Electrodeposition of granular Cu-Co alloys

Article abstract of DOI:10.1149/1.1527938

Electrodeposition characteristics of Cu-Co films were studied for the formation of heterogeneous alloys for giant magnetoresistance applications. In situ scanning tunneling microscopy, Auger electron spectroscopy (AES), and high resolution scanning electron microscopy studies showed that rough films with a low concentration of cobalt [Cu92.5-Co7.5] (atom %) were deposited mainly due to a higher deposition rate of copper than of cobalt toward the end of the deposition process, and due to the formation of copper grains after the electrodeposition process by chemical exchange between copper ions in the solution and with the cobalt in the deposit. AES analysis revealed that the Cu-Co film is not homogeneous; the bulk of the film is richer in Co while the surface and the bottom of the film are Co-poor. X-ray diffraction showed that the electrodeposition is a topotaxial crystallization process and that the as-deposited film is composed of two phases, a solid solution of face centered cubic Cu-Co with preferred orientation of {111} planes, and a hexagonal close packed Co phase. Scanning electron microscopy micrographs and energy dispersive spectroscopy indicated the segregation of cobalt grains resulting from thermal treatments, according to the phase diagram of Cu-Co.

Full text of DOI:10.1149/1.1527938

C28
Journal of The Electrochemical Society, 150 ͑1͒ C28-C35 ͑2003͒
0013-4651/2003/150͑1͒/C28/8/$7.00 © The Electrochemical Society, Inc.
Electrodeposition of Granular Cu-Co Alloys
b
T. Cohen-Hyams,a,z W. D. Kaplan,a D. Aurbach,b, Y. S. Cohen,
*
and J. Yahaloma,
*
aDepartment of Materials Engineering, Technion-Israel Institute of Technology, 32000 Haifa, Israel
bDepartment of Chemistry, Bar Ilan University, 52900 Ramat-Gan, Israel
Electrodeposition characteristics of Cu-Co films were studied for the formation of heterogeneous alloys for giant magnetoresis-
tance applications. In situ scanning tunneling microscopy, Auger electron spectroscopy ͑AES͒, and high resolution scanning
electron microscopy studies showed that rough films with a low concentration of cobalt ͓Cu92.5-Co7.5͔ ͑atom %͒ were deposited
mainly due to a higher deposition rate of copper than of cobalt toward the end of the deposition process, and due to the formation
of copper grains after the electrodeposition process by chemical exchange between copper ions in the solution and with the cobalt
in the deposit. AES analysis revealed that the Cu-Co film is not homogeneous; the bulk of the film is richer in Co while the surface
and the bottom of the film are Co-poor. X-ray diffraction showed that the electrodeposition is a topotaxial crystallization process
and that the as-deposited film is composed of two phases, a solid solution of face centered cubic Cu-Co with preferred orientation
of ͕111͖ planes, and a hexagonal close packed Co phase. Scanning electron microscopy micrographs and energy dispersive
spectroscopy indicated the segregation of cobalt grains resulting from thermal treatments, according to the phase diagram of
Cu-Co.
© 2002 The Electrochemical Society. ͓DOI: 10.1149/1.1527938͔ All rights reserved.
Manuscript submitted February 1, 2002; revised manuscript received June 20, 2002. Available electronically December 3, 2002.
One of the most exciting and startling properties exhibited by
forming continuous series of solid solutions. Thus, the formation of
supersaturated two-phase Co-Cu is possible depending on the depo-
sition conditions.
In this paper, the electrodeposition and properties of a Cu-Co
alloy are described. This alloy should be a metastable phase because
the Cu-Co system is immiscible under equilibrium conditions at
room temperature.
some magnetic multilayer systems is the giant magnetoresistance
͑GMR͒ effect. GMR refers to a significant change in the electrical
resistance of a film or a device when an external magnetic field is
applied. The GMR effect which was first discovered in magnetic
multilayers1 also exists in heterogeneous alloys with ferromagnetic
granules ͑i.e., Fe or Co͒ embedded in a nonmagnetic metal ͑i.e., Cu
or Ag͒.2
Heterogeneous alloy films are a potentially useful alternative for
GMR applications, especially for magnetic sensor applications,
where the sensitivity is less important than the magnitude of the
response.2 It is generally simpler to prepare a heterogeneous alloy
film rather than a multilayer system. The heterogeneous alloy films
are immiscible combinations usually prepared by physical deposi-
tion methods.3,4
Experimental
Electrodeposition.—The electrodeposition experiments were
conducted on 2 cm2 substrates cleaved from Si wafers, which had
been previously coated by physical vapor deposition ͑PVD͒ with a
copper layer ͑75 nm͒ on top of a TaN layer ͑30 nm͒. Some of the
polarization measurements were performed on a platinum ͑Pt͒ sheet
of 4.8 cm2, in order to eliminate the reduction of an air-formed
copper oxide. The substrates were pretreated before the deposition
of the alloy coating by degreasing in subsequent baths of methanol
and acetone, rinsed in distilled water, dipped into a concentrated
H2SO4 solution in order to remove the surface oxide, and finally
rinsed in distilled water.
The electrochemical apparatus used for the alloy deposition was
EG&G ͑PARC͒ model 273 potentiostat-galvanostat. The deposition
was done in a Pyrex glass cell with a standard three-electrode con-
figuration with a platinum wire as a counter electrode and a satu-
rated calomel reference electrode ͑SCE͒, which was positioned close
to the working electrode. The following blank solution was used: 0.2
M H3BO3 and 0.13 M citric acid (C6H8O7). Copper and cobalt ions
were added to the blank solution as CuSO4 and CoSO4 . The adjust-
ment of pH was done by additions KOH or H2SO4 . The deposition
bath composition and the electrochemical parameters were deter-
mined by the experiments, described in the next section.
An alternative processing technique is electrodeposition.5-13
Electrodeposition has several advantages over dry processes. There
are a large number of possible alloy combinations. Electrodeposition
does not require vacuum technology and consequently is less expen-
sive. It can easily be scaled up for use in large size areas, and it is
capable of depositing uniform films on complex surfaces without
shadowing effects often encountered in other deposition methods.
The experimental systems used are much simpler than evaporation
or sputtering apparatus, and electrodeposition can be a room-
temperature technology. There are some drawbacks associated with
the process such as the need for a conducting/semiconducting sub-
strate, the limited number of elements that can be deposited, and the
large number of variables that control this process ͑composition, pH,
concentration, current density, temperature, agitation, etc.͒.
Electrodeposited binary alloys may or may not have the crystal-
lographic structure expected from the equilibrium phase diagram
formed by other methods.14 Gelchinski et al.15 electroplated at room
temperature cobalt alloys containing the A-15 structure phase that is
stable only at high temperature, according to the equilibrium phase
diagram. Usually it is assumed14 that when metals are codeposited at
low polarization values, the formation of solid solutions or super-
saturated solid solution results. This was found to be so even when
the metals are not mutually soluble in the solid state according to the
bulk phase diagram. Codeposition at high polarization values was
reported14 to result in two-phase alloys even for systems capable of
Anodic polarization measurements were conducted on a Pt sub-
strate at room temperature (ϳ25°C) at a scan rate of 0.5 mV/s. The
bath composition used for the anodic polarization consisted of 100
g/L K2SO4 and 20 g/L H3BO3 . The pH of the solution was 4. The
alloy coatings were deposited at room temperature. Following depo-
sition the specimens were annealed at 450°C for 0.5, 1.5, 10, and 20
h in Ar with 7% H2 atmosphere.
Characterization methods.—The microstructural characteristics
of the granular Cu-Co films were investigated by in situ scanning
tunneling microscopy ͑STM͒, a scanning Auger microprobe ͑Auger
electron spectroscopy, AES͒, high resolution scanning electron mi-
croscopy ͑HRSEM͒, and X-ray diffraction ͑XRD͒.
* Electrochemical Society Active Member.
z E-mail: mtcohen@tx.technion.ac.il

Journal of The Electrochemical Society, 150 ͑1͒ C28-C35 ͑2003͒
In situ STM experiments were conducted using a PicoSPM from
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Molecular Imaging. The tunneling tip was a platinum-iridium ͑9:1͒
wire ͑250 ␮m thick͒ coated with polyethylene and etched electro-
chemically using a self-constructed control unit, which uses ac
current.16 The etched tip radius was between 1 and 2 ␮m. The tun-
neling scanning current was set to 0.3 nA. The experiments were
performed in the same solution used for Cu-Co alloy deposition. The
potentials of the specimens and the tip were controlled potentiostati-
cally against a copper electrode as a reference electrode. The poten-
tial of the copper reference electrode in the solution was found to be
0 mV vs. SCE and remained constant throughout the experiment.
A scanning Auger microprobe ͑Physical Electronics PHI 590A͒
system was used to measure the surface composition and the depth
profile of Cu and Co atoms in the Cu-Co film. The incident electron
beam energy was 5 keV, the current was 2 ␮A, and the scanned area
of the beam was 40 ϫ 40 ␮m. Auger transitions of the elements
used were C-KLL, 272; N-KLL, 390; O-KLL, 503; Cu-LMM, 918;
Co-LMM-775; Si-KLL, 1619; and Ta-MNN, 1675 eV. Arϩ ions ac-
celerated at 4 kV were used to sputter an area of ϳ2 ϫ 2 mm. The
incident angles of the Arϩ bombardment and that of the electron
beam relative to the samples were ϳ60 and 30°, respectively. The
approximate sputter rate was 20-100 nm/min, as estimated by a
Ta2O5 standard. Depth profiles were obtained by measuring the
peak-to-peak heights in the first derivative mode of the correspond-
ing Auger transitions.
The composition ratio of the Cu-Co alloy was determined by
energy dispersive spectroscopy ͑EDS͒, using a Pentafet EDS detec-
tor and Oxford Link ISIS Software attached to a Philips XL30 scan-
ning electron microscope ͑SEM͒, operated at 20 kV at a working
distance of 10.5 mm ͑penetration radius of ϳ1 ␮m for Cu K␣ at 20
kV͒. The EDS results were obtained after calibration with a cobalt
standard. HRSEM was conducted using a Leo 982 Gemini equipped
with a Schottky electron source. Micrographs of Cu-Co films before
and after thermal treatments were recorded at an accelerating volt-
age of 4 kV and a working distance of 2 mm.
The XRD system used in this study ͑Philips PW-3020 goniom-
eter͒ included a long fine-focus copper X-ray tube (␭ ϭ 0.15406
nm͒, operated routinely at 40 kV and 40 mA with 1° divergent and
antiscattering slits coupled with 0.2 mm receiving slits. A curved
graphite monochromator (2␣ ϭ 26.4°) preceded the detector. Dif-
fraction patterns were acquired at steps of 0.01° 2␪ and 25 s/step
exposure.
Results
Polarization measurements.—Preliminary experiments were
conducted to find a suitable electrolyte for the deposition of Cu and
Co.17 Generally, since copper is the more noble metal in the system,
the concentration ratio of Co2ϩ to Cu2ϩ ions in the electrolyte has to
be much higher than the ratio required in the alloy system. The
standard electrode potential18 of Cu is ϩ0.337 V ͑vs. standard hy-
drogen electrode, SHE͒ and that of Co is Ϫ0.277 V ͑SHE͒, which
means that their deposition potentials are far apart for noncom-
plexed ions, making the codeposition difficult from a simple sulfate
solution of Cu and Co ions. In order to bring these deposition po-
tentials closer, a complex forming agent of citric acid is added to the
Cu-Co solution.17 The solution also contains boric acid, which acts
as a pH buffer agent. The concentrations of Co2ϩ and Cu2ϩ varied
from 0.08 M Co2ϩ and 0.1 M Cu2ϩ in which no cobalt was detected
in the film, to 0.3 M Co2ϩ and 0.15 M Cu2ϩ in which a high content
of cobalt was obtained in the deposit (Ͼ42 atom %). In order to
understand the influence of each component of the solution, polar-
ization measurements were performed on a Pt substrate at a scan
rate of 5 mV/s in acidic ͑pH 2͒ and neutral ͑pH 7͒ blank solutions
which did not contain Co and Cu ion additions. The polarization was
then repeated after the addition of each of the ions, and a combina-
tion of both. These measurements are shown in Fig. 1 at two values
of pH.
Figure 1. Cathodic potentiodynamic curves in the blank plating electrolyte
with and without additions of 5 mM CuSO4 and/or 0.25 M CoSO4 . Scan rate
was 5 mV/s. ͑a͒ pH 2 and ͑b͒ pH 7. The substrate used was a Pt sheet. The
curves are presented in a linear scale in order to elucidate the influence of the
different components in the solution.
It can be seen in Fig. 1 that at the lower pH the evolution of
hydrogen begins at a less negative potential than at the higher pH
͑Fig. 1b͒, as would be expected. Since copper has a high overvoltage
for hydrogen evolution, a significant H2 evolution starts only at ca.
Ϫ0.8 V ͑SCE͒ for solutions that deposit Cu ions, and at ca. Ϫ0.5 V
͑SCE͒ for the other solutions at pH 2. A similar effect is noticed at
pH 7 where a significant H2 evolution starts at around Ϫ0.7 V
͑SCE͒ for the blank solution while the addition of copper ions in-
hibits hydrogen evolution to ca. Ϫ0.9 V ͑SCE͒.
When 5 mM CuSO4 was added to the blank solution at pH 2
͑Fig. 1a͒, the deposition of copper started at about the equilibrium
value of ϳ0 V ͑SCE͒, with a diffusion-limited process of copper

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Journal of The Electrochemical Society, 150 ͑1͒ C28-C35 ͑2003͒
Figure 2. Potentiodynamic curves of a neutral blank electrolyte ͑pH 7͒ con-
taining 0.125 M CuSO4 and 0.25 M CoSO4 . Scan rate was 5 mV/s. The
substrate used was a PVD Cu platelet.
Figure 3. Anodic potentiodynamic curves at pH 6.5 of Co deposited at
Ϫ1.05 V ͑SCE͒, Cu deposited at Ϫ0.5 V ͑SCE͒, and ϳ1 ␮m Cu-Co alloy
Cu-Co deposited at Ϫ1.2 V ͑SCE͒. Scan rate was 0.5 mV/s. The substrate
used was a Pt sheet.
deposition at ca. Ϫ0.1 V ͑SCE͒ and the current started rising again
with hydrogen evolution at ca. Ϫ0.7 V ͑SCE͒. At pH 7, the deposi-
tion of copper began at a lower potential, Ϫ0.2 V ͑SCE͒, and the
current density reached a limiting value, as a result of the low con-
centration of copper ions in the electrolyte. Subsequently, hydrogen
evolution starts as indicated by a rise in the current above Ϫ0.9 V
͑SCE͒ ͑Fig. 1b͒. From the polarization curves, it is obvious that (i)
the copper reduction potential changes to more negative values by
raising the pH, and (ii) the process is diffusion limited due to the
low concentration of copper ions in the solution.
When 0.25 M CoSO4 was added to the blank solution, at pH 2
the current remained very low until it rose sharply at ca. Ϫ0.5 V
͑SCE͒, which is close to the reduction potential of cobalt. In a so-
lution that contained both Co and Cu ions at pH 2, the characteris-
tics of the polarization curve were similar to those with only Cu
additions, except that the current densities were somewhat lower.
Blockage of the current was observed again at pH 2 after the initia-
tion of copper deposition ͑ca. Ϫ0.2 V SCE͒ until a final current rise
that began at Ϫ0.8 V ͑SCE͒. The final current at the high cathodic
polarization would be compatible with both hydrogen evolution and
cobalt deposition.
this solution. At the potential range of copper reduction there is a
deflection in the slope of the current density vs. the cathodic poten-
tial. This deflection is attributed to side reactions, such as the reduc-
tion of an air-formed copper oxide, a transition of copper divalent to
copper monovalent, or reduction of dissolved oxygen in the
solution.
Anodic polarization curves were performed on ϳ800 nm Co,
ϳ100 nm Cu, and ϳ1 ␮m Cu-Co films, which were electrodepos-
ited on a Pt substrate from a solution that contained 0.2 M H3BO3 ,
0.13 M citric acid (C6H8O7), and the ions of either copper or cobalt
at pH 6.5 ͑Fig. 3͒. Cobalt dissolution started at ca. Ϫ0.5 V ͑SCE͒
and copper dissolution at ϳ0 V ͑SCE͒. When the anodic polariza-
tion measurements of the Cu-Co film were performed at a scan rate
higher than 1 mV/s, only one broad dissolution transition was ob-
served. However, three dissolution transitions were observed at a
scan rate of 0.5 mV/s. The first two transitions represent cobalt
dissolution, and the third transition is attributed to copper dissolu-
13
´
tion. Gomez et al. have also examined the stripping behavior of
Cu-Co alloys. Only one dissolution transition was observed which
was attributed to the dissolution of a Cu-Co solid solution. The
´
stripping analysis of Gomez et al. was performed at a higher scan
The solution chosen for Cu-Co electrodeposition was 0.25 M
CoSO4 , 0.125 M CuSO4 , 0.2 M H3BO3 , and 0.13 M citric acid
(C6H8O7).
rate than in our experiments, 10 mV/s vs. 0.5 mV/s. The appearance
of only one transition does not necessarily mean the formation of a
Cu-Co solid solution, and may be the result of convolution of two
transitions.
From the anodic polarization curves one can conclude that both
Co and Cu are deposited. The alloy concentration was estimated
according to the integrated Cu and Co dissolution areas. The alloy
concentration was 89.5% atom % Cu and 10.5% atom % Co. The
cathodic current efficiency of the Cu-Co alloy deposition process
was calculated by comparing the deposit weight and the calculated
charge according to the Faraday law. The cathodic current efficiency
was 89%.
While the polarization curve indicated that a high pH is favor-
able, the film quality was also examined as a function of the solution
pH. The pH of the solution was varied between 2 and 6.5. In an acid
solution ͑pH 2͒ a dark, spongy, and powdery film was obtained. As
the pH was raised, a sound film was formed, which was reddish in
color and more reflective. At a pH of 6.5, an optimum film was
obtained. Consequently, the pH of the solution was set between 6
and 6.5.
From the polarization curve on a copper substrate in the solution
that was chosen for the alloy deposition ͑Fig. 2͒, it seems that cop-
per deposition starts at ca. Ϫ0.1 V ͑SCE͒ and cobalt deposition
starts at ca. Ϫ0.4 V ͑SCE͒. These values correspond to a scan rate of
5 mV/s. The difference between the deposition potentials is small,
meaning that the electrodeposition of a Cu-Co alloy is possible from
STM.—The PVD copper on wafer substrate was first scanned in
the deposition solution without polarization under an open circuit
potential ͑OCP͒ of 0 V ͑SCE͒, then in situ electrodeposition of only
copper at Ϫ0.5 V ͑SCE͒ was performed ͑Fig. 4͒. The copper layer

Journal of The Electrochemical Society, 150 ͑1͒ C28-C35 ͑2003͒
C31
Figure 4. 3-D in situ STM image of a Cu layer deposited in a 0.25 M
CoSO4 and 0.125 M CuSO4 electrolyte at Ϫ0.5 V ͑SCE͒, scanned at
5 ϫ 5 ␮m. The copper layer exhibited a polycrystalline morphology, and
coalescence of copper grains in clusters of 200-750 nm.
exhibited polycrystalline morphology. At the beginning of the depo-
sition process, a fine microstructure was formed, but with increasing
deposition time a rough film was formed by coalescence of copper
grains in clusters of less than 100 nm and 200-750 nm ͑Fig. 4͒.
When the potential was shifted to Ϫ0.92 V ͑SCE͒ ͑Fig. 5͒, in which
both copper and cobalt are deposited with a higher content of copper
in the film, a finer microstructure was formed, although there were
still clusters of small grains of more than 1 ␮m in size. This is
probably due to the rough underlayer of copper. Only when the
potential was shifted to Ϫ1.05 V ͑SCE͒ ͑Fig. 6a͒ did the large cop-
per clusters disappear and the film was found to contain small
grains, less than 50 nm in size. As expected, a finer microstructure
was obtained by lowering the deposition potential. In addition, it
seems that faceting occurred in the films, which were deposited at
Ϫ0.92 and Ϫ1.05 V ͑SCE͒.
Figure 6. In situ STM images of a Cu-Co alloy layer deposited from a 0.25
M CoSO4 and 0.125 M CuSO4 electrolyte at Ϫ1.05 V ͑SCE͒. Three-
dimensional scans at ͑a͒ 5 ϫ 5 ␮m and ͑b͒ 1 ϫ 1 ␮m. A finer microstruc-
ture was obtained by lowering the deposition potential; the film contains
small grains, less than 50 nm in size. Faceting still exists in the Cu-Co film.
EDS.—The composition ratio in the Cu-Co film was checked at
several potentials from Ϫ0.8 to Ϫ1.2 V ͑SCE͒, and a ϳ1 ␮m
͓Cu92.5-Co7.5͔ ͑atom %͒ film was obtained at Ϫ1.2 V ͑SCE͒. The
concentration ratio of cobalt decreased with increasing film thick-
ness, low cobalt content ͑7.5 atom %͒ was obtained in thicker films
above ϳ0.5 ␮m compared to ϳ20 atom % at thin films below ϳ300
nm. Films of ͓Cu92.5-Co7.5͔ ͑atom %͒ with a thickness of ϳ1 ␮m
were prepared according to the experimental description.
AES.—Auger surface composition profile analysis was per-
formed on the as-formed Cu-Co film. Almost no cobalt was detected
at the surface of the Cu-Co film. AES depth profiling was conducted
to study the concentration gradients of the elements in the as-formed
film. Carbon and oxygen signals rapidly decreased to a low and
constant value after a few seconds of sputtering. The oxygen signal
was below the quantificial limit of our AES system (ϳ0.5%), and
one cannot rule out the incorporation of cobalt hydroxide in the
Cu-Co films from the AES measurements. The Auger peaks were
computer differentiated and the peak-to- peak height values for Cu
and Co are presented in Fig. 7. The AES data are used to show a
trend rather than the precise concentration of Cu and Co since no
suitable calibration was performed for the normalization of the Co
peak-to-peak.
Figure 5. 3-D in situ STM images of a Cu-Co alloy layer deposited from a
0.25 M CoSO4 and 0.125 M CuSO4 electrolyte at Ϫ0.92 V ͑SCE͒, scanned
at 5 ϫ 5 ␮m. A finer microstructure of Cu-Co alloy was formed with clus-
ters of small grains of more than 1 ␮m in size. Appearance of faceting
occurred in the film when cobalt is incorporated in the Cu-Co alloy.

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Journal of The Electrochemical Society, 150 ͑1͒ C28-C35 ͑2003͒
Figure 9. HRSEM micrographs ͑secondary electrons͒ of 1500 nm electrode-
posited Cu-Co alloy ͑top view͒ recorded at 4.0 kV annealed at 450°C for 90
min.
Figure 7. AES depth profile of Cu-Co film as formed. Co and Cu present
opposite trends, the relative cobalt concentration was negligible at the sur-
face and went through a peak concentration before gradually decreasing
close to the copper substrate, whereas the deposition rate of copper decreased
slightly during the middle stage of the electrodeposition process.
and in the particles was the same. Some of the thick specimens (Ͼ1
␮m͒ were cracked and spalled off after a thermal treatment, prob-
ably due to the evolution of hydrogen from the film since the elec-
trodeposition of the Cu-Co films are accompanied by the reduction
of hydrogen, which was partially trapped in the deposit.
After thermal treatments the number of the particles on top of the
surfaced increased compared to the as-deposited state. The size of
these particles was 2-5 ␮m ͑Fig. 9 and 10͒. The cobalt content in
these particles was reduced (Ͻ0.5 atom %) compared to its content
in the deposit ͑8-10 atom %͒. In addition, as a consequence of the
annealing procedures, holes were formed ͑Fig. 9 and 10͒ indicating
diffusion of cobalt away from the Cu-Co alloy, according to the
EDS results.
The concentration profile of Cu and Co was not uniform along
the film depth. Co and Cu presented opposite trends, where the
relative cobalt concentration was negligible at the surface and went
through a peak concentration before gradually decreasing close to
the copper substrate, whereas the deposition rate of copper de-
creased slightly during the middle stage of the electrodeposition
process.
HRSEM.—HRSEM micrographs of Cu-Co films before and after
thermal treatments are presented in Fig. 8-10. The film was rather
dense at thicknesses up to ϳ300 nm. As the deposition time in-
creased ͑meaning a thicker deposit͒, the roughness of the surface
increased. Spherical particles of 0.5-1 ␮m in diam were found on
the surface ͑Fig. 8͒. The content ratio of copper/cobalt in the film
XRD.—XRD analysis was conducted on samples before and af-
ter thermal treatment ͑Fig. 11 and 12͒. No clear evolution of a cobalt
phase after thermal treatment was detected. The XRD spectrum of
the 75 nm PVD copper substrate is an fcc structure with a preferred
orientation of ͕111͖ planes ͑inset in Fig. 11͒. Another weak reflection
at 2␪ ϭ 41.6° appeared. This reflection corresponds to an hcp co-
¯
balt phase of (1010) planes.
Figure 8. HRSEM micrographs ͑secondary electrons͒ of 1350 nm electrode-
posited Cu-Co alloy ͑top view͒ recorded at 4.0 kV. The roughness of the film
increased with its thickness, and spherical particles of 0.5-1.0 ␮m in diam
were found on the surface.
Figure 10. HRSEM micrograph ͑secondary electrons͒ of 1500 nm elec-
trodeposited Cu-Co alloy ͑top view͒ recorded at 4.0 kV annealed at 450°C
for 10 h. Holes were formed, indicating diffusion of cobalt away from the
Cu-Co alloy according to the EDS results.

Journal of The Electrochemical Society, 150 ͑1͒ C28-C35 ͑2003͒
Discussion
C33
The results presented in this paper indicate that a high pH is
favorable for the electrodeposition of Cu-Co alloys. In a solution
that contained both Co and Cu ions at pH 2, the limited cathodic
process between Ϫ0.1 and Ϫ0.8 V ͑SCE͒ ͑Fig. 1a͒ may indicate that
the cobalt content in the alloy will be limited. At high overpoten-
tials, more negative than Ϫ0.8 V ͑SCE͒, both cobalt and copper may
be deposited together with hydrogen evolution, but the likelihood of
the formation of Cu-Co solid solution will decrease due to high
polarization values.14 However, at a higher pH, no blockage of the
current densities occurred ͑Fig. 1b, 2͒ and the formation of a Cu-Co
solid solution is possible with a higher content of cobalt in the
deposit.
The anodic polarization curve indicates that both Cu and Co are
deposited, although three dissolution transitions are observed. The
first and third transitions are related to cobalt dissolution at ca.
Ϫ0.5 V ͑SCE͒ and copper dissolution at ca. Ϫ0 V ͑SCE͒, respec-
tively. The second dissolution transition may be related to dealloy-
ing of cobalt from a metastable state of Cu-Co phase which accord-
ing to the XRD results exists in the Cu-Co films, or due to the
formation of an overlayer of copper as found by AES studies. Shima
et al.19 have studied the anodic dissolution of an ultrathin cobalt
film buried beneath a copper overlayer. It was found that the cobalt
underlayer can dissolve even though it is covered by a thin copper
layer, and that the blocking nature of Cu overlayer was reflected in
the displacement of the cobalt stripping potential to a more positive
potential. The dissolution of an artificial Co/Cu structure is de-
scribed by two stages: breakdown of the copper overlayer and sub-
sequent dissolution of the buried cobalt layer. From the anodic po-
larization curve, it seems that in either of the two possibilities, the
stripping potential of cobalt should move to more positive values.
The two dissolution transitions of cobalt may indicate a nonhomo-
geneous Cu-Co alloy as was observed in the AES and XRD results.
According to the AES results, the bulk part of the film is richer in
Co while the surface and bottom of the film are Co-poor. It was
reported20,21 that in a vacuum environment copper segregates to the
surface of Cu/Co films, in order to reduce the free surface energy.
Thus, it might be that some degree of stability is established when
the present films are covered by copper.
Figure 11. XRD of the Cu-Co films before and after thermal treatment.
Inset: XRD of a Si substrate which was deposited by 75 nm PVD copper on
top of 30 nm TaN. The Cu-Co film as deposited is composed of two phases,
an hcp Co phase and a solid solution of fcc rich copper Cu-Co phase with a
preferred orientation of ͕111͖ planes, similar to the copper substrate.
From XRD of the Cu-Co alloy after different sequences of ther-
mal treatments there is no indication of cobalt precipitation or evo-
lution of a new cobalt phase ͑Fig. 11-12͒. Only one peak is observed
at 2␪ ϭ 43.7°, which correlates mainly to ͑111͒ face centered cubic
͑fcc͒ copper. However, there is a slight shift toward high angles for
the ͑111͒ reflections with increasing annealing time ͑Fig. 12͒ and
consequently a slight decrease in the d-spacing.
The diffusion coefficient of Co22 and Cu23 ions in aqueous solu-
tions are 5.8 ϫ 10Ϫ6 and 3 ϫ 10Ϫ6 cm2/s, respectively. The higher
rate of copper deposition at the first stage of the film growth may be
due to the higher polarization of copper compared to cobalt, because
copper is much nobler than cobalt. A possible reason for the appear-
ance of a peak in the rate of cobalt deposition and a small decay in
the copper deposition rate may be that after the electrodeposition
process is started, copper deposition process becomes more diffu-
sion limited than that of cobalt due to the higher concentration of
cobalt ions in the electrolyte and a higher diffusion coefficient.
Another possible reason for the opposite trends in the deposition
rate of each element is that at the beginning copper grains are
formed on a copper substrate, while as the deposition process
progresses the substrate changes from pure copper to the Cu-Co
mixture, either as a metastable solid solution and/or as nanoclusters
of cobalt inside a matrix of copper. This can lead to different growth
modes. Bulk Co and Cu are considered immiscible below
ϳ900 K.24 According to Coyle et al.,24 the Cu-Co bond energy is
probably greater than 0.190 eV for the nearest neighbors and it is
0.271 and 0.190 eV for Co-Co and Cu-Cu, respectively. The depo-
sition rate of cobalt on cobalt will be faster than cobalt on Cu and
vice versa for the deposition rate of copper. The different modes of
cobalt and copper deposition will be the focus of a future study. This
may be also the reason for the appearance of faceting that was
observed by STM. Faceting appeared in films that were deposited at
Ϫ0.92 and Ϫ1.05 V ͑SCE͒ only after cobalt was incorporated into
the deposit ͑Fig. 5, 6b͒. A recent in situ STM study of Cu growth on
Ni͑001͒ reports a novel model for internal faceting as a consequence
Figure 12. XRD of Cu-Co films before and after sequences of thermal treat-
ments. There is no indication of evolution of a new cobalt phase, only a
slight shift toward high angles for the ͑111͒ reflections with increasing an-
nealing time.

C34
Journal of The Electrochemical Society, 150 ͑1͒ C28-C35 ͑2003͒
of strain relief at the interface, where the lattice misfit in the Cu and
Ni is 2.6%.25 Such a strain relief mechanism, in principle, can occur
in the Cu-Co system as well. In the fcc Cu-Co system, the lattice
misfit is less than 2%.26 The growth of the Cu-Co alloy may be
accompanied by strain relief, resulting in the formation of faceting.
It seems from HRSEM results that the existence of a copper
overlayer increases the roughness of the Cu-Co films. There may be
a number of reasons for the formation of a rough film in the rela-
tively thick layer ͑larger than 1 ␮m͒. As the deposition time in-
creases, the content of Co2ϩ and Cu2ϩ ions near the cathode surface
is diminished, and the overpotential reduces, there are fewer nucle-
ation events and more grain growth with hydrogen voids in the
films. Since the concentration of both metal ions is high and since
copper is nobler than cobalt, its deposition rate does not change
considerably, while that of cobalt reduces with increasing time. As
seen in the AES studies the deposition rate of copper is higher than
that of cobalt toward the end of the deposition process, and an
overlayer of copper is formed on top of the alloy. This also may be
the reason for the low cobalt content ͑7.5 atom %͒ in thicker films
above ϳ0.5 ␮m compared to 20 atom % at thin films ͑less than
ϳ300 nm). In addition, according to the polarization curve of the
solution containing only Co2ϩ ͑Fig. 1͒, the cobalt deposition rate
remains constant until hydrogen evolution begins. Nakahara and
Mahajan27 claim that during cobalt deposition with hydrogen evo-
lution, cobalt hydride or hydroxide are formed in the films. This may
be the reason for a decrease in the cobalt deposition rate. Nakahara
and Mahajan have examined the content of O2 and H2 in electrode-
posited cobalt by using a vacuum fusion technique and mass spec-
trometry. They found 230 ppm H2 and 890 ppm O2 in the films
deposited at a high pH (ϳ5.7). According to Nakahara and Ma-
hajan, the solubility of H2 in cobalt is 0.041 ppm at room tempera-
ture ͑20°C͒ and its diffusivity is 1.71 ϫ 10Ϫ11 cm2/s for fcc cobalt.
Hydrogen has a high diffusivity in cobalt at room temperature and it
could escape from the film very rapidly. Another option for the
decrease in the deposition rate of cobalt is that when cobalt hydride/
hydroxide forms, some of the reduction current may be consumed
by the reduction of these species, and hence the deposition rate of
cobalt is reduced. These may be the reasons why phases of cobalt
precipitates were not detected in the XRD reflections.
In addition, the OCP becomes more anodic with deposition time.
This behavior is similar to the results of Bradley et al.9 At thick-
ness up to ϳ300 nm the OCP remains constant ca. Ϫ0 mV ͑SCE͒
while at a thickness higher than 500 nm, the OCP decreases to ca.
Ϫ500 mV ͑SCE͒. At least 1 min is required for the OCP to reach
Ϫ100 mV ͑SCE͒. According to the polarization curve ͑Fig. 2͒ this
OCP is enough for copper deposition but not for cobalt deposition.
This means that at the end of the electrodeposition process copper
grains may form by chemical exchange of Cu ions with the more
active Co atoms in the alloy. Since the OCP is close to the equilib-
rium potential of copper, coarse grains of copper are deposited, as
seen in the STM measurements. This also may be the reason for the
particles on the surface ͑Fig. 7͒. The reason for detecting cobalt in
these particles probably stems from the large penetration depth in
EDS ͑penetration radius of ϳ1 ␮m for Cu K␣ at 20 kV͒, due to
columnar growth of copper which causes an incomplete coverage of
the alloy surface9 and since the volume fraction of the copper grains
is small compared to that in the alloy.
process. The fcc structure remains even in thick films ͑1 ␮m͒, in
which the role of the substrate ͑75 nm physical vapor deposited
PVD copper͒ is negligible. The main reason for the formation of an
alloy with an fcc structure is probably the high copper content in the
film. A shift in the d-spacing of ͕111͖ copper planes from 2.088 to
2.07 Å was detected in the as-deposited films. Since the atomic
volume of Co atom is smaller than that of Cu, the shift may indicate
the substitution of Cu by Co atoms and the formation of a meta-
stable phase. It may also indicate the merging of two reflections;
͑111͒ fcc Cu with ͑111͒ fcc Co.
Peter et al.30 reported that the structure of a dc-plated Co rich
Co-Cu alloy consisted of both hcp and fcc Co-Cu crystallites and it
contained more fcc components at the solution side than at the sub-
strate side ͑titanium sheet͒. The fraction of the fcc crystallites in-
creased gradually during deposition from about 50 to 67%. Since
both titanium and cobalt have hcp structures, the expected structure
of the Co-rich alloy is hcp. According to Peter et al.,30 a possible
reason for obtaining an fcc structure is the incorporation of a small
amount of Cu into the deposit. In a Cu-rich film, the reverse situa-
tion may occur, where incorporation of small amount of Co into the
¯
deposit may form an hcp structure. If the (1010) reflection was
correlated to a Cu-Co alloy, the d-spacing would decrease. How-
¯
ever, the d-spacing of the (1010) reflection is slightly higher than
the d-spacing in a pure cobalt film. Therefore one can deduce that
¯
the (1010) reflection corresponds to a pure cobalt phase and the
change in the d-spacing is related to a strained Co phase.
It is thus concluded that the as deposited Cu-Co film is composed
of two phases, a solid solution of fcc Cu-Co phase and an hcp Co
phase. The mechanism by which the Co hcp phase forms and grows
is not yet known.
The formation of two phases, pure hcp phase and a rich Cu phase
can result in a GMR effect in the as-deposited film without heat-
treatment. The magnetic properties of the Cu-Co film will be re-
ported in a following paper.
The HRSEM and EDS results of the Cu-Co films indicate the
precipitation of cobalt grains during the thermal treatments. How-
ever, from XRD of the Cu-Co alloy after different sequences of
thermal treatments there is no indication of cobalt precipitation or
evolution of a new cobalt phase ͑Fig. 11 and 12͒. The reflections
from fcc ͕111͖ planes of copper and cobalt are at 2␪ ϭ 43.297° and
2␪ ϭ 44.216°, respectively. The difference is almost 1°, which
should be large enough to distinguish between the two peaks. Only
one peak is observed at 2␪ ϭ 43.7°, which correlates mainly to
͑111͒ fcc copper. However, there is a slight shift toward high angles
for the ͑111͒ reflections with increasing annealing time ͑Fig. 12͒ and
consequently a slight decrease in the d-spacing. The shift may indi-
cate the progressive decomposition of Co atoms from the Cu rich
matrix. The absence of an fcc cobalt reflection is probably related to
the relatively low content of cobalt ͑7.5 atom %͒ and to the small
grain size of cobalt ͑nanoscale͒. The annealing temperature of the
Co-Cu deposit ͑450°C͒ is higher than the hcp-fcc transition tempera-
ture ͑422°C͒. Thus it is expected that the Co particles precipitate in
the fcc form at 450°C and that the hcp Co phase, which was formed
during the deposition process, transformed to the fcc phase. There is
also a minor possibility that the fcc Co phase transformed to hcp Co
while cooling down to room temperature after annealing, although
the cooling process was quite rapid ͑15-20 min͒. The mechanism
that inhibits the phase transformation in the Cu-Co film is not yet
clear.
The deposition was conducted under a high polarization potential
of Ϫ1.2 V ͑SCE͒ ͑current density of ϳ8 mA/cm2). Several
authors28,29 have reported that hcp cobalt is formed by electrodepo-
Conclusions
Ͼ
sition from solution at pH
2.9. Thus it was expected that a two-
phase structure would be formed,14,29 i.e., hexagonal close packed
͑hcp͒ cobalt and fcc copper. However, the XRD results indicate the
formation of a supersaturated solution of Cu-Co solid solution and
an hcp cobalt phase. The Cu-Co alloy, which is supposed to be a
metastable phase at room temperature, is a copper rich Cu-Co fcc
phase with a preferred orientation of ͕111͖ planes, similar to the
copper substrate ͑Fig. 11͒. This indicates a topotaxial crystallization
Electrodeposition from a neutral solution with a high concentra-
tion of Co and Cu ions results in a topotaxial crystallization process
and a deposit ͓Cu92.5-Co7.5͔ ͑atom %͒ of two phases; a solid solu-
tion of fcc Cu-Co with preferred orientation of ͕111͖ planes and an
hcp Co phase. Electrochemical analysis and AES showed that cobalt
deposition is limited toward the film surface, possibly due to a
higher deposition rate of copper than cobalt toward the end of the

Journal of The Electrochemical Society, 150 ͑1͒ C28-C35 ͑2003͒
C35
7. O. F. Bakkaloglu, I. H. Karahan, H. Efeoglu, M. Yildirim, U. Cevik, and Y. K.
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8. H. Zaman, A. Yamada, H. Fukuda, and Y. Ueda, J. Electrochem. Soc., 145, 565
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9. P. E. Bradley and D. Landlot, Electrochim. Acta, 45, 1077 ͑1999͒.
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5905 ͑2000͒.
deposition process and by the formation of copper grains after the
electrodeposition process by chemical exchange between the copper
ions in the solution and with the cobalt in the Cu-Co film. This is the
main reason for the lower cobalt content in the film, and for the
rough Cu-Co films revealed by in situ STM and HRSEM studies.
In situ STM studies revealed that when cobalt is incorporated
into the deposit, faceting develops which is believed to be due to
strain relief in the growth of the metastable Cu-Co alloy. AES depth
profiling analysis revealed that the Cu-Co film is not homogeneous;
the bulk of the film is richer in Co while the surface and the bottom
of the film are Co-poor.
SEM micrographs and EDS results indicate the precipitation of
cobalt grains resulting from thermal treatments, according to the
phase diagram of Cu-Co. From XRD studies there is no definite
indication of evolution of an fcc cobalt phase after thermal treat-
ments, probably due to the low content of Co ͑atom % 7.5͒ and due
to its small grain size.
12. G. R. Pattanaik, S. C. Kashyap, and D. K. Pandya, J. Magn. Magn. Mater., 219,
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Both of the phases ͑hcp and fcc͒ remain stable after annealing
above the hcp-fcc transition temperature.
20. S.-K. Kim, J.-S. Kim, J. Y. Han, J. M. Seo, C. K. Lee, and S. C. Hong, Surf. Sci.,
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Acknowledgments
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The authors wish to thank theTechnion Research Funds for sup-
porting this project, and Applied Materials Inc. for supplying the
substrate wafers.
¨
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Technion-Israel Institute of Technology assisted in meeting the publica-
tion costs of this article.
25. B. Muller, B. Fischer, L. Nedelmann, A. Fricke, and K. Kern, Phys. Rev. Lett., 76,
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