Structural and magnetic characterization of electrodeposited, high moment FeCoNi films

Article abstract of DOI:10.1149/1.1545462

We investigated the electrochemical deposition process of Fe-Co-Ni alloys from sulfate electrolytes containing no organic additives and the structural, magnetic, and electrical properties of the resulting films. The effect of pH and Co and Fe concentrations in the electrolyte on alloy composition is examined. The pH of the solutions is kept at 2.8 to minimize the incorporation of nonmagnetic matter. With increasing Ni content, a transition of the film structure from body-centered cubic (bcc) to face-centered cubic (fcc) is observed. X-ray diffraction and transmission electron microscopy results show that the smallest grain size of about 10 nm is obtained when the bcc and fcc phases coexist. Films with biphasic structure exhibit soft magnetic properties (coercivity about 1 Oe), high magnetic moment (about 2.1 T), narrow dispersion of the easy axis (less than 3°), and good high frequency response, with a resonance frequency above 1.8 GHz. Stress anisotropy has an important role in determining soft magnetic properties.

Full text of DOI:10.1149/1.1545462

Journal of The Electrochemical Society, 150 ͑3͒ C159-C168 ͑2003͒
C159
0013-4651/2003/150͑3͒/C159/10/$7.00 © The Electrochemical Society, Inc.
Structural and Magnetic Characterization of Electrodeposited,
High Moment FeCoNi Films
Xiaomin Liu,a,b, Giovanni Zangari,
and Mohammed Shamsuzzohad
*
a,b,c,**,z
b
aCenter for Materials for Information Technology, Department of Metallurgical and Materials Engineering, and
dSchool of Mines and Energy Development, University of Alabama, Tuscaloosa, Alabama 35487-0209, USA
We investigated the electrochemical deposition process of Fe-Co-Ni alloys from sulfate electrolytes containing no organic addi-
tives and the structural, magnetic, and electrical properties of the resulting films. The effect of pH and Co and Fe concentrations
in the electrolyte on alloy composition is examined. The pH of the solutions is kept at 2.8 to minimize the incorporation of
nonmagnetic matter. With increasing Ni content, a transition of the film structure from body-centered cubic ͑bcc͒ to face-centered
cubic ͑fcc͒ is observed. X-ray diffraction and transmission electron microscopy results show that the smallest grain size of about
10 nm is obtained when the bcc and fcc phases coexist. Films with biphasic structure exhibit soft magnetic properties ͑coercivity
about 1 Oe͒, high magnetic moment ͑about 2.1 T͒, narrow dispersion of the easy axis ͑less than 3°͒, and good high frequency
response, with a resonance frequency above 1.8 GHz. Stress anisotropy has an important role in determining soft magnetic
properties.
© 2003 The Electrochemical Society. ͓DOI: 10.1149/1.1545462͔ All rights reserved.
Manuscript submitted June 6, 2002; revised manuscript received September 11, 2002. Available electronically February 5, 2003.
anisotropy on the stress state. However, the presence of additives
adsorbed on the electrode surface induces incorporation of sulfur or
carbon-containing impurities into the growing film, which influences
its microstructure and magnetic properties18 and degrades its corro-
sion resistance.17 As a consequence, various groups have reported
different properties for Fe-Co-Ni alloy films ͑for example, the com-
position of alloys with zero magnetostriction or high moment de-
pends on the additive used͒.8,12,18 We present here a study of the
electrodeposition process and magnetic properties of Fe-Co-Ni films
grown from sulfate solutions and containing the lowest possible
amount of foreign matter. By this procedure, we hope to achieve a
better understanding of the dependence of the properties of elec-
trodeposited ternary alloys on composition and microstructure, thus
determining the conditions at which optimum magnetic properties
are obtained. To achieve high purity, the films are grown from elec-
trolytes not containing any type of additives, and deposition condi-
tions, in particular the pH, are optimized to minimize the incorpo-
ration of nonmagnetic matter ͑metal hydrides or hydroxides͒.
Soft magnetic films with high saturation induction Bs , low co-
ercivity Hc , and high permeability ␮ are essential components of
write heads in high density recording systems, as well as in inte-
grated thin film inductors and transformers. In particular, Fe20Ni80
͑Permalloy͒ and later Fe55Ni45 alloys have been almost exclusively
used as pole materials for thin film heads.1 Currently, however, the
Bs of these materials, 1 and 1.6 T ͑10 and 16 kG͒, respectively, is
insufficient to provide the writing field to saturate the magnetic me-
Ͼ
dia with high coercivity H
320 kA/m ͑ϳ4000 Oe͔͒ required to
͓
c
store information at the bit density of 100 Gb/in.2 recording and
higher.2,3 Fe-Co alloys on the other hand are ideal candidates for
these applications, as in the bulk they exhibit Bs values up to 24 kG,
2.4 times the value of Permalloy. In addition, these materials poten-
tially provide higher roll-off frequencies of the permeability,4 nec-
essary to enable device operation in the gigahertz range.
Electrodeposition offers several advantages in the fabrication of
poles for write heads or magnetic components in integrated devices
or microsystems. For example, soft magnetic properties ͑usually bet-
ter than the corresponding alloys made by sputtering or evaporation͒
and low internal stress can be obtained at high rates of deposition.1
More importantly, the purely additive process of plating through a
patterned mask allows easier and better definition and control of
small features. Subtractive processes on the contrary must follow
blanket deposition by gas-phase methods, and these are usually dif-
ficult to implement due to the lack of practical methods for selective
etching.5 Most recently, focused ion beam etching ͑FIBE͒ has shown
improvements in this respect, allowing the reliable definition of
poles with submicrometer width, at the expense, however, of using a
complicated and expensive instrument.
Experimental
Composition of the electroplating bath is reported in Table I. All
the reagents were analytical grade, and the solutions were made with
triply distilled water. Simple sulfate salts of the metals were used,
and boric acid was added as a buffer. The choice of an all-sulfate
bath allows minimization of the internal stresses.19 pH was varied in
the range 2 to 3.5 to minimize the incorporation of impurities, as
discussed in the Results section.
Electroplating was performed in a prismatic Pyrex cell ͑300 mL
volume͒ using facing vertical electrodes flush to the side walls, to
achieve a uniform current distribution. The anode was a Co sheet
99.9% pure. The electrolyte was heavily stirred using a magnetic
stirrer. A magnetic field of 600 Oe was applied along the plane of the
film during film growth to achieve a uniaxial in-plane magnetic
anisotropy. FeCoNi alloy films were electrodeposited under current
control. The alloy films were deposited onto either Cu sheets which
were first electropolished and then electroless plated with 2 ␮m of
amorphous, nonmagnetic NiP or 500 ␮m thick n-type Si͑100͒ sub-
strates coated with sputtered Ti ͑10 nm͒ and Cu ͑200 nm͒ layers,
with Cu being the substrate for plating. Electroless plated amor-
phous NiP and sputtered Cu were chosen to provide flat, inert sub-
strates having limited influence on the nucleation and growth of the
magnetic film. Magnetic properties were independent of the
substrate.20
High moment Fe-Co bulk alloys exhibit a high magnetostriction,
6
up to about 10Ϫ4
.
To avoid stress-induced degradation of the soft
magnetic properties, the first Fe-Co system to be investigated was
7
the zero-magnetostrictive composition Co90Fe10
. More recently,
Co-Fe-Ni alloys have been developed with a higher value of Bs, and
low coercivities.8-15 These alloys also exhibit improved corrosion
resistance with respect to Fe-Co alloys.16,17
Current research and development on electrodeposited transition
metal alloys for write-head applications is conducted using electro-
lytes containing additives to decrease the internal stresses of the
films, favor adhesion, and minimize the dependence of the magnetic
Potentiodynamic measurements were performed in the plating
cell using an EG&G PAR potentiostat/galvanostat ͑A273͒. The elec-
trolytes used for Fe, Co, Ni single-metal deposition and FeCoNi
deposition ͑Table II͒ had a constant ionic strength, adjusted with the
addition of Na2SO4 , and were not stirred during the experiments.
* Electrochemical Society Student Member.
** Electrochemical Society Active Member.
c Present address: Department of Materials Science and Engineering and CESE,
University of Virginia, Charlottesville, VA 22904-4745, USA.
z E-mail: gz3e@virginia.ed
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Journal of The Electrochemical Society, 150 ͑3͒ C159-C168 ͑2003͒
Table I. Plating conditions and electrolyte composition for the
electrodeposition of Fe-Co-Ni films.
Plating conditions
Temperature
pH
Current density
Counter electrode
25°C
2-3.5
5-50 mA/cm2
Pure Co ͑99.9%͒
Plating solution
0.01-0.05 mol/L
0.2 mol/L
0.01-0.05 mol/L
0.4
CoSO4•7H2O
NiSO4•6H2O
FeSO4•7H2O
H3BO3
For these measurements Si͑100͒ substrates as above sputtered with
Ta ͑10 nm͒ and Au ͑10 nm͒ layers were used as the working elec-
trodes, a Pt sheet as the counter electrode, and saturated calomel as
the reference electrode. Cathodic polarization curves were measured
by ramping the potential at a sweep rate of 2 or 10 mV/s.
Figure 1. Polarization curves for Fe, Co, Ni single metals and FeCoNi alloy
from the solutions shown in Table II. Sweep rate, 2 mV/s; no N2 purging
before experiment.
Film composition was determined by energy dispersive X-ray
spectrometry ͓EDS, Philips XL30 scanning electron microscope
͑SEM͒ attachment͔. In some cases, compositional profiles of the
films and the amount of impurities were determined using quantita-
tive X-ray photoelectron spectroscopy ͑XPS, Kratos AXIS 165͒. The
two methods were in good agreement. Crystalline structure, grain
size, and preferential orientation of the films were evaluated by
X-ray diffraction ͑XRD, Philips X’pert͒ using Cu K␣ radiation, and
transmission electron microscopy ͑TEM, Hitachi H-8000, 200 kV͒.
Static magnetic properties were determined by vibrating sample
magnetometry ͑VSM, DMS880͒, ferromagnetic resonance ͑FMR͒,
and low field susceptibility. In the FMR method, the FMR frequency
is measured as a function of the angle between the easy axis and the
applied field in the plane of the film, and Bs is calculated using the
method reported in Ref. 21 and 22. In the low-field susceptibility
method, the sample is magnetized perpendicular to the film plane
and the linear behavior of 4 ␲M vs. the applied field thus obtained is
extrapolated to the measured in-plane saturation magnetization. The
field at which the in-plane and perpendicular loops intersect is Bs .
Magnetostriction was measured with an in-house built tester which
employs the optical cantilever method under a rotating field of up to
60 Oe.23 In a few cases, this field was insufficient to saturate the
films and the reported magnetostriction values are the maximum
values measured. Permeability was evaluated with an in-house built
system24 up to a frequency of 3 GHz. Internal stresses in the films
were determined by comparison of the substrate curvature before
and after electroplating, using a Flexus F2320 instrument.25 Film
resistivity was measured using a four-point probe at room tempera-
ture.
single metals and the alloy of about 10 to 20 mV ͑not shown͒. In
addition, the small current due to reduction of dissolved oxygen is
decreased significantly between 0 and Ϫ500 mVSCE . Figure 2
shows cathodic potential sweep curves taken at a scan rate of 10
mV/s. In all cases, the cathodic currents increase rapidly with volt-
age below an onset potential, associated with nucleation and growth
of the film. The nucleation potentials En for single metals and the
Ͼ
Ͼ
Ͼ
En,FeCoNi En,Fe , which
alloy follow the order, En,Ni
En,Co
differs from the results obtained by Tabakovic et al.18 In Ref. 18
however, the electrolytes used contained organic additives, and dif-
ferences in the nucleation and growth behavior should be expected.
Saturating the solutions with N2 gas ͑Fig. 2b͒ does not vary En ,
indicating that the oxygen dissolved in solution has a limited influ-
ence on the metal reduction reactions, although it has a definite
influence on the equilibrium for metal hydroxide formation. A small
reduction peak at Ϫ0.68 VSCE due to hydrogen evolution was ob-
served for Fe and Co deposition. Such a peak was not observed for
Ni or alloy deposition. This effect may be explained by the recent
observation26 that the addition of metal ions to an acid chloride
solution inhibited hydrogen reduction, with a stronger inhibiting ef-
fect for Ni than for Fe, or may be due simply to the different con-
centration of the various metal ions in solution. Large deposition
currents were observed below Ϫ0.9 VSCE for Ni and FeCoNi depo-
sition. On the contrary, Co and Fe solutions exhibit a diffusion-
limiting current at these potentials due to the lower metal concen-
tration.
Alloy composition and impurities.—The operating pH of the
FeCoNi plating solution investigated here was varied between 2 and
3.5. Below pH 2, in fact, the current efficiency at room temperature
was unacceptably low, due to the increased hydrogen evolution,
while above pH 3.5 it was possible to observe formation of precipi-
tates with time, probably due to the hydrolysis of Fe3ϩ, which led to
irreversible degradation of the solution. Figure 3 shows film com-
position as a function of applied current density for the lowest,
highest, and optimum pH values investigated. At constant current
density, the Fe and Co content in the alloy decreased with increasing
pH, while the Ni content increased. This is in agreement with the
Results and Discussion
Potentiodynamic studies.—Figure 1 shows the polarization
curves at low cathodic overpotential for Fe, Co, Ni single metals and
the alloy obtained from the solutions shown in Table II. The zero
current potential E for the single metals follows the same trend as
Ͼ
Ͼ
E
the standard equilibrium potentials, ENi
ECo
Fe , and the
mixed potential for the FeCoNi alloy is between those of the single
metals. Deaerating the solutions by sparging N2 gas for 30 min
before the measurement results in a cathodic shift of E for both the
Table II. Composition of the solutions utilized for electrochemical studies.
Fe
Co
Ni
FeCoNi
0.025 mol/L FeSO4•7H2O ϩ 0.25 mol/L Na2SO4 ϩ 0.4 mol/L H3BO3 , pH 2.8
0.05 mol/L CoSO4•7H2O ϩ 0.225 mol/L Na2SO4 ϩ 0.4 mol/L H3BO3 , pH 2.8
0.2 mol/L NiSO4•6H2O ϩ 0.075 mol/L Na2SO4 ϩ 0.4 mol/L H3BO3 , pH 2.8
0.025 mol/L FeSO4•7H2O ϩ 0.05 mol/L CoSO4•7H2O ϩ 0.2 mol/L NiSO4•6H2O ϩ 0.4 mol/L H3BO3 ,
pH 2.8
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Journal of The Electrochemical Society, 150 ͑3͒ C159-C168 ͑2003͒
C161
Figure 2. Cathodic linear potential sweep curves for single metals and alloy.
Scan rate 10 mV/s; composition of the electrolytes is shown in Table II. ͑a͒
Without bubbling N2 , ͑b͒ with N2 .
patent by Castellani et al.,27 which reports a decrease in Fe content
with increasing pH in the range 1.5 to 3.5. At a given pH, the Fe and
Co content in the alloy decrease with increasing current density,
while Ni follows the opposite trend. The ternary alloy system clearly
follows an anomalous deposition behavior,28 with enhanced Fe and
Co reduction at low current densities due to the presence of Ni. At
high current densities, Fe and Co deposition are under diffusion
control due to the lower ion concentration in solution and conse-
quently the alloys are enriched in Ni. In Fig. 3b, pH 2.8, the cathodic
current efficiency is also reported, showing an increase with current
density from 80 to 94% in the range investigated. This is due to the
decreasing importance of hydrogen reduction at high current den-
sity.
Figure 3. Variation of alloy composition with current density for FeCoNi
films electrodeposited at ͑a͒ pH 2.0; ͑b͒ 2.8, and ͑c͒ 3.5. Fe2ϩ
The concentrations of metal ions in the plating solution are im-
portant variables in defining the composition of the plated film. This
is a consequence of the anomalous nature of the FeCoNi system,
whereby the preferential deposition of Fe and Co over Ni translates
small variations of the Fe2ϩ and Co2ϩ concentrations in solution
͓
͔
ϭ 0.025 M, Co2ϩ ϭ 0.05 M, and Ni2ϩ ϭ 0.2 M.
͓
͔
͓
͔
͓
͔
͓
͔
into larger variations of their content in the alloy film. Ternary dia-
increase of that metal in the alloy, yielding approximately parallel
lines on the ternary diagram.
grams ͑Fig. 4͒ show the dependence of alloy composition on Fe2ϩ
͓
͔
͑Fig. 4a͒, Co2ϩ ͑Fig. 4b͒, and current density. In each figure the
͓
͔
Electrodeposited Permalloy films are known to exhibit a compo-
sitional gradient along their thickness, with enrichment of Fe in the
first 100 nm deposited.29 This may cause undesirable inhomogeneity
of the magnetic properties, but can be avoided by a suitable choice
of the deposition conditions.30 Our group, using EDX on films with
different thicknesses, reported measuring a compositional gradient
concentration of two metals in solution is held constant while that of
the third one is varied. The current density increases along the arrow
direction, and it changes in the sequence 5, 10, 15, 20, and 50
mA/cm2, corresponding to each marker. It is observed that increas-
ing Fe2ϩ or Co2ϩ in the solution results in a corresponding
͓
͔
͓
͔
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Journal of The Electrochemical Society, 150 ͑3͒ C159-C168 ͑2003͒
Figure 4. Effect of ͑a͒ Fe2ϩ and ͑b͒ Co2ϩ on film composition. In ͑a͒
͓
͔
͓
͔
Co2ϩ ϭ 0.05 M and Ni2ϩ ϭ 0.2 M; while in ͑b͒ Fe2ϩ ϭ 0.05 M and
͓
͓
͔
͓
͔
͓
͔
Ni2ϩ ϭ 0.2 M.
͔
in FeCoNi alloys,10 while more recently Tabakovic et al.31 using
XPS depth profiling showed no such effect. EDX and XPS depth
profiling are complementary techniques; EDX, in fact, probes sev-
eral micrometers inside the alloy film, and for films about 1 ␮m
thick it measures a composition averaged over the entire film thick-
ness. XPS, on the contrary, yields information on the composition of
a surface layer only a few nanometers thick, and sputter etching is
necessary to probe film composition at different depths. Both tech-
niques were utilized in this work, and the results are shown in Fig. 5.
Figure 5a shows the composition of FeCoNi films, grown under the
same deposition conditions vs. film thickness as determined by
EDX. A very small compositional gradient was observed, which
however may be due to experimental errors. Figure 5b depicts the
composition profile obtained by sputter-etching a 70 nm thick film.
XPS data reveal a limited decrease of Fe content and a correspond-
ing increase in Ni near the film surface; on the other hand alloy
composition is nearly uniform near the substrate interface. We con-
clude that compositional gradients, if present, are of limited extent
and likely to have a limited influence on the films’ magnetic prop-
erties.
Figure 5. FeCoNi film composition as a function of film thickness. Film
composition determined by ͑a͒ EDX on different films; ͑b͒ XPS depth profile
of a 70 nm thick film. FeCoNi films were grown on Si/Ti/Cu substrates, with
Cu being the substrate for plating. ͑c͒ Magnification of the oxygen content
vs. film thickness shown in ͑b͒. Electrodeposition condition: Fe2ϩ
͓
͔
ϭ 0.025 M, Co2ϩ ϭ 0.05 M, and Ni2ϩ ϭ 0.2 M, pH 2.8. Current
͓
͔
͓
͔
density ϭ 10 mA/cm2.
In order to quantify the amount of impurities incorporated in
films grown at different pH, quantitative XPS was used to detect
sulfur ͑S͒, carbon ͑C͒, and oxygen ͑O͒. After the contaminated sur-
face film ͑about 3 nm͒ was eliminated by sputtering, the XPS signal
from S and C was below the detection limit of our instrument ͑Ͻ1
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Journal of The Electrochemical Society, 150 ͑3͒ C159-C168 ͑2003͒
atom%͒. Oxygen, however was observed, and its content increased
C163
with increasing pH. At pH 3.5, the O content was around 3 atom %.
For pH values below 2.8, the O content was at most 2 atom %.
Figure 5c is a magnification of the depth profile for oxygen shown in
Fig. 5b, and shows that O content increases slightly with thickness.
The small intensity of the O 1s peak does not allow determination of
the oxygen chemical state. At lower pH, the hydrogen reduction
reaction becomes more important and the probability for incorporat-
ing interstitial hydrogen into the deposit is higher. In order to avoid
this effect and thus further limit the incorporation of nonmagnetic
matter, the pH of the solution was kept at 2.8.
Crystalline structure.—X-ray diffraction spectra of FeCoNi films
deposited at various current densities are presented in Fig. 6a. A
transition of the crystalline structure from bcc to fcc was observed
with increasing current density ͑increasing Ni content͒, as expected
on the basis of the bulk phase diagram. Figure 6b shows the alloy
composition for which the bcc and fcc phases were seen to coexist,
superposed to the fcc/bcc equilibrium boundary for the bulk alloys.
Films with high Fe content, low Co content, and Ni content below
20% showed a predominantly bcc structure. Films with high Co and
Ni content were found to be predominantly fcc. XRD indicates that
the fcc phase grows with the ͑111͒ and ͑200͒ planes parallel to the
film surface, while the bcc phase is textured with a ͑110͒ preferential
orientation. Lattice constants were calculated from the XRD spectra
in Fig. 6a by subtracting the background and fitting the diffraction
peaks with a Lorentzian line shape; the results are displayed in Fig.
6c. The lattice parameters did not appreciably vary with the average
alloy composition, which suggests that the composition of each
phase is approximately constant. The measured average composi-
tional change may thus be attributed to a variation of the relative
amount of the two phases. In this figure are also shown the lattice
constants of bulk fcc Co, fcc Ni, and bcc Fe. The lattice constant of
the fcc phase is around 3.55 Å, larger than that of fcc Co and fcc Ni.
The lattice parameter of the bcc phase is about 2.83 Å, smaller than
in the bcc Fe. These results can be rationalized assuming that two
substitutional solid solutions nucleate and grow in parallel during
electrocrystallization of the films. In fact, the atomic radii of cobalt
͑1.67 Å͒ and nickel ͑1.62 Å͒ are smaller than that of iron ͑1.72 Å͒;
therefore, the incorporation of cobalt and nickel into the bcc iron
structure can result in a smaller lattice parameter. Similarly, the in-
corporation of Fe should increase the lattice constant of the fcc
phase above that of Ni and Co, as observed.
The average grain size in the direction normal to the film plane
was estimated from the XRD spectra by using the Debye-Scherrer
formula, neglecting the instrumental linewidth ͑which is acceptable
for a nanocrystalline material͒ and stress-induced line broadening32
L ϭ 0.9␭/B cos ␪
͓1͔
c
where L is the grain size ͑Å͒, ␭ is the wavelength of the X-ray
radiation ͑Å͒, B is the full width at half maximum ͑rad͒ of the
diffraction peak at 2␪, and ␪ is the angle of the peak center. The
c
estimated grain size of single-phase bcc films was the largest, about
35 nm, while fcc films showed a smaller value ͑about 25 nm͒. For
films with mixed fcc ϩ bcc phases, the grain size was around
10 nm.
Figure 6. ͑a͒ XRD spectra for electrodeposited FeCoNi films grown at dif-
ferent current densities with Fe2ϩ ϭ 0.025 M, Co2ϩ ϭ 0.05 M, and
͓
͔
͓
͔
Ni2ϩ ϭ 0.2 M, pH 2.8. A transition from bcc to coexisting fcc and bcc
͓
͔
phases is shown; film compositions are Fe37Co48Ni15
, Fe33Co48Ni19 ,
Fe30Co47Ni23 , Fe24Co46Ni30 for current density 5, 10, 15, 20 mA/cm2, re-
spectively. ͑b͒ Ternary diagram showing compositions of Fe-Co-Ni films
with coexisting bcc and fcc phases. ͑c͒ Lattice constants of the different
phases calculated from the XRD spectra shown in ͑a͒.
The microstructure of FeCoNi alloys was also examined by
TEM. Figures 7-9 show bright field images and high energy electron
diffraction patterns for representative fcc ͑Fig. 7͒, bcc ͑Fig. 8͒ and
mixed fcc ϩ bcc ͑Fig. 9͒ films. The fcc film with a composition of
Fe24Co46Ni30 has an average grain size of 25-50 nm. Stacking faults
due to twinning can be seen in Fig. 7a, in the grains showing lamel-
lar contrast, as also observed by Nam et al.33 Twinning in fcc metals
requires low energy, which can be easily provided during growth by
the driving force for crystallization. The estimated grain size for the
bcc film (Fe36Co48Ni16) is about 40-90 nm, with few relatively large
grains. Some contrast is observed inside single grains, indicative of
an unresolved subgrain structure. For the biphasic film with a com-
position of Fe33Co48Ni19 shown in Fig. 9, the grain size is around
10-15 nm and some twinned grains can still be observed. Film mi-
crostructure is much less resolved in the latter case, with grain
boundaries clearly visible only in limited regions. The grain size
estimated from TEM images is larger than that calculated from
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Journal of The Electrochemical Society, 150 ͑3͒ C159-C168 ͑2003͒
Figure 7. ͑a͒ TEM bright field image and ͑b͒ electron diffraction pattern of
a film with single fcc phase. Film composition is Fe24Co46Ni30 . Current
density 20 mA/cm2, and solution composition Fe2ϩ ϭ 0.025 M, Co2ϩ
Figure 8. ͑a͒ TEM bright field image and ͑b͒ electron diffraction pattern of
a film with single bcc phase. Film composition is Fe37Co48Ni15 . Current
density 5 mA/cm2, and solution composition Fe2ϩ ϭ 0.025 M, Co2ϩ
͓
͔
͓
͔
͓
͔
͓
͔
ϭ 0.05 M, and Ni2ϩ ϭ 0.2 M. pH 2.8.
ϭ 0.05 M, and Ni2ϩ ϭ 0.2 M. pH 2.8.
͓
͔
͓
͔
electrodeposited films than in bulk alloys,36 and/or with the ex-
tremely small grain size, which may enhance magnetic moment due
to the high fraction of atoms at interfaces, characterized by reduced
atomic coordination.37
XRD, and this is probably due to a relevant contribution of lattice
strain to the peak width in the XRD spectra.
The observation that films with coexisting fcc and bcc phases
exhibit the smallest grain size can be explained in terms of competi-
tive nucleation and growth of the two phases. In fact, when the two
crystal structures grow in parallel it can be assumed that the prob-
ability of formation of bcc and fcc clusters is similar. Under such
conditions, stable bcc clusters can nucleate on fcc surfaces and vice
versa. Atoms at these sites will consequently be in irregular crystal-
lographic positions and will necessarily form grain boundaries,34
decreasing the average grain size of the films.
Coercivity, internal stresses, magnetostriction, and easy axis
dispersion.—Magnetic films for write-head poles should exhibit low
coercivity and well-defined uniaxial anisotropy, with a square loop
in the easy direction and a closed, linear loop in the hard direction
up to saturation. Figure 11a shows the composition of FeCoNi alloys
with coercivity less than 2 Oe, and Fig. 11b presents typical hyster-
esis loops of these films in the easy and hard directions. The hyster-
esis loops indicate that films with a well-defined uniaxial anisotropy
are obtained upon application of a magnetic field during growth. The
anisotropy field is in the range 20-24 Oe. Narrow easy axis disper-
Magnetic properties.— Saturation magnetization.—A wide range
of alloys exhibit Bs values in excess of 1.8 T ͑18 kG͒, as shown in
the ternary diagram of Fig. 10. Most of the points are located in the
region 40-60% Co, 25-40% Fe, and 10-30% Ni, although there are
also some points at higher Fe content. The highest values of Bs are
in the range 21-21.5 kG, obtained for compositions in the range
48-52% Co, 34-36% Fe, and 14-18% Ni. The latter values are sig-
nificantly larger than those reported for bulk alloys of the same
composition.35 Since the largest source of error in the determination
of Bs by VSM is in the estimate of sample volume, two volume-
independent methods were used to independently confirm VSM val-
ues. Both FMR and low-field susceptibility agreed with the VSM
results, thus confirming the unusually high values of Bs . These val-
ues are always observed in biphasic samples and are tentatively
associated with the higher relative percentage of the bcc phase in
Ͻ
sion (␣90
3°) was also observed in these films.
The coercivity Hc along the easy axis is determined by the field
at which a domain wall is able to move through the sample and
reverse its magnetization, which is mainly related to inhomogene-
ities in the film. On the other hand, if the film has a narrow easy axis
dispersion, magnetization reversal in the hard axis direction pro-
ceeds by rotation. In this case the hard axis coercivity is zero and
magnetization reversal is completed at the anisotropy field Hk
ϭ 2Keff /Ms, where the effective anisotropy constant is the sum of
contributions from magnetocrystalline anisotropy K, stress K , and
␴
shape effects Ks,
Keff ϭ K ϩ K ϩ Ks and Ms ϭ Bs/4␲ is the mag-
␴
netic moment. For thin films with in-plane anisotropy and macro-
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Journal of The Electrochemical Society, 150 ͑3͒ C159-C168 ͑2003͒
C165
Figure 9. ͑a͒ TEM bright field image and ͑b͒ electron diffraction pattern of
film with bcc ϩ fcc two-phases coexistence. Film composition is
Fe33Co48Ni19
Current density 10 mA/cm2, and solution composition
Fe2ϩ ϭ 0.025 M, Co2ϩ ϭ 0.05 M, and Ni2ϩ ϭ 0.2 M. pH 2.8.
a
.
Figure 11. ͑a͒ Ternary diagram showing compositions of Fe-Co-Ni films
with coercivity less than 2 Oe. ͑b͒ Typical in-plane hysteresis loops along
both easy and hard axis for films with coercivity less than 2 Oe.
͓
͔
͓
͔
͓
͔
scopic sample size, Ks does not influence the reversal. Magnetocrys-
talline and stress effect are determined, respectively, by the
magnetic properties are thus favored in homogeneous films with
crystalline structure of low anisotropy ͑typically cubic͒. Exchange
interactions in a magnetic material tend to couple the magnetic mo-
crystalline structure of the film and by the product ␭ ␴, where ␭s is
s
the saturation magnetostriction constant and ␴ the local stress. Soft
ͱ
ments over an exchange length Lex ϰ A/Keff, where A is the ex-
change stiffness.38 If the crystalline grains have dimensions smaller
than Lex , the magnetic moments in each grain interact strongly over
Lex and do not necessarily follow the local magnetocrystalline an-
isotropy. As a consequence, an effective anisotropy results which is
averaged over several grains and is smaller than the anisotropy in
each grain. On the basis of this random anisotropy model Herzer39
predicts that Hc ϳ D6, with D the average grain size. On similar
grounds Hoffman40 predicts for polycrystalline films with random
crystallographic orientation Hc ϳ D2. Although the exponents for
the two cases differ, both theories predict a decrease in coercivity
with decreasing grain size. Comparison of Fig. 11 and Fig. 6b shows
that indeed there is a strong correlation between films with low
coercivity and those exhibiting fcc ϩ bcc mixed phases, which are
also those with the smallest grain size.
Stress anisotropy, on the other hand, varies on a scale character-
istic of the stress state in the films, and may not follow the same
scaling rules as the magnetocrystalline anisotropy. To determine the
influence of stress effects on the films’ magnetic properties we de-
termined both ␭ and internal stresses in a representative series of
s
films obtained at increasing current density. Figure 12 shows the
Figure 10. Ternary diagram showing compositions of Fe-Co-Ni films with
saturation magnetization larger than 18 kG.
dependence of the saturation magnetostriction ␭ on current density
s
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C166
Journal of The Electrochemical Society, 150 ͑3͒ C159-C168 ͑2003͒
Figure 14. The variation of coercivity and easy axis dispersion angle with
current density for
series of films obtained at pH 2.8, Fe2ϩ
ϭ 0.025 M, Co2ϩ ϭ 0.05 M, and Ni2ϩ ϭ 0.2 M. Film thickness, 1 ␮m.
a
͓
͔
͓
͔
͓
͔
Figure 12. The dependence of magnetostriction on deposition current den-
sity ͑alloy composition͒ for a set of FeCoNi films. Film composition is also
shown. Solution composition: Fe2ϩ ϭ 0.03 M, Co2ϩ ϭ 0.05 M, and
Film compositions are shown in Fig. 3b.
͓
͔
͓
͔
Ni2ϩ ϭ 0.2 M. pH 2.8.
͓
͔
negative. The variation in stress-induced anisotropy causes impor-
tant changes in the micromagnetic structure and magnetic properties
of the films, which are discussed in detail in Ref. 14. In summary,
films obtained at low current density exhibit large easy axis disper-
sion originated by the large in-plane magnetoelastic energy, while
films grown at high current density show a relevant perpendicular
anisotropy. In both cases, large coercivities and dispersion in the
easy axis are observed, as shown in Fig. 14 for a series of samples
grown at increasing current density. While ␭s is mainly a function of
alloy composition, internal stresses depend on a number of deposi-
tion parameters, including current density.42,43 As a consequence,
soft magnetic properties cannot only be predicted in terms of the
alloy composition, but will heavily depend on deposition conditions.
for selected samples, with the composition also reported in this fig-
ure. With increasing current density ͑increasing Ni content in the
alloy͒, ␭ decreases and changes sign, similar to the case of Fe-Ni
s
alloys around the Permalloy composition.41 At a current density of
35 mA/cm2,
␭ is near zero and the film composition is
Fe14Co31Ni55 . Other authors12 observed ␭s Х 0 for a similar com-
position when saccharin was used as an additive. At higher Ni con-
tent the films cannot be saturated at the maximum applied field of 60
Oe, and the reported values should be interpreted as ␭͑60 Oe͒.
Figure 13 shows the variation of internal stresses with current
density for the same samples described in Fig. 12. Internal stresses
are always positive ͑i.e., tensile͒ and show a minimum at 190 MPa
with increasing current density. This value is comparable to the in-
ternal stress measured on CoNiFe films electrodeposited from
chloride-based electrolytes with additives.31
Anisotropy and permeability.—In materials for write heads, induc-
tors, and transformers a high permeability ␮ is necessary to maxi-
mize the sensitivity of the material magnetization to the external
field and the ability of the material to concentrate external flux lines.
The operating frequency of write-heads and thin-film inductors for
microwave applications is currently in the gigahertz range; it is thus
necessary that the permeability of suitable materials remains at suf-
ficiently high levels at these frequencies.
The magnetoelastic energy can be expressed by
3
Eme
ϭ
␭ ␴ sin2 ␪
͓2͔
s
2
Assuming pure magnetization rotation ͑i.e., applied field along
the hard axis and narrow easy axis dispersion͒, at low frequencies
the permeability is given by the relation
where ␪ is the angle between Ms and ␴. Eme is proportional to the
product ␭ ␴, and thus it is positive and decreases with increasing
s
4␲Ms
current density in our samples, changing sign when ␭ becomes
s
␮ ϭ
͓3͔
Hk
Ms and Hk were determined from the quasi-static hysteresis loops
along the hard axis. Before measuring the hard axis loop, the sample
was saturated along the easy direction to eliminate possible fine
domain structures and to point the magnetization along the easy
direction. The dependence of Hk and 4␲Ms on current density is
shown in Fig. 15 for a representative set of films. Hk increases with
the current density in correspondence with the microstructural tran-
sition bcc → bcc ϩ fcc → fcc. The decrease in Hk observed for
films obtained at the highest current density is due to a different
magnetic microstructure.14 Maximum 4 ␲Ms is obtained when the
two phases bcc and fcc coexist. In the same figure the permeability
calculated from Eq. 3 is also reported. It can be seen that the low
frequency permeability decreases with increasing current density,
the films with lowest Hk showing the highest value.
Various loss mechanisms contribute, however, to the decrease of
␮ at high frequency. For instance, time-varying magnetic fields in-
duce eddy currents, which screen the external magnetic field from
penetrating the material, thus decreasing the effective permeability.
The length of penetration of the magnetic field is the skin depth ␦
Figure 13. The dependence of internal stress on current density for FeCoNi
films. Film composition and deposition parameters are the same as in
Fig. 12.
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Journal of The Electrochemical Society, 150 ͑3͒ C159-C168 ͑2003͒
C167
Figure 15. The dependence of anisotropy field Hk , saturation magnetization
4 ␲Ms and calculated low frequency permeability ␮ on current density for a
series of films obtained in the same conditions as Fig. 14.
Figure 17. Real and imaginary components of the permeability of a
Fe33Co48Ni19 film with bcc ϩ fcc phases, easy axis coercivity is 1.5 Oe, and
easy axis dispersion ϳ3°. Film thickness ϳ300 nm.
ͱ
ϰ
␳/f␮␮0, where ␳ is the resistivity of the material and f the
frequency of the driving field. Figure 16 reports ␳ for a representa-
tive set of samples obtained by varying the deposition current den-
sity. ␳ values are in the range 30 to 50 ␮⍀ cm, comparable to those
of bulk alloys, but much lower than values obtained by others on
electrodeposited films.31,44 This is due to the fact that foreign inclu-
sions in our samples are kept to a minimum, while high resistivity
͑Ͼ100 ␮⍀ cm͒ is usually achieved by incorporation of S- or
C-containing compounds,8,44 amorphization of the film,45 or
codeposition of nano-oxides.46
In thin films ͑Ͻ1 ␮m͒ the dominant loss mechanism is ferromag-
netic resonance ͑FMR͒, wherein permeability starts to roll off when
the drive field frequency approaches the precession frequency of the
spin in a magnetic field. The FMR frequency can be calculated by
the narrow dispersion of the easy axis as discussed above.14 On the
contrary, nonoptimized films exhibit wide easy axis dispersion and
smaller permeability at high frequency. For example, films with a
bcc structure exhibit a strong in-plane stress-induced anisotropy,
which contributes to local fluctuations in the effective anisotropy. As
a consequence, the permeability at 100 MHz is only 500. Films
obtained at high current density exhibit perpendicular stress-induced
anisotropy, and their permeability is low ͑130 at 100 MHz͒. In the
latter case, the low high-frequency permeability is related to domain
wall motion being involved in magnetization switching. Magnetic
films with high magnetic moment and low coercivity thus also ex-
hibit the best high frequency response.
ͱ
f
FMR ϭ ␥ 4 ␲Ms ϫ Hk
͓4͔
Conclusions
where ␥ is the gyromagnetic ratio.47 Figure 17 displays the real and
imaginary components of the permeability of a FeCoNi film with
biphasic structure, small grain size, Bs ϭ 20.7 kG and Hk
ϭ 21 Oe. The real component of ␮ is 1000-1250 at frequencies up
to 1.5 GHz. The imaginary component of ␮ is at a maximum at 1.85
GHz, almost equal to the theoretical value of 1.82 GHz that can be
calculated from the measured static properties and Eq. 4. In this
calculation, ␥ was taken as 2.8 MHz/Oe. This agreement is due to
We studied the electrochemical deposition process of Fe-Co-Ni
alloys containing a limited amount of impurities and their structural,
magnetic, and electrical properties. The alloy system exhibits
anomalous deposition behavior, with both Fe and Co reduction en-
hanced at low current densities. The compositional gradient reported
for electrodeposited Ni-Fe alloys is observed only to a very limited
extent in Fe-Co-Ni films. A transition of the film crystalline structure
from bcc to fcc is seen with increasing Ni content. Mixed-phase
films are obtained for a wide range of compositions and exhibit
small grain size down to 10 nm. These films also exhibit large mag-
netic moments up to 21.5 kG, higher than reported for bulk alloys of
corresponding composition. This behavior is tentatively related to
the growth of supersaturated bcc solid solutions and/or the small
grain size of such films. Small internal stresses and small positive
magnetostriction are important for achieving a narrow easy axis dis-
tribution and a small coercivity ͑ϳ1 Oe͒. As a result of the narrow
easy axis distribution, the high frequency permeability of these films
reaches high values ͑200 at 1 GHz͒.
The magnetic properties of Fe-Co-Ni alloys are determined to a
large extent by the deposition conditions. As a consequence, even
when using electrolytes containing no additives, it is not possible to
define a range of alloy composition yielding optimum magnetic
properties. Optimization of the properties must necessarily be car-
ried out experimentally for any given set of experimental conditions.
Acknowledgments
The authors thank F. Huang for his help in stress measurement.
This work was supported in part by the MINT Center through an
industrial sponsorship by Seagate Technology and made use of the
NSF-MRSEC shared facilities, grant no. DMR 98-09423.
Figure 16. Resistivity as a function of deposition current density for a series
of FeCoNi films obtained at pH 2.8, Fe2ϩ ϭ 0.05 M, Co2ϩ ϭ 0.02 M,
͓
͔
͓
͔
and Ni2ϩ ϭ 0.2 M.
͓
͔
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C168
Journal of The Electrochemical Society, 150 ͑3͒ C159-C168 ͑2003͒
25. W. D. Nix, Metall. Trans. A, 20A, 2217 ͑1989͒.
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