To submit to PRB (2013)
Chemical local order signature on magnetism in CoPt and FePt
nanoalloys
A. Hillion1†, G. Khadra1, A. Cavallin2, A. Tamion1, J. Tuaillon-Combes1, F. Tournus1, E. Otero3, P. Ohresser3, S. Rusponi2,
H. Brune2, A. Rogalev4, F. Wilhelm4 and V. Dupuis1*,
1
ILM, UMR 5306 CNRS/Université de Lyon, F-69622 Villeurbanne cedex
Institute of Condensed Matter Physics, EPFL, 1015 Lausanne, Switzerland
3
Synchrotron SOLEIL, L’Orme des Merisiers, Saint-Aubin BP 48, F-91192 Gif-sur-Yvette Cedex
4
European Synchrotron Radiation Facility, BP 220, F-38043 Grenoble Cedex
2
Abstract. By combining the high photon flux and the chemical selectivity of several synchrotron radiation
facilities, X-ray absorption spectroscopy has been intensively used to study finite size effect on the local
structure and magnetism of CoPt and FePt clusters upon annealing-driven transition to chemically ordered L10like phase. Reversely to the bulk alloys upon chemical ordering, X- ray magnetic circular dichroism (XMCD)
investigations at each L2, 3 edge, revealed a significant increase of both spin and orbital of Co (resp. Fe) and Pt
magnetic moments after thermal ordering transition. Moreover these values largely overpass all the reported
values known in the literature. Nevertheless, limited Magnetic Anisotropy Energy (MAE) enhancements have
been obtained from SQUID magnetometry measurements on the same chemically ordered CoPt and FePt
nanomagnets samples. In view to try to relate such magnetic behaviour to local structures in nanoalloys, we
performed extended x-ray absorption fine structure (EXAFS) experiments and simulations at Co, Fe and Pt
edges. In agreement with previous Vienna ab-initio simulation package (VASP) calculations, we
experimentally evidenced a systematic element-specific dependence of the local atomic relaxations in asprepared and annealed CoPt and FePt clusters. We claim that this strong distortion in pure magnetic planes
which do not match the underlying Pt layer, could explain the deviation to bulk alloys magnetic behaviours.
*
Electronic address : Veronique.Dupuis@univ-lyon1.fr
1 Introduction
Magnetic nanoalloys attract a lot of attention because
they offer the possibility to tune the magnetic moments
and the magnetic anisotropy energy (MAE) probably up
to the ultimate density storage limit. In particular, an
extremely high magnetocrystalline anisotropy is expected
from the stacking of pure Co (or Fe) and Pt atomic planes
in the (001) direction for CoPt (or FePt) bulk alloys in the
chemically ordered L10 phase. Nevertheless, even if a
well chemical order can be observed at nanosize, the
achievement of a consequent MAE enhancement without
any coalescence after annealing-driven transition,
remains so far absent [1, 2, 3, 4]. The atomic structure and
magnetic moment of CoPt and FePt nanoparticles have
been experimentally observed to differ from the
corresponding bulk materials in a favourable way or not,
due to small size effects as peculiar symmetry [5], partial
chemical ordering [6], surface segregation [7],.. . Numerous
theoretical works have been performed to try to explain
exotic structure [8, 9, 10, 11] or electronic properties [12,13]
observed in nanoalloys by integrating a great number of
parameters. Recently, we evidenced element-specifc
relaxation in size-selected CoPt clusters with 2-4 nm
diameter range, from experiment to theory confrontation
[14]
.
In this paper, in one hand, by using powerful X-ray
magnetic circular dichroism (XMCD) investigations on
FePt and CoPt clusters embedded in amorphous carbon
matrix upon annealing, we found a significant increase of
Fe, Co and Pt magnetic moments for the chemically
ordered L10-like clusters structures, reversely to the bulk
values. In the other hand, the magnetic anisotropy energy
(MAE) of annealed samples determined from SQUID
magnetometry is only twice that of the as-prepared
samples, i.e. one order of magnitude smaller to what is
expected for the L10 bulk phase. So we decided to
quantify the Co-Co (resp. Fe-Fe), Co-Pt (resp. Fe-Pt) and
Pt-Pt nearest neighbour interatomic distances from
EXAFS simulations and compared to theoretical
calculations obtained from Vienna ab-initio simulation
package (VASP) in the same size range. From
complementary
x-ray
absorption
spectroscopy
experiments at each specific Co (resp. Fe) and Pt
environment, we have been able to relate structural and
magnetic local order in bi-metallic CoPt (resp. FePt)
nanoclusters in the 3 nm range.
2 Experimental procedures
1) Sample Synthesis and Transmission electron
Microscopy characterization
Bi-metallic clusters are preformed in the gas phase thanks
to a laser vaporization source working in the Low Energy
Clusters Beam Deposition (LECBD) regime. Briefly, a
YAG laser ( = 532 nm, pulse duration 8 ns, frequency 
30 Hz) is used to vaporize a mixed equiatomic (CoPt or
FePt) target rod. Simoustaneously, a continuous flow of
inert gas (He, 30 mBar), is injected to rapidly cool the
generated plasma and to nucleate clusters submitted to a
supersonic expansion under vacuum. Moreover our
apparatus is equipped with a quadrupolar electrostatic
mass-deviator allowing us when necessary to deposit
mass-selected clusters in an UHV deposition chamber [5,
6]
. The matrix is evaporated with an electron gun working
under UHV conditions (base pressure of 5.10-10 Torr).
Clusters and atomic matrix beams are simultaneously codeposited in UHV on 45°-tilted substrate front of both
independent arriving beams.
We have focused our attention on CoPt and FePt
nanoparticles embedded in an inert carbon matrix, in
order to preserve and investigate the intrinsic cluster
surface properties. Using LECBD, we have been able to
synthesize mixed FePt nanoparticles assemblies with a
log normal size distribution centered on 3 nm in diameter
which cover the range of mass-selected CoPt clusters
present in this paper. The nanoparticles are 1%-diluted in
volume to quite avoid magnetic interaction among
nanoparticles.
The idea was to study structure and magnetism
of both FePt and CoPt nanoparticles (NPs) embedded in
amorphous carbon matrix before and after an optimum
annealing of 2 hours at 750 K under vacuum without any
coalescence.
First of all, structural characterizations have been
performed using Transmission Electron Microscopy
(TEM) on isolated CoPt and FePt clusters deposited on
an amorphous carbon coated grid (then protected by a
carbon thin film), before and after annealing under High
Vacuum. The clusters' size distributions are
experimentally determined from transmission electron
microscopy (TEM) observations. For FePt samples
prepared without mass selection, the size-histogram has
been fitted with a log normal size distribution centered on
~ 2.6 nm in diameter and relative size dispersion around
40% [1]. While for mass-selected CoPt sample obtained
thanks to the electrostatic quadrupole deflector, we used a
Gaussian shape centered on 3nm with a size-dispersion
lower than about 8% [15]. As upon 2 hours annealing
under vacuum at 750 K, the size histogram fitted with a
similar Gaussian curve, we claimed that there is no
coalescence [15].
High Resolution TEM-observations and simulations
on both CoPt and Fept NPs revealed that upon annealing
a transition occurred from as-prepared A1 fcc-structure to
a tetragonal chemically ordered L10 phase with a quasi
perfect order parameter [15]. Notice that a perfect fcctruncated CoPt or FePt octahedrons predicted from the
Wulff theory in our range of interest with a diameter of
Dm = 1.8 nm, (resp. Dm = 2.7 nm) contains 201 atoms
(resp. 586 atoms) and a number of atoms in the first
surface-monolayer equal to 60.7 % (resp. 46.4 %).
(a)
(c)
(b)
(d)
Fig. 1. HRTEM images of FePt cluster as-prepared (a) and
annealed (c) FePt clusters coated wirh thin carbon film. In
the fast Fourier Transform (b) and (c ), we clearly evidenced
peaks and angles corresponding respectively to A1 fcc and
L10 FePt phase transition upon annealing without
coalescence.
Representative HRTEM micrographs are presented in
Fig.1 even if a few percent of decahedron and multitwined shape has also been detected [5].
2) Magnetic Results – XMCD experiments
Then, x ray magnetic circular dichroism (XMCD) signal
has been measured on as-prepared and annealed
bimetallic nanomagnets CoPt and FePt samples,
respectively at the Co L-edge on SOLEIL – DEIMOS
and at the Fe -Ledge on X’Treme – SLS beamline (see
resp. Figure 2 and 3). Moreover, taking into account that
induced magnetic moments resulting from 3d-5d
proximity effects are also expected in non magnetic
metal, we performed (XMCD) measurements at Pt-L
edge on ID12 - ESRF beam line on the previous samples
(see resp. Figure 4 and 5). In table I, the mean orbital and
spin magnetic moments per Co, Fe and Pt atom, L and
S have been determined using the wellknown sum
rules16, 17 and the numbers of hole per Co, Fe and Pt
atoms estimated from theoretical band structure
calculations for L10 CoPt and FePt by applying the sum
rules to theoretical XMCD spectra (n h Co = 2.628, n h Fe
= 3.705 and n h Pt = 2.369) 18.
TABLE I. Magnetic size diameter (dm and ), MAE (Keff) and maximum susceptibility (Tmax) determined from
SQUID measurements. Co, (resp. Fe) and Pt spin and orbital magnetic moments from experimental XMCD
measurements where the values L and S were calculated using the numbers of hole per Co, Fe and Pt atoms
estimated from theoretical band structure calculations for L10 CoPt and FePt by applying the sum rules to
XMCD spectra (n h Co = 2.628, n h Fe = 3.705 and n h Pt = 2.369). Average interatomic distances (R
P.theoretical
Ohresser
in Å) and obtained on as-prepared) and annealed CoPt (resp. FePt) clusters (3nm) from EXAFS simulations at
both Co (resp. Fe)-K and Pt-L edges. The site specific L10 lattice parameter ratio [c/a] is given for the annealed
samples. Notice that the best fit has been obtained for relaxed geometry chemically ordered CoPt and FePt NPs
with 201 and 586 atoms respectively from VASP calculations by assuming truncated octahedral structures in
agreement with TEM observations.
3
SQUID
dm
(nm)

EXAFS
Co-(resp. Fe)
K edge
R (Å)
EXAFS
Pt-L edge
XMCD
Pt-L2. 3
SB/at.)
LB/at.)
L/S
SQUID
Keff
(KJ/m )
K
SQUI
D
Tmax
(K)
Å
XMCD
Co- (resp. Fe)
L2. 3
SB/at.)
LB/at.)
L/S
[c/a]
[c/a]
Å
R (Å)
CoPt asprepared
TM
3.-- nm
?
Co-Co: 2.52
Co-Pt: 2.58
0.012
Pt-Pt: 2.70
Co-Pt: 2.57
0.0045
1.67
0.13
0.077
0.47
0.07
0.150
218
37%
15 ?
CoPt annealed
TM
3.-- nm
?
Co-Co: 2.57
Co-Pt: 2.62
Pt-Pt: 2.71
Co-Pt: 2.63
22 ?
[0.92]
0.52
0.10
0.192
293
28%
[1.03]
1.98
0.20
0.101
0.009
0.004
Fe-Fe: 2.57
Fe-Pt: 2.63
0.012
Pt-Pt: 2.71
Fe-Pt: 2.63
0.008
1.71
0.11
0.06
?
15 ?
_
Fe-Fe: 2.63
Fe-Pt: 2.63
Pt-Pt: 2.70
Fe-Pt: 2.64
?
[0.93]
0.57
0.07
0.13
?
[1.00]
2.58
0.18
0.07
0.012
0.008
FePt asprepared
TM
3 nm
FePt annealed
NT
3 nm
40 %
?
?
a)
a)
b)
b)
Fig. 2. Comparison between the experimental XMCD signal
at the Co-L edge for the mass-selected CoPt sample before (a)
and after annealing (b).
Fig. 3. Comparison between the experimental XMCD signal at
the Fe-L edge for the FePt sample before (a) and after annealing
(b).
Fig. 4. Comparison between the experimental XMCD signal at
the Pt-L edge for the mass-selected CoPt sample before (a) and
after annealing (b).
Note that, even if the samples are made of randomly
oriented crystallized nanoparticles, the magnetic dipole
term T, which reflects the asphericity of the spin
moment distribution around the absorbing atom, averages
to zero only for 3d metal sites19. Thus, in table 1,  S
represents an effective spin moment at the Pt site (where
T, can be significant) while is the true spin magnetic
moment at the Co and Fe edges.
We first verified that the mass-selected CoPt sample at
the Co-site, present an analogous evolution compared to
the previous samples of ref 1 (see figure 2). In the margin
of error, estimations of the spin (1.7/2.0 B/at.) and
orbital (0.13/ 0.20 B/at.) magnetic moment for the (as
prepared/annealed) CoPt sample, lead to the same
magnetic moment increase upon annealing. While the
chemically disordered to the L10 phase has been clearly
evidence in CoPt nanoparticles, this anomalous net Co
magnetic moment enhancement is in contradiction with
the bulk behaviour where the magnetization is decreasing
when going from A1 to L10 phase20. To verify that this
effect could not be due to an extensive Pt magnetic
moment reduction which could compensate the Co
magnetic moment variation, XMCD signal has been
measured on the same sample at the Pt sites (see Figure
4). In this case, induced magnetic moments originate
from the hybridization of the Pt 5d orbitals with the Co
spin-polarized 3d states, we estimated the effective spin
(0.47/0.52 B/at.) and orbital (0.07/ 0.10 B/at.) induced
magnetic moment for the same (as prepared /annealed)
CoPt sample. Once more, there is a striking change in the
induced moment upon annealing: the S value is
enhanced by 11%. This effect definitively invalidates the
previous assumption on the Co magnetic moment
compensation. Moreover, it is worth noting that spin
momentum at both Co and Pt sites in annealed CoPt
clusters (resp. equal to 2.00 and 0.52 B/at.) are larger
than the calculated values for L10 CoPt bulk phase (1.85
and 0.37 B/at.)18. In addition, due to the fairly large spinorbit coupling of the Pt 5d electrons, the orbital moment
largely contributes to the total moment on the Pt site and
increase significantly as L/S ratio at Pt edge upon
Fig. 5. Experimental XMCD signal at the Pt-L edge for the FePt sample
after annealing.
annealing. Because of the largely directional chemical
bonding of Co atoms to Pt, anisotropic chemical local
order, L10 like phase, plays a key role in enhancing the
volume anisotropy as compared to the isotropic
chemically disordered FCC structure.
Then, we also performed XMCD measurements on FePt
samples (see Figure 3 and 5). In figure 3a, one can notice
a small multiplet in the Fe-L3 absorption edge for the asprepared FePt sample, which transforms into a small
shoulder upon annealing (see Fig; 3b). It is due to the
different chemical environment of Fe atoms at the NP-C
interface as we will see in details later. By regarding the
evolution of the mean spin magnetic moments per Fe
atom, reported in table I, one can notice that the initial
spin iron moment is smaller than in pure Fe-bulk site
(2.12B/at.) but larger at both Fe and Pt sites in annealed
FePt clusters (resp. equal to 2.58 and 0.57 B/at.) than the
calculated values for L10 FePt bulk phase (2.50 and 0.47
B/at.)18. This must be undoubtedly related to finite-size
effect in nanoalloys.
According to Stern and Gerlach experiments on pure free
Co and Fe clusters, the deviation from the bulk behaviour
occurs around 400 atoms21. In our case, the mean CoPt
and FePt cluster size corresponds to 200-600 atoms. So
we can claim that finite size effects in deposited bimetallic clusters have been clearly evidenced for the first
time without any in-situ removing native oxide shell7. In
addition, the carbon matrix provides a efficient external
degradation and oxidation protection of nanoparticles, in
particular one of the best thermal resistant due to a
preferential carbon graphitization upon high temperature
annealing under vacuum 6, 22, 23 .
SQUID measurements
Finally, the magnetic behavior of the cluster-assembled
films has also been studied using SQUID magnetometry
(Superconducting Quantum Interference Device). An
accurate analysis has been reached using our recently
developed “triple fit” method where the ZFC/FC curves
and a room temperature magnetization loop are
simultaneously adjusted using a semi-analytical model
[24]
. Our method is based on an improved description of
the susceptibility curves with a progressive crossover
between the blocked regime and the superparamagnetic
regime instead of the usual abrupt transition model, and
takes into account the experimental temperature sweeping
rate. By fitting simultaneously the entire magnetic curves,
we can reach a reliable and accurate determination of the
most important characteristics of a sample, namely the
particle magnetic size distribution and MAE [10].
In Fig. 6, we present these 3 experimental
SQUID curves and their fits obtained on both as-prepared
and annealed CoPt samples. Notice that in the
superparamagnetic regime, the magnetization curves at
300 K are found identical for both samples.
Fig. 6: Experimental and simulated magnetic susceptibility
curves from Field Cooled, Zero-Field-Cooled protocole (top)
and hysteresis loop at 300 K (bottom) of as-prepared and 750 K
annealed CoPt nanoparticles embedded in carbon matrix. The
effective anisotropy constant and the clusters magnetic size
distribution are obtained by a simultaneous triple fit procedure
of the FC, ZFC and a room temperature magnetization loop [10].
From systematic “triple fit” simulations [10], we have
been able to deduce the average magnetic diameter (Dm)
taking into account the TEM distribution, the mean
anisotropy constant (Keff) and their corresponding
standard deviation and(see TAB. I).
One can note that the median magnetic diameter Dm, is
equal to the observed TEM one and remains constant
upon annealing. This means that there is no magnetically
dead interface between the cluster surface and the matrix
and that the cluster size is conserved at least up to 750 K.
The main difference comes from the MAE evolution
between the as-prepared and annealed CoPt sample
where the median effective anisotropy constant (Keff)
increased and its corresponding dispersion  slightly
decreased. This last dispersion decrease is well explained
by the appearance of the chemical order at nanosize [ 15].
But, the MAE value is one order of magnitude smaller
that what is expected for the L10 CoPt bulk.
As-prepared FePt sample in carbon matrix has the same
Tmax as CoPt… A suivre
3) Structural characterization
Then, we have studied the local environment in these
magnetic nanoalloys on the CRG-FAME and ID12
beamlines of the European Synchrotron Radiation
Facilities at Grenoble. Extended x-ray absorption fine
structure (EXAFS) at both Co (resp. Fe)- K and Pt-L
edge has been performed in the fluorescence mode on asprepared and annealed on CoPt (resp. FePt) samples.
From FEFF Fit simulations, we clearly show that the CoCo/Co-Pt (resp. Fe-Fe/Fe-Pt) coordination ratio changes
upon annealing from 1 to 2 as expected for the A1 to L1 0
transition with contracted lattice parameters related to the
bulk phase.
After annealing, EXAFS experiments performed at both
Co (resp. Fe) K and Pt L edges, on the same annealed
CoPt (resp. FePt) sample are presented in figures 7 and 8.
The best simulation has been obtained for a perfect
chemically ordered CoPt (resp. fePt) packing (L10 phase)
with a unique Co-Pt (resp. Fe-Pt) first neighbour distance
at both edges. But surprisingly, we found a reverse
apparent c/a ratio at both edges, namely c/a = 1.03 (resp.
c/a = 1) at the Co-K (resp. Fe-K) edge and c/a = 0.92
(resp. 0.93) at the Pt-L edge for annealed CoPt sample
(resp. FePt sample) see Table I.
From previous performed spin-polarized densityfunctionnal calculations using the Vienna ab initio
simulation package (VASP) with first principles magnetic
and structural optimizations [25], applied to perfect CoPt
clusters in the same size range, we found 3 different
Gaussian relaxed distance distributions for Co-Co, Co-Pt
and Pt-Pt both in chemically disordered and L1 0 ordered
phases. In particular, while in the bulk phases, an equal
Co-Co (resp. Fe-Fe) and Pt-Pt Dirac distance distribution
is expected, a strong Co-Co distance dispersion has been
found for relaxed clusters even in the chemically ordered
phase. The tendency of Co nearest neighbour bonds to be
shorter than the Pt ones results in structural stress, which
in a finite CoPt cluster can be more easily distorted by
moving Co atoms (than Pt atoms) in the surface faces.
As seen in Table I, one can notice the very good
agreement between EXAFS simulation and VASP
predictions. As expected from the respective TEM sizedistribution, the FePt samples presents a larger DebyeWaller value compared to the size selected CoPt sample.
An other important difference existing in the FePt
samples compared to CoPt, is the fact that it is necessary
to consider C atoms neighboring in the Fe-environment,
especially in as-prepared FePt phase. The contribution of
interstitials carbon chemical environment at the FePt
clusters surface is noticeable by the smoothing of the
EXAFS oscillations compared to the CoPt curves (see
Figures 7 and 8). In fact, the enthalpy of carbide
formation is close to zero in Fe fcc and increases by
Fig.7. Comparison between the experimental EXAFS signal
(dots, contribution of the NN peak only) and simulated
curves (solid lines) with c/a = 1.03 at the Co-K edge (a) and
with c/a = 0.92 at the Pt-L edge (b) obtained on L10 3 nm in
diameter CoPt clusters in very good agreement with VASP
calculations.
reducing the particle size26, 27. This metastable carbide is
probably at the origin of the magnetically dead interface
layer found in the as-prepared FePt sample from XMCD
at Fe edge. Annealing the samples at 750 K favors the
graphitization of the matrix and removes the dead layer
without deteriorating the nanoparticle size distribution 23.
Moreover, we have to notice that this demixing has not a
damage effect on the magnetic moment of annealed
sample. Discussions concerning MAE and magnetism
from SQUID magnetometry measurements on FePt
sample, are in progress …
3 Conclusion
The VASP ab-initio calculations fully confirm the
experimental structural trends in the size-range of our
nanoalloys, providing a detailed account of the element
specific local relaxations which bring back together the
EXAFS experimental results at both element edges.
Indeed, this element-specific dependence of the local
atomic relaxations in CoPt and FePt clusters lead to a
strong distortion in pure magnetic planes. The fact that
Fig. 8. Comparison between the experimental EXAFS signal
(dots, contribution of the NN peak only) and simulated curves
(solid lines) with c/a = 1.00 at the Fe-K edge (a) and with c/a
= 0.95 at the Pt-L edge (b) obtained on L10 FePt clusters in
very good agreement with VASP calculations.
pure Co and Fe layers do not match the underlying Pt
layer is an specifity of finite size effects in nanaoalloy,
certainly relate to the low MAE results obtained on
chemically ordered L10-like nanoalloys.
We claim that we obtain new insights on the correlation
between magnetic properties and short- or long-range
chemical order in nanoalloys. We clearly demonstrate
that the magnetic moment increase while the MAE for
nearly perfect chemical ordered nanoparticles assemblies
is one order lower to what is expected for the bulk.
In a previous paper [28], we have shown that the MAE
dispersion in pure Co clusters essentially comes from the
effect of additional facets and is relatively small. But in
nanoalloys, the size and shape dispersion is not the
unique source of MAE dispersion. Indeed, since the
anisotropy enhancement in as-prepared CoPt compared to
pure Co clusters is due to the presence of Pt atoms, the
dispersion of the magnetocrystalline anisotropy (which
depends on the neighborhood of each Co atom) increases
with the number of possible chemical arrangements. It is
the reason why the MAE of chemically disordered CoPt
particles is quite large even if mass-selected clusters have
a small size dispersion (8%) and a highly symmetrical
shape (regular truncated octahedron) [15]. As long as a
well-defined and high enough degree of chemical order
can be reached, the multiplicity of atomic configurations
is strongly reduced and the MAE dispersion decreases
while its median value increases29.
Discussions concerning MAE and magnetism in FePt
sample, are in progress …
To go further, non-collinear calculations, including the
spin-orbit coupling in order to quantitatively access to the
MAE values are in progress for such relaxed nanoalloys
(Ab-initio+ tight-binding calculations).
Acknowledgment
The authors are grateful to A. Ramos, H. Tolentino, M.
de Santis and Olivier Proux for stimulating discussions
and for their help during experiments on the French
CRG-D2AM and BM30b-FAME beamlines at ESRF and
to J. Dreiser, C. Piamonteze and F. Nolting from Swiss
Light Source for their investment on the X’Treme
beamline. Support is acknowledged from both GDR
CNRS 3182 and COST-STSM-MP0903-7318 on
Nanoalloys
†
Current address: Laboratory of Solid State Physics and Magnetism,
K.U. Leuven, Belgium
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