Formation of pure Cu nanocrystals upon post-growth annealing of Cu–C material obtained from focused electron beam induced deposition: comparison of different methods

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In this paper we study in detail the post-growth annealing of a copper-containing material deposited with focused electron beam induced deposition (FEBID). The organometallic precursor Cu(II)(hfac) 2 was used for deposition and the results were compared to that of compared to earlier experiments with (hfac)Cu(I)(VTMS) and (hfac)Cu(I)(DMB). Transmission electron microscopy revealed the deposition of amorphous material from Cu(II)(hfac) 2 . In contrast, as-deposited material from (hfac)Cu(I)(VTMS) and (hfac)Cu(I)(DMB) was nano-composite with Cu nanocrystals dispersed in a carbonaceous matrix. After annealing at around 150–200 °C all deposits showed the formation of pure Cu nanocrystals at the outer surface of the initial deposit due to the migration of Cu atoms from the carbonaceous matrix containing the elements carbon, oxygen, and fluorine. Post-irradiation of deposits with 200 keV electrons in a transmission electron microscope favored the formation of Cu nanocrystals within the carbonaceous matrix of freestanding rods and suppressed the formation on their surface. Electrical four-point measurements on FEBID lines from Cu(hfac) 2 showed five orders of magnitude improvement in conductivity when being annealed conventionally and by laser-induced heating in the scanning electron microscope chamber.

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# Formation of pure Cu nanocrystals upon post-growth annealing of Cu–C material obtained from focused electron beam induced deposition: comparison of different methods

Aleksandra Szkudlarek 1,2, Alfredo Rodrigues Vaz 1,3, Yucheng Zhang 4,
Andrzej Rudkowski 5, Czesław Kapusta 5, Rolf Erni 4, Stanislav Moshkalev 3 and Ivo Utke *1

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

Address:
1Empa, Swiss Federal Laboratories for Materials Science and

Technology, Laboratory for Mechanics of Materials and

Nanostructures, Feuerwerkerstrasse 39, 3602 Thun, Switzerland,
2AGH University of Science and Technology, Academic Centre for

Materials and Nanotechnology, al. A. Mickiewicza 30, 30-059 Krakow,
Poland, 3 Center for Semiconductor Components, State University of
Campinas, 13083-870, Campinas, SP, Brazil, 4 Empa, Swiss Federal

Laboratories for Materials Science and Technology, Electron

Microscopy Center, Überlandstrasse 129, 8600 Dübendorf,
Switzerland and 5 AGH University of Science and Technology, Faculty

of Physics and Applied Computer Science, Department of Solid State

Physics, al. A. Mickiewicza 30, 30-059 Krakow, Poland

Email:
Ivo Utke * [- ivo.utke@empa.ch](mailto:ivo.utke@empa.ch)

- Corresponding author

Keywords:
Cu(hfac)2; Cu nanocrystals; focused electron beam induced
deposition (FEBID); post-growth annealing of Cu–C material

_Beilstein J. Nanotechnol._ 2015, _6,_ 1508–1517.

[doi:10.3762/bjnano.6.156](http://dx.doi.org/10.3762%2Fbjnano.6.156)

Received: 16 March 2015

Accepted: 24 June 2015

Published: 13 July 2015

This article is part of the Thematic Series "Focused electron beam

induced processing".

Guest Editor: M. Huth

License and terms: see end of document.

## Abstract

In this paper we study in detail the post-growth annealing of a copper-containing material deposited with focused electron beam

induced deposition (FEBID). The organometallic precursor Cu(II)(hfac)2 was used for deposition and the results were compared to
that of compared to earlier experiments with (hfac)Cu(I)(VTMS) and (hfac)Cu(I)(DMB). Transmission electron microscopy

revealed the deposition of amorphous material from Cu(II)(hfac)2. In contrast, as-deposited material from (hfac)Cu(I)(VTMS) and
(hfac)Cu(I)(DMB) was nano-composite with Cu nanocrystals dispersed in a carbonaceous matrix. After annealing at around

150–200 °C all deposits showed the formation of pure Cu nanocrystals at the outer surface of the initial deposit due to the migra
tion of Cu atoms from the carbonaceous matrix containing the elements carbon, oxygen, and fluorine. Post-irradiation of deposits

with 200 keV electrons in a transmission electron microscope favored the formation of Cu nanocrystals within the carbonaceous

matrix of freestanding rods and suppressed the formation on their surface. Electrical four-point measurements on FEBID lines from

Cu(hfac)2 showed five orders of magnitude improvement in conductivity when being annealed conventionally and by laser-induced
heating in the scanning electron microscope chamber.

1508

## Introduction

Focused electron beam induced deposition (FEBID) is a direct

maskless nanolithography technique, based on the local dissoci
ation of adsorbates upon the irradiation with electrons [1]. The

molecules are delivered into the microscope chamber by a gas

injection system (GIS) where they reversibly physisorb onto the

substrate surface. Part of the energy from the primary electron

beam or from the secondary electrons generated in the vicinity

of the impinging primary beam is transferred to the adsorbates

and breaks their chemical bonds. The non-volatile fragments

stick to the substrate surface whereas the volatile fragments are

removed from the chamber by the pumping system. By control
ling the beam scanning three dimensional structures of a com
plex shape can be created in a single direct-write deposition step

onto planar or non-planar surfaces [2].

Nanodevices with various functionalities have been deposited

comprising gas sensors [3,4], magnetic sensors [5,6] strain

sensors [7], thermal sensors [8], photodetectors [9], and mode

stabilizers for vertical surface emitting lasers [10]. Other

deposits were used as ferromagnetic wires [11,12], supercon
ducting wires [13], plasmonic structures [14], or as electrode

nanocontacts [15,16]. The feasibility of obtaining 3D nanostruc
tures with a high aspect ratio makes FEBID suitable for fabrica
tion of high resolution probes to scanning magnetic force

microscopy (MFM) [17-19].

Purification methods of FEBID structures

For FEBID direct-write nanostructures lateral resolution can be

well-controlled by adjusting the beam and gas flow settings [20]

as well as by optimizing scanning strategies [21]. However, the

purity of FEBID materials obtained with organic precursors still

remains an issue. Recently, post-growth purification methods

using electron beam irradiation in combination with thermal

annealing and co-injection of reactive gases/ions were devel
oped. In the case of Pt–C deposits, the catalytic properties of Pt

nanoparticles facilitate the process of molecular oxygen dissoci
ation, thereby increasing the efficiency of removing the carbo
naceous matrix [22]. Pure Pt material was obtained with a post

deposition treatment using O2 gas and a) laser pulsing [23] or
b) low-temperature substrate annealing (up to 50 °C) [24]. The

presence of H2O during electron irradiation performed at rt
allowed for a total elimination of carbon from Pt–C deposits

without affecting the shape [25]. The combination of thermal

heating to 300 °C, injection of H2 gas and simultaneous electron irradiation led to pure Co deposits [26]. Microstructural

changes were observed upon simple 5 keV electron beam

curing of FEBID structures. The Pt–C deposits exhibited an

increased conductivity by three to four orders of magnitude

[27,28]. For W–C deposits an improvement of one order of

magnitude was found [29].

_Beilstein J. Nanotechnol._ 2015, _6,_ 1508–1517.

Conventional post-growth annealing of FEBID material in

vacuum was summarized in a review by Botman et al. [30]: The

thermal energy which is delivered to the sample can cause a de
sorption of carbonaceous fragments increasing the metal

concentration from 15 atom % of Au (rt) to 24 atom % of Au (at

100 °C). Increasing the substrate temperature during FEBID

also favors the desorption of non-metallic dissociation by-prod
ucts as it was observed by Mulders et al. [31] for various

precursors: TEOS (tetraethylorthosilicate), Co(CO)3NO,
Co2(CO)8, and Me2Au(acac) with the best purity enhancement
for W(CO)6 (from 37 atom % at 25 °C to 59 atom % of W at
280 °C). However, the temperature rise during the deposition

may not be favorable as it also decreases the residence time of

adsorbates, significantly lowering the growth rate. Furthermore,

the high temperature may also cause the diffusion of deposit

atoms into or from the substrate.

A compromise approach is based on pulsed heating with an IR

laser as a heat source which can generate abrupt temperature

peaks in the microseconds range only (sufficient for

desorption), allowing the substrate surface to equilibrate quickly

and to replenish with new adsorbates before the next electron
beam scan frame. FEBID together with synchronized pulsed IR

laser heating helped to increase the metal concentration of Au,

W, Pt in FEBID deposits, however, did not fully remove the

carbon. For deposits obtained from Me2Au(acac) the initial
atomic ratio of C to Au decreased from 4 to 0.5 with the laser

assistance [32]. For W(CO)6 FEBID the atomic ratio of W to C
was improved from 1:4 to 2:1 [33]. For MePtCpMe3 FEBID the
Pt concentration increased from about 15 atom % to 35 atom %

[34]. Such an improvement was not observed when using a

conventional heating stage during the deposition process with

MePtCpMe3. The deposit obtained at 350 °C did not exhibit a
different Pt/C ratio than the deposit obtained at room tempera
ture, where 15 atom % Pt were measured [31].

Copper purification

The very high conductivity of copper makes the localized depo
sition of this metal very attractive for applications in nano-elec
tronics. The organometallic hexafluoroacetylacetonate (hfac)
based Cu(I) and Cu(II) precursors are widely used in chemical

vapor deposition (CVD) methods due to their stability and high

vapor pressure. They allow to obtain pure metal CVD films

with the same resistivity as in a bulk material at deposition

temperatures below 300 °C [35-37].

Recently, it was shown for condensed monolayers of Cu(hfac)2
(also Pt(hfac)2 and Pd(hfac)2) that electron-beam irradiation
results in about 80 atom % of carbon content [38]. The metal

content could be then increased by two sequential purification

1509

steps: 1) deposit bombardment with atomic oxygen 2) deposit

bombardment with atomic hydrogen. In the first step the carbo
naceous material was fully removed from the material and in

the next step the metal oxide was reduced to the metal.

Although this method was successfully applied to obtain a

deposit with high metal purity, the total exposure time was

rather long: 40 h for oxygen and 2 h for hydrogen. The effi
ciency of atomic hydrogen only for purification of Cu–C ma
terial obtained by an ion-induced deposition process at room

temperature was shown by Chiang et al. [36]. It led to

99 atom % pure Cu films. H2/Ar microplasma-assisted FEBID
increased the Cu content from 12 atom % to 41 atom % but also
caused extended halo deposits [39]. Ga + ion beam deposition

showed that heating the substrate surface has a crucial influ
ence on the properties of the deposit structure, from small

isolated nanocrystals of Cu (ca. 20 nm) at 25 °C towards contin
uous thin films of pure copper at 100 °C, using (hfac)CuVTMS

[40].

In this paper we will show results obtained by SEM, TEM, and

electrical resistance monitoring during post-growth annealing of

Cu–C FEBID material from Cu(II) and Cu(I) precursors with

respect to nanostructural changes and conductivity showing the

potential of fabricating pure copper nanodots, from the

as-grown amorphous Cu–C deposit. The thermal energy input

favors the migration of Cu atoms to coalesce to pure Cu

nanocrystals being dispersed inside and on top of the carbona
ceous matrix.

## Experimental

FEBID

The experiments were performed using a Hitachi 3600 scan
ning electron microscope (SEM) with a tungsten filament. The

precursors were filled into their reservoirs inside a glove box in

argon or dry nitrogen atmosphere. The deposition process has

been carried out at room temperature on two types of substrates:

Si with a 200 nm top layer of SiO2 and copper TEM grids with
holey carbon films. The beam energy was fixed to 25 keV.

_Beilstein J. Nanotechnol._ 2015, _6,_ 1508–1517.

In this study bis(hexafluoroacetylacetonato)copper(II)

[Cu(hfac)2, Cu(HC5O2F6)2] was used as a precursor and
compared to earlier experiments with the precursors

vinyltrimethylsilane copper(I) hexafluoroacetylacetonate,

[(hfac)CuVTMS, (C6H12Si)Cu(HC5O2F6)], and dimethylbutene copper(I) hexafluoroacetylacetonate [(hfac)CuDMB,

(C5H12)Cu(HC5O2F6)] [41].

The precursor flux was estimated to be about 10 monolayers per

second for Cu(hfac)2. The exposure parameters for Cu(hfac)2
for 1 µm × 1 µm square deposits were: dwell time of 1 µs, pixel

distance of 0.4 nm, and frame repetitions varying from 100 to

1000 with refreshment times of 0.625 s. The beam current was
0.4 nA. This corresponds to doses of 0.25 nC/µm 2 (100 repetitions) and 2.5 nC/µm 2 (1000 repetitions). For the 15 μm long

lines we used 100 µs dwell time per pixel, 0.5 nm pixel dis
tance, and 300 line repetitions with a refreshment time of 3 s.

The beam current was 1 nA, which corresponds to the dose of
9 nC/µm 2 and exposure time of 900 s. Tip deposits were

obtained by the stationary dot exposure mode exposing a pixel

for two minutes. Line and freestanding-rod deposits with

(hfac)CuDMB and (hfac)CuVTMS from earlier FEBID experi
ments [39,40] were performed with a single scan at around

30 nm/s with 600 pA.

Annealing

Post-growth annealing in vacuum was achieved by various

setups shown in Figure 1. The conventional heating using a hot

plate was performed without breaking the vacuum after the

deposition process. The experiments were performed inside the

SEM for a temperature range from room temperature up to

220 °C. The sample was supported in a custom-made massive

copper block on a resistive heater (Boraelectric). The control of

temperature was done with a thermocouple directly coupled to

the base of the substrate. During the heating we observed an
increase of the pressure inside the SEM chamber from 5 × 10 −6

to 5 × 10 −5 mbar. We cannot exclude possible reactions of the

deposits with the residual gas (most likely water and residual

Figure 1: Sketch of post-growth annealing experiments: a) conventional heating using a hot plate in an SEM, b) SEM integrated laser heating; the
substrate is heated by the IR laser pulses, c) post-growth thermal annealing during in situ TEM analysis. Conventional and IR laser annealing experiments were combined with in situ four-point probe resistance measurements.

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hydrocarbons) but we judge the influence as negligible as TEM
annealing experiments at pressures better than 1 × 10 −6 mbar

also resulted in nanocrystal formation (see below in Figure 6).

An infrared laser system was integrated into the SEM chamber

for in situ post-growth annealing directly after the deposition.

The system was composed of a SvetWheel’s MU40-960-01

laser diode module and a fast diode current modulator VFM

20-25s (Messtec Power Converter GmbH) triggered by a signal

function generator. The wavelength of emitted photons was

equal to 960 nm (corresponding to an energy of 1.29 eV) with a

total maximum laser power of 29 W. The tip of the laser optical

fiber was placed in close vicinity of the FEBID deposit. The

laser power, pulse duration and its frequency were adjustable by

the signal generator. The size of the laser spot on the substrate

was estimated to 65 µm × 130 µm, based on the molten area in

SiO2 using the maximum laser power and a high pulse
frequency.

The substrate surface temperatures generated by the laser during

post-deposition annealing were determined experimentally and

theoretically. Numerical finite element simulations with the

annealing conditions (peak laser power of 13.6 W, pulse dura
tion of 10 μs and frequency of 10 kHz) gave a stationary

substrate surface temperature at the center of the laser beam

of 158 °C which was reached after approximately 30 s.

This was in good agreement with the value of 150 °C deduced

from in situ four-point-probe resistance measurements of a

gold wire deposited by e-beam lithography and PVD metal

lift-off together with the pre-structure of the four-point

electrodes.

TEM in situ annealing of FEBID deposits from Cu(hfac)2 was
performed with a Gatan double tilt heating holder (Model 652)

in a JEOL-2200FS microscope. The samples were heated up to

220 °C with a ramp rate of 20 K/min.

Characterization

For SEM observation we used a Hitachi S4800 electron micro
scope and for the EDX measurements a Brooker/Oxford device

mounted on a TESCAN LYRA microscope. EDX measure
ments were performed at 5 kV and 2 nA over 30 s with a

300 nm × 300 nm scan area on the squares to account for in

homogeneities in the lateral copper nanocrystal precipitation.

Standard EDX software was used to calculate the composition

from the spectra.

High-resolution TEM (HR-TEM) images were taken at

200 keV. Selected area diffraction (SAD) was taken using the

second smallest selected area aperture corresponding to an area

of 400 nm in diameter on the sample. Chemical mapping was

_Beilstein J. Nanotechnol._ 2015, _6,_ 1508–1517.

obtained using electron energy loss spectroscopy (EELS) oper
ated in the scanning TEM (STEM) mode. The Cu K edge

(928 eV) and a signal energy window of 40 eV (920–960 eV)

after background subtraction were used to map the composition

distribution of Cu, see Figure S1 in Supporting Information

File 1.

The changes of resistance during the annealing were monitored

by four-point probe measurements using an SEM-integrated

15-stage nanomanipulator from SmarAct. Conductive micro
probes were connected via an SEM feedthrough to a Keithley

2400 Sourcemeter, with a source current of 100 µA and a

voltage compliance level of 500 mV.

## Results and Discussion

As-deposited material

Energy dispersive X-ray (EDX) analysis showed that at room

temperature deposited Cu–C lines and squares obtained

from Cu(II)(hfac)2 had an atomic ratio of approximately
Cu/C/O/F = 10:64:25:1 with standard deviations of ±1 atom %

for Cu, ±2 atom % for C, ±1 atom % for O, and ±0.3 atom %

for F on eight deposits. This amounts to 10 ± 2 atom % of Cu.

With respect to the stoichiometric copper content in the

Cu(hfac)2 precursor of 3.7 atom % (disregarding the hydrogen)
this was 2.7 times more copper in the deposit. The Cu content in

deposits obtained from the (hfac)Cu(I)VTMS and

(hfac)Cu(I)DMB precursors was about twice as high [41].

Transmission electron microscopy (TEM) showed that dot,

square, and line deposits from Cu(II)(hfac)2 on an amorphous
carbon membrane were amorphous (Figure 2).

In contrast, as-deposited freestanding rods from earlier FEBID

experiments with (hfac)Cu(I)VTMS showed small Cu

nanocrystals homogeneously dispersed in a polymeric carbona
ceous matrix (see below in Figure 7 taken with a Philips

EM-430 TEM at 300 keV). This matrix contains all the ligand

elements: carbon, oxygen, fluorine, and silicon as well as prob
ably some hydrogen (not detectable by EDX) [42]. The differ
ence between amorphous and nano-composite materials

obtained for the Cu(II)(hfac)2 and (hfac)Cu(I)VTMS precursor,
respectively, can be attributed to the lower thermal stability of

(hfac)CuVTMS which is 63 °C compared to 250 °C for

Cu(hfac)2.

Electrical measurements showed that the as-deposited FEBID

lines from Cu(hfac)2 were highly resistive with a value of a few
gigaohms being around the measurement limit of the Keithley

device. This is in line with our former experiments which

showed for all three copper precursors nonconductive behavior

for the room temperature as-deposited material [41].

1511

_Beilstein J. Nanotechnol._ 2015, _6,_ 1508–1517.

Figure 2: TEM of as-deposited lines and squares from Cu(hfac)2 on an amorphous carbon membrane on a TEM grid. a) An overview of deposits,
b) zoom into a line deposit, c) a SAD pattern showing amorphous nature of the deposit.

Annealed material
SEM observations

Upon annealing conventionally or with a laser the flat

morphology of square, line, or tip deposits on the pre-patterned

SiO2/Si substrate changes (Figure 3 and Figure 4). While the
laser allows for local heating, the conventional hotplate ap
proach allows for more accurate temperature measurements.

The visible onset of Cu nanocrystal precipitation on the deposit

surface starts at around 150 °C for the Cu(hfac)2 deposits on the
pre-patterned SiO2/Si substrate. Further heating to about 200 °C
for 30 min did not visibly change the appearance of the Cu

nanocrystal precipitation.

EDX analysis after conventional heating to 200 °C

for 30 min showed an atomic ratio of approximately

Cu/C/O/F = 12:75:13:0 with standard deviations of ±2 atom %

for Cu, ±2 atom % for C, and ±1 atom % for O on the same

deposits. The 12 ± 2 atom % Cu content in the deposit is thus

rather constant compared to the as-grown sample value of

10 ± 2 atom % Cu within the error limits and may have even

been slightly overestimated systematically by the standard EDX

software due to the surface-precipitated Cu nanocrystals

corrected for absorption. Compared to the as-deposited ma
terial the fluorine completely disappeared and the oxygen

content was halved. An estimation of the amount of precipi
tated copper nanocrystals that are visible in SEM gave only

about 9 to 17 wt % (see Supporting Information File 1).

Compared to the average Cu content in the as-grown and

deposited samples of roughly 11 atom % (corresponding to

an average 37 wt % Cu) this means that 20 to 28 wt % or

5 to 7 atom % of the copper dispersed in the matrix did not

precipitate to be visible in SEM or segregated during the

annealing. The matrix deposit volume shrunk during annealing

to about 70% of its initial volume pointing to reticulation of the

carbon network after release of fluorine and oxygen.

Figure 3: Post-growth annealing of FEBID line from Cu(hfac)2
between four gold electrodes. SEM tilt images (60°) of a) as-deposited
line, b) after annealing at 200 °C for 30 min, and c) top-view zoom into
central part of the annealed line showing the Cu nanocrystals on the
line deposit and inside as well as on the halo deposit.

Interestingly, the size of the nanocrystals is much smaller in the

thin halo region of the deposited line, yet the nanocrystal

density is much larger. This may be due to the small amount of

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_Beilstein J. Nanotechnol._ 2015, _6,_ 1508–1517.

Figure 4: Post-growth laser annealing of FEBID deposits from Cu(hfac)2. SEM top view images of a) as-deposited square, b) Cu nanocrystals
forming due to laser annealing at a power of 13.6 W over 1 min, c) laser-annealed FEBID line between two gold electrodes.

deposit material available to form the Cu nanocrystals but could

be also due to the low irradiation dose of these regions and

hence a slightly less reticulated carbon network which facili
tates segregation at many places.

For laser-induced heating we found the same phenomenon of

surface precipitation of Cu nanocrystals although the annealing

time was only one minute. Also the same trend in the variation

of the atomic ratio with annealing was found. Energy

dispersive X-ray measurements of the deposit shown in

Figure 4a and Figure 4bgave an atomic ratio of Cu/C/O/F =

9:52:32:6 for the as-grown FEBID material which changed to

Cu/C/O/F = 10:66:23:0.6 after laser annealing. Considering the

error limits of 2 atom % there is no fundamental difference to

the trends and absolute values for Cu stated for conventional

annealing. The Cu/C = 1:6.6 ratio after laser annealing of 1 min

at around 158 °C (see section Experimental) is smaller than for

thermal annealing, however, time and annealing temperature

were lower for the laser than for thermal annealing. Figure 4 c

shows that precipitation of Cu nanocrystals is not fully uniform

across the line length in contrast to the conventionally annealed

lines in Figure 3b and Figure 3c (see also Figure S5 in

Supporting Information File 1). At the vicinity to the gold elec
trodes there are fewer nanocrystals visible which might be due

to a varying distribution of the laser-induced temperature.

The precipitation of Cu nanocrystals on the initial deposit

surface was observed also in former experiments with the Cu(I)

precursors. Figure 5 shows SEM tilt views of a periodic three

dimensional line deposit obtained from (hfac)Cu(DMB). Such

periodicity can arise when the vertical deposition rate is compa
rable to the scan speed of the focused electron beam. For the

annealing experiments this is not of importance (for more

details we refer to Bret et al. [43]). Upon annealing the same

precipitation at deposit surfaces and at halo regions due to

forward and backscattered electrons can be seen (Figure 5a–c).

For a tip deposit the same features develop upon heating

pointing to the fact that the dwell time per pixel during FEBID

is not a very sensitive parameter for Cu nanocrystal precipita
tion (Figure 5d). The prolonged irradiation of a few minutes in

the spot mode during FEBID with 20 keV electrons does not

seem to change the reticulation of the matrix in such a way that

formation of Cu nanocrystals remains contained inside the

matrix. On the other hand, post-irradiation experiments with

electrons energy of 5 keV were shown to be already effective to

change the electron transport mechanisms in Pt–C and W–C

FEBID deposits [27-29].

TEM observations

In Figure 6, the results of in situ TEM annealing experiments

performed on the line and square deposit material from

Figure 5: Periodic 3D FEBID line deposits from (hfac)Cu(DMB) between gold electrodes on SiO2/Si. SEM tilt images: a) before annealing, b) after in
situ vacuum annealing, c) zoom on Cu nanocrystals, d) tip deposit showing also Cu nanocrystal formation.

1513

Cu(hfac)2 FEBID (shown in Figure 2) are presented. During
annealing the deposits were not observed except for one to

capture the temperature of crystal formation in a video.

The formation of nanocrystals took place at a temperature of

200 °C. The time for Cu nanocrystal formation was below

500 ms – the time range of the CCD camera to take a video

frame image. After this rapid transition the nanocrystal arrange
ment did not change anymore. We compared it to regions that

were not observed in TEM and found no differences so that the

given temperature for crystal formation was not subject to ir
radiation during the video. We have tried different annealing

rates and taken TEM images at different beam currents (which

can be easily done by spreading the electron beam) and the

results remained the same. In Figure 6 it can be seen that the

amorphous material from Cu(hfac)2 turns to nano-composite
with Cu nanocrystals embedded in an amorphous carbonaceous

matrix. The size ranges from 2 nm to 20 nm for the face

centered cubic (fcc) copper nanocrystals. This underlines that

post-growth annealing has the potential to achieve one-digit
nanometer sized copper nanocrystals when using a high-resolu
tion electron microscope for smallest dot deposition [44].

Theoretically, annealing a hemispherical FEBID deposit having

a radius _r_ met and containing a given weight percentage _w_ met to
pure metal would result in a metal hemisphere of radius

(1)

with ρmet and ρdep being the densities of the metal and
deposit, respectively. For our deposits from Cu(hfac)2
Equation 1 yields _r_ met = 0.46· _r_ dep with _w_ met = 0.35 (corresponding to the 10 atom % Cu in the as-grown deposit, see

_Beilstein J. Nanotechnol._ 2015, _6,_ 1508–1517.

section “As-deposited material”), ρdep = 2.5 g/cm 3 (from [45])
and ρCu = 8.9 g/cm 3 . This means that 1 nm Cu dots could be
annealed out of a 2 nm sized hemispherical FEBID material

from Cu(hfac)2.

Selected results of earlier in situ TEM annealing experiments

within a Philips EM-430 TEM on freestanding FEBID rods

obtained from (hfac)Cu(VTMS) are shown in Fsigure 7. Inter
estingly, the comparison of Figure 7b and Figure 7c shows that

the place of Cu nanocrystal precipitation can be controlled by

post-irradiation with the high-energy electrons of a TEM in

contrast to the experiments with Cu(hfac)2. The Cu crystals
segregate around 140 °C inside the carbonaceous matrix when

the deposit was exposed to the 300 keV electrons of the TEM

during the annealing experiment, probably due to an electron
triggered reticulation of the polymeric carbonaceous matrix.

The stronger reticulation of the carbon matrix seems to suppress

the long-range mobility of Cu atoms to diffuse to the outside

surface. In contrast, irradiation with only 20 keV is not effi
cient in reticulating the matrix in such a way that out-diffusion

is hampered as was observed for (hfac)CuDMB deposits in

Figure 5 and for Cu(hfac)2 deposits in Figure 3 and Figure 4.

In contrast, purification by electron beam induced heating of

freestanding rods obtained from FEBID with (hfac)CuVTMS

[42] gave much larger pure copper crystals (up to 150 nm in

size). In comparison to the above results the differences may be

explained by invoking an electron stimulated desorption effect

which was suggested for FEBID by van Dorp et al. [46] which

would facilitate the desorption of carbonaceous fragments and

thus the formation of larger copper nanocrystals.

Electrical measurements

Figure 8 shows the typical behavior of a Cu–C FEBID line for a

heating/cooling cycle. The as-deposited lines were non-conduc

Figure 6: In situ TEM annealing for 10 min at 220 °C on a line deposit from Cu(hfac)2 shown in Figure 2. a) STEM high angle annular dark field
(STEM-HAADF) image of Cu nanocrystals forming from the deposit material. b) HAADF image showing distribution of Cu nanocrystals.
c) High-resolution TEM (HR-TEM) image of 15 nm sized polycrystalline Cu precipitate. The inset shows the SAD pattern of Cu fcc nanocrystals.

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_Beilstein J. Nanotechnol._ 2015, _6,_ 1508–1517.

Figure 7: TEM in situ annealing of FEBID rods grown from (hfac)Cu(VTMS). a) Dark field image of an as-deposited freestanding rod. Inset: rod apex
with small Cu nanocrystals in carbonaceous matrix. b) Bright-field image of same rod after 270 °C annealing and continuous TEM observation
(200 keV). Large Cu nanocrystals form inside the rod. c) Bright field image of another rod not observed during the same annealing process. Cu
nanocrystals form at the outside surface of the rod.

Figure 8: Calculated resistivity from the resistance measurement of a Cu–C line during in situ post-growth heating with a hot plate (red dots) and
cooling down (blue dots) inside the SEM chamber. The resistance did not change when opening the chamber. The top SEM images show the
morphology changes of an adjacent FEBID line which was observed simultaneously during the in situ resistance measurement.

tive, showing an electrical resistance of few gigaohms at room

temperature. After thermal annealing the resistance dropped by

four to five orders of magnitude to hundreds of kiloohms. At

the current densities, used in the experiments (less than
0.1 MA/cm 2 ) the migration process of Cu atoms is rather slow.

Gazzadi and Frabboni [47] reported grain formation and

1515

electromigration in Pt–C material at current densities
approaching 10 MA/cm 2 which is about a factor 100 below our

maximum current densities applied. A large resistance drop was

observed after nanocrystals precipitated on the surface around

150 °C.

An increase in resistivity of one order of magnitude can be

observed upon the cooling cycle from 180 °C to 25 °C. The
temperature coefficient for bulk Cu is 0.00386 K −1 and would

amount to an increase of the resistance by a factor of 1.6 only

for the temperature difference of 155 K. The thermally acti
vated transport observed in this case, can be due to the variable

range hopping mechanism [48], corresponding to the insulating

transport regime for granular materials, as it was observed

previously for various FEBID deposits, composed of metallic

grains embedded in carbonaceous matrix [49]. More detailed

studies are planned to characterize the electronic transport of

this material including the question whether the resistance-vs
temperature behavior of the samples would be stable. Here we

focused on a proof of concept study.

## Conclusion

We have shown that nanostructural changes were induced in

Cu–C FEBID material from Cu(I) and Cu(II) precursors upon

post-growth annealing causing the segregation and precipita
tion of Cu nanocrystals. No fundamental differences between

laser induced heating and conventional heating in the SEM and

TEM with respect to trends and composition values were

observed. As deposited the Cu–C FEBID deposits obtained

from Cu(hfac)2 were amorphous while nanocomposite deposits
were obtained from (hfac)CuVTMS and (hfac)CuDMB. The

as-deposited materials were non-conductive. The transition into

conductive material as well as the segregation and precipitation

of the copper atoms occur upon conventional or laser heating.

The surface precipitation of copper nanocrystals upon annealing

opens a route for depositing pure Cu nanodot patterns using

highly focused electron beams.

## Supporting Information

Supporting Information features additional information

about the chemical mapping with electron energy loss

spectroscopy, the estimation of Cu precipitation on deposit,

and the distribution of Cu nanocrystals along the Cu–C

lines after conventional and IR laser thermal annealing.

Supporting Information File 1

Additional experimental data.

[[http://www.beilstein-journals.org/bjnano/content/](http://www.beilstein-journals.org/bjnano/content/supplementary/2190-4286-6-156-S1.pdf)

[supplementary/2190-4286-6-156-S1.pdf]](http://www.beilstein-journals.org/bjnano/content/supplementary/2190-4286-6-156-S1.pdf)

_Beilstein J. Nanotechnol._ 2015, _6,_ 1508–1517.

## Acknowledgements

Metal deposition for gold electrode lift-off by Vipin Chawla is

acknowledged. I.U. would like to thank Philippe Buffat for the

TEM heating experiments at CIME at EPFL shown in Figure 7.

A.S and I.U acknowledge support by the COST Action

CM1301 (CELINA). A.R.V. acknowledges the support of the

program ‘Ciência sem Fronteiras’ from CNPq(Brazil) and State

University of Campinas (UNICAMP).

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