Electrochimica Acta 56 (2011) 5142–5150
Contents lists available at ScienceDirect
Electrochimica Acta
journal homepage: www.elsevier.com/locate/electacta
Electrodeposition of cobalt–yttrium hydroxide/oxide nanocomposite films from
particle-free aqueous baths containing chloride salts
K.M. Sivaraman a , O. Ergeneman a , S. Pané a,∗ , E. Pellicer b,∗∗ , J. Sort c , K. Shou a , S. Suriñach b ,
M.D. Baró b , B.J. Nelson a
a
b
c
Institute of Robotics and Intelligent Systems (IRIS), ETH Zürich, CH-8092 Zürich, Switzerland
Departament de Física, Facultat de Ciències, Universitat Autònoma de Barcelona, E-08193 Bellaterra, Spain
Institució Catalana de Recerca i Estudis Avançats (ICREA) and Departament de Física, Universitat Autònoma de Barcelona, E-08193 Bellaterra, Spain
a r t i c l e
i n f o
Article history:
Received 11 December 2010
Received in revised form 13 March 2011
Accepted 14 March 2011
Available online 22 March 2011
Keywords:
Nanocomposite
Cobalt
Yttrium hydroxide/oxide
Pulse deposition
X-ray photoelectron spectroscopy
a b s t r a c t
The feasibility of growing nanostructured films composed of cobalt and yttrium hydroxide/oxide phases
by electrodeposition is demonstrated. Particle-free aqueous solutions containing YCl3 and CoCl2 salts and
glycine were used. The incorporation of yttrium compounds into the cobalt deposit was achieved using
pulse deposition (ton = 0.1 ms, toff = 0.9 ms) and for cathodic pulses higher than −500 mA cm−2 . Deposits
obtained were crack-free, typically with 1–5 wt% yttrium, and exhibited morphologies markedly different from the ones shown by pure cobalt deposits. Moreover, yttrium-rich films (up to 30 wt% Y) could be
deposited under certain conditions, though incipient cracking developed in this case. X-ray photoelectron spectroscopy analyses revealed that Y(OH)3 /Y2 O3 compounds were present in the films. From the
structural viewpoint, the composites exhibited a partially amorphous/nanocrystalline character, with
the crystalline fractions originating from the hexagonal-close packed structure of ␣-Co. A refinement of
the ␣-Co crystallite size was observed in deposits containing higher weight percentage of yttrium compounds. Nanoindentation tests revealed that hardness increased with the yttrium content. This result can
be explained by taking into account both the presence of intrinsically hard oxide phases and the effects
promoted by incorporation of yttrium hydroxides/oxides on the ␣-Co matrix (namely, grain-refining and
higher concentration of stacking faults).
© 2011 Elsevier Ltd. All rights reserved.
1. Introduction
Nanocomposite coatings consisting of ultra-fine ceramic particles (e.g., Al2 O3 , SiO2 , SiC, TiN, AlN) embedded in a metal matrix
have become the focus of widespread research in recent years
due to their superior properties compared to purely metallic films.
Benefits include high specific heat, optical non-linearity, novel
magnetic properties, enhanced mechanical behavior (large hardness and wear resistance) and good corrosion resistance, amongst
others [1–5].
Nanocomposite coatings can be produced by various methods, including electrodeposition. Electrodeposited nanocomposite
coatings are generally obtained by suspending charged ceramic
nanoparticles in the electrolyte and co-depositing them with the
metal [1–5]. However, this method suffers from some drawbacks.
Nanoparticles can easily agglomerate due to the compressive effect
caused by the high ionic strength on the diffuse double layer
∗ Corresponding author. Tel.: +41 44 632 33 12.
∗∗ Corresponding author. Tel.: +34 93 581 14 01; fax: +34 93 581 21 55.
E-mail addresses: [email protected] (S. Pané), [email protected] (E. Pellicer).
0013-4686/$ – see front matter © 2011 Elsevier Ltd. All rights reserved.
doi:10.1016/j.electacta.2011.03.058
surrounding the particles [6]. As a consequence, the anticipated
advantageous chemical and/or physical properties of the composite coatings are often not obtained. Furthermore, hydration
forces can also hinder particle co-deposition [7]. Suitable surfactants can improve the stability of the suspension by increasing the
wettability of suspended particles, while enhancing the electrostatic adsorption of the dispersed particles on the cathode surface
by increasing their positive charge [8]. Nevertheless, there are
also drawbacks associated with the use of surface-active agents.
Surfactants can get adsorbed on the cathode surface leading to
unfavorable changes in the mechanical properties of the electrodeposit, such as high internal residual stress or brittleness [9]. Due
to these reasons, the search for new approaches for synthesizing
heterogeneous, multi-phase deposits with enhanced and tunable
properties is highly desirable.
Cathodic electrodeposition of certain ceramic materials from
their metal salts, including metal oxides/hydroxides or complex
oxide compounds, has been demonstrated over the last few years
[10–14]. In particular, deposition of yttrium hydroxide has received
particular attention because it enhances the corrosion resistance of
several metals and alloys such as carbon steel, stainless steel, zinc,
bronze, aluminium and magnesium alloys [15]. Research on Y(OH)3
K.M. Sivaraman et al. / Electrochimica Acta 56 (2011) 5142–5150
deposition has been primarily motivated by the need to replace
Cr(VI) based corrosion inhibitors by eco-friendly protective layers,
given that Cr(VI) is extremely toxic and carcinogenic [16]. In addition to the corrosion inhibiting performance, electrodeposition of
Y(OH)3 has also been a topic of research because it can be converted into Y2 O3 (yttria) after an appropriate thermal treatment.
Yttria is highly resistant to most chemicals and has a melting point
of approximately 2400 ◦ C. It also exhibits a low vapor pressure, optical transparency in the infrared region, high corrosion and thermal
resistance, and high strength and fracture toughness [17–20]. This
has prompted the use of yttria in a variety of industrial applications, including optical sensors, cutting tools, protective coatings,
and solid oxide fuel cell (SOFC) interconnects [20–23].
Cathodic electrodeposition of Y(OH)3 can be performed from
either organic (typically, water–alcohol mixtures) or fully aqueous
solvents. According to the early work by Switzer, the formation of
an electrogenerated base leads to the precipitation of Y(III) from
the electrolyte as a hydroxide film on the surface of the cathode
[24]. From ecological and health perspectives, fully aqueous electrolytes are certainly more desirable. However, there is a dearth
of literature concerning the aqueous electrodeposition of Y(OH)3
films. In [25,26], Y(NO3 )3 and YCl3 salts dissolved in water were
used to cathodically deposit Y(OH)3 thin films and nanospheres. In
both works, Y(OH)3 was thermally treated at 600 ◦ C in air to obtain
crystalline yttria.
To date, attempts to incorporate yttrium oxides/hydroxides
into metallic electrodeposits have proven difficult. The production
of nickel–yttria composite coatings by direct current electroplating from a Watts nickel plating bath containing variable amounts
of micron-sized yttria particles suspended in the electrolyte was
explored by McCormack et al. [27]. However, there was little or no
yttria inclusion in the deposits for the investigated range of yttria
loadings (1–5 g dm−3 ). The authors reported only on the effects
caused by the presence of yttria dispersed in the electrolytic bath
on the Ni deposit texture and morphology. Hence, incorporation
of a useful amount of yttrium-based compounds into electroplated
metallic films remains a challenge and requires investigation.
The aim of the present work is to demonstrate the feasibility
of producing yttrium-containing cobalt films by electrodeposition from aqueous solutions. The novelty of the present synthetic
approach relies on incorporating yttrium oxides/hydroxides into
nanocrystalline Co films from a fully aqueous electrolyte containing both YCl3 and CoCl2 salts. Glycine was chosen as a complexing
agent since it shows high buffering properties and favors the codeposition of iron-group (Fe, Co and Ni) with rare earth (RE) metals
[28]. In order to induce both nanocrystallinity of the Co matrix and
an effective inclusion of yttrium oxide/hydroxide into the deposits,
the pulse plating technique was utilized.
5143
Table 1
Detailed composition of the employed baths, in mol dm−3 .
Bath
YCl3 ·6H2 O
CoCl2 ·6H2 O
Glycine (C2 H5 NO2 )
KCl
H3 BO3
1
2
3
4
0
1
1
1.5
0.1
0.05
0.1
0.1
0.15
1
0.16
wherein the [Co(II)] was much smaller than the [Y(III)] (Table 1). KCl
was used as a supporting electrolyte and boric acid as a buffering
agent. The working temperature was 60 ◦ C. This temperature was
selected on the basis that yttrium chemically behaves as a RE element. RE metal deposition is generally more favorable above room
temperature [28]. The pH was kept at 4 for all the electrodeposition
experiments.
The working electrode consisted of 5 mm × 6 mm silicon chips
(crystal orientation 1 0 0), on top of which a titanium adhesion
layer of 50 nm and a copper seed-layer of 500 nm had been successively deposited through e-beam evaporation. Evaporation was
performed using an E306A coating apparatus controlled by an eVap CVS-6 system. Before being used for electrodeposition, the
copper surface was degreased by first dipping it in acetone followed
by iso-propanol and then stripped of oxide and organic residues
by dipping in 10% H2 SO4 solution. The three electrodes were connected to a Gamry Reference 3000 potentiostat/galvanostat which
was controlled by the Gamry Framework software during galvanostatic and linear sweep voltammetry (LSV) experiments and
by the Gamry VFP-600 software during pulsed electrodeposition.
The charge density was kept at 20 C cm−2 across all depositions.
Deposits were obtained by galvanostatic pulse plating with a
cathodic pulse (CP) (from −50 to −1500 mA cm−2 ) with a duration
of ton = 0.1 ms followed by an open circuit potential (OCP) with a
duration of toff = 0.9 ms (Fig. 1). The data acquisition frequency was
10 kHz which provided 10 data points for each cycle. The deposition
was carried out under quiescent conditions.
The morphology of the deposits was observed using Zeiss NVision 40 and Zeiss Merlin scanning electron microscopes (SEM).
The chemical composition of the deposits was analyzed by energy
dispersive X-ray spectroscopy (EDX). Data reported are based on
the averaging of four measurements on three replicas per sample. The yttrium content is calculated by taking into account the
weight percentages of cobalt and yttrium elements (without considering the oxygen content) normalized to 100%. For cross-section
SEM images, the films were embedded in epoxy resin and polished to a mirror-like appearance using diamond paste. X-ray
2. Experimental
The electrodeposition was performed in a one-compartment
three-electrode cell. A 100 ml of bath was used in the electrochemical cell. Distilled water from a Julabo ED-5 thermostat-circulator
system was circulated in the outer jacket of the cell to maintain
the required working temperature. A double junction Ag|AgCl reference electrode was used with 3 mol dm−3 KCl inner solution and
an outer solution made of 1 mol dm−3 NaCl. An insoluble platinum
wire mesh acted as counter electrode. An inert atmosphere was
ensured inside the electrochemical cell by maintaining a blanket of
nitrogen gas on top of the solution. The level of water in the cell
and the pH of the bath were checked at the beginning and end of
each deposition. The required acidity was maintained using sodium
hydroxide and hydrochloric acid. To study the feasibility of electrodepositing yttrium-containing films, a set of baths was prepared
Fig. 1. Plot of the galvanostatic pulse with a CP of 0.1 ms and an OCP of 0.9 ms.
5144
K.M. Sivaraman et al. / Electrochimica Acta 56 (2011) 5142–5150
Table 2
Y content in the deposits as a function of applied CP for Baths 2–4. Note that the Y
content is given taking into account the weight percentages of Co and Y (without
considering O), normalized to 100%.
−j (mA cm−2 )
750
1000
1200
1500
Fig. 2. Linear sweep voltammetries of Baths 1–4 recorded under quiescent conditions, at a scan rate of 50 mV s−1 .
photoelectron spectroscopy (XPS) analyses were carried out on
a PHI equipment 5500 Multitechnique using the Al K␣ radiation
(1486.6 eV), after sputtering the sample surface with Ar ions for
1 min. All spectral positions have been corrected taking C 1s peak
at 284.5 eV. X-ray diffraction (XRD) patterns were obtained with
a Philips X’Pert diffractometer using the Cu K␣ radiation in the
35–100◦ 2 range (0.03◦ step size, 10 s holding time). Nanoindentation tests were carried out in the load control mode using
a UMIS device from Fischer-Cripps Laboratories equipped with
Berkovich pyramid-shaped diamond tip. The load–unload curves
were taken on the films’ cross-section applying a maximum force
of 20 mN. Hardness (H) values were obtained applying the method
of Oliver and Pharr [29] at the beginning of the unloading curve.
Data reported corresponds to the averaging of 20 indentations per
sample.
3. Results and discussion
3.1. Linear sweep voltammetry and pulse galvanostatic transients
During the design and development of electrolytic solutions for
incorporating yttrium into cobalt deposits, considering bath stability was essential. In this context, at 25 ◦ C all glycine-containing
baths were stable from pH 3 through pH 5. However, at 60 ◦ C,
the baths at pH 5 were not stable. Turbidity was observed but
was reversible on bringing the pH to 4. Hence, pH 4 was chosen to synthesize the cobalt electrodeposits containing yttrium
oxides/hydroxides. It can be conjectured that, the formation of
insoluble species of Y3+ is promoted at pH 5. Tkachenko et al. [30]
point out that, in the presence of boric acid and ammonium hydroxide, there is a possibility of formation of yttrium orthoborate salts
around pH 5. Therefore, the observed turbidity might be due to the
presence of both boric acid and the ammonium group in glycine. At
room temperature, the baths exhibited a pink coloration characteristic of octahedral aqua complexes of Co2+ . When the temperature
of the solution was increased, the bath became violet-blue in color.
This can be attributed to the fact that a fraction of the Co(II) ions
form tetrahedral chloride [CoCl4 ]2− complexes.
The characteristics of cobalt–yttrium electrodeposition in the
presence of glycine at pH 4 were investigated by LSV. On comparing the response of Baths 1, 3 and 4, it can be seen that, for a
given Co(II) concentration, the onset of deposition remains almost
the same when Y(III) is added to the bath (cf. Baths 1 and 3), but
slightly shifts towards more positive potentials when Y(III) concentration is further increased (Bath 4) (Fig. 2). Simultaneously, the
diffusion-controlled reduction peak is also shifted towards more
wt% Y
Bath 2
Bath 3
3
2
5
3
2
3
3
2
Bath 4
2
2
3
2
positive potentials and its height decreases, suggesting changes in
the diffusion/charge controlling factors. As expected, for a constant
Y(III) concentration, a decrease in the concentration of Co(II) (Baths
2 and 3) causes a delay in the onset of deposition and also a decrease
in the height of the reduction peak (Fig. 2). These trends suggest
that addition of Y(III) species to cobalt(II)–glycine solutions causes
changes in the electrolyte solution which ultimately might affect
the characteristics of the reduction process.
Fig. 3 shows the galvanostatic cathodic pulses for the initial
and the final cycles at each current density. Especially for the final
cycles, the potentials corresponding to the pulse are clearly differentiated by the applied current density. The system reaches more
negative potentials as the cathodic pulse is increased. For the last
cycles, the potentials corresponding to the relaxation time seem
to converge to a similar potential value. The regular pattern of
the CP–OCP cycles suggests that deposition proceeds in a wellcontrolled manner.
3.2. Morphology and composition of yttrium-containing cobalt
deposits
Table 2 shows the composition of the deposits obtained from
Baths 2 to 4. Yttrium content varied between 1 and 5 wt%, and Y
was detected only in deposits obtained using galvanostatic pulses
more negative than −500 mA cm−2 . Upon unaided visual inspection, all the films appeared matte dark and well adherent to the
substrate. Most importantly, yttrium incorporation was almost
negligible under direct current electroplating conditions. The pulse
plating mode allowed the application of very high peak current densities, which could be the key factor in enabling Y incorporation.
The results show random variability in the yttrium content, with
no clear tendency observed with respect to both the applied current
density and the [Y(III)]/[Co(II)] ratio in solution, though maximum
Y incorporation was usually attained at CP of −1200 mA cm−2 .
Though levels of Y incorporation were not very high, clear
changes in morphology were observed with respect to pure cobalt
deposits (Fig. 4), indicating a clear influence of Y on the morphology
of the deposits. Yttrium-free cobalt films displayed rounded grains
which looked highly facetted at higher magnification (Fig. 4(a) and
(b)), whereas elongated grains were observed when Y was incorporated in the deposits (Fig. 4(c) and (e)). The distinctive feature of
these grains was the presence of ridge-like structures on their surface, which gave them a seashell-like appearance (Fig. 4(d) and (f)).
The width of these ridge-like structures was in the range of a few
nanometers (roughly around 10–20 nm). Also, the trends observed
in the surface morphology of these deposits were independent of
the [Co(II)] in their deposition baths. Significantly, as opposed to the
deposits mentioned in [31], these deposits were completely free of
micro-cracks. This can also be attributed to the use of the pulse
plating technique [32]. For a given bath, the grains became slightly
finer with an increase in the applied current density. On the other
hand, an increase of Y(III) concentration in solution (Bath 4) made
the grains evolve to more rounded shapes, but they still displayed
characteristic ridge-like features on their surface as in the deposits
K.M. Sivaraman et al. / Electrochimica Acta 56 (2011) 5142–5150
5145
Fig. 3. Galvanostatic cathodic pulses for the initial and final cycles for Bath 3 ([Co(II)] = 0.1 mol dm−3 , [Y(III)] = 1 mol dm−3 ) and quiescent conditions.
prepared from Bath 3. The pure cobalt deposits were always free of
such ridge-like features.
Yttrium-rich deposits (with Y percentages larger than 5 wt% on
average) could be obtained after a number of successive depositions
from a given bath, typically after depositing 10–15 films, and for
CPs more negative than j = −750 mA cm−2 . The mechanism behind
the growth of such yttrium-richer deposits is not yet fully understood. Since the depletion of Co(II) ions in solution could intuitively
Fig. 4. SEM images (secondary electrons) of deposits: (a and b) from Bath 1 ([Co(II)] = 0.1 mol dm−3 , [Y(III)] = 0 mol dm−3 ) at j = −1500 mA cm−2 ; (c and d) from Bath 3
([Co(II)] = 0.1 mol dm−3 , [Y(III)] = 1 mol dm−3 ) at j = −750 mA cm−2 ; (e and f) from Bath 3 at j = −1500 mA cm−2 .
5146
K.M. Sivaraman et al. / Electrochimica Acta 56 (2011) 5142–5150
Fig. 5. SEM images (secondary electrons) of yttrium-rich films: (a and b) 13 wt% Y deposit obtained from Bath 2 ([Co(II)] = 0.05 mol dm−3 , [Y(III)] = 1 mol dm−3 ) at
j = −1500 mA cm−2 ; (c and d) 19 wt% Y deposit obtained from Bath 3 ([Co(II)] = 0.1 mol dm−3 , [Y(III)] = 1 mol dm−3 ) at j = −1000 mA cm−2 ; (e and f) 30 wt% Y deposit obtained
from Bath 3 at j = −1200 mA cm−2 .
explain such enrichment, freshly prepared baths with a lower CoCl2
concentration were tested. However, no clear enrichment in Y was
observed. Thus, it is likely that concurrence of several factors during
successive depositions (e.g., changes in the distribution of chloro
and glycine–metal complexes, local pH variations, etc.) takes place,
which ultimately lead to the formation of Y-rich deposits. Despite
the strong buffering properties of glycine and boric acid, the bulk
electrolyte became slightly acidic (pH around 3.5) after each deposition. This decrease in pH can be correlated to an increase in the
concentration of protons that are generated at the cathode as a
replacement for the metal ions that get reduced. On the other hand,
local pH changes near the cathode surface are likely to take place
due to the electrogeneration of hydroxide ions. Such local changes
of pH could affect the morphology and chemical composition of
deposits as well. Despite this, there was no turbidity or precipitation observed at the end of deposition, nor was there any change in
the color of the solution. Fig. 5 shows selected SEM images of these
yttrium-rich deposits. At first glance, it is clear that the morphology of these deposits differs from the ones previously described.
Namely, nodular grains developed with an increase in yttrium
percentage (Fig. 5(a), (c) and (e)). Simultaneously, incipient microcracking appeared, which can be attributed to the stress caused
by high levels of yttrium inclusion. Such a change in morphology
towards hemispherical grains was also observed in nickel electrodeposits obtained in the presence of yttria particles suspended
in the electrolyte [27]. When the films were imaged at higher magnification, cauliflower-like features were observed for intermediate
yttrium percentages (Fig. 5(b)) while the nodules found for larger
yttrium contents consisted of nanometric rod-like subgrains (see
Fig. 5(d) and (f)).
As expected, the measured thickness of the deposits was lesser
than the theoretical value due to the intense evolution of hydrogen
gas during the deposition. Evolution of hydrogen gas bubbles was
noticed during the deposition process, but the bubbles did not stick
to the substrate. In all cases, the deposit thickness was rather regular across the substrate (Fig. 6(a)). A careful examination of their
cross-sections revealed the existence of differences in chemical
composition as detected qualitatively by back-scattered electron
(BSE) imaging. Fig. 6(b) shows a typical secondary-electron SEM
image (right) and its corresponding BSE image (left) of a zoomed
K.M. Sivaraman et al. / Electrochimica Acta 56 (2011) 5142–5150
5147
Fig. 6. (a) Cross-sectional SEM image of a 5 ␮m thick deposit with 13 wt% Y (Bath 2, j = −1500 mA cm−2 ). (b) Backscattered-electrons (BSE) (left) and secondary-electrons
SEM (right) images of a zoomed area. (c) Detail of two nanoindentations on deposit’s cross-section. The profile of one of the indentations is indicated with a dotted line.
region. Notice that the BSE image is not affected by topological features (e.g., changes in surface flatness due to polishing scratches).
Hence, the phase contrast is linked to variations in composition. In
order to gain a better understanding on such differences, EDX-spot
analyses were carried out (see Fig. 7(a) and Table 3). It was observed
that yttrium is distributed across the entire thickness of the films.
Interestingly, the regions enriched in yttrium coincide with depletion in cobalt content. Conversely, larger cobalt contents coincide
with a decrease in yttrium percentage. If only an alloy or a mixed
cobalt–yttrium oxide had formed throughout the film, then homogeneous distribution of Co, Y (and O) elements would have been
Table 3
Yttrium and cobalt weight percentages (normalized to 100%) determined by EDXspot analyses at the regions indicated in Fig. 7(a).
Region
wt% Co
1
2
3
4
97
87
81
94
wt% Y
3
13
19
6
Fig. 7. (a) Cross-sectional secondary-electrons SEM image of a deposit with 13 wt% Y (Bath 2, j = −1500 mA cm−2 ). EDX mappings taken on this area for (b) Co K␣1, (c) Y
L␣1 and (d) O K energies. Cobalt- and yttrium-rich zones are indicated by green and red arrows, respectively. Notice that yttrium-enriched zones coincide with depletion of
cobalt. Conversely, yttrium-depleted regions coincide with cobalt enrichment. (For interpretation of the references to color in this figure caption, the reader is referred to
the web version of the article.)
5148
K.M. Sivaraman et al. / Electrochimica Acta 56 (2011) 5142–5150
Fig. 8. XPS survey spectrum of an yttrium-containing Co deposit obtained at CP of
j = −750 mA cm−2 from Bath 3 ([Co(II)] = 0.1 mol dm−3 , [Y(III)] = 1 mol dm−3 ).
observed. Complementary EDX mappings confirmed this observation (see Fig. 7(b)–(d)). The oxygen content as determined by EDX
was found to be significant and increased in yttrium-rich films.
3.3. XPS analyses
Fig. 8 shows a typical XPS survey spectrum of a deposit containing low amounts of yttrium, featuring the characteristic peaks of
cobalt and yttrium elements. Copper and carbon peaks, belonging
to the copper seed-layer and surface contaminants, respectively,
were also identified. Oxygen was also detected, in agreement with
EDX analyses, suggesting that either oxide or hydroxide groups
were present in the deposits. The corresponding core-level Y3d,
Co2p and O1s spectra are displayed in Fig. 9. The Co2p spectrum shows a complex structure broadened by multiplet splitting
effects, the sharpest peaks at 778.5 and 793.8 eV being characteristic of metallic cobalt [33,34]. On the other hand, less intense and
broadened peaks denoted by an asterisk are attributed to Co2+ compounds. According to the literature, the Co2+ binding energies in
CoOOH and Co(OH)2 are about ∼780 and ∼781 eV, respectively
[35]. Likewise, CoO displays characteristic reflections at around 781
and 797 eV [36]. Hence, it can be concluded that metallic cobalt is
accompanied by small amounts of oxidized cobalt. This is indeed
corroborated on analyzing the O1s signal. Namely, the dominant
feature in the O1s spectrum at around 532.4 eV can be assigned to
structure water, which is commonly detected in electrodeposited
oxides and hydroxides [37]. Also, the shoulder located at lower
binding energies can be deconvoluted, according to the literature,
into the contributions of hydroxide OH− (531.5 eV) and lattice oxide
O2− ions (530 eV). Concerning the Y3d spectrum, the two peaks
centered at 158.6 eV and 156.7 eV match with the binding energies expected for Y–OH and Y–O bindings [38,39]. This leads to
the hypothesis that yttrium incorporation into cobalt films chiefly
proceeds via hydrolysis of Y(III) species in solution by an electrogenerated base such as OH− (Eq. (1)) [26]. As a result, colloidal
Y(OH)3 forms (Eq. (2)), which tend to accumulate at the cathode
vicinity and become entrapped into the growing cobalt deposit
during electrodeposition, rendering a nanocomposite structure (Eq.
(3)):
−
−
2H2 O + 2e → 2 OH (aq) + H2 (g)
(1)
Y 3+ (aq) + 3OH− (aq) → Y(OH)3 (ads)
(2)
Fig. 9. Co2p, O1s and Y3d spectra of same deposit in Fig. 8. Note that the symbol *
in cobalt spectrum indicates weak peaks corresponding to Co2+ compounds.
K.M. Sivaraman et al. / Electrochimica Acta 56 (2011) 5142–5150
5149
Table 5
Hardness values, evaluated by nanoindentation on the cross-section of various
deposits, as a function of yttrium content.
Fig. 10. XRD spectra of deposits with 0 wt% (black curve), 3 wt% (dark grey curve)
and 13 wt% (light grey curve) of yttrium. The Miller indices corresponding to the hcp
planes of the ␣-Co phase are indicated. Sharp reflections belong to the Cu seed-layer.
Co2+ (aq) + Y(OH)3 (ads) + 2e− → Co(s) + Y(OH)3 (s)
(3)
The detection of Y–O binding also suggests the deposition of
some Y2 O3 . The concurrent presence of glycine and Co(II) ions in
solution could make the deposition of Y2 O3 feasible. In fact, samarium oxides were also detected in cobalt–samarium electrodeposits
obtained from chloride salts in glycine-containing electrolytes [40].
Similar XPS results were obtained for deposits containing larger
amounts of yttrium. For these films, the higher oxygen signal
detected by XPS can be explained on the basis of larger amounts
of Y(OH)3 /Y2 O3 being incorporated into the composite. The darker
regions in the BSE image of Fig. 6(b) probably correspond to
Y(OH)3 /Y2 O3 compounds embedded in the metallic cobalt matrix,
which appears brighter.
3.4. Structure of the deposits
The structure of the deposits was analyzed by XRD. Peaks were
rather broad, indicating the nanocrystalline character of the films
(Fig. 10). The main reflections could be indexed on the basis of
a hexagonal close-packed (hcp) unit cell (␣-Co phase) independent of the baths used. Small amounts of CoO and Co(OH)2 were
also detected. These compounds would be likely formed during
the electrodeposition process. Moreover, the surface of cobalt is
prone to oxidation in air, leading to the formation of a native oxide
layer as well. In addition, the broad background between 35◦ and
55◦ 2␪ range suggests the presence of amorphous material. In fact,
electrodeposited Y(OH)3 from chloride baths was reported to be
fully amorphous [26]. The angular positions of the ␣-Co XRD peaks
did not vary as a function of the Y content, indicating that no Y
was incorporated as a solid solution within the hcp Co structure.
Furthermore, no additional XRD peaks appeared for large Y percent-
wt% Y
H (GPa) (±0.2 GPa)
0
3
5
13
19
30
1.9
2.2
2.6
2.7
3.4
2.2
ages. This suggests that there is no formation of Co–Y alloys during
film growth. Table 4 lists the full width at half maximum (FWHM) of
(1 0 0), (1 0 1) and (1 1 0) hcp reflections for a series of deposits with
progressively larger Y content. The FWHM values from films containing low amounts of yttrium do not vary significantly from those
of pure Co films. However, in films with a relatively high Y content,
the (1 0 0) and (1 0 1) peaks clearly broaden, suggesting that the
incorporation of yttrium hydroxide/oxide into the films promotes
grain-refining of the ␣-Co phase. The crystallite size estimated by
applying the Scherrer’s formula on the (1 0 0) peak width gives values of ∼13 nm for almost pure cobalt films and of ∼9 nm for films
with the highest yttrium percentages. Most importantly, the (1 0 1)
reflection becomes significantly wider in comparison with (1 0 0)
and (1 1 0) reflections, which indicates a higher density of stacking
faults [41]. This is evidenced in the ratio between the peaks widths
given in Table 4. Hence, the effect of yttrium hydroxide/oxide incorporation into cobalt deposits is twofold: it induces crystallite size
refinement of the hcp phase, and it increases the amount of stacking
faults. Overall, since the background broadens with yttrium incorporation, it can be concluded that the resulting composites exhibit a
mixed nanocrystalline/amorphous character. The growth of amorphous Y(OH)3 and Y2 O3 films by physical/chemical methods is well
documented [26,42]. However, the direct growth of metal–metal
oxide or metal–metal hydroxide composite films from particle-free
electrolytes has not yet been reported.
3.5. Mechanical properties
The hardness (H) of the nanocomposites was evaluated by
nanoidentation. It is known that rough surfaces tend to increase
the scatter in the measured hardness, leading to underestimation
[43]. For this reason, given the rough finish of the surface of the
deposits, analyses were made on the cross-section of the deposits
after proper mechanical polishing (Fig. 6(a)). Fig. 6(c) shows a detail
of two nanoidentations. Table 5 lists the H values as a function of
yttrium content in the films. Note that in Fig. 6 the lateral size of the
indent impressions is about 2.5–3 ␮m. Since the variations in composition within the films are much smaller in size (see Fig. 7), each
indentation covers both cobalt and yttrium hydroxide/oxide phases
and are thus, representative of the mechanical strength of the entire
film. It can be observed that H increases with the incorporation of
Table 4
FWHM values of the (1 0 0), (1 0 1) and (1 1 0) hcp reflections and their ratios as a function of the yttrium content in the deposits.
Bath
1
1500
500
750
1000
1500
1000
1200
2
3
a
−j (mA cm−2 )
Y-rich deposits.
wt% Y
0
1
2
3
13a
19a
30a
FWHM (◦ )
Ratio of XRD
peak widths
(1 0 0)
(1 0 1)
(1 1 0)
FWHM(1 0 0)
FWHM(1 0 1)
FWHM(1 1 0)
FWHM(1 0 1)
0.48
0.63
0.64
0.53
0.69
0.71
0.88
0.97
1.03
0.98
1.02
1.43
1.62
2.01
0.85
1.28
1.27
1.20
1.07
1.08
1.34
0.495
0.612
0.703
0.519
0.482
0.438
0.438
0.876
1.243
1.296
1.176
0.748
0.666
0.666
5150
K.M. Sivaraman et al. / Electrochimica Acta 56 (2011) 5142–5150
yttrium oxide/hydroxide. In particular, H increases from 1.9 GPa (in
films without Y incorporation) to 3.4 GPa in films containing 19 wt%
Y. This could be due, at least in part, to the intrinsically large hardness of yttrium oxide and the mechanical interactions occurring
between the yttrium hydroxide/oxide and hcp-Co phase during the
course of nanoindentation experiments. The Vickers hardness of
pure Co is around 1 GPa, whereas that of Y2 O3 exceeds 9 GPa [21].
Therefore, it is likely that the incorporation of Y(OH)3 and Y2 O3
causes an increase of mechanical strength. In addition, such an
increase in hardness can be explained by taking into account both
the grain refining effect and the higher density of stacking faults
caused by yttrium hydroxide/oxide incorporation into the films.
In crystallites with small sizes, the motion of dislocations is hindered to a large extent by grain boundaries. This leads to piling up
of dislocations, and, consequently to an increase in stress concentration and hardness. Apart from the presence of grain boundaries,
the increase in hardness can be due to the presence of stacking
faults as well. In terms of disrupting the motion of dislocations,
stacking faults play a role similar to that of high-angle grain boundaries [44,45]. The decrease in hardness in deposits with a higher Y
wt% (∼30%) is probably due to an increase in the density of microcracks. Indeed, if one compares the SEM images of samples with
13 wt% Y (Fig. 5 panels (a) and (b)), 19 wt% Y [panels (c) and (d)]
and 30 wt% Y [panels (e) and (f)] it is clear that the most dense film
(with less porosity and lower density of micro-cracks) is the one
with 19 wt%, which corresponds to the sample with largest value of
hardness. The sample with 13 wt% Y shows enhanced porosity with
respect to the one with 19 wt% Y (and, concomitantly, the hardness
decreases from 3.4 GPa to 2.7 GPa). In turn, the film with 30 wt%
Y shows a higher density of micro-cracks. Such micro-cracks are
known to have a detrimental effect on the mechanical properties
of films (e.g., hardness), particularly when the indentation impression is performed near or on top of a micro-crack. This is probably
the reason for the lower value of hardness (2.2 GPa) obtained
in this film. Hence, a combination of both compositional and
microstructural effects needs to be taken into account to explain the
hardness variations observed in the electroplated nanocomposite
films.
4. Conclusions
The synthesis of cobalt–yttrium hydroxide/oxide nanocomposites has been accomplished by galvanostatic pulse plating directly
from chloride salts. Incorporation of these yttrium compounds into
the cobalt films brings about significant changes at morphological
and structural levels, even for small yttrium percentages. In spite
of the high current densities applied, the composites were mostly
crack-free and adhered well to the substrate. XRD analyses revealed
that the films displayed mixed nanocrystalline/amorphous character. Incorporating yttrium oxides/hydroxides into the electrodeposits promoted grain size refinement of the ␣-Co matrix phase,
and increased the quantity of stacking faults. Both effects play
a key role on the mechanical behavior of the composites, which
displayed larger hardness compared to pure cobalt coatings. This
particle-free electrodeposition paves the way for synthesizing new
metal–ceramic nanocomposites.
Acknowledgements
This work is supported by the NCCR Co-Me of the Swiss National
Science Foundation. E.P., J.S., S.S. and M.D.B. acknowledge fund-
ing from the Generalitat de Catalunya through the 2009-SGR-1292
project and from the Spanish Ministry of Science and Innovation (MICINN) through MAT2007-61629. The authors sincerely
acknowledge the staff from the Servei de Microscòpia of the Universitat Autònoma de Barcelona for their assistance with SEM
characterization. E.P. is indebted to the Generalitat de Catalunya
for the Beatriu de Pinós postdoctoral fellowship. M.D.B. was partially supported by an ICREA Academia award. S.P. acknowledges a
postdoctoral fellowship from MICINN.
References
[1] A. Hovestad, L.J.J. Janssen, J. Appl. Electrochem. 25 (1995) 519.
[2] M. Musiani, Electrochim. Acta 45 (2000) 3397.
[3] L. Benea, P.L. Bonora, A. Borello, S. Martelli, F. Wenger, P. Ponthiaux, J. Galland,
J. Electrochem. Soc. 148 (2001) C461.
[4] S. Arai, T. Saito, M. Endo, J. Electrochem. Soc. 157 (2010) D147.
[5] D. Thiemig, A. Bund, J.B. Talbot, Electrochim. Acta 54 (2009) 2491.
[6] J. Fransaer, E. Leunis, T. Hirato, J.-P. Celis, J. Appl. Electrochem. 32 (2002) 123.
[7] J. Fransaer, J.P. Celis, J.R. Roos, J. Electrochem. Soc. 139 (1992) 413.
[8] M.D. Ger, Mater. Chem. Phys. 87 (2004) 67.
[9] T.J. Tuaweri, G.D. Wilcox, Surf. Coat. Technol. 200 (2006) 5921.
[10] A.R. Boccaccini, I. Zhitomirsky, Curr. Opin. Solid State Mater. Sci. 6 (2002) 251.
[11] S.H. Baeck, T. Jaramillo, G.D. Stucky, E.W. McFarland, Nano Lett. 2 (2002) 831.
[12] E. Gómez, E. Pellicer, E. Vallés, J. Electroanal. Chem. 580 (2005) 238.
[13] C. Lepiller, S. Poissonnet, P. Bonnaillie, G. Giunchi, F. Legendre, J. Electrochem.
Soc. 151 (2004) D13.
[14] X. Xia, I. Zhitomirsky, J.R. McDermid, J. Mater. Process. Technol. 209 (2009)
2632.
[15] M. Tran, D. Mohammedi, C. Fiaud, E.M.M. Sutter, Corros. Sci. 48 (2006) 4257.
[16] L. Nylén, J. Gustavsson, A. Cornell, J. Electrochem. Soc. 155 (2008) E136.
[17] R. Siab, G. Bonnet, J.M. Brossard, J. Balmain, J.-F. Dinhut, Appl. Surf. Sci. 253
(2007) 3425.
[18] T. Ikegami, J.-G. Li, T. Mori, J. Am. Ceram. Soc. 85 (2002) 1725.
[19] J. Al-Haidary, M. Al-Haidari, S. Qrunfuleh, Biomed. Mater. 3 (2008) 015005.
[20] G.Y. Chen, G. Somesfalean, Z.G. Zhang, Q. Sun, F.P. Wang, Opt. Lett. 32 (2007)
87.
[21] A.S. Kumar, A.R. Durai, T. Sornakumar, Mater. Lett. 58 (2004) 1808.
[22] J.-M. Brossard, J. Balmain, J. Creus, G. Bonnet, Surf. Coat. Technol. 185 (2004)
275.
[23] E. Tondo, M. Boniardi, D. Cannoletta, M.F. De Riccardis, B. Bozzini, J. Power
Sources 195 (2010) 4772.
[24] J.A. Switzer, Am. Ceram. Soc. Bull. 66 (1987) 1521.
[25] I. Zhitomirsky, A. Petric, J. Mater. Chem. 10 (2000) 1215.
[26] M. Aghazadeh, A. Nozad, H. Adelkhani, M. Ghaemi, J. Electrochem. Soc. 157
(2010) D519.
[27] A.G. McCormack, M.J. Pomeroy, V.J. Cunnane, J. Electrochem. Soc. 150 (2003)
C356.
[28] M. Schwartz, N.V. Myung, K. Nobe, J. Electrochem. Soc. 151 (2004) C468.
[29] W.C. Oliver, G.M. Pharr, J. Mater. Res. 7 (1992) 1564.
[30] E.A. Tkachenko, R. Mahiou, G. Chadeyron, D. Boyer, P.P. Fedorov, S.V. Kuznetsov,
Russ. J. Inorg. Chem. 52 (2007) 829.
[31] S. Pané, O. Ergeneman, K. Sivaraman, E. Pellicer, M.D. Baró, B.J. Nelson, Meet.
Abstr. - Electrochem. Soc. 902 (2009) 3132.
[32] E. Pellicer, E. Gómez, E. Vallés, Surf. Coat. Technol. 201 (2006) 2351.
[33] F. Allegretti, G. Parteder, M.G. Ramsey, S. Surnev, F.P. Netzer, Surf. Sci. 601 (2007)
L73.
[34] P.R. Sajanlal, T. Pradeep, J. Phys. Chem. C 114 (2010) 16051.
[35] I.G. Casella, M.R. Guascito, J. Electroanal. Chem. 476 (1999) 54.
[36] W. Chu, P.A. Chernavskii, L. Gengembre, G.A. Pankina, P. Fongarland, A.Y. Khodakov, J. Catal. 252 (2007) 215.
[37] J.-K. Chang, M.-T. Lee, C.-H. Huang, W.-T. Tsai, Mater. Chem. Phys. 108 (2008)
124.
[38] V. Bondarenka, S. Grebinskij, S. Kaciulis, G. Mattogno, S. Mickevicius, H. Tvardauskas, V. Volkov, G. Zakharova, J. Electron. Spectrosc. Relat. Phenom. 120
(2001) 131.
[39] NIST X-ray Photoelectron Spectroscopy Database: http://srdata.nist.gov/
xps/Default.aspx.
[40] J.C. Wei, M. Schwartz, K. Nobe, J. Electrochem. Soc. 155 (2008) D660.
[41] G.B. Mitra, N.C. Halder, Acta Cryst. 17 (1964) 817.
[42] S.N. Mukherjee, C.R. Aita, J. Vac. Sci. Technol. A 10 (1992) 2723.
[43] M.S. Bobji, S.K. Biswas, J. Mater. Res. 14 (1999) 2259.
[44] J. Sort, A. Zhilyaev, M. Zielinska, J. Nogués, S. Suriñach, J. Thibault, M.D. Baró,
Acta Mater. 51 (2003) 6385.
[45] L. Lu, Y. Shen, X. Chen, L. Qian, K. Lu, Science 304 (2004) 422.