Simultaneous magnetic and chemical imaging of Nd-Fe-B thin films
by means of XMCD-PEEM technique
R. Goto1,*, S. Okamoto1, 2, T. Ohkochi3, N. Kikuchi1, O. Kitakami1, and T. Nakamura2, 3

1 Institute of Multidisciplinary Research for Advanced Materials (IMRAM), Tohoku University,

980-8577, Japan

2-1-1 Katahira, Aoba-ku, Sendai

1-2-1 Sengen,Tsukuba, Ibaraki 305-0047, Japan
3 Japan Synchrotron Radiation Research Institute/SPring-8, 1-1-1, Kouto, Sayo-cho, Sayo-gun, Hyogo 679-5198, Japan

2 Elements Strategy Initiative Center for Magnetic Materials (ESICMM), National Institute for Materials Science,

The coercivity mechanism of Nd-Fe-B thin film samples with and without Nd deposition was investigated bysimultaneous chemical and magnetic imaging by means of X-ray magnetic circular dichroism-photoemission
electron microscopy. The two Nd-Fe-B thin film samples exhibited almost the same surface morphology andmagnetic domain patterns, but their coercive fields were quite different, being 0.98 T for the sample with Nd
deposition and 0.55 T for sample without Nd deposition. The chemical distribution of Nd revealed that Nd
aggregates with the average size of 130 nm exist, and the density in the Nd-deposited sample is much higher thanthat in the not deposited one. A comparison of the magnetic and chemical contrast images implied that the Nd
aggregates act as domain wall pinning sites, indicating that the large coercivity difference in the Nd-Fe-B thinfilms would be attributable to the different density of wall pinning sites.
Key words: Neodymium-Iron-Boron magnets, thin films, magnetic reversal, X-ray magnetic circular dichroism,
photoelectron emission microscopy
1. Introduction
Nd-Fe-B magnets are a key material for highefficiency electric generators and motors. For theseapplications, numerous attempts have been made toenhance the coercivity while maintaining highsaturation magnetization. Although heavy rare earth(RE) elements such as Dy or Tb are usually added toincrease the coercivity of Nd-Fe-B magnets, these RE
elements degrade their saturation magnetization due
to antiferromagnetic coupling between 3d and 4f
magnetic moments. Moreover, the availability andprice of the heavy RE elements are currently serious
resource problems. Recently, the coercivity μ0Hc of
Nd-Fe-B sintered magnets has been successfullyenhanced to about 2 T even without addition of heavy
RE elements by various methods such as refinement of
the constituent crystal grains of the Nd2Fe14B phase 1)
and/or grain boundary diffusion of RE elements torestore magnetic anisotropy near the grain boundaries
or to suppress intergranular exchange coupling 2), 3).
However, the μ0Hc obtained by these techniques
remains much smaller than the anisotropy field of
μ0Hk = 7.5 T 4).
The large discrepancy between the μ0Hc and μ0Hk
observed in a wide variety of permanent magnets,
which is known as Brown’s paradox 5), is the subject ofa long-standing debate in magnetics. To solve thisproblem, a detailed analysis of the magnetization
reversal process is indispensable. So far it is widely
accepted that nucleation of reversed domains
dominates the magnetization reversal in Nd-Fe-B

magnets, judging from the initial and minor
magnetization curve behaviors 6). On the other hand,the angular dependence and c-axis dispersion
dependence of μ0Hc 7), 8) strongly suggest that domain
wall depinning is the dominant process in Nd-Fe-B
magnets.
Recently, we proposed a new approach to examine
the magnetization reversal process 9). At a finite
temperature, magnetization reversal proceeds under
the influence of thermal activation, which enables themagnetization to overcome a finite energy barrier
separating local energy minima. The energy barrier isa function of a magnetic field H, and its function form
depends on the magnetization reversal process. For
example, the energy barrier is a quadratic function of
H for coherent rotation and/or nucleation 10) and alinear function of H for weak pinning of a domain wall11). The energy barrier function can be determined
experimentally by analysis of magnetic viscosity
measurements. We applied this type of analysis to two
kinds of (001) oriented Nd-Fe-B thin film samples withand without a Nd overlayer which is preferentiallydiffused along the Nd-Fe-B grain boundaries during
the deposition 9). In spite of the large difference in theircoercivities, both of their barriers were found to belinear functions of H, suggesting that weak domain
wall pinning is dominant in both samples. It should be
stressed here that the above two Nd-Fe-B thin film
samples have almost the same structures, such as
c-axis orientation, grain size, and so on. In order tounderstand the reason for the large coercivity

difference mentioned above, simultaneous magnetic
and chemical distribution imaging were performed on
the two Nd-Fe-B film samples by means of X-ray
magnetic circular dichroism-photoemission electron
microscopy (XMCD-PEEM).
2. Experiments
Nd-Fe-B thin films examined in the present study
are grown on MgO (001) substrates with the layer
structure of MgO/Mo(30 nm)/Nd(0.25 nm)/Nd-Fe-B(50nm)/Nd(tNd = 0, 1 nm)/Mo(10 nm). For brevity, the
Nd-Fe-B film samples with and without a Nd
deposition are denoted as Nd-Fe-B/Nd and Nd-Fe-B,
respectively. The composition of the Nd-Fe-B layer isadjusted to Nd12.8Fe78.7B8.5 in at% by controlling the
sputtering rates of Nd, Fe, and B. The substrate is
kept at Ts = 500oC during deposition of both Nd-Fe-B
and Nd. The detailed sample fabrication conditions andmagnetic and structural characterization methods aredescribed elsewhere 9). Since both Nd-Fe-B/Nd andNd-Fe-B are prepared under the same condition, they
have the same crystal structure with good c-axis
orientation. The surface morphologies and magnetic
domain structures were observed by atomic force
microscopy (AFM) and magnetic force microscopy(MFM),
respectively.
X-ray magnetic
circulardichroism-photoemission
electronmicroscopy(XMCD-PEEM) was performed with Nd M4, 5 and Fe L2,
3 edges at the BL25SU beamline of SPring-8. Soft
X-rays were incident on the sample at 30 degree withrespect to the sample plane, enabling us to detect
perpendicular component of the magnetizations. Prior
to the XMCD-PEEM experiment, the Mo capping
layers were etched ex situ down to about 5 nm
considering the very short probing depth of PEEM. The
samples were observed at room temperature without
application of a magnetic field.
3. Results and Discussions
Figures 1(a) and 1(b) show the perpendicularmagnetization curves of Nd-Fe-B and Nd-Fe-B/Nd,
respectively. While they both exhibit high remanence
ratios, their coercive fields are quite different, being
μ0Hc = 0.55 T for Nd-Fe-B and 0.98 T for Nd-Fe-B/Nd.
According to the AFM images in Figs. 1(c) and 1(d),both samples have a partially-connected island
structure with the mean size of 100±50 nm. Figures
1(e) and 1(f) show the MFM images of thermallydemagnetized Nd-Fe-B and Nd-Fe-B/Nd, respectively.
Their domain structures are very similar with theaverage size of 200 nm. This is roughly two times
larger than the grain size in Figs.1 (c) and 1(d),indicating that each Nd-Fe-B grain is more or lessexchange-coupled to neighboring grains. In spite of thevery similar micro- and domain-structures, the
Nd-Fe-B with and without Nd deposition exhibit quite

Fig. 1 (a) and (b) perpendicular magnetization
curves, (c) and (d) AFM images, and (e)and (f) MFM images of thermal
demagnetization states. (a), (c), and (e)show Nd-Fe-B, and (b), (d), and (f) show
Nd-Fe-B/Nd.
different coercivities.
In order to explore the reason for the difference, we
performed simultaneous chemical and magnetic
imaging by means of XMCD-PEEM. Figure 2 shows the
X-ray absorption spectra (XAS) of the Nd-Fe-B
obtained at the Fe L2, 3 and Nd M4, 5 edges. In addition
to the clear Fe and Nd absorption edges, a small
shoulder is observed at the slightly higher energy side
near the Fe L3 edge owing to oxidation of Fe. This
oxidation may be caused by partial over-etching of the
Mo capping layer. We believe this small amount ofoxidation does not affect the following discussion.
Figures 3(a) and 3(b) show the chemical contrast
maps of Nd for Nd-Fe-B and Nd-Fe-B/Nd, respectively,

Fig. 2 XAS spectra for (a) Fe L2, 3 and (b) Nd M4, 5
edges of Nd-Fe-B and Nd-FeB/Nd.

Fig. 3(a) and (b) chemical contrast images of Nd,(c) and (d) magnetic contrast images, and(e) and (f) contours of magnetic contrast
shown in (c) and (d) overlapped with Ndchemical images. The red-colored contrast
of contours corresponds to the Nd content
under themselves. (a), (c), and (e) show
Nd-Fe-B, and (b), (e), and (f) show
Nd-Fe-B/Nd.

which were displayed as the spatial distributions of thephotoemission intensities at the Nd M5-edge (978.5 eV)
normalized to those at a pre-absorption-edge (972.0 eV).
Since Nd is one of the main constituents of the samples,
Nd is distributed over the whole area of the samples,
as indicated by the blue-colored areas in Figs. 3(a) and3(b). In addition to this uniform Nd distribution, many
tiny Nd-rich aggregates with the size of 130 nm can beclearly observed as indicated by green areas in thesame figures. Obviously, the density of the Nd
aggregates in Nd-Fe-B/Nd is much higher than that in
Nd-Fe-B. According to cross-sectional transmission
electron microscopy (TEM) observation of the Nd-Fe-B
film prepared under the same conditions, Ndaggregates exist laterally beside the Nd-Fe-B particles
on the Mo buffer layer 12). Figures 3(c) and 3(d) show
XMCD-PEEM images of Nd-Fe-B and Nd-Fe-B/Nd,
respectively, which were taken to be the XMCD
amplitude defined as the difference of XMCD intensity
between the Fe L3 edge (707.6 eV) and Fe L2 edge
(720.8 eV) to highlight the very weak XMCD contrast.
In spite of this data processing, the XMCD contrast is
still very weak, probably due to the presence of the Mo
capping layer. However, XMCD contrasts reflecting themagnetic domain states can be clearly noted, and thepatterns are similar to the MFM images in Figs. 1(e)and 1(f). The average domain width of Nd-Fe-B in fig.

3(c) is about 200 nm, being comparable with the value
from MFM images in Fig. 1(e). For Nd-Fe-B/Nd,
however, the domain width roughly reaches 400 nm,
which is about 2 times larger than that from MFM in
Fig. 1(f). This discrepancy is probably brought about by
weak magnetic contrast in the XMCD measurements.
The contours of these XMCD contrasts are overlaid on
the chemical mapping of Nd in Figs. 3(e) and 3(f). Notethat the magnetic contours which correspond to
domain walls partially run across Nd aggregates as
indicated by the red-colored portions on the contours.
It suggests that Nd-rich aggregates act as pinning sitesof magnetic domain walls. This tendency is more
obvious in the Nd-Fe-B/Nd sample shown in Fig. 3(f).
Each magnetic contour overlaps multiple Nd
aggregates, and in some cases lies along the aggregates.
The above observation implies that Nd aggregates act
as pinning sites of magnetic domain walls. When amagnetic domain wall is pinned at multiple defects, Ndaggregates in this study, the effective pinning force inthe two-dimensional system as in the present study
may be proportional to the density of defects 13) if theso-called weak pinning process is dominant 14). Thus
based on the present experimental results, we atpresent consider that the larger coercivity for
Nd-Fe-B/Nd is due to higher density of pinning sites, inthis case Nd aggregates. This conclusion isqualitatively consistent with our previous report, inwhich we found that domain wall pinning was
dominant for both Nd-Fe-B and the Nd-Fe-B/Ndsamples 9).
4. Conclusion
To discuss the origin of the different coercive fields of
Nd-Fe-B thin film samples with and without Nddeposition, simultaneous observations of magnetic andchemical contrast images were carried out by the
XMCD-PEEM technique. These two Nd-Fe-B films
exhibit almost the same surface morphologies andmagnetic domain patterns, but their coercive fields arequite different. From the Nd chemical contrast imaging,it was found that a large number of Nd aggregateswith the average size of 130 nm exist, and the densityof the Nd aggregates for the Nd deposited sample ismuch higher than that for not deposited one. Although
the XMCD contrasts of these thin film samples werevery weak, patterns similar to the MFM images wereobtained. A comparison of the chemical and magnetic
contrast images implied that the Nd aggregates act aswall pinning sites. This result indicates that the large
difference in coercive fields of the Nd-Fe-B thin films
with and without Nd deposition is attributable to thedifferent density of wall pinning sites.
Acknowledgements This work was partially supportedby the Elements Strategy Initiative Center forMagnetic Materials from MEXT, the Management
Expenses
Grantsfor
National
Universities
Corporations from MEXT, and JSPS KAKENHI Grant
Number 24360261. The XMCD measurements were
performed with the approval of the Japan Synchrotron
Radiation Research Institute (JASRI) (Proposal No.
2014B1272).
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Present address*Santoku Corporation, 4-14-34 Fukaekitamachi,
Higashinada-ku, Kobe, Hyogo 658-0013, Japan
Received Sep. 25, 2015; Revised Jan. 29, 2016; Accepted Apr.
11, 2016