In the next sections of this work we

present the main results of our comprehensive study of the in
fluence of electron beam parameters on the properties of Au

nanostructures fabricated by FEBID.

Au nanopillars

Generally, the full experimental parameter space of a typical

FEBID system is multidimensional – it encompasses electron

beam parameters (such as acceleration voltage, i.e., high

tension, electron beam current, electron dose, i.e., deposition

time), precursor parameters (such as precursor type, flux, and

possibly addition of other gases), and substrate parameters

(such as substrate type and temperature). To reduce the level of

complexity, here we restrict the full parameter space to a

subspace spanned by the axes of high tension, electron beam

current, and deposition time.

The deposition of Au NPs was carried out at 4 different HT

values: 1 kV, 3 kV, 5 kV, and 10 kV, using a 60 μm condenser

aperture, which provided a current of 1.0 nA, as was set by the

condenser lens and measured by a Faraday cup placed inside the

SEM chamber. Additionally, for the fixed HT of 3 kV, the elec

2533

tron beam current was varied by using the 30 μm, 60 μm, and

120 μm condenser apertures. For every HT and current, we set

the deposition time to 20 s, 40 s, and 60 s and produced at least

6 NPs using identical experimental conditions. In total we

deposited over 120 samples and conducted a broad SEM and

EDX study to elucidate the effect of electron beam parameters

on the deposition rate, morphology, and elemental composition

of NPs. For the sake of clarity we here concentrate on the

results obtained for 60 s deposition at the current of 1.0 nA,

with varying HT values. These samples are chosen as they

provide the clearest picture of the behaviour generally observed

across all Au NP samples.

All NPs fabricated under the same experimental conditions

appeared indistinguishable in terms of both the morphology (as

evidenced by SEM imaging) and chemical composition (as evi
denced by SEM EDX linescans); see Figure S1, Supporting

Information File 1. Therefore, the results shown here are highly

representative for the whole group of NPs produced under the

same conditions. The top section of the pillar slightly below

(50–150 nm) the tip was analyzed, as this section was more

similar in diameter to the bottom part. The chemical analysis of

the topmost tip of the nanopillar is plagued by the small self
absorption of the carbon K-line in the tip region and the depth
dependence of the absorption in tapered tip geometries. We

would like to point out that a thickness-dependent error for the

given results still exists, so the measured data might slightly

differ from the actual purity levels.

Commonly, for the same deposition time and current (aperture),

the height of the NPs was approximately equal regardless of the

HT used, although the appearance of the NPs changed from

more rugged and wider, produced using the lowest HT value of

1 kV, to smoother and thinner, deposited at higher HT values

(Figure 2), in line with previous observations for W [72]. SEM

EDX studies showed that the NPs are in reality mostly made up

of C, having a smaller fraction of Au present in the deposits

(see Figure S2, Supporting Information File 1). A small O frac
tion was also detected (<10 atom %), but it remains unclear

whether it partially originates from decomposed precursor mol
ecules (C7H10AuF3O2) or partially from the atmospheric O2
and H2O that had been adsorbed to the NPs during a short
transfer of the samples under ambient conditions from the depo
sition system to the analytical SEM. The amount of F in the

NPs was found to be below the detection limit of EDX mea
surements, which is a bit surprising considering the fact that the

precursor molecules contained more F than Au or O atoms.

Detailed linescan EDX investigations of produced samples

revealed a compositional gradient in literally every NP that was

sufficiently tall for multiple probing: as a rule of thumb, the

_Beilstein J. Nanotechnol._ 2017, _8,_ 2530–2543.

gold fraction gradually increased from the top to the bottom of

the NPs, by as much as 57% (relative) in the case of 10 keV

beam (from ≈24 atom % at the top to ≈38 atom % at the bottom;

see Table S1, Supporting Information File 1). At the same time,

the carbon fraction showed the opposite trend, notably decreas
ing towards the bottom of NPs. It is worth noting that the O

content did not exhibit any evident variation between the top

and bottom parts of NPs, which indicates the dominant pres
ence of uniformly adsorbed O atoms. It was commonly ob
served that the Au gradients in NPs deposited for shorter

periods were not as evident as in NPs that were deposited for

longer periods (Figure S2, Supporting Information File 1).

These results were obtained repeatedly and were independent of

the direction of EDX scanning and the relative position of the

sample to the EDX detector. We note that very similar composi
tion values and gradients were found by STEM EDX per
formed on long straight (uniform diameter) sections of NPs

deposited on silicon nitride TEM grids in our previous study

[65]. More accurate compositional quantification could proba
bly be achieved by performing atom probe tomography or

thickness-corrected STEM electron energy loss spectroscopy

(EELS).

Planar Au nanostructures

The main goal here was to elucidate the effects of each experi
mental parameter on the properties of planar nanostructures; in

particular, we wanted to see which parameter had the strongest

influence on the elemental composition of the deposited struc
tures, in order to find the experimental recipe that will maxi
mize the Au content in pristine planar FEBID structures. As

such, this investigation represents a major step towards fabrica
tion of exceptionally pure, patterned planar Au nanostructures

by giving crucial pointers for their practical straightforward re
alization.

As an archetype for all planar FEBID experiments in this study,
we deposited rectangular structures measuring 2 × 0.5 μm 2 in

area, using a single pass, with PPS set to either 3 nm or 0.7 nm,

and DT set to either 0.2048 ms or 2.0480 ms. These specific

values are selected on the basis of results of an initial set of experiments performed on smaller 0.5 × 0.5 μm 2 structures which

suggested that sufficiently small PPS and/or sufficiently long

DT lead to an increase of the Au content in the deposits. How
ever, setting too small PPS and too long DT drastically in
creased the total deposition time to a few hours, even for a rela
tively small planar nanostructure having an area of a few square

micrometers. Thus the PPS and DT values had to be chosen

wisely in order to keep the whole study experimentally feasible.

The overall strategy was to firstly produce a planar structure by

setting PPS to 3 nm and DT to 0.2048 ms, which represents the

smallest electron dose per unit area. In the second step, keeping

2534

_Beilstein J. Nanotechnol._ 2017, _8,_ 2530–2543.

Figure 2: (a)–(d) SEM images (45° tilt) and the corresponding EDX line profiles taken along the height of the Au nanopillars (NPs), as indicated by the
dotted line (to the scale). These vertical NPs were deposited for 60 s using an electron beam current of 1 nA for various acceleration voltage values,
as indicated in the images. By definition, the top section relates to the top quarter of the NP, whilst the bottom section relates to the bottom quarter, as
indicated by the red and blue squares, respectively. The scale bar applies to all SEM images in this Figure. (e) Elemental composition in the top and
bottom sections of samples presented in (a–d), expressed in terms of Au/C atom % ratio. The EDX measurements were performed on 100 × 100 nm 2

areas with an acquisition period of 60 s. For detailed EDX data see Table S1 in Supporting Information File 1.

PPS at 3 nm, DT was increased to 2.048 ms. In the final step,

PPS was decreased to 0.7 nm, with DT set back to 0.2048 ms.

By following such a procedure, the total electron dose was

varied relatively as approximately 1:10:18, for a given electron

beam current. The procedure was performed for HT values of

1 kV, 3 kV, 5 kV, and 10 kV for the fixed current of 1.0 nA

(condenser aperture of 60 µm). In total more than 80 samples

were deposited and subsequently analysed. The main results of

the study are presented in Figure 3; for the sake of clarity and

brevity, we only show results for a limited set of samples that

still provide clear insight into a general trend.

It is clear that more material was deposited when the electron

dose was higher (compare Figure 3c,e to Figure 3d,f). The

growth rate of the samples produced at HT = 1 kV and
HT = 5 kV was around 0.43 μm 3 /μC. The volume of the

deposited material was found to be linearly dependent of the

electron dose which, in combination with relatively long dwell

times and high currents used in these experiments, indicates that

we were most likely working in the mass transport regime [73].

There were notable asymmetric, disjointed deposits (“halo”)

present around the main planar nanostructure, with more mate
rial deposited in the SSA direction, clearly indicating a shad
owing effect. This likely occurred due to scattering of primary

electrons on already deposited material, which resulted with

favoured generation of secondary electrons in front of the

deposit (as seen from the deposition direction). On the other

hand, one might speculate that the body of already deposited

main structure may have obstructed the supply of adsorbed pre
cursor molecules and/or secondary electrons to the back side.

These observations are in agreement with the previous reports

on FEBID of Cu [74] and Pt [75]. In spite of the “halo”, the

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_Beilstein J. Nanotechnol._ 2017, _8,_ 2530–2543.

Figure 3: (a) Schematics of the scanning strategy used in FEBID of planar structures. (b) Definition of “beginning” and “end” of the structures, with
respect to the scanning strategy. Note that both areas are ≈50 nm away from the respective edges of the planar structure. (c–f) SEM images (tilted
view) of planar Au structures measuring 2 × 0.5 μm 2 in area, deposited using various experimental conditions. Plots containing EDX line profiles are
given in the insets. Note the change of scale for Au/C count ratio among the images. The scale bar indicated in (f) applies to all SEM images in this
figure. (g) Elemental composition of the planar structure presented in (c–f), as determined by EDX measurements performed on 100 × 100 nm 2 areas
for 60 s, on the “beginning” and “end” of the deposition, marked with a red and blue square, respectively. Note a pronounced Au content on the
“beginning” side in the samples produced using a higher electron beam dose. For detailed EDX data see Table S2 in Supporting Information File 1.

main structure remained fairly well defined for all HT values

used. Interestingly, in the case of maximum electron dose, the

planar structures at one point start to lift up and deposit

under an angle relative to the substrate plane – a 2D analogue

of 1D lines from [74]. According to our experimental results,

this “disconnection electron beam dose” lies between
2.3 × 10 −7 C·μm −2 and 4.2 × 10 −7 C·μm −2 for electron beam

energy ranging from 1 keV to 10 keV.

2536

A comprehensive EDX investigation of the planar structures

followed and detailed line EDX profiles were acquired, bearing

in mind that EDX profiling inevitably induces some unwanted

C deposition from the residual hydrocarbons present in the ana
lytical SEM chamber. To minimize this effect which may have

consequently lead to misleading results, the chamber contain
ing the samples was always pumped overnight prior to SEM

characterization. The chemical composition was probed on two

sides of the structures, defined as “deposition beginning” side

and “deposition end” side, with respect to the electron beam/

deposition advancement. It was found that the chemical composition was almost uniform across the whole 2 × 0.5 μm 2 struc
ture for the samples produced with the lowest electron beam

dose (see Figure 3c,e), regardless of the acceleration voltage

used during the deposition. Notably, the composition of these

planar structures was quite similar to the composition of the top

sections of vertical Au NPs.

In the case of deposition with higher electron beam exposure

doses (Figure 3d,e), the results clearly reveal that the composi
tion is not consistent across the whole area of the structures: al
though the lateral line profiles (i.e., along FSA) showed no sig
nificant compositional variation in that direction, the profiles

taken in the SSA direction exhibited large differences from one

side of the deposit to the other. Evidently the “deposition begin
ning” side was systematically richer in Au than the “deposition

end” side. For instance, in the case of the sample produced by

setting HT = 1 kV, PPS = 0.7 nm, and DT = 0.2048 ms

(Figure 3c), the Au/C atom % ratio was nearly 7 times higher

on the “deposition beginning” side than on the “deposition end”

side. The same behaviour of Au and C content was usually ob
served for other highest electron dose samples with no excep
tion, albeit at smaller differences between the opposite sides of

planar structures. This behaviour in planar Au nanostructures

seems fundamentally equal to the one observed earlier in the

study of Au NPs, where the bottom sections of NPs were

consistently Au-richer than the top sections. In particular, the

long dwell time “planar” nanostructures that lift-off from the

substrate exhibit compositional profiles very similar to those

seen in the nanopillars. This indicates that the optimization of

pixel spacing may enable fabrication of planar structures of a

high gold content, which is of interest in the context of nanotech

applications. These findings strongly indicate that the elemental

composition of FEBID Au nanostructures significantly depends

on the total electron dose that certain sections of Au nanostruc
tures experience during deposition. Therefore, one can con
clude that the electron beam, primarily used for generating sec
ondary electrons that cause the cracking of precursor molecules

and their initial deposition, plays an extra role in FEBID experi
ments – it cures the pre-deposited material at the same time,

likely by additionally breaking the solid carbonaceous frag

_Beilstein J. Nanotechnol._ 2017, _8,_ 2530–2543.

ments of the non-decomposed precursor molecules still present

in the deposited material. As gaseous C-containing fragments

are pumped out, sections of a relatively higher Au content are

left behind. This curing mechanism during FEBID of 2D and

3D structures from the Me2-Au-tfac precursor is in essence
analogous to the mechanism that governs the deposition from

the Me2-Au-acac precursor, studied by Wnuk et al. [69,70].
These results show that the Au content in FEBID nanostruc
tures can indeed be significantly increased by setting the proper

values of the experimental parameters used for deposition. In

general, our experiments indicate that the deposition with

higher electron beam doses yields structures of a higher metallic

content, in line with the findings of postdeposition Pt curing

using high currents and oxygen injection [56]. Looking at a

wider perspective, our study may point to a more general

strategy for direct realization of exceptionally pure FEBID ma
terials. Moreover, bearing in mind that this deposition tech
nique allows for extremely precise patterning of 2D and 3D

structures on the nano- and microscale, these findings may

facilitate a widespread usage of FEBID in a range of applica
tions.

Oxygen-plasma cleaning of FEBID Au
nanostructures

As was demonstrated in the previous sections, the experimental

results show that, despite significant improvements of the

process, a large part of the deposited structures is still carbona
ceous material, which is generally undesired. Hence, in the

next part of our study of Au nanostructures, we adopted a

postdeposition cleaning procedure that would further decrease

the C content in the structures. Building up on the work

from Miyazoe et al. [76] we here take a similar approach by

subjecting the samples to O-plasma, imposing a crucial require
ment that the morphology of the structures had to be preserved

to a high degree. This requirement is set in order to fully exploit

the initial advantage of FEBID over other deposition tech
niques – direct writing of patterned structures on the nanoscale.

Taking this constraint into consideration, it is immediately clear

that a substantial removal of carbonaceous material from the

deposits represents a major threat to their structural stability and

morphology. Thus, the study was focused on finding just the

right set of parameters for the O-plasma cleaning procedure that

would maximally increase the Au content in the structures,

while maintaining the structures’ shape at the same time. In this

sense, initial optimization of the Au content during deposition

itself seems even more important, because otherwise any post
deposition purification would be significantly limited by the

imposed condition of maintaining the structures’ morphology.

As a test sample, a patterned set of Au NPs was deposited by

FEBID (see Figure 4a). We should note that the composition of

2537

_Beilstein J. Nanotechnol._ 2017, _8,_ 2530–2543.

Figure 4: (a) SEM image (tilted view) of the test configuration of patterned pristine Au nanopillars (NPs). Note the variation in the number of neighbouring NPs, i.e., adjacent NPs within 1 μm distance; (b) Close-up view on a single pristine NP from the pattern having 2 neighbouring NPs, with the
EDX line profile given in the inset; (c) SEM image (tilted view) of the same configuration after the O-plasma treatment; (d) Close-up view on the same
NP as in (b), with the EDX line profile given in the inset. (e) Elemental composition of the top and bottom parts of Au NPs presented in (b,d) acquired
before and after O-plasma treatment, expressed as Au/C atom % ratio. The EDX measurements were performed on 100 × 100 nm 2 areas marked
with the red and blue squares in b,d, for a period of 60 s. For detailed EDX data see Table S3 in Supporting Information File 1.

each NP in this set may slightly differ from the previously dis
cussed stand-alone NPs, since proximity effects are known to

come into play when a dense array of FEBID structures is

deposited [77-79]. Nevertheless, we think that such a compact

configuration might be more similar to the envisaged patterns

that will eventually be used in real-life applications, so these NP

assemblies may be of greater practical importance.

The NPs were deposited using an acceleration voltage of 3 kV,

a beam current of 1 nA, and a deposition time of 60 s. The

overall appearance of the complete configuration and a close-up

view of a single pristine NP are presented in Figure 4a,b, re
spectively. To demonstrate the design capabilities of FEBID the

NPs form the letters "fke" which is the abbreviated name of the

Institut für Festkörperelektronik. It can be seen that the shape of

the NPs was not fully uniform across the set: those NPs having

more neighbouring NPs appear to have a slightly wider diame
ter than the more isolated NPs. This is strong evidence of

FEBID proximity effects, where deposition of neighbouring

NPs leaves some residue on the previously deposited NPs. The

EDX study performed on a single pristine NP (conveniently

chosen to have two neighbouring NPs, which seems interesting

from the application point of view) showed a gradient of Au

along the height of the NP (see inset plot in Figure 4b), which

was consistent with the results obtained for stand-alone NPs

deposited under the same experimental conditions in the first

part of the study. Here one has to be careful when deriving

conclusions because of a sizeable “halo” deposit present in the

2538

surroundings of each NP that may have given a certain back
ground contribution to the EDX line profiles. Such an effect is

not pronounced when NPs are deposited far from each other, as
described earlier. The EDX probing done on 100 × 100 nm 2

areas in the top and bottom sections of the NP showed that the

NP mostly consisted of a carbonaceous material, having only a

minor Au fraction, which was in accordance with the EDX

results obtained for stand-alone NPs. These findings seem to be

representative for all those NPs in our configuration having two

neighbouring NPs. On the other hand, the NPs with three or

more neighbouring NPs generally showed slightly higher levels

of C content, which can be attributed to more pronounced prox
imity effects.

Following the initial SEM investigation of pristine Au NPs, the

samples were then exposed to O-plasma and examined again

using the same SEM parameters. It was immediately evident

that some material was indeed removed from the sample,

leading to a significant volume reduction (approximately

40–45%) of the O-plasma treated NPs compared to the pristine

NPs (see Figure S3, Supporting Information File 1 for detailed

analysis). In spite of the loss of material, the stability of indi
vidual NPs was largely preserved; consequently, the overall

pattern of the NP configuration was conserved as well. The

EDX line profiles indicate a prominent increase of the Au

content in the NPs: the area-averaged EDX results show

increase of the Au fraction, reaching the values of over 70 atom

% at the bottom of NPs; at the same time, the C content dropped

to below 25 atom %, resulting in a six-fold enhancement of the

Au/C atom % ratio due to O-plasma treatment (Figure 4e).

Interestingly, this enrichment factor applies almost uniformly to

all parts of the NPs, so that the compositional gradient still

existed even after the O-plasma treatment. Since the treatment

is assumed to have removed carbonaceous species from the sur
face of the structures, this result implies that the Au gradient in

the pristine structures was not a surface-bound feature – it

existed throughout the body of the NPs.

To confirm that O-plasma cleaning is generally applicable, the

same experimental procedure was used on a set of planar sam
ples. We note that, although the underlying cleaning mecha
nism is assumed the be the same in both planar (2D) and

vertical NPs (3D), some differences in the cleaning rate might

be expected due to variation in the geometry of exposed sur
faces. Consequently, slight modifications in the experimental

cleaning protocols between 2D and 3D structures might arise.

Again though, a common observation was that, upon O-plasma

treatment, the planar nanostructures to a small extent reduced in

volume, presumably due to partial removal of carbonaceous

material, leaving the original morphology largely preserved. As

a rule, the Au fraction increased in all the samples, as revealed

_Beilstein J. Nanotechnol._ 2017, _8,_ 2530–2543.

by EDX analyses. The enhancement factor tended to be greater

for the structures (or their parts) that initially had a lower Au

content. A typical example presented in Figure 5 demonstrates

this effect quite clearly: a compositionally inhomogeneous pris
tine planar structure (i.e., structure exhibiting a large Au

gradient across the opposite sides) was subjected to O-plasma

treatment which resulted in an evident loss of material,

rendering its initially smooth surface quite rough. Notably, the

shape of the deposit was largely preserved, but more important
ly, the EDX analysis showed an increase in Au content throug
hout the nanostructure. The enhancement factor (EF) was

considerably larger on the initially Au-poor, “end” side

(EF ≈ 4), while the initially Au-rich “beginning” side saw only

a very modest increase (EF ≈ 1.1). We note that, in this way,

postdeposition O-plasma treatment may work as a leveller of

compositional gradient in inhomogeneous planar FEBID struc
tures.

Figure 5: (a) SEM image (tilted view) of a pristine planar FEBID Au
nanostructure of a large compositional gradient (see Figure 3d) before
O-plasma treatment. (b) SEM image (tilted view) of the same structure
after O-plasma treatment. (c) Elemental composition of the “beginning”
and “end” of the planar nanostructure, acquired before and after
O-plasma treatment. The EDX measurements were performed on
100 × 100 nm 2 areas for 60 s. For detailed EDX data see Table S4 in
Supporting Information File 1.

Furthermore, the increased roughness and SEM-brightness of

the nanostructures’ surfaces, seen in Figure 4d and Figure 5b,

indicate that the O-plasma cleaning exposed some of the em

2539

bedded Au nanocrystallites to the exterior. Such purified

patterned nanostructures are deemed very suitable for further

functionalization, e.g., for immobilization of biomolecules or

similar applications where having clean Au surfaces is essential

for successful binding of ligands, such as thiols [80]. From that

point of view, our cleaning procedure presents a valuable ap
proach for experimental realization of structurally sound,

precisely patterned nanostructures of a high Au content, with

exposed Au surfaces available for functionalization.

## Conclusion

In summary, we investigated a FEBID experimental parameter

subspace, aiming to fabricate 2D and 3D nanostructures with a

high Au content, starting from a common metal-organic precur
sor, Me2-Au-tfac. In the first instance, we realized vertical freestanding FEBID Au nanopillars having a compositional

gradient along the height of the structure, where the bottom part

was systematically Au-richer than the top part, with up to

45 atom % of Au. These results suggest that the electron beam

dose sustained during deposition plays a key role in defining the

purity of the FEBID Au material. In a similar fashion, a range

of planar nanostructures were produced, aiming to identify the

optimal set of experimental parameters that will yield deposits

of the highest Au content. It was found that, in line with the

results obtained for vertical NPs, higher electron beam doses

generally leave a purer FEBID Au material, which in the

optimal case exceeded 70 atom %. We note that this straightfor
ward approach is feasible with any standard SEM and does not

require specialized instrumentation.

The second approach towards higher gold purity is postdeposi
tion purification by oxidation. As molecular oxygen is not reac
tive enough, oxygen radicals were used. The oxidation was per
formed in a plasma reactor that is common equipment in most

clean rooms. This second purification approach can quickly

improve the gold purity of a number of FEBID structures, in
cluding larger ones. In this work we developed a process that

employs an oxygen-plasma treatment to further increase the

purity of the deposited material. It was found that such an ap
proach can significantly increase the Au content in the FEBID

nanostructures, while preserving their structural stability. The

enhancement factor of Au content was found to be dependent

on the initial composition in such a way that the structures (or

their parts) that initially had a lower Au fraction experienced a

larger boost, levelling the original compositional gradient. We

conclude that a combination of optimized deposition followed

by a suitable oxygen-plasma cleaning protocol may lead to

exceptionally pure, structurally stable, patterned 2D and 3D Au

nanostructures. In addition, the oxygen-plasma cleaning might

potentially expose the Au surfaces to the exterior, making such

freshly processed patterned FEBID Au nanostructures very suit

_Beilstein J. Nanotechnol._ 2017, _8,_ 2530–2543.

able for prospective surface functionalization, which is highly

desired for a range of applications in nanotechnology.

## Experimental

The samples were prepared by FEBID of Au inside a custom
modified Zeiss Leo 1530VP SEM. The microscope was

equipped with a home-made gas injection system containing

dimethylgold(III)-trifluoroacetylacetonate precursor (Me2-Autfac, chemical formula C7H10AuF3O2). The base pressure in
the microscope chamber was around 1.6 × 10 −6 mbar, while the

working pressure during the introduction of precursor gas
was maintained at 6.0 × 10 −6 mbar (except in the case of a

large planar structure presented in Figure 1e, where it was
1.2 × 10 −5 mbar). The deposition took place in a non-reactive

environment at room temperature. Customized macros were

executed by the SmartSEM user interface to exactly control the

electron beam exposure, ensuring complete reproducibility of

experimental parameters. When depositing planar structures we

made use of a Raith Elphy Plus pattern generator. The samples

for SEM analyses were deposited onto ultrasonically cleaned

p-type silicon substrates having a thin native oxide layer of

around 2 nm, while the samples for TEM investigation were

deposited onto a perforated silicon nitride TEM membrane (SPI

Supplies). The oxygen-plasma treatment was carried out in a

Tepla 100 plasma system using 300 W of power for 180 s, at

the gas pressure of 0.9 mbar. At these conditions the main reac
tive species in the plasma are typically oxygen atoms [81] with
an approximate concentration of 10 12 to 10 13 atoms/cm 2 [82].

SEM and EDX characterization was performed in a Zeiss Neon

40EsB cross-beam microscope equipped with an Oxford Instru
ments EDS 7427 detector, operating at an acceleration voltage
of 5 kV and a base pressure of 1.2 × 10 −6 mbar. The samples

were briefly exposed (<2 min) to cleanroom air during the

transfer from the Zeiss Leo 1530VP SEM or Tepla 100 plasma

system to the Zeiss Neon SEM. TEM characterization was per
formed in a FEI Tecnai F20 microscope equipped with an

EDAX detector and a Fischione 3000 HAADF detector, at an

acceleration voltage of 200 kV. The TEM samples were

exposed to atmosphere for a few hours during the transfer from

the fabrication facility to the TEM lab. AFM investigations

were carried out on a Veeco/Bruker Dimension 3000 atomic

force microscope.

## Supporting Information

Supporting Information File 1

Additional experimental data.

[[http://www.beilstein-journals.org/bjnano/content/](http://www.beilstein-journals.org/bjnano/content/supplementary/2190-4286-8-253-S1.pdf)

[supplementary/2190-4286-8-253-S1.pdf]](http://www.beilstein-journals.org/bjnano/content/supplementary/2190-4286-8-253-S1.pdf)

2540

## Acknowledgements

The research leading to these results has received funding from

the European Community's Seventh Framework Programme

(FP7/2007-2013) under grant number ENHANCE-238409 and

the Austrian Science Fund (FWF) under grant number P24093.

A part of this work was conducted within the framework of the

COST Action CM1301 (CELINA). D.B. acknowledges funding

from the European Research Council (ERC) via the Advanced

Grant “PANDORA”. Authors greatly acknowledge the support

of the Center for Micro- and Nanostructures (ZMNS) at Vienna

University of Technology. TEM characterization was carried

out using facilities at the University Service Centre for Trans
mission Electron Microscopy (USTEM) at Vienna University of

Technology.

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