high-energy spectroscopic study of mn-based magnetic...
TRANSCRIPT
High-energy spectroscopic study
of Mn-based magnetic
semiconductors
Master Thesis
Yoshitaka Osafune
Department of Physics, University of Tokyo
January, 2006
Contents
1 Introduction 1
2 Principles of high-energy spectroscopy 9
2.1 Photoemission spectroscopy . . . . . . . . . . . . . . . . . . . . . 9
2.2 Resonant photoemission spectroscopy . . . . . . . . . . . . . . . . 12
2.3 X-ray absorption spectroscopy . . . . . . . . . . . . . . . . . . . . 13
2.4 X-ray magnetic circular dichroism . . . . . . . . . . . . . . . . . . 13
3 Experimental 15
4 Chalcopyrite-type magnetic semiconductor MnGeP2 thin films 17
4.1 Physical properties of MnGeP2 . . . . . . . . . . . . . . . . . . . 17
4.2 Sample preparation . . . . . . . . . . . . . . . . . . . . . . . . . . 20
4.3 Results and discussion . . . . . . . . . . . . . . . . . . . . . . . . 21
5 Thermally diffused Mn/GaAs (001) thin films in depth profile 27
5.1 Physical properties of thermally diffused
Mn/GaAs (001) . . . . . . . . . . . . . . . . . . . . . . . . . . . . 27
5.2 Sample preparation . . . . . . . . . . . . . . . . . . . . . . . . . . 31
5.3 Results and discussion . . . . . . . . . . . . . . . . . . . . . . . . 33
6 Summary 39
i
Chapter 1
Introduction
It is well known that electrons have the charge degrees of freedom correspond-
ing to the electric properties and the spin degrees of freedom corresponding to
the magnetic ones. These two characteristics have been separately applied to
various devices so far. For instance, the charges of electrons have been utilized
in semiconductor electronics such as integrated circuits, transistors, and lasers.
On the other hand, the spins of electrons have been utilized in magnetics such
as hard disks. Recently significant movement toward the coupling between the
charge and spin of electron has occurred and a new field, called “spintronics”
(spin electronics). The fascination of spintronics is that spintronics devices hold
the possibilities to show properties which have never been achieved by the in-
dependent use of devices in semiconductor electronics or magnetics. Diluted
magnetic semiconductors (DMS’s), where magnetic ions are doped into the non-
magnetic host semiconductors, are promising candidates for the application of
spintronics. In the early 1980s, II-VI-based DMS’s such as Cd1−xMnxTe have
been extensively studied, because Mn2+ ions substitute isovalently the II-site ions
in II-VI-based DMS’s, which enables high concentration of Mn doping. However,
most of II-VI-based DMS’s have few carriers, thus superexchange interaction
between the magnetic moments becomes predominant and show paramagnetic,
anti-ferromagnetic, or spin glass behavior. Although Cd1−xMnxTe is now used as
optical isolators [1], this has not been sufficient to demonstrate the potentiality
of spintronics.
The prosperity of spintronics has started by the breakthrough of crystal
growth technique in the late 1980s, that is, low temperature molecular beam
epitaxy (MBE), which is a sophisticated version of vacuum deposition and en-
ables us to grow single crystal sample epitaxially on a substrate in an ultrahigh
vacuum (UHV) by evaporating independently each element from each crucible.
1
Note that since this method is a non-equilibrium thermal growth technique, one
can incorporate transition metals into host semiconductors beyond its solubility
limit. Thus, III-V-based DMS’s, where solubility limit in equilibrium state is low
because heterovalent substitution of divalent ions for III-site ions takes place, has
attracted considerable interest. Using MBE method, Munekata et al. succeeded
in the synthesis of In1−xMnxAs [2], and Ohno et al. Ga1−xMnxAs [3]. Different
from II-VI-based DMS’s, they showed ferromagnetism, which made a big impact
in this research field. Ferromagnetism in III-V-based DMS’s is thought to be
caused by coupling between the local magnetic moments of magnetic ions medi-
ated by the charge carriers of the host semiconductors, that is, “carrier-induced
ferromagnetism”. Such ferromagnetism in III-V-based DMS’s initiates various
interesting properties. As for Ga1−xMnxAs, it was reported that magnetotrans-
port properties such as large negative magnetoresistance [4] and anomalous Hall
effect [5] appeared as shown in Fig. 1.1.
Here, we would like to introduce several proposed heterostructure devices
based on ferromagnetic III-V-based DMS’s. Figure 1.2 shows a light-emitting
p-i-n diode [6]. Under forward bias, the device emits circularly polarized light
through the recombination of spin-polarized holes and spin-unpolarized electrons,
which is the significant evidence of electrical spin injection. Figure 1.3 shows a
(In,Mn)As/GaSb heterostructure where ferromagnetism can be induced by pho-
togenerated carriers. The electron-hole pairs produced by the light irradiation
are split by the internal electric field, causing the holes to accumulate in the
(In,Mn)As layer at the surface. Since (In,Mn)As exhibits hole-induced ferromag-
netism, phase transition from paramagnetic state to ferromagnetic state occurs
in the device after light irradiation, as depicted in Fig. 1.3 (b), (c).
Search for high-TC ferromagnetic DMS’s, in particular, room-temperature fer-
romagnetic DMS’s have been heated up because the the range of application is
expected to be widely extended. Dietl et al. predicted the T C’s of various DMS’s
using mean-field approximation [8] as shown in Fig. 1.5(a). Inspired by this pre-
diction, many room-temperature ferromagnetic DMS’s have been fabricated such
as ZnO:Co [9], ZnV:O [10], GaN:Mn [11], GaN:Cr [12].
Recently, Mn-based II-IV-V2 chalcopyrite-type DMS’s showed ferromagnetism
above the room-temperature, and have attracted much attention of many re-
searchers. For instance, it has been reported that the T C’s of ZnGeP2:Mn and
CdGeP2:Mn are ∼350 K and ∼320 K, respectively, well above the room tem-
perature, as shown in Fig. 1.4(a)-(c) [13, 14]. In the II-IV-V2 DMS’s, isovalent
substitution of Mn2+ for the II-site ions become possible, which enables high
concentration Mn-doping. Carriers are expected to be doped by various kinds of
2
(a) (b)
(c)
(d)
Figure 1.1: Magnetotransport properties of Ga1−xMnxAs. (a) Temperature dependence ofresistivity for samples with x ranging from 0.015 to 0.071 [4]. (b) Magnetoresistance for asample with x = 0.071 at different temperatures ranging from 1.4 to 250 K [4]. (c) Magne-toresistance for a sample with x = 0.035 at different temperatures from 1.7 to 200 K [4]. (d)Dominant contribution of the anomalous Hall effect to the Hall coefficient RH (ρ/RH), whichis proportional to (T − TC), for a sample with x = 0.06 [5].
3
(a)
(b)
Figure 1.2: Electrical spin injection [6]. (a) Schematic diagram of device structure. Underforward bias, spin-polarized holes (h+) from p-(Ga,Mn)As recombine with spin-unpolarizedelectrons from n-GaAs substrate, and as a result, circularly polarized light (σ+) is emittedfrom the edge of the (In,Ga)As quantum well. (b) Dependence of the polarization ∆P of theemitted light on the magnetic field at each temperature.
(a)(b)
(c)
Figure 1.3: Photoinduced ferromagnetism [7]. (a) Device structure of (In,Mn)As/GaSb. (b)M -H curves at 5 K observed before (open circles) and after (solid circles) light irradiation.Solid line shows a theoretical curve. (c) Hall resistivity observed at 5 K before (dashed line)and after (solid line) light irradiation.
4
(a) (b)
(c) (d)
Figure 1.4: Magnetization curves of Mn-based II-IV-V2 chalcopyrite-type magnetic semicon-ductors. (a) M -H curves of CdGeP2:Mn at T = 300 K for magnetic field in two orientations(solid curves: in plane, open circles: perpendicular) [13]. (b) M -T curve of CdGeP2:Mn at H= 0 T [13]. (c) M -H curve of ZnGeP2:Mn at T = 350 K. The inset shows an enlarged plotindicating a hysteresis behavior [14]. (d) M -T curves of MnGeP2 and MnGeAs2 at H = 0.01T [15].
5
defects such as antisite IIIV ions and vacancies VII, VIV, and VV. Recent stud-
ies have revealed that 100 % Mn substitution in ZnGeP2, i.e., MnGeP2 can be
synthesized and shows p-type semiconducting properties (see Section 4.1) and
ferromagnetism at room-temperature as shown in Fig. 1.4(d) [15]. MnGeP2 also
shows magneto-optical effect such as Kerr effect [16] and application to various
magneto-optical devices can be considered.
Figure 1.5: Theoretical predictions of TC for various p-type semiconductors containing 5 %of Mn and 3.4 × 1020 holes per cm3 [8].
As the fabrication technique of DMS’s, thermal diffusion method as well as
MBE method has been in use. As a matter of fact, ZnGeP2:Mn and CdGeP2:Mn
mentioned above have been fabricated by thermal diffusion method [13, 14].
Schematic illustration of thermal diffusion is described in Section 5.1. Taking Mn
on ZnGeP2 substrate as an example, Mn is thermally diffused and reacted with
ZnGeP2 substrate by annealing the substrate. Then, by removing the surface
metallic Mn layer in some way such as noble gas ion sputtering, one can fabri-
cate the DMS ZnGeP2:Mn. In order to elucidate the origin of the ferromagnetism
of ZnGeP2:Mn, detailed electronic structure along the depth direction has been
provided by depth profile study using photoemission spectroscopy (PES) [17].
Mn doping into GaN thin film has also been achieved by the thermal diffusion
method [18, 19]. As noted above, Ga1−xMxAs (synthesized by MBE method)
is a remarkable DMS, and shows various magnetotransport properties. We at-
tempted to fabricate thermally diffused Mn/GaAs and compare the electronic
structure of it with that of Ga1−xMnxAs.
In this thesis, we have studied MnGeP2 thin films and thermally diffused
Mn/GaAs (001) thin films using high-energy spectroscopic method. As for
MnGeP2 thin films, we have focused on the elucidation of the electronic struc-
6
ture and magnetic properties. Photoemission spectroscopy (PES) is a powerful
technique to elucidate the core-level and valence-band electronic structure. X-
ray absorption spectroscopy (XAS) is also a useful technique to investigate the
electronic structure, and x-ray magnetic circular dichroism (XMCD) measure-
ments provide us with the information about element-specific spin and orbital
magnetic moments. We have adopted combined PES, XAS, and XMCD. As for
thermally diffused Mn/GaAs (001) thin films, we have combined PES with Ar+-
ion sputtering, which enabled us to perform depth profile analysis, which is a
suitable method for studying the electronic structure along the depth direction
and diffusion effect.
The present thesis is organized as follows. First, the principles of high-energy
spectroscopic method are explained in Chapter 2. The experimental setup is
described in Chapter 3. Detailed physical properties and experimental results
of MnGeP2 thin films and thermally diffused Mn/GaAs (001) thin films are
presented in Chapters 4 and 5, respectively. Finally, Chapter 6 is devoted to
summary.
7
Chapter 2
Principles of high-energy
spectroscopy
In this chapter, we describe the principles of high-energy spectroscopy, that is,
photoemission spectroscopy, resonant photoemission spectroscopy, x-ray absorp-
tion spectroscopy, and x-ray magnetic circular dichroism.
2.1 Photoemission spectroscopy
Photoemission spectroscopy (PES) is a useful technique to investigate the occu-
pied electronic structure in solids. Photoelectric effect, where an illuminated solid
emits electrons (photo-electrons), is utilized in PES. PES spectra are provided by
the measuring kinetic energy distribution of photo-electrons which escape from
the solid through the surface and overcome the vacuum level EV . Assuming that
when light with photon energy hν is illuminated on a solid, an electron with the
binding energy EB, which is referenced to the Fermi level EF, is emited with the
kinetic energy EVkin, which is referenced to EV . In this case, we can describe the
relationship between EB and EVkin using the energy conservation law:
EVkin = hν − φ −EB , (2.1)
where φ = EV − EF is the work function of the solid. The kinetic energy Ekin
which is reference to EF is practically observed in PES measurements, thus the
notation is rather simplified by
Ekin = hν − EB. (2.2)
Schematic diagram of PES is shown in Fig. 2.1. The density of states (DOS)
9
EB
Ekin hνphoton
photo-electron
EF
EV
E
DOS (N(E))
φ
EkinV
valence band
core levelIntensity (I(E))
Ekin
EF
PES spectra
Figure 2.1: Schematic diagram of photoemission spectroscopy. The density of states N(E) isobtained by measuring the photoemission spectra I(E).
N(E) is obtained by measuring the photoemission spectra I(E), which are broad-
ened by the resolution of light source and electron energy analyzer. In one-
electron approximation, the binding energy is equal to the negative Hartree-Fock
orbital energy with Bloch wave number k,
EB = −εk. (2.3)
Here, Koopmans’ theorem is used [20]. This assumption is valid when the wave
functions of both the initial and final states can be expressed by the single Slater
determinants of the n- and (n − 1)-electron systems, respectively, and the one-
electron wave functions do not change by the removal of the electron. If we apply
this approximation, the photoemission spectrum I(EB) can be expressed as:
I(EB) ∝∑
k
δ(EB + εk) ∝ N(−EB). (2.4)
Thus, when the one-electron approximation is valid, the photoemission spectrum
is proportional to the density of states of the occupied one-electron states N(E).
If the electron correlation effect is taken into account, one can no more con-
sider the electron system within the one-electron picture, because the relaxation
influences the photoemission final state such as screening of photo-holes by va-
lence electron. Thus, the energy difference between the n-electrons initial state
energy Eni and the (n− 1)-electrons final state energy En−1
f provides the binding
energy EB, that is,
10
Figure 2.2: Mean free path of electrons in solids as a function of electron energy. Dashedcurves indicate the approximate range of distribution [21].
EB = En−1f −En
i . (2.5)
Using Fermi’s golden rules, the PES spectrum, which now corresponds to the
single-particle excitation spectrum of the electron system, is expressed as:
I(EB) ∝∑
k
|〈Ψn−1f |ak|Ψn
i 〉|2δ[EB − (En−1f −En
i )], (2.6)
where Ψn−1f and Ψn
i denote the final and initial states, respectively, ak is the
annihilation operator of the electron occupying orbital k. Considering electron
correlation effect, the finite lifetime of quasi-particle also contributes to the spec-
tral broadening.
Figure 2.2 shows the energy dependence of the mean free path of electrons in
solids. The escape depth of electrons is described by the universal curve, roughly
independent of the material. Around the electron energy of 20-1000 eV, where
we perform the PES measurements, the escape depth is 5-10 A. This suggests
that PES measurement is quite surface-sensitive. Therefore, we should always
keep it in mind to eliminate the surface effect.
11
E
EF
final state
E
EF
final state
E
EF
direct photoemission
Fanoresonance
Auger decay
p6dn + hν p5dn+1 p5dn+1 p6dn-1 + e-
p6dn + hν p6dn-1 + e-
3p 3d absorption
3pvalence
EF
E
hν
hν
initial state
Figure 2.3: Schematic diagram of resonant photoemission spectroscopy. The case of 3p → 3dabsorption is taken as an example.
2.2 Resonant photoemission spectroscopy
Resonant photoemission spectroscopy (RPES) is an effective approach to extract
the PES spectrum for an impurity atom from the entire spectrum in the valence
band. RPES measurement is achieved by synchrotron radiation, where photon
energy hν can be continuously varied. Let us take 3p → 3d absorption as an
example. Schematic diagram of RPES is shown in Fig. 2.3. The direct PES
process of a valence 3d electron is described as:
p6dn + hν → p6dn−1 + e−. (2.7)
When the photon energy is equal to the absorption energy from the 3p core level
to the valence 3d state, 3p → 3d absorption and subsequent Auger decay, called
super Coster-Kronig decay occur.
p6dn + hν → p5dn+1 → p6dn−1 + e−. (2.8)
The final states of these two processes are the same electronic configurations,
and quantum-mechanically interfere with each other in consequence. Thus, the
photoemission intensity is resonantly enhanced and shows a so-called Fano profile
[22]. This enhancement helps detecting weak signals such as photoemission from
transition metal impurities in the valence band, which is difficult to obtain by
normal PES.
12
2.3 X-ray absorption spectroscopy
X-ray absorption spectroscopy (XAS) is a powerful technique to investigate the
unoccupied electronic structure in solids. The photo-absorption intensity by ex-
citation of a core-level electron into unoccupied states as a function of photon
energy hν is given by
Iµ(hν) ∝∑
f
|〈Ψf |Tµ|Ψi〉|2δ(Ef −Ei − hν), (2.9)
where T is the dipole transition operator, µ is the index of light polarization, and
Ei and Ef are the energies of the initial and final state, respectively. In the 3d
transition-metal compounds, transition-metal 2p (L2,3-edge) XAS spectra reflect
the electronic structure of the 3d states such as the spin state and the crystal-
field splitting. In order to interpret experimental spectra, various theoretical
calculations have been applied.
The measurement modes for XAS can be classified broadly into the trans-
mission mode and the total electron-yield mode. In the transmission mode, the
intensity of the x-ray is measure before and after the sample and the ratio of the
transmitted x-rays is counted. Transmission-mode experiments are standard for
hard x-rays, while for soft x-rays, they are difficult to perform because of the
strong interaction of soft x-rays with the sample and hence strong absorption.
In the present work, we have adopted the total electron-yield mode, because all
measurements have been performed in the region of soft x-rays.
2.4 X-ray magnetic circular dichroism
If circularly polarized light is used in XAS, the absorption intensity of magnetic
materials depends on the helicity of the incident light. This phenomenon has
been utilized in x-ray magnetic circular dichroism (XMCD), which is defined
as the difference in absorption spectra between right- and left-handed circularly
polarized x-rays when the helicity of x-rays are parallel and antiparallel to the
magnetization direction of the magnetic materials with a magnetic field.
The characteristic features of XMCD measurements are as follows.
• If the absorption region of an element does not overlap with other ab-
sorption region, we can study element-specific magnetic moments, which
enables us to investigate precisely the magnetism of particular orbitals of
each element.
13
• XMCD reflects the orbital and spin polarization of local electronic states.
Thus, using integrated intensity of the L2,3-edge XAS and XMCD spectra
of a transition-metal atom, one can separately estimate the values of orbital
[23] and spin [24] magnetic moments by applying XMCD sum rules.
XMCD sum rules provide the values of orbital and spin magnetic moments using
the following formulae,
Morb = −4∫
L3+L2(µ+ − µ−)dω
3∫
L3+L2(µ+ + µ−)dω
(10 − nd). (2.10)
M spin + 7MT = −6∫
L3(µ+ − µ−)dω − 4
∫L3+L2
(µ+ − µ−)dω∫
L3+L2(µ+ + µ−)dω
(10 − nd), (2.11)
where Morb and M spin are the orbital and spin magnetic moments in units of
µB/atom, respectively, µ+(µ−) is the absorption intensity for the positive (neg-
ative) helicity, nd is the d electron occupation number of the specific transition-
metal atom. L3 and L2 denote the integration range. MT is the expectation
value of the magnetic dipole operator, which is small when the local symmetry
of the transition-metal atomic site is high and is neglected here with respect to
M spin.
14
Chapter 3
Experimental
Ultraviolet photoemission spectroscopy (UPS) measurements were performed at
beamline BL-18A of Photon Factory (PF), Institute for Material Structure Sci-
ence, High Energy Accelerator Research Organization (KEK) using a VG CLAM
hemispherical analyzer. All the photoemission spectra were taken at room tem-
perature under an ultra high vacuum of 5.0 × 10−10 Torr. The Fermi level (EF)
was calibrated by the Fermi edge of a Cu metal in electrical contact with the sam-
ple. X-ray photoemission spectroscopy (XPS) measurements were performed at
University of Tokyo using a Gammadata-Scienta SES-100 hemispherical analyzer
for MnGeP2 and VSW125 analyzer for thermally diffused Mn/GaAs (001). The
total energy resolution was estimated to be ∼200 meV for UPS and ∼800 meV for
XPS including temperature broadening. In both UPS and XPS measurements,
photoelectrons were collected in the angle-integrated mode. XAS and XMCD
measurements at the Mn 2p → 3d (Mn L2,3) edge were performed at BL-11A of
PF. The XAS spectra were taken by the total electron yield mode with magnetic
fields applied perpendicular to the sample plane under an ultra high vacuum of
2.0 × 10−9 Torr. X-ray absorption spectra for right-handed (µ+) and left-handed
(µ−) circularly polarized x-rays were obtained by reversing the direction of mag-
netization. The difference between the µ+ and µ− spectra yields XMCD spectra.
Depth profile studies were performed by repeated cycles of sputter-etching and
subsequent PES measurement. Sputter-etching was done with Ar+-ion at 1.0 kV.
Here, it is necessary to eliminate possible surface effects such as oxidation of
the surface in ex situ treatment of the sample.
MnGeP2 thin films for the XAS, XMCD, and XPS measurements were coated
with a Ge-cap layer (∼3 nm). As for the UPS measurements, which are quite
surface-sensitive, the uncapped sample was used and cleaned repeatedly by Ar+-
ion sputtering at 0.8 kV with a vacuum of 4.0 × 10−5 Torr. Thermally diffused
15
Mn/GaAs (001) thin films were coated with an As-cap layer.
(a)
(b)
Figure 3.1: Schematic optical layout of (a) BL-18A at PF and (b) BL-11A at PF.
Manipulator
Photon source
Electron energy analyzer
Sample
Ion gun
Photo-electron
Ar+-ion
Figure 3.2: Schematic drawing of the experimental apparatus used for the photoemissionstudy.
16
Chapter 4
Chalcopyrite-type magnetic
semiconductor MnGeP2 thin
films
4.1 Physical properties of MnGeP2
MnGeP2 is a II-IV-V2 chalcopyrite-type room-temperature ferromagnetic semi-
conductor. The crystal structure of MnGeP2 is shown in Fig. 4.1(a). The lat-
tice parameters for MnGeP2 were estimated as a = 5.693 A, and c = 11.303
A and α = β = γ = 90 by reciprocal lattice mapping of x-ray diffraction (XRD)
shown in Fig. 4.1(b) [16], which is consistent with the experimental value of poly-
crystalline MnGeP2 [15]. Recently, Cho et al. have succeeded in synthesizing
MnGeP2 and MnGeAs2 thin films using MBE method and measured resistance
and Hall resistance as shown in Fig. 4.2(a), (b) [15]. The resistivity of MnGeP2
at room-temperature is ∼10−3 Ωcm, which is rather small for a semiconductor.
Hall measurements revealed that the carrier type of MnGeP2 and MnGeAs2 is
p- and n-type, respectively, with carrier concentration ∼1019 cm−3. Based on
the fact that various native defects such as group II vacancies and antisite de-
fects are present in II-IV-V2 chalcopyrites with carrier concentrations up to 1019
cm−3 [25, 26], p-type behavior of MnGeP2 may arise from point defects such as
cation Mn and Ge vacancies (VMn, VGe) and antisite defects MnGe. Anomalous
Hall effect was observed as shown in Fig. 4.2(b), suggesting the spin polarized
hole carriers exist in MnGeP2. In order to confirm the electrical properties, a p-n
junction has been fabricated by use of 100 A p-MnGeP2 and 100 A n-MnGeAs2
and measured I-V characteristics as shown in Fig. 4.2(c). Clear diode character-
istic was observed, indicating that MnGeP2 is certainly a p-type semiconductor.
17
Theoretical studies have also been carried out using full-potential linearized
augmented plane wave (FLAPW) method [27] in the local density approximation
(LDA). For instance, Zhao et al. derived the lattice constants of MnGeP2 as a =
5.673 A, c = 10.716 A [28], which is consistent with the value deduced by Sato
et al. [16]. Cho et al. applied total energy calculation to MnGeP2, revealing that
MnGeP2 is a semiconductor with an energy gap of ∼0.24 eV.
It is reported that MnP cluster was observed in polycrystalline Zn1−xMnxGeP2
by nuclear magnetic resonance measurement [29], and therefore we must be con-
cerned about the formation of MnP secondary phase in MnGeP2. MnP, which
is distorted from the NiAs-type structure with a = 5.916 A, b = 5.260 A, and c
= 3.173 A, is a ferromagnetic metal with T C ∼293 K [30, 31]. Hence, MnP also
contributes to the room-temperature ferromagnetism, however, a phase transi-
tion from the ferromagnetic to an antiferromagnetic state occurs at T N ∼47 K.
Therefore, the presence or the absence of MnP can be checked by investigating
the magnetization around T N.
18
Mn Ge P
a
c
a = 5.693 [A]
c = 11.303 [A]
(a) (b)
Figure 4.1: Chalcopyrite-type MnGeP2. (a) Crystal structure. (b) Reciprocal lattice mappingof x-ray diffraction for MnGeP2 grown on an InP (001) substrate [16], which provide the latticeconstants.
(c)
(a) (b)
(d)
Figure 4.2: Several physical properties and theoretical calculation for MnGeP2 and MnGeAs2thin films [15]. (a) Temperature dependence of the electrical resistances of MnGeP2 andMnGeAs2 thin films grown on GaAs (001) substrates. (b) Anomalous Hall resistances ofp-MnGeP2 and n-MnGeAs2 thin films. (c) I-V diode characteristics of a junction between 100A p-MnGeP2 and 100 A n-MnGeAs2 thin films grown on a Si (001) substrate. (d) Theoreticalcalculation of density of states (DOS) by the full-potential linearized augmented plane wave(FLAPW) method [27] in the local density approximation (LDA).
19
4.2 Sample preparation
Two single-crystal MnGeP2 thin films were fabricated by the molecular beam
epitaxy (MBE) method. One was a Ge-capped sample, and the other was without
capping. These samples were epitaxially grown on a Ge buffer layer at 435 C
which had been deposited on a GaAs (001) substrate at 380 C. It has been
reported that the Ge buffer layer enables a two-dimentional growth of MnGeP2
and improves the crystallinity of the sample [32]. We have done magnetization
measurements using a SQUID magnetometer (MPMS, Quantum Design, Co.,
Ltd.) on the MnGeP2 films prior to the spectroscopic measurements. Figure
5.9 shows magnetization curves for a magnetic field perpendicular to the sample
plane. The sample showed T C ∼320 K and thus was confirmed to be a room-
temperature ferromagnet as in the previous report [15]. Additionally the M-T
curve indicated that there was a Curie-Weiss component in the sample. This fact
suggested that a paramagnetic component existed in the sample. Details of the
sample fabrication are given in Ref. [32].
-1.5
-1.0
-0.5
0
0.5
1.0
1.5
Mag
netiz
atio
n ( µ
B/ M
n)
-2 -1 0 1 2Magnetic Field (T)
10K 100K 310K 330K
(a)
MnGeP2
0.8
0.6
0.4
0.2
0Mag
netiz
atio
n ( µ
B/ M
n)
300200100Temperature (K)
H = 0.01 T
(b)
Figure 4.3: Magnetization data of MnGeP2 thin film. (a) M -H curves at T = 10 K, 100 K,310 K, and 330 K. The linear component in the high magnetic field region has been subtracted.(b) M -T curve at H = 0.01 T. We obtained the M -H curve as follows. The sample was firstcooled down to 5 K without applying the magnetic field. Then the magnetic field was appliedperpendicular to the sample plane up to 1.0 T, and set it at 0.01 T and we started to measurethe sample with increasing temperature.
20
4.3 Results and discussion
Inte
nsit
y (a
rb. u
nits
)
15 10 5 0 -5Binding Energy (eV)
Mn Auger
h =60
57
55
54
53
52
515049484746 eV
difference (51 – 48 eV)
(a)MnGeP2
CIS
Int
ensi
ty (
arb.
uni
ts)
70656055504540Photon Energy (eV)
EB =15 eV
9.5 eV
7.5 eV
5.3 eV
2.6 eV
0.7 eV
(b)
Inte
nsit
y (a
rb. u
nits
)
12 8 4 0Binding Energy (eV)
MnGeP2
MnP
Ga0.931Mn0.069As
difference (51 eV – 48 eV)(c)
-1.0-0.500.51.0
Ga0.931Mn0.069As
MnP
MnGeP2
Figure 4.4: A series of valence-band spectra taken at various photon energies in the Mn 3p→ 3d core-excitation region. (a) Energy distribution curves. Vertical bars indicate a constantkinetic energy and represent the threshold (lower EB’s) and the peak position (higher EB’s) ofthe Mn Auger signals. The difference between the on-resonant (hν = 51 eV) and off-resonant(hν = 48 eV) spectra, which approximately represents the Mn 3d partial density of states, isshown at the bottom. (b) Constant-initial-states spectra at various binding energies. Arrowsshow the threshold and peak position of the Mn Auger signals. (c) Comparison of the difference(hν = 51−48 eV) spectra between MnP [33], Ga0.931Mn0.069As [34], and MnGeP2. Inset showsan enlarged plot around the Fermi level.
Figure 4.4(a) shows the valence-band spectra of the uncapped MnGeP2 thin
film taken at various photon energies in the Mn 3p→ 3d core-excitation region.
The intensities have been normalized to the photon flux. The vertical bars rep-
resent the threshold and the peak position of the Mn M2,3L4,5L4,5 Auger signals,
which is quite strong, indicating that there exist the itinerant nature of the Mn
3d states in MnGeP2. The intensity near EF was high as shown in Fig. 4.4(a),
but there was no Fermi edge in MnGeP2 as shown in Fig. 4.4(a). In contrast, a
clear Fermi edge has been observed in single crystal MnP [35]. Therefore, one can
conclude that MnP is not a dominant component of the sample. The resonantly
enhanced Mn 3d partial density of states (PDOS) was obtained by subtracting
the off-resonant (hν = 48 eV) spectrum from the on-resonant (hν = 51 eV) one as
shown at the bottom of Fig. 4.4(a). In order to properly carry out the subtraction,
we consider the photon energy dependence of the photoionization cross-section
of Ge 4p and P 3p. The difference spectrum showed a peak at binding energy
(EB) ∼2.6 eV. Another broad peak which appeared around EB = 6-15 eV can
be attributed to a charge-transfer satellite as in the case of Ga1−xMnxAs [34].
21
The existence of the charge-transfer satellite indicates that there exist strong
Coulomb interaction between the Mn 3d electrons and strong hybridization be-
tween the Mn 3d and other valence orbitals. However, it should be noted that
Mn M2,3L4,5L4,5 Auger signals exist in this region and overlaps with the satellite.
Figure 4.4(b) shows the constant-initial-state (CIS) spectra for various EB ’s in
the range of hν = 40-70 eV. Arrows represent the threshold and the peak posi-
tion of the Mn M2,3L4,5L4,5 Auger signals. All the CIS’s showed a peak at hν
∼51 eV (shown by a dashed line), whereas Mn M2,3L4,5L4,5 Auger overlaped with
the 51-eV peak for EB = 5-11 eV. Figure 4.4(c) shows comparison of the differ-
ence (hν = 51 − 48 eV) spectra between MnP [33], Ga0.931Mn0.069As [34], and
MnGeP2. The difference spectra of MnP show broader structure at EB = 1-4 eV
than that of MnGeP2, and there is a clear Fermi edge in the difference spectrum
of MnP, indicating that the Mn 3d states in MnP exhibit more itinerant behavior
than those in MnGeP2. On the other hand, the main peak of Ga0.931Mn0.069As
is sharper than that of MnGeP2, reflecting more localized nature of the Mn 3d
states in Ga0.931Mn0.069As than those in MnGeP2.
Inte
nsit
y (a
rb. u
nits
)
655 650 645 640 635Binding Energy (eV)
Mn 2p
2p3/22p1/2
(a)A
BA'B'
h = 1486.6 eV
MnGeP2
645 640 635Binding Energy (eV)
Mn 2p3/2
20º 60º
(b)A
B
Figure 4.5: XPS spectrum of the Mn 2p core level for MnGeP2. (a) Entire spectrum. (b)Mn 2p3/2 spectra for the take-off angle of 20 and 60 relative to the sample surface. Theintensity has been normalized to the peak height of structure B. The 20 spectrum was moresurface-sensitive than the 60 one.
Figure 4.5(a) shows the Mn 2p core level XPS spectra taken with the Al
Kα source. The spin-orbit doublet was observed and each spin-orbit component
has a main peak (A, A′) and a strong shoulder structure (B, B′), which is a
charge-transfer satellite located at a higher EB of the main peak. On the other
hand, structure A is fairly sharp, suggesting the itinerant nature of the Mn 3d
states. Figure 4.5(b) shows Mn 2p3/2 spectra for the measurement condition that
22
the sample surface forms an angle of 20 and 60 with the analyzer axis. The
20 measurement was more surface-sensitive than the 60 measurement. The
relative intensity of structure A to structure B in the 20 spectrum was weaker
than that in the 60 spectrum. This indicates that structure A derived not from
the surface region, but from the intrinsic origin of the bulk, meaning that the
itinerant character of the Mn 3d electrons is that of bulk MnGeP2 and not that
of its surface.
Figure 4.6(a), (b), and (c) show circularly polarized XAS spectra (µ+, µ−)
and XMCD spectra (µ+ − µ−) at the Mn L2,3 edge under various experimental
conditions. Circularly polarized XAS spectra have been normalized to the Mn
L3 peak height of the unpolarized XAS spectra [(µ+ + µ−)/2]. XMCD signal
for the remanence (H = 0 T), which represents the ferromagnetic component,
was clearly observed at 30 K, thus one can exclude contributions from MnP,
because a ferromagnetic-to-antiferromagnetic transition occurs at T N ∼47 K in
MnP [30, 31]. In Fig. 4.6(d), we compared the XAS spectra between MnAs [36],
Ga0.94Mn0.06As [36], and MnGeP2. Note that the XAS line shape of MnAs is
broad, which reflects the itinerant nature of the Mn 3d states in MnAs, while that
of Ga0.94Mn0.06As is narrower and exhibits multiplet structures, which reflects the
localized nature of the Mn 3d states. The XAS spectrum of MnGeP2 exhibits a
broad asymmetric line shape, which is similar to that of MnAs, and has multiplet
structures, suggesting that the Mn 3d states have both the itinerant and localized
character in MnGeP2.
In order to decompose the XMCD signals into the paramagnetic and ferro-
magnetic component, we subtracted the XMCD spectra at H = 0 T and T = 30
K (multiplied by the ratio of the saturation magnetization to the residual mag-
netization ∼1.77 at T = 30 K and ∼1.90 at T = 200 K) from those at H = 5.0
T and T = 30 K and 200 K in Fig. 4.7(a). It turns out that the spectral features
of the paramagnetic component at T = 30 K and 200 K [para (30 K), para (200
K)] are quite similar to each other, and are clearly different from those of the
ferromagnetic component at T = 30 K [ferro (30 K)]. The orbital (Morb), spin
(M spin), and total (M tot = Morb + M spin) magnetic moments of the paramag-
netic and ferromagnetic component were estimated by applying the XMCD sum
rules, and plotted as functions of temperature in Fig. 4.7(b). While Morb/M spin
for the paramagnetic component is ∼0.3-0.5, Morb/M spin for the ferromagnetic
component at 30 K is tiny. Therefore, a large orbital magnetic moment is thought
to be realized by the paramagnetic component.
23
-0.20-0.15-0.10-0.0500.05
XM
CD
(arb. units)
660640Photon Energy (eV)
1.00.80.60.40.2
0
Abs
orpt
ion
(arb
. uni
ts)
660640 660640
T = 200 KH = 5.0 T
+
–
+
–
T = 30 KH = 5.0 T
–
+
T = 30 KH = 0 T
+–
–
(a) (b) (c) Mn L2,3 edgeMnGeP2A
bsor
ptio
n (a
rb. u
nits
)
660655650645640635Photon Energy (eV)
MnGeP2
Ga0.96Mn0.04As
MnAs
Mn L2,3 edge(d)
Figure 4.6: Circularly polarized x-ray absorption spectra (top panels) and XMCD spectra(bottom panels) at the Mn L2,3 edge under various experimental conditions. (a) T = 200 K,H = 5.0 T. (b) T = 30 K, H = 5.0 T. (c) T = 30 K, H = 0 T (remanence). (d) Comparisonof the XAS spectra between MnAs [36], Ga0.94Mn0.06As [36], and MnGeP2.
24
Inte
nsit
y (a
rb. u
nits
)
665660655650645640635630Photon Energy (eV)
(a) MnGeP2
Mn L2,3 edge
ferro (30 K)
para (30 K)
para (200 K)
0.6
0.4
0.2
0
Mag
netic
Mom
ent (
B/M
n)
200150100500Temperature (K)
Morb(para)
Mspin(para)
Mtot(para)
(b)
Morb(ferro, residual)
Mspin(ferro, residual)
Mtot(ferro, residual)
Figure 4.7: Analysis of the XMCD spectra of MnGeP2. (a) Ferromagnetic component at 30K, and paramagnetic component at 30 K, 200 K of the XMCD spectra. The paramagneticcomponent has been deduced by subtracting the ferromagnetic component (H = 0 T, T = 30K) multiplied by the ratio of the saturation magnetization (M sat) to the residual magnetization(M res) from the XMCD spectra at H = 5.0 T, T = 30 K and 200 K. (b) Temperature depen-dence of the magnetic moments M spin, Morb, and M tot = Morb + M spin of the ferromagneticcomponent and the paramagnetic component estimated using the XMCD sum rules. The XASsignals were not decomposed into the paramagnetic and ferromagnetic component, thus thevalue of the vertical axis is rather underestimated.
25
Chapter 5
Thermally diffused Mn/GaAs
(001) thin films in depth profile
5.1 Physical properties of thermally diffused
Mn/GaAs (001)
Figure 5.1 shows a schematic illustration of the Mn incorporation into GaAs (001)
substrate by thermal diffusion. The thermal diffusion procedures can be classified
into “pre-annealing” procedure [Fig. 5.1(a)] and “post-annealing” [Fig. 5.1(b)]
one. We define “pre-annealing” as the procedure where Mn metal is deposited on
the annealed substrate. In contrast, “post-annealing” procedure means that Mn
metal is deposited on the unannealed substrate and then annealed. I would like
to introduce several reports for the thermally diffused Mn/GaAs (001) prepared
by both “pre-annealing” and “post-annealing” procedure.
Dong et al. have performed a depth profile study of thermally diffused Mn/GaAs
(001) thin films prepared by “pre-annealing” procedure using XPS [37]. Figure
5.2 shows relative core-level intensities as functions of sputtering time under var-
ious conditions. Note that there existed Mn diffusion layer, where percentage of
Mn decreased with increasing sputtering time at TS = 300 K, namely, the diffu-
sion occurred even at TS = 300 K (room temperature) as shown in Fig. 5.2(a),
suggesting a large diffusion coefficient. It was found that the metallic Mn region,
where the percentage of Mn was ∼100 % and those of Ga and As were ∼0 %,
that is, deposited Mn was not reacted with the GaAs substrate, existed at TS =
300 and 400 K [Fig. 5.2(a), Fig. 5.2(b)], while metallic Mn region did not exist
at TS = 450 K [Fig. 5.2(c)], indicating that deposited Mn was fully reacted with
the GaAs substrate.
The diffusion coefficients of the thermally diffused Mn(200 nm)/GaAs (001)
27
GaAs (001) substrate
Annealing temperature (T [oC])
deposited Mn
Mn diffusion into bulk
d [A]
GaAs (001) substrate
Substrate temperature (TS [oC])
Mn
Mn diffusion into bulk
d [A]
(b)
(a)
Annealing time (t [min])
Figure 5.1: Schematic illustration of the Mn incorporation into the GaAs (001) substrate bythermal diffusion. The procedure of the thermal diffusion can be classified broadly into twocategories. (a) Depositing Mn metal on the annealed substrate. (b) Depositing Mn metal onthe unannealed substrate and then annealing.
(c)
(b)(a)
Figure 5.2: Relative core-level intensities (Mn 2p, Ga 2p3/2, As 2p3/2) as functions of sput-tering time under various conditions [37]. (a) the thickness of Mn (d) is 12.5 nm, substratetemperature (TS) is 300 K. (b) d = 14.5 nm, TS = 400 K. (c) d = 8.5 nm, TS = 450 K.
28
(e)
(f)
(200 nm)
Figure 5.3: RBS spectra with a 2.3 MeV He+ ion beam for Mn(200 nm)/GaAs (001) thick filmsof as-grown sample(a), and sample with post-annealing up to 300 C for 1(b), 8(c), and 16 h(d)[38]. The simulations of the RBS spectra were done using Rutherford universal manipulationprogram (RUMP) [39]. The dashed curves are calculated spectra, which was derived on theassumption that the composition of the Mn diffusion layer is uniform (Mn0.6Ga0.2As0.2). Theindividual contributions of Mn, Ga, and As of the calculated spectra are also shown. Thevertical dashed lines represent the energies where each element reaches the sample surface. (e)Thickness of the Mn diffusion layer estimated by Fig. 5.3(a)-(d) as a function of square rootof annealing time at 275, 300, and 325 C. The horizontal dashed line represents the point ofcomplete Mn reaction [40]. (f) Temperature dependence of the diffusion coefficients determinedfrom the slope of (e) [40].
29
thick films were obtained by Rutherford backscattering spectroscopy (RBS) [38,
40]. Figure 5.3 shows the RBS spectra for Mn/GaAs (001) of as-grown sample(a),
and sample with post-annealing up to 300 C for 1(b), 8(c), and 16 hours(d).
The calculated spectra can be obtained using Rutherford universal manipulation
program (RUMP) [39] and decomposed into Mn-, Ga-, and As-derived spectra,
and one can see the time evolution of the Mn diffusion from the decomposed
calculated spectra. The thickness of the Mn diffusion layer was estimated by
Fig. 5.3(a)-(d) as a function of square root of annealing time at 275, 300, and
325 C as shown in Fig. 5.3(e). One can determine the temperature dependence
of the diffusion coefficients from the slope of Fig. 5.3(e) as shown in Fig. 5.3(f),
indicating that the diffusion coefficient exponentially increased with increasing
temperature.
30
5.2 Sample preparation
Thermally diffused Mn/GaAs (001) thin films were fabricated by the MBE method
as follows. First, an ordered GaAs buffer layer was grown on an epi-ready Si-
doped n+-GaAs (001) substrate with carrier concentration ∼1020 cm−3. The
GaAs buffer layer was obtained by heating the GaAs substrate up to 600 C
during the Ga and As depositions. Then, Mn metal was deposited on the GaAs
buffer layer at the rate of 0.029 nm/s as thick as ∼2 nm. The growth temper-
ature of Mn on the GaAs substrate was set at 50 C. After the Mn deposition,
Mn diffusion into the GaAs substrate was achieved by post-annealing the sam-
ple up to 600 C for 10 min. The growth procedure is illustrated in Fig. 5.4(a).
Surface reconstruction was monitored during the growth and the post-annealing
with reflection high-energy electron diffraction (RHEED), which showed (2×2)
pattern after the post-annealing as shown in Fig. 5.4(b). We prepared uncapped
samples of the same growth condition for PES to study the surface morphology
of the Mn/GaAs (001) thin film. Figure 5.5(a)-(c) shows the surface morphology
of the as-grown and thermally diffused (300 C, 600 C) Mn/GaAs (001) thin
films measured by atomic force microscopy (AFM). The AFM images suggested
that the level of the surface roughness decreased with increasing temperature. In
particular, a fairly flat surface was obtained for the thermally diffused (600 C)
Mn/GaAs (001) thin film.
31
(a)
(b)
Figure 5.4: The growth of thermally diffused Mn/GaAs (001) thin films. (a) The growthprocedure. (b) RHHED patterns in various growth steps. (2×2) surface reconstruction wasobserved down to 150 C.
(a) as-grown (b) 300 oC for 10 min (c) 600 oC for 10 min
Mn/GaAs (001)
Figure 5.5: Atomic force microscopy images for (a) as-grown sample. (b) Sample with post-annealing up to 300 C for 10 min. (c) Sample with post-annealing up to 600 C for 10 min.
32
5.3 Results and discussion
Figure 5.6(a) shows the core-level XPS spectra in a sputter-etching series, and
Fig. 5.6(b) shows their relative core-level intensities as functions of sputtering
time. The Mn 2p peak completely disappeared at t = 210 min, thus we consider
that Mn diffusion layer disappeared and the GaAs substrate appeared at t = 210
min. At this stage we set the ratio of Ga 2p3/2 to As 2p3/2 at 1.0. Throughout
the sputter-etching series, the Mn 2p spectra showed non-metallic signals, which
indicated that the Mn layer entirely reacted with the GaAs substrate and that
Mn was not metallic even in the first surface layer. In the early stage of the
sputtering (t = 0-10 min), one can see As-excess composition and a high EB of
the Mn 2p3/2 peak positions, which may be both attributed to the existence of
As-cap layer. One can also see that the As intensity rapidly decreased while the
Mn 2p3/2 position rapidly shifted to lower EB ’s for t = 0-20 min, indicating the
removal of the As-cap layer. We suspect that some kind of unexpected reaction
between As-cap layer and Mn layer took place in the surface region. We therefore
conclude that the peak energy shift of the Mn 2p3/2 core-level can be explained as
a change of the chemical species from the Mn-As compound found in the surface
region to the Mn diffusion layer. The ratio of the Ga 2p3/2 to As 2p3/2 intensities
became ∼1 after 20 min sputtering, and therefore the As-cap layer was thought
to be removed and the Mn diffusion layer appeared from this point. Except for
the surface region, the intensities of the Mn 2p3/2 changed so slowly that Mn
diffused into deep region where a dilute Mn phase existed. This is consistent
with the large diffusion coefficient for thermally diffused Mn/GaAs reported by
the Rutherford backscattering studies (see Section 5.1) [38,40].
Figure 5.6(c) shows the Mn 2p spectrum of Mn/GaAs (001) sputtered for
40 min, compared to that of Ga0.926Mn0.074As [41]. As for Mn/GaAs (001), the
spin-orbit doublet was observed and each spin-orbit component had a main peak
and a strong shoulder structure, which is a charge-transfer satellite located at
a higher EB of the main peak. The existence of the charge-transfer satellite
indicates that there are strong Coulomb interaction between the Mn 3d electrons
and strong hybridization between the Mn 3d and other valence orbitals. This
structure is quite similar to that of Ga0.926Mn0.074As (shown by vertical bars in
Fig. 5.6(c)), and therefore the Mn 3d state is expected to be basically localized
in the dilute Mn phase.
Figure 5.7 shows a series of valence-band spectra for various sputtering time
taken at various photon energies in the Mn 3p→ 3d core-excitation region. The
intensities have been normalized to the photon flux. The vertical bars represent
33
Inte
nsit
y (a
rb. u
nits
)
660 655 650 645 640 635
Mn 2p
t(min) = 0
5
10152030406090150210
(a)Mn/GaAs (001)
45 42 39
As 3d
1326 1322
As 2p3/2
1120 1115
Ga 2p3/2
21 19 17
Ga 3d
Inte
nsit
y (a
rb. u
nits
)
660 650 640Binding Energy (eV)
Ga0.926Mn0.074As
Mn/GaAs (001) (t = 40 min)
Mn 2p(c) 2p3/2
2p1/2Mn Auger1.0
0.8
0.6
0.4
0.2
0Rel
ativ
e C
hem
ical
Com
posi
tion
200150100500Sputtering Time (min)
(b)As 2p3/2
Ga 2p3/2
Mn 2p3/2×5
Mn/GaAs (001)
1.0
0.5
020151050
As 2p3/2
Ga 2p3/2
Mn 2p3/2×5
Binding Energy (eV)
Figure 5.6: Core-level XPS spectra in the sputter-etching series. (a) Mn 2p, Ga 2p3/2, As2p3/2, Ga 3d, and As 3d spectra. t(min) denotes the sputtering time. (b) Relative core-levelintensities as functions of sputtering time. Inset shows an enlarged plot for t = 0-20 min. Itshould be noted that the relative chemical composition of Mn 2p3/2 is rather approximate. (c)Comparison of the Mn 2p core-level spectra between Mn/GaAs (001) sputtered for 40 min andGa0.926Mn0.074As [41].
34
Inte
nsity
(ar
b. u
nits
)
15 10 5 0
difference (50 – 48 eV)
t = 0 (min)
47
48
495051
52
55 eV
Mn/GaAs (001)
h =
(a)
15 10 5 0
difference(50 – 48 eV)
t = 5 (min)
h =
47
484950
51
52
55 eV
As Auger
(b)
15 10 5 0
difference(50 – 48 eV)
t = 10 (min)
h =
47
48
49
50
51
52
55 eV
As Auger
(c)
15 10 5 0
t = 85 (min)
difference(50 – 48 eV)
h =
47
48
49
50
51
52
55 eV
As Auger
(d)
15 10 5 0
(e)
h =
55 eV53
51
5049
48
46
As Auger
difference(50 – 48 eV)
Ga0.931Mn0.069As
Binding Energy (eV)
Figure 5.7: A series of valence-band spectra in the sputter-etching series of as-grown(a), 5 min-sputtered(b), 10 min-sputtered(c), 85 min-sputtered Mn/GaAs (001)(d), andGa0.931Mn0.069As(e) [34] taken at various photon energies in the Mn 3p→ 3d core-excitationregion. Vertical bars represent the As Auger signal. The difference between the on-resonant(hν = 50 eV) and off-resonant (hν = 48 eV) spectra, which approximately represents the Mn3d partial density of states, is shown at the bottom of each panel.
Inte
nsity
(ar
b. u
nits
)
15 10 5 0
0
t(min)
5
=
10152535455585
h = 51 eV
(a)
Mn/GaAs (001)
12 8 4 0
difference@10 min (50 – 48 eV)
difference@51 eV (t = 10 – 85 min)
(b)
difference@51 eV (t = 15 – 85 min)
Binding Energy (eV)
Figure 5.8: Comparison of Mn 3d PDOS. (a) Valence-band spectra for hν = 51 eV in thesputter-etching series. The photon energy is fixed at hν = 51 eV. t(min) denotes the sputteringtime. (b) Comparison of Mn 3d PDOS deduced from the difference between the spectra at t =10 min and 85 min, the spectra at t = 15 min and 85 min, and the spectra at hν = 48 eV and50 eV.
35
the expected energy position of the As-derived Auger signals. Throughout the
depth profile study, the Fermi edge was not observed, reflecting the non-metallic
character of Mn 3d electrons, consistent with the XPS result. Although the
valence-band spectra drastically changed during the first 10 minutes, there was
little change after that. Hence, we conclude that the spectra of the first 10
minutes correspond to those of As-cap layer and the subsequent spectra represent
those of the dilute Mn phase. The resonantly enhanced Mn 3d partial density of
states (PDOS) has been obtained by subtracting the off-resonant (hν = 48 eV)
spectrum from the on-resonant (hν = 50 eV) one (shown at the bottom of each
panel). In order to properly carry out the subtraction, we considered the photon
energy dependence of the photoionization cross-section of As 4p. Throughout
the measurement, Mn-derived Auger was not observed and the Mn 3d PDOS
was suppressed near EF, indicating the localized nature of the Mn 3d states.
There was a main peak at EB ∼3.7 eV and a broad and strong charge transfer
satellite at EB = 5-13 eV in the Mn 3d PDOS at t = 10 min, whose feature was
similar to that of Ga1−xMnxAs [34]. One can see a slight energy difference in the
main peak between Mn/GaAs (001) (EB = 3.7 eV) and Ga1−xMnxAs (EB = 4.5
eV), possibly due to the different positions of the EF caused by the sputtering.
The Mn 3d PDOS completely disappeared at t = 85 min, and thus we considered
that the Mn diffusion layer disappeared and the GaAs substrate appeared at t =
85 min. The absence of Mn at t = 85 min was also confirmed by the absence of
Mn 2p core-level signal (not shown). We have also attempted to obtain the Mn
3d PDOS by subtracting the spectrum at t = 85 min (GaAs) from the one at t =
10 min, 15 min (Mn/GaAs in the dilute Mn phase) for the fixed photon energy
of hν = 51 eV as shown in Fig. 5.8(b). These difference spectra are similar to
those obtained from the RPES measurement.
We performed a magnetization measurement, using a SQUID magnetometer
(MPMS, Quantum Design, Co., Ltd.) for the sample which had been sputtered
for 30 min by the same ion gun as that used for the XPS measurements, that is,
the sputtered Mn/GaAs (001) in the dilute Mn phase. A tiny hysteresis was ob-
served as shown in Fig. 5.9(a), indicating that the sample exhibits ferromagnetic
behavior in the dilute Mn phase.
36
-1.0
0
1.0
Mag
netiz
atio
n (1
0-6em
u/m
m2 )
-1.0 -0.5 0 0.5 1.0Magnetic Field (T)
Mn/GaAs (001)
T = 5 K
(a)t = 30 (min)
-0.4
-0.2
0
0.2
0.4
.
0.01-0.01 0
0.05
0.04
0.03
0.02
0.01
0
Mag
netiz
atio
n (1
0-6em
u/m
m2 )
300200100Temperature (K)
Mn/GaAs (001)
H = 0.005 T
(b)
t = 30 (min)
Figure 5.9: Magnetization curves of thermally diffused Mn/GaAs (001) thin film sputteredfor 30 min. (a) M -H curve at T = 5 K. The linear component in the high magnetic field regionhas been subtracted. The inset shows an enlarged plot for H = −0.01-0.01 T. (b) M -T curvefor H = 0.005 T (zero-field cooled).
37
Chapter 6
Summary
In this thesis, we have studied the Mn-based magnetic semiconductors MnGeP2,
and thermally diffused Mn/GaAs (001) using high-energy spectroscopic tech-
nique.
In Chapter 4, we have investigated the electronic structure and the mag-
netic properties of the II-IV-V2 chalcopyrite-type room-temperature ferromag-
netic semiconductor MnGeP2 thin films by PES, XAS, and XMCD measure-
ments. All the spectra indicated intermediate electronic states between the itin-
erant and localized character of the Mn 3d states in MnGeP2. It has become
clear from the XMCD measurement that a paramagnetic component coexisted
with a ferromagnetic component and had a large orbital magnetic moment.
In Chapter 5, we have performed the depth profile study of the thermally
diffused Mn/GaAs (001) thin films using PES combined with Ar+-ion sputtering
to investigate the electronic structure of the sample along the depth direction.
We confirmed that Mn was thermally diffused into the GaAs substrate into the
deep region, and completely reacted with the the GaAs, consistent with the large
diffusion coefficient predicted by previous report [40]. The Mn 2p core-level and
Mn 3d valence-band spectra of thermally diffused Mn/GaAs (001) in the dilute
Mn phase are similar to those of Ga1−xMnxAs, indicating that the Mn 3d states
are well localized. Ferromagnetism was observed even in the dilute Mn phase.
39
Acknowledgements
It is my great pleasure to express my special gratitude to the following people
for their help concerning my master thesis.
First of all, I would like to express my heartfelt gratitude to Prof. Atsushi Fuji-
mori, who has given me a lot of attentive guidance and valuable advice throughout
this work. I have always got an impression by his knowledge and foresight in the
field of condensed matter physics. Thanks to his pertinent advice, I have been
able to go ahead with my investigation in an efficient manner so far. I also thank
to Prof. Takashi Mizokawa for his instructive advice about the interpretation of
the experimental results.
The experiments at Photon Factory were supported by a number of people.
I am indebted to the members of Kinoshita group, Dr. Taichi Okuda, Dr. Ayumi
Harasawa, Dr. T. Takanori Wakita, and Prof. Toyohiko Kinoshita, for their valu-
able technical support during the beamtimes at BL-18A of Photon Factory. I
also acknowledge Dr. Kazutoshi Mamiya, and Prof. Tsuneharu Koide for their
vital technical support and fruitful discussions about XMCD spectra during the
beamtimes at BL-11A of Photon Factory.
I am very grateful to the members of Sato Group. Mr. Kazuyuki Minami,
Dr. Takayuki Ishibashi, and Prof.Katsuaki Sato willingly provided me with the
high-quality samples of the MnGeP2 thin films.
I am deeply thankful to the members of Oshima Group. Mr. Ken Kanai,
Mr. Kotaro Kubo, Dr. Jun Okabayashi, and Prof.Masaharu Oshima, gave me
the invaluable opportunities to grow the thermally diffused Mn/GaAs (001) thin
films using MBE and perform AFM measurements.
I like to thank the members of Uchida Group for giving me helpful advice to
perform MPMS measurements. In particular, Dr. Kenji Kojima gave me a lot of
useful guidance for the maintenance of the SQUID magnetometer.
I would like to thank the members of Fujimori-Mizokawa Group. Mr. Yukiaki
Ishida always gave me suggestive advice about the interpretation for experimental
results. His attitude for research has inspired me to greater efforts. Mr. Jong-Il
41
Hwang taught me how to use and maintain photoemission instruments. I had
much to learn from his deep insight about DMS’s. Mr. Kazuaki Ebata gave me a
helping hand with my work when I was in trouble. Mr. Masaki Kobayashi gave
me valuable comments on my experimental results and helped me at Photon
Factory. Mr. Masaru Takizawa gave me a lot of useful advice about the neces-
sary things to analyze the experimental data. Mr. Yasuhiro Ooki helped me with
the maintenance of photoemission instruments and experiments. I also wish to
thank Dr. Teppei Yoshida, Dr. Kiyohisa Tanaka, Mr. Hajime Yagi, Mr. Hiroki Wa-
dati, Mr. Makoto Hashimoto, Mr. Masaki Ikeda, Mr. Takashi Maekawa, Dr. James
Quilty, Dr. Jin-Yong Son, Mr. Daishuke Asakura, Mr. ThangTrung Tran, Mr. Akira
Shibata, Mr. Yasuhiro Fujii, Mr. Kou Takubo, and Ms. Ayako Fukuya for their
cordial supports.
Finally, I would like to express my special thanks to my friends and family.
My friends have always encouraged me to bring my master thesis to completion.
My family has supported me mentally as well as financially all the time. I cannot
be too appreciative of their kindness.
Kashiwa, Chiba
January 2006
Yoshitaka Osafune
42
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