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Magnetoelectric Fe 2TeO 6 thin films
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Journal of Physics: Condensed Matter J. Phys.: Condens. Matter 26 (2014) 055012 (9pp)
doi:10.1088/0953-8984/26/5/055012
Magnetoelectric Fe2TeO6 thin films Junlei Wang1 , Juan A Colón Santana2 , Ning Wu1 , Chithra Karunakaran3 , Jian Wang3 , Peter A Dowben1 and Christian Binek1 1
Department of Physics and Astronomy, University of Nebraska, Theodore Jorgensen Hall, 855 North 16th Street, PO Box 880299, Lincoln, NE 68588-0299, USA 2 Department of Electrical Engineering, W Scott Engineering Center, University of Nebraska, North 16th Street, Lincoln, NE 68588-0656, USA 3 Canadian Light Source, University of Saskatchewan, 44 Innovation Boulevard, Saskatoon, SK, S7N 2V3, Canada E-mail: [email protected] Received 22 August 2013, in final form 5 November 2013 Published 17 January 2014
Abstract
We demonstrate that Fe2 TeO6 is a magnetoelectric antiferromagnet with voltage-controllable boundary magnetization. This provides experimental evidence of the theoretical prediction that boundary magnetization is a universal property of magnetoelectric antiferromagnets including boundary magnetization at a surface orthogonal to the polar direction. Highly (110) textured Fe2 TeO6 thin films were grown by pulsed laser deposition, terminating in Te–O at the (110) surface due to a surface reconstruction. Magnetic dc susceptibility measurements of both Fe2 TeO6 powder and thin film samples confirm antiferromagnetic long-range order. Finally, measurements of x-ray magnetic circular dichroism combined with photoemission electron microscopy (XMCD–PEEM) provide a lower bound to the spin and angular magnetic moment of the surface Fe ions. Keywords: boundary polarization, magnetoelectric, XMCD–PEEM, antiferromagnetic domains S Online supplementary data available from stacks.iop.org/JPhysCM/26/055012/mmedia (Some figures may appear in colour only in the online journal)
1. Introduction
magnetization is scarce and microscopic evidence has only been provided for the Cr2 O3 (0001) surface using x-ray magnetic circular dichroism–photoemission electron microscopy (XMCD–PEEM) [5], spin-polarized inverse photoemission [5] and spin-polarized photoemission [1]. In order to bring the concept of voltage-controlled boundary magnetization into a broader experimental context, and demonstrate general applicability, evidence for voltage-controllable boundary magnetization is needed from additional magnetoelectric antiferromagnetic insulators. Fe2 TeO6 is another material that has been demonstrated to be both antiferromagnetic and magnetoelectric [6], and therefore is expected to show boundary magnetization at surfaces and interfaces, based upon general theoretical considerations [3, 4]. Motivated by the prospect of largely dissipationless voltage-controlled interface magnetization in future ultralow power post-CMOS technology, we report on thin film
Voltage-controlled boundary magnetization of dielectric materials, such as in Cr2 O3 based voltage-controlled exchange bias systems [1, 2], has the potential to revolutionize spintronics. Boundary magnetization is a roughness insensitive net magnetization which emerges at the surface or interface of a magnetoelectric antiferromagnet in a single-domain state [1, 3]. The bulk linear magnetoelectric effect is quantified by the magnetoelectric susceptibility tensor, αi j . This tensor determines the magnetization induced in linear response due to the presence of an applied electric field. Symmetry arguments reveal that equilibrium boundary magnetization is a generic property of magnetoelectric antiferromagnets [3, 4], that may be obtained by magnetoelectric field cooling, with simultaneously applied electric and magnetic fields, thus utilizing the magnetoelectric effect to select an antiferromagnetic singledomain state. However, experimental evidence of boundary 0953-8984/14/055012+09$33.00
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Figure 2. Temperature dependence of the magnetic moment (open
circles) of the powder sample and deduced temperature dependence of the parallel component of the magnetic susceptibility (solid circles). Vertical arrow indicates N´eel temperature.
Figure 1. X-ray diffraction pattern (dots) with the calculated profile
(line) of Fe2 TeO6 . The peak positions and difference pattern are shown at the bottom of the chart. Major peaks are labeled.
general trirutile input structure of A2 BO6 compounds [9], provides the lattice parameters indexed for a tetragonal cell, with space group P42 /mnm, and atomic positions, as summarized in table 1. The temperature dependence (5 K < T < 300 K) of the field-induced magnetization was measured for approximately 0.1 cm3 of the Fe2 TeO6 powder sample in an applied magnetic field of 0.6 T, after zero-field cooling using a superconducting quantum interference device, SQUID (Quantum Design, MPMS XL). Figure 2 shows the field-induced magnetic moment (open circles), m, versus temperature, T. A decrease of the moment sets in at T < 230 K, indicating the onset of antiferromagnetic order [10, 11]. The N´eel temperature, TN , is determined from the inflection point in m versus T close to the maximum at 230 K. The antiferromagnetic long-range order establishes with mutually compensating sublattice magnetization. Our data suggest that TN ≈ 210 K, which is in good agreement with the prior literature value [6]. Due to the random orientation of the Fe2 TeO6 crystallites in the powder sample, the magnetization is isotropic and the susceptibility, hχi = ∂ M/∂ H ≈ M/H , is a weighted superposition of the parallel and perpendicular susceptibilities, χk , and χ⊥ , according to hχi = χk /3 + 2χ⊥ /3. Using χ⊥ (T < TN ) ≈ const [12] and χk (T TN ) χ⊥ (T < TN ), we extracted χk for T < TN (the susceptibility χk is plotted as the solid circles in figure 2) from hχ i versus T. One could use the assumption of low anisotropy, which would allow one to continue extrapolation of the parallel component of the susceptibility, χk to above TN using the measured hχ i(T ) for T > TN based on the mean-field approximation χ⊥ (T > TN ) = χk (T > TN ). This assumption is not adopted as we prefer to keep the number of assumptions at a minimum and rather not to deduce the χk values for T > TN . Close to TN , the T-dependence of the parallel magnetoelectric susceptibility αk is given by αk = 8λvµ0 Bη χk . Here η is the antiferromagnetic order parameter, v0 is the volume of the unit cell, µB is the Bohr magneton, and λ is the coupling strength which is determined by the compound’s spin
growth of the magnetoelectric antiferromagnet Fe2 TeO6 , a representative in a class of materials where dissipationless switching of boundary magnetization is feasible through voltage-control. In this work, we discuss the growth of highly (110) textured Fe2 TeO6 thin films utilizing pulsed laser deposition (PLD), and compare some of the magnetic properties in both powder and thin film samples. In addition, we provide direct evidence of voltage-controlled boundary magnetization in Fe2 TeO6 via XMCD–PEEM. Our results strongly support the generality of the concept of boundary magnetization in antiferromagnetic magnetoelectrics and demonstrate that chromia is not an exclusive material for voltage-controlled boundary magnetization. Indeed, we show that the boundary magnetization for a magnetoelectric can occur at a surface orthogonal to the polar direction. 2. Synthesis and characterization of Fe2 TeO6 powder
The starting point for thin film growth is the synthesis of fine Fe2 TeO6 powders, which we used to prepare Fe2 TeO6 targets by successive sintering and annealing. The synthesis of fine Fe2 TeO6 powders was done via a solid state reaction routine [6]. The synthesized powder was characterized by x-ray diffraction confirming the correct chemical structure of the sample. A Rigaku D/Max-B diffractometer using monochromatized Cu Kα radiation was used to record diffraction intensity in the 2θ range from 15◦ to 100◦ in steps of 0.02◦ scanning with speed of 0.2◦ min−1 . Figure 1 shows the x-ray diffraction pattern of Fe2 TeO6 powder and is in good agreement with the x-ray diffraction data reported in the literature [7]4 . In order to derive the structure, Rietveld refinement of the x-ray diffraction pattern was conducted utilizing the Bruker AXS DIFFRAC TOPAS 4.2 software [8]. The Rietveld least squares profile refinement, based on the 4
International Centre for Diffraction Data (ICDD) Powder Diffraction File (PDF) 04-009-3444, and 01-070-7443. 2
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Table 1. Lattice parameters and crystal parameters describing atomic positions within the unit cell obtained by Rietveld refinement of the
x-ray diffraction pattern. Wavelength (nm) a (Å) c (Å) Space group Crystal density (g cm−3 ) Rexp (%) Rwp (%) Goodness of fit Crystal parameters
0.154 06 4.605 98(95) 9.091 58(74) P42 /mnm (#136) 5.773 8.20 10.24 1.25
Atom Fe Te O1 O2
y 0 0 0.303 61(68) 0.305 43(46)
Site 4e 2a 4f 8j
x 0 0 0.303 61(68) 0.305 43(46)
z 0.334 94(12) 0 0 0.336 26(37)
B (Å2 ) 0.368(69) 0.624(67) 0.71(17) 0.58(12)
structure [13–16]. Since the T-dependence of αk is known for bulk Fe2 TeO6 [6] and we measured the T-dependence of χk we are able to deduce η ∝ (TN − T )β with β = 0.37. The value of β, the critical exponent of the order parameter, indicates 3D Heisenberg type criticality. Certainly, additional accurate data in the critical regime are needed to ultimately conclude on the critical behavior. It is tantalizing to correlate the potential similarity in the criticality of chromia and Fe2 TeO6 since the measurements of the critical exponent for Fe2 TeO6 are remarkably similar to the critical behavior of the archetypical magnetoelectric antiferromagnet chromia [17, 18]. 3. Growth of Fe2 TeO6 thin films and structure characterization below the surface
The growth of the Fe2 TeO6 thin films starts with preparation of targets, which are not commercially available at present. The target was prepared by pressing synthesized Fe2 TeO6 powder into a pellet with a hydraulic and cold isostatic press (CIP KJY s150-300). The CIP processed pellets were taken back to the furnace for annealing at 975 K using a ramping speed of 2 K min−1 . Our targets serve for both RF sputtering and pulsed laser deposition. Attempts to grow Fe2 TeO6 thin films by sputtering (ATC 2000-F, AJA International) resulted in only amorphous films. Some improvement towards formation of a polycrystalline material could be achieved through post-growth annealing, as indicated by (110) and (103) peaks in the XRD θ –2θ scan (not shown). The alternative, pulsed laser deposition (PVD MBE/PLD 2000), was employed to grow thin films of Fe2 TeO6 , and led to the growth of superior films. The energy of the KrF excimer laser was set to 130 mJ per pulse with 10 Hz pulse rate in order to prevent the target surface from overheating. During growth the target was kept at a distance of 10 cm relative to the substrate. The Al2 O3 (0001) substrate was maintained at 300 ◦ C with oxygen partial pressure environment, set to 10 mbar to keep the material stable. A deposition time of 1 h was used to achieve relatively thick films (nominally several hundred nanometers) for characterization and further analysis.
Figure 3. X-ray diffraction of the PLD grown sample (thick line)
with the sapphire substrate pattern (thin line). The prominent (110) Bragg peak of the Fe2 TeO6 thin film has been labeled.
Figure 3 shows the XRD pattern of the PLD grown Fe2 TeO6 /Al2 O3 (0001) sample, and in sharp contrast to the sputtered thin films, there is a single, pronounced, and narrow (110) peak present in the wide-angle diffraction pattern. The presence of the single sharp Bragg peak indicates that the thin film is highly (110) textured. The PLD grown thin films were used for all further analysis. 4. Magnetic characterization of (110) Fe2 TeO6 thin film
Magnetic susceptibility data for Fe2 TeO6 are scarce and to the best of our knowledge only available for powder samples [11, 19, 20] and this limits the comparisons available. There is a further complication in that susceptibilities of these antiferromagnetic thin films are challenging to measure due to the unfavorable film-to-substrate signal ratio. The sapphire substrate (0.5 mm × 3.9 mm × 6.8 mm) creates a diamagnetic background of the order of 10−4 emu = 10−7 A m2 in an 3
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Figure 4. (a) Linear combination 1D versus T (squares) of two m versus T data sets measured in orthogonal in-plane magnetic fields of µ0 H = 0.9 T applied along the respective sample edges. (b) Circles (left axis) show the sum of squared residuals S versus Tmin resulting from least squares fits of a Curie–Weiss law to 1D versus T data when systematically varying the fitting interval (see text). Upper lines are 1 versus T linear interpolations of the S versus Tmin data. Their intercept is used as a marker for TN (vertical arrow). Triangles show 1D (right axis). A corresponding linear fit (lower line, see text) is extrapolated towards interception with the T-axis. Onset of deviation of the data from linear behavior is used as a marker for TN .
applied magnetic field of µ0 H = 0.9 T or about two orders of magnitude larger than the signal for a Fe2 TeO6 thin film with a thickness of a few hundred nanometers. In addition, the temperature dependence of the combined substrate and sample holder signal complicates the analysis. The standard approach to address this problem is to measure the magnetic behavior of the substrate separately and then subtract the background signal from the overall sample signal. Note, however, that with this methodology it is not possible to measure the genuine substrate on which the film has been deposited. In cases where the ratio of film-to-substrate signal is 10−2 at most, it is more appropriate to utilize a method which allows subtraction of the background signal of the genuine substrate and holder in the presence of the magnetic film. In order to overcome the problems outlined above, we measured the temperature dependence of the magnetic moment of our sample in two different geometries. The measurements took place with an in-plane applied magnetic field of µ0 H = 0.9 T parallel and orthogonal to the (001) axis (as established by reflection high-energy electron diffraction: see supplementary material available at stacks.iop.org/JPhysCM/26/055012/ mmedia). Hence, the two orthogonal m versus T measurements probe the parallel and perpendicular susceptibilities of the antiferromagnet. The measurements lead to data sets D1,2 , where D1 = m k + bs + bother and D2 = m ⊥ + bs + αbother . Here bs (T ) is the diamagnetic background of the sapphire substrate and bother (T ) is a background signal, which depends on the respective conditions of sample mounting involving different amounts of diamagnetic polymer materials in each configuration. The factor α takes into account that the magnitude of this background contribution varies between the two configurations.
A least squares minimization algorithm was used to derive a constant factor β which allowed us to remove the background when linearly combining the two data sets according to 1D = D2 − β D1 . The algorithm providing β relies on a universal property of antiferromagnets. Above TN , any linear combination of m k and m ⊥ has a temperature dependence which follows the functional form of a Curie–Weiss behavior f (T ) = T C−θ . Here C is a constant and θ is the Curie–Weiss PN temperature. By minimizing S(β, C, θ ) = i=1 (1D(Ti ) − f (Ti ))2 with Tmin = 230 K ≤ Ti ≤ 300 K and Ti+1 − Ti ≈ 5 K, we obtain the parameter β = 0.42, in addition to the auxiliary parameters C and θ . Obviously one can always find a parameter β nulling the background difference bs + αbother − β (bs + bother ). In general, β has the distinct potential to be temperature dependent; however, by systematically increasing Tmin from 230 to 280 K we found that in our case, β remains virtually constant within an error of less than 2.5%. Figure 4(a) shows 1D versus T (squares) calculated from the raw data sets D1,2 using β = 0.42. Indeed, the temperature dependence resembles a linear combination of parallel and perpendicular magnetic responses and, in particular, shows the expected Curie–Weiss type behavior in the high temperature limit. Using the susceptibility data from [20] and, assuming that our film has a magnetic response similar to the powder, we make an order of magnitude estimate of the thickness of our film. It yields a thickness of 200 nm in agreement with the target nominal thickness sought in the PLD growth. Specifically, we utilized the mean-field approximation for the perpendicular susceptibility of antiferromagnets, which implies that the susceptibility remains at the constant value χ⊥ (T < TN ) = χ (TN ). With this and the fact that the parallel susceptibility approaches zero with decreasing temperature 4
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below TN we use the low temperature value of 1D as an estimate for the perpendicular magnetic moment at TN . With the definition of susceptibility as moment per sample volume and applied magnetic field we are able to estimate the sample thickness by comparison with the powder data from [20]. Note that, when applying the approximation from above, the low temperature value of 1D multiplied by 1 − β = 0.58 provides 1D(TN ). Indeed, the data in figure 4(a) satisfy this relation remarkably well, thus providing independent confirmation of the validity of our analysis and the determination of TN outlined below. In order to extract the antiferromagnetic long-range ordering temperature of the film from the SQUID data we follow two complementary approaches. The triangles in figure 4(b) 1 show the 1D versus T dependence, which, in mean-field approximation, is expected to have linear temperature dependence for antiferromagnets for temperatures T ≥ TN [21]. Linear temperature dependence is indeed observed for T > 201 K. For T < 201 K the data deviate increasingly from 1 the linear function, obtained from a linear regression of 1D versus T in the temperature interval 245 K < T < 300 K. The deviation from the Curie–Weiss behavior below TN is clearly evident from a comparison with the extrapolation of the fit towards the intercept with the T-axis, i.e. the lower line in figure 4(b). A virtually identical value of TN is determined by a more sophisticated independent analysis using 1D versus T and fitting the Curie–Weiss law f (T ) in various temperature intervals Tmin ≤ T ≤ 300 K where Tmin has been systematically decreased from 250 K down to 160 K. Figure 4(b) (circles) shows the sum of squared residuals S versus Tmin resulting from the individual least squares fits. The intercepting lines mark a prominent increase of S versus Tmin for temperatures Tmin < 201 K. The Curie–Weiss behavior is strictly speaking an approximation which is asymptotically exact for temperatures T TN . Nevertheless, one expects a very pronounced deviation from the qualitative Curie–Weiss type functional form for antiferromagnets in the temperature region T < TN . Here the parallel susceptibility decreases with decreasing temperature and the perpendicular susceptibility remains virtually constant in accordance with mean-field approximation. From the temperature dependence of S, shown in figure 4(b) (circles), and the T-dependence of 1/1D (triangles in figure 4(b)), we conclude that TN ≈ 201 K or a value slightly higher. This is based on the fact that S can only increase above the noise level if some sizable temperature interval below TN is included in the sequence of fits [22]. Nevertheless, a possible ordering temperature of 201 K, slightly reduced from the powder value of 210 K (see figure 2), is expected in the thin film geometry, as observed elsewhere as well, e.g., in thin films of the magnetoelectric chromia [23].
in this complication. Therefore, angle-resolved x-ray photoemission spectroscopy (ARXPES) [24–29] was employed to determine the surface composition of the PLD grown Fe2 TeO6 thin films. ARXPES studies were performed using an x-ray source with a Mg Kα (1253.6 eV) line from a fixed anode. X-ray photoemission experiments were performed over a wide range of temperature, but the core level binding energies were obtained from data taken at a sample temperature of 600 K to suppress any photovoltaic charging effects in this insulating material. The photoemission spectra, taken in normal emission, provided core level binding energies of 576.4 ± 0.1 eV (figure 5(a)) and 711.6 ± 0.2 eV (figure 5(c)) for the Te (3d5/2 ) and Fe (2p3/2 ) core levels respectively and 531 ± 0.1 eV (figure 5(b)) for the O (1s) core level. The increase of the binding energy, by 3.4 eV for the Te (3d5/2 ) core level with respect to the metallic state (573 eV [30–33]), indicates that the Te underwent oxidation. Specifically, the Te (3d5/2 ) binding energy 576.4 ± 0.1 eV suggests the formation of an oxide TeOx with x = 2 (binding energy 575.9 [34] to 576.4 [30]) or even possibly x = 3 resembling TeO3 (binding energy 576.2 [35] to 576.4 [36]). Likewise, the increase in the Fe (2p3/2 ) binding energy (measured value of 711.6 ± 0.2 eV) from the metallic iron value of 706.8 eV, by about 4.8 eV suggests the formation of an oxidation state akin to Fe2 O3 (binding energy 710.9 eV [37] to 711 eV [38]) leading again to the conclusion that the oxidization states of the metals are as expected for the Fe2 TeO6 compound. Using angle-resolved core level photoemission [24–29], the relative surface concentration of a binary alloy of the form Ax B1−x can also be determined via the expression A R(θ ) = ( IIAB (θ)/σ (θ)/σB )(
p
E kin (A)−C ), p E kin (B)−C
where IA (θ ), IB (θ ) represent
the intensities of the respective element’s core level as a function of emission angle θ, normalized by the photoemission cross sections σA , σB , and the analyzer transmisp sion function E kin (A) − C at the kinetic energy of the core level of interest (see supplementary material available at s tacks.iop.org/JPhysCM/26/055012/mmedia). The oxygen to tellurium core level photoemission intensity ratios, normalized for cross section and transmission function, are roughly 6.2:1, indicating that the chemical stoichiometry of the film is preserved during deposition. As we probe more of the surface region at increased emission angles, the O (1s) to Te (3d3/2 ) ratio decreases slightly (supplementary material available at stacks.iop.org/JPhysCM/26/055012/mmedia) while the Fe (2p3/2 ) signal remains significantly smaller. This suggests a Te–O surface termination and that the tellurium stands slightly out of the plane of the surface compared to oxygen; the latter being consistent with the measured variation in relative intensity with emission angle variation is expected, as tellurium is slightly larger than oxygen, although complications arising from surface oxygen vacancies cannot be excluded. The surface photovoltaic effects are commonly observed in oxides and were also evident in the Fe2 TeO6 film as shown in figures 5(d)–(f). In general, surface photovoltage refers to the non-zero net charge induced at the surface of semiconductor (or insulator) material. In the case of photoemission spectroscopy,uncompensated charges are likely due to the
5. Surface characterization via angle-resolved x-ray photoemission spectroscopy
For many oxides, the chemical composition at the surface differs from that of the bulk [24–27] and Fe2 TeO6 films share 5
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Figure 5. Left panel shows x-ray photoemission spectra obtained for (a) Te (3d), (b) O (1s) and (c) Fe (2p) core levels. The spectra were taken at a temperature of 600 K. Photoelectrons were collected along the surface normal. Right panel shows the core level binding energy corresponding to (d) Te (3d3/2 ), (e) O (1s) and (f) Fe (2p3/2 ) photoemission peaks as a function of temperature. The dashed lines indicate temperatures at which the binding energies of the photoemission peaks begin to shift.
6. XMCD–PEEM investigation of Fe2 TeO6 thin film
creation of a hole after the ejection of the photoelectron from the films. In order to characterize any surface photovoltaic effects, binding energies were also measured as a function of temperature, with an accuracy of ±5 K. Here the binding energies of Te (3d3/2 ) (figure 5(d)), O (1s) (figure 5(e)) and Fe (2p3/2 ) (figure 5(f)) core levels were determined as a function of temperature. The data indicate a surface with a higher conductivity than that of the bulk. This is indicated in figure 5(d) by the shift in the binding energy of the Te (3d3/2 ) core level at a lower temperature (315 K) in comparison to the other core level peaks. This is indeed consistent with a Te–O terminated surface. The Te–O termination of the (110) surface might appear to be at odds with the XRD results, because this is not the expected termination of the bulk structure of this trirutile. However, significant surface reconstructions are well known. Examples with similarities to our system are the insulating, magnetoelectric antiferromagnet chromia [39–43] and the manganites [25–27, 27, 28]. In chromia, the details of the surface reconstructions are still under investigation due to their complicated temperature dependence with reentrant phenomena, and sensitivity to small energy changes. Note that complexity of the Fe2 TeO6 (110) reconstruction can be expected when considering related complex examples such as the well-known non-planar rutile (110) reconstruction [44]. With this in mind, it is no longer surprising that surface reconstructions, with significant deviations from the unrelaxed bulk structure, take place in the insulating, antiferromagnetic magnetoelectric Fe2 TeO6 .
6.1. Evidence for magnetoelectricity
X-ray magnetic circular dichroism combined with photoemission electron microscopy (XMCD–PEEM) [45–47] studies were carried out on our (110) textured Fe2 TeO6 thin film utilizing the spectromicroscopy (SM) beamline at the Canadian Light Source [48]. The SM beamline is capable of producing linearly and circularly polarized photons with energies ranging from 130 to 2500 eV. The photon source is an elliptically polarized APPLE II type undulator (EPU). The spatial resolution of the X-PEEM (Elmitec Gmbh) microscope is better than 30 nm for an ideal flat sample. Left and right circularly polarized light comes in at an angle of 74◦ relative to the surface normal. The X-PEEM measures, as used here, the local difference of intensities, 1I , in electron yield from left and right circularly polarized light. A piezo shutter was used in the data acquisition and the photon flux was reduced by closing the exit slits to 100 µm × 100 µm to keep charging effects at a minimum while enabling focused imaging. Our XMCD–PEEM investigation replicates the strategy of our recent experimental study on chromia [5]. This includes the concept of magnetoelectrically annealing the system to discriminate a single-domain state with virtually uniform boundary magnetization from a multi-domain state with zero net boundary magnetization. The domain pattern can be strongly affected by an electric field applied during cooling when a small but symmetry breaking magnetic field such as the Earth’s magnetic field is present. 6
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Figure 6. (a) Reference XMCD–PEEM image of the PLD grown (110) Fe2 TeO6 film measured at T = 300 K. The inset is a qualitative
depiction of the XMCD experiment indicating the spin polarization with respect to positively circularly polarized incident light. (b) XMCD–PEEM image after cooling to below TN in the absence of an applied electric field. The two black ellipses indicate regions where the spectra in figure 7 are taken. Each ellipse encloses roughly an area of 200 µm2 . (c) XMCD–PEEM image after 13.3 kV mm−1 electric field cooling to below TN in the simultaneous presence of Earth’s magnetic field. The bar to the right of the figures provides identical color codes used in all three images. Zero spin polarization is marked in the center of the color code while the upper and lower ends of the bar correspond to positive and negative spin polarization of the surface.
the degeneracy between the two antiferromagnetic domains is lifted thanks to the transverse magnetoelectric component, as long as the projection of the magnetic field on the surface normal does not vanish. In order to determine the components of the magnetization allowed for a thermodynamic equilibrium surface of a certain orientation, we need to determine how the existence of a polar vector designating the macroscopic orientation of the surface reduces the magnetic point group of the crystal. First, the magnetic point group of Fe2 TeO6 contains a plane m 0 that is normal to the tetragonal axis. This plane is not removed by any surface orthogonal to the c-axis, and antireflection in this plane, which is reflection combined with time reversal, forbids the magnetization component along c. Therefore, for any such surface orientation the equilibrium surface magnetization cannot have a component parallel to the c-axis. Moreover, for the (110) surface the magnetic point group becomes 2m 0 m 0 with the 2-fold axis normal to the surface, which allows only the out-of-plane component of the magnetization. Since the 2-fold axis interchanges the two antiferromagnetic sublattices, the out-of-plane component of the magnetization appears due to relativistic effects [3]. In the geometry of our experiment, the photon helicity has a non-zero projection onto the surface normal, and a spinpolarized signal is expected from the out-of-plane component of the magnetization. However, the spin contrast should also have a significant in-plane contribution, because the oblique incidence of the photon beam further reduces the symmetry. Indeed, in the absence of macroscopic time reversal symmetry, vanishing of the in-plane components of the spin-polarized response of a magnetoelectric surface is only protected by reflections and/or rotations retained by that surface. In the experimental setup used here, the beam lies in the plane containing the tetragonal axis and the surface normal. The symmetry of the combined beam–surface system is reduced from 2m 0 m 0 to m 0 . Here m 0 is the plane of incidence, which allows an axial vector component parallel to the tetragonal axis. The structurally equivalent antiferromagnetic sublattices are not oriented in the same way with respect to the incident x-ray beam giving rise to a microscopic origin of the associated XMCD signal. This may be dominant because both the
Figure 6(a) displays the reference XMCD image at T = 300 K > TN ≈ 201 K. Here virtually no contrast is visible due to the absence of antiferromagnetic long-range order. Figure 6(b) is recorded after cooling to slightly below the N´eel temperature in the absence of an applied electric field. The image shows that long-range order has been established within individual domains forming collectively a mesoscopic multi-domain state. Each domain carries boundary magnetization at the surface with an orientation determined by the registration of the antiferromagnetic order parameter of the respective domain. As a result a very prominent contrast in the XMCD–PEEM signal evolves. Interestingly, this domain contrast is far more pronounced when compared to the related XMCD–PEEM images of the magnetoelectric chromia [5]. Figure 6(c) shows the XMCD–PEEM image recorded after a 13.3 kV mm−1 electric field cooling across the N´eel temperature, obviously in the presence of the Earth’s magnetic field. Such a magnetoelectric field cooling procedure is known to lift the degeneracy of the antiferromagnetic domains [1, 49] and brings the system closer to a single-domain state. The observed XMCD–PEEM image implies that after magnetoelectric annealing, the overwhelming majority of surface spins are uniformly aligned. The XMCD intensity is proportional to the projection of the magnetization on the x-ray polarization direction. The XMCD absorption reflects electronic transitions from the 2p core level to the unoccupied majority-spin (minority-spin) 3d states. These transitions are induced predominantly (but not exclusively) by photons with positive (negative) helicity [50]. When magnetization is aligned parallel (antiparallel) to the photon angular momentum, there is a maximum (minimum) intensity in the absorption yield spectra. The contrast at the L3 edge is used to image domains in the boundary magnetization via XMCD–PEEM as shown in figure 6(b). The magnetic structure of Fe2 TeO6 [10] corresponds to the magnetic point group 4/m 0 m 0 m 0 , for which all the diagonal components of α are non-zero [6]. Therefore, magnetoelectric annealing is possible for almost any geometry and direction of the electric and magnetic fields. In particular, for the (110) surface with the electric field normal to the surface, 7
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Figure 7. (a) Right (circles) and left (line) circular polarized spectra taken from the blue area as shown in figure 6(b). (b) XMCD signal deduced from (a). (c) Right (circles) and left (line) circular polarized spectra taken from the red area as shown in figure 6(b). (d) XMCD signal deduced from (b). Green lower curves are the two-step background fits to the data, respectively. The spin and orbital magnetic moments are denoted accordingly, as derived using standard sum rules [54] from the data.
out-of-plane surface magnetization is related to spin–orbit coupling and the projection of angular momentum, along the surface normal, may be small. In order to distinguish these sources of contrast, it is necessary to analyze the dependence of the spin contrast on the angle of incidence. We can conclude that the XMCD–PEEM measurements provide clear evidence that our Fe2 TeO6 thin film is magnetoelectric. In addition, symmetry analysis reveals that boundary magnetization associated with the (110) surface must be normal to the (110) surface. Although XMCD–PEEM is sensitive to the contrast in perpendicular spin polarization one needs to keep in mind that interpretation of the contrast is based on a combination of two effects inherent to the single-domain magnetoelectric state: the out-of-plane surface magnetization, and the nonequivalence of the antiferromagnetic sublattices due to the oblique incidence of the photon beam. The relative magnitude of these effects was not determined in this work.
Fe2 TeO6 (110) as discussed above. The two absorption peaks at the shoulder of the L2 edge, observed at photon energy of 726.9 and 729.6 eV, are likely the result of the octahedral ligand field of the O2− ions, in accordance with multiplet ligand field calculations [51]. They indicate that the atomic multiplet spectrum of Fe3+ ions has two peaks at the L2 edge and the effect of the octahedral ligand field is to increase the splitting of these two peaks. Once the ligand field is larger than about 1 eV, the degeneracy of the second multiplet peak is lifted and thus the spectrum of Fe3+ at the L2 edge should have three distinguishable peaks. This is in very good agreement with our experiment where three peaks can be resolved. Different from γ -Fe2 O3 [52] and magnetite [53], Fe3+ ions in Fe2 TeO6 actually all reside in the octahedral sites. Figures 7(a) and (c) show the XMCD spectra taken from the different domains: the dark blue and red areas shown in figure 6(b), respectively. In figure 7(a), at the Fe L3 edge, the spectrum taken under left circular polarized light has a higher intensity than that taken under right circular polarized light. This intensity contrast is opposite at the Fe L2 edge. Such an XMCD contrast is characteristic of a ferromagnetic material but obtained here from the surface of an antiferromagnet. On the other hand, as shown in figure 7(c), the right circular polarized spectrum has a higher intensity at the Fe L3 edge and a lower intensity at the Fe L2 edge. Therefore, figures 7(a) and (c) exhibit XMCD spectra taken from opposite magnetic domains. A two-step background signal (green curves figures 7(a) and (c)) was subtracted, in the standard way [54], and subtracted from the difference of left and right circular polarized spectra respectively. Using regular sum rules [54], we deduce a lower bound for the local spin magnetic moment of 0.52 µB per
6.2. Lower bound to the spin and angular magnetic moment of the surface Fe ions
The XMCD spectra, recorded for the two different magnetic domains, are shown in figure 7. The two regions where the spectra have been taken are indicated by black ellipses in figure 6(b). Each ellipse encloses roughly an area of 200 µm2 . The Fe L3 edge and L2 edge are located at 711.8 eV and 724 eV, respectively. Both L3 and L2 edges in our measurement are at higher photon energy compared with spectra of Fe2+ ions (L3 at 710 eV and L2 at 723 eV) [51], indicating that the iron ions in Fe2 TeO6 should be in a higher oxidation state. In fact the Fe3+ configuration suggested by the L3 and L2 edge energies is consistent with the x-ray photoemission spectra for 8
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Fe atom in the blue area (figure 6(b)). This value is likely underestimated due to the possible misalignment of the Fe 3d spin moments and the photon polarization direction. The orbital magnetic moment in the blue area is 0.17 µB . In the red area (figure 6(b)), the spin magnetic moment and orbital magnetic moment contributions are estimated to be about 0.88 µB and 0.33 µB respectively. These large values rule out the possibility that the Fe moment measured is the result of the bulk magnetoelectric effect: the magnetoelectric effect contribution to the XMCD contrast would be very small as the bulk magnetization from the magnetoelectric effect is also small. Since XMCD– PEEM is not completely surface sensitive (it is a signal that is only predominantly from the surface region, not strictly a top surface layer measurement) and because of misalignment of the light angular momentum and the surface magnetization direction(s), our estimates of apparent spin magnetic moment must represent a lower bound and not the true spin magnetic moment and orbital magnetic moment contributions.
[2] Binek Ch and Doudin B 2005 J. Phys.: Condens. Matter 17 L39 [3] Belashchenko K D 2010 Phys. Rev. Lett. 105 147204 [4] Andreev F 1996 JETP Lett. 63 758 [5] Wu N, He X, Wysocki A L, Lanke U, Komesu T, Belashchenko K D, Binek Ch and Dowben P A 2011 Phys. Rev. Lett. 106 087202 [6] Buksphan S, Fischer E and Hornreich R 1972 Solid State Commun. 10 657 [7] Krishnan K, Mudher K S and Rao G R 2001 J. Alloys Compounds 316 264 [8] Bruker A X S 2008 TOPAS V4: General Profile and Structure Analysis Software for Powder Diffraction Data—User’s Manual (Karlsruhe: Bruker AXS) [9] Saes M, Raju N P and Greedan J E 1998 J. Solid State Chem. 140 7 [10] Kunnmann W, La Placa S, Corliss L, Hastings J and Banks E 1968 J. Phys. Chem. Solids 29 1359 [11] Yamaguchi M and Ishikawa M 1997 J. Phys. Soc. Japan 63 1666 [12] Morrish H 2001 Physical Principles of Magnetism (Piscataway, NJ: IEEE Press) p 447 [13] Domb C 1965 Magnetism vol IIA, ed G T Rado and H Suhl (New York: Academic) [14] Mostovoy M, Scaramucci A, Spaldin N A and Delaney K T 2010 Phys. Rev. Lett. 105 87202 [15] Rado G T 1962 Phys. Rev. 128 2546 [16] Hornreich R and Shtrikman S 1967 Phys. Rev. 161 506 [17] Pisarev R, Krichevtsov B and Pavlov V V 1991 Phase Transit.: Multinational J. 37 63 [18] Fischer E, Gorodetsky G and Shtrikman S 1971 J. Physique Suppl. 32 C1 650 [19] Mudher K D S, Krishnan K, Gopalakrishnan I K, Kundaliya D C and Malik S K 2005 J. Alloys Compunds 392 40 [20] Dehn J T 1968 J. Chem. Phys. 49 3201 [21] Binek Ch 2003 Ising-Type Antiferromagnets: Model Systems in Statistical Physics and the Magnetism of Exchange Bias (Springer Tracts in Modern Physics vol 196) (Berlin: Springer) [22] Binek Ch and Kleemann W 1995 Phys. Rev. B 51 12888 [23] He X, Echtenkamp W and Binek Ch 2012 Ferroelectrics 426 81 [24] Cheng R, Xu B, Borca C N, Sokolov A, Yang C-S, Yuan L, Liou S H, Doudin B and Dowben P A 2001 Appl. Phys. Lett. 79 3122 [25] Borca C N, Xu Bo, Komesu T, Jeong H-K, Liu M T, Liou S-H and Dowben P A 2002 Surf. Sci. Lett. 512 L346 [26] Dulli H, Plummer E W, Dowben P A, Choi J and Liou S-H 2000 Appl. Phys. Lett. 77 570 [27] Dulli H, Dowben P A, Liou S-H and Plummer E W 2000 Phys. Rev. B 62 14629 (R) [28] Chambers S A et al Surf. Sci. Rep. 65 317 [29] Dowben P A 1990 Core level spectroscopy and the characterization of surface segregation Surface Segregation Phenomena ed P A Dowben and A Miller (Boca Raton, FL: CRC Press) p 145 [30] Bensch W, Helmer O, Muhler M, Ebert H and Knecht M 1995 J. Phys. Chem. 99 3326 [31] Nyholm R and Martensson N 1980 J. Phys. C: Solid State Phys. 13 L279 [32] Shalvoy R B, Fisher G B and Stiles P J 1977 Phys. Rev. B 15 1680 [33] Pessa M, Vuoristo A, Vulli M, Aksela S, Vayrynen J, Rantala T and Aksela H 1979 Phys. Rev. B 20 3115 [34] Briggs D and Seah M P 1993 Practical Surface Analysis 2nd edn, vol 1 (Hoboken, NJ: Wiley) [35] Kang C K et al 2000 J. Appl. Phys. 88 2013 [36] Moulder J F, Stickle W F, Sobol P E and Bomben K D 1992 Handbook of X-Ray Photoelectron Spectroscopy ed J Chastain (Minnesota: Perkin-Elmer Corp.) p 131
7. Summary and conclusions
For the first time, highly (110) textured thin films of the ME antiferromagnet Fe2 TeO6 are fabricated with the help of pulsed laser deposition methodology. The raw material is synthesized from a solid state reaction and used to sinter targets for thin film deposition. Magnetometry of the Fe2 TeO6 powder suggests 3D Heisenberg universality. Magnetometry on the film sample confirms in-plane magnetic anisotropy and the presence of long-range antiferromagnetic order below the N´eel temperature of approximately 201 K. Photoemission supports correct chemical stoichiometry and bulk electrically insulating properties with an increased conductivity at the reconstructed Te–O terminated (110) surface of the Fe2 TeO6 film. Further investigations of Fe2 TeO6 films are necessary to clarify details of the reconstruction. XMCD–PEEM investigations give evidence of magnetoelectricity and voltage-controlled boundary magnetization in thin film geometry. The finding shows that boundary magnetization is not a unique functional property of chromia but can be found in other ME antiferromagnets supporting the theoretical expectation. This result puts voltage-controlled boundary magnetization in a broader context allocating it a key role in future voltage-controlled spintronic devices [55]. Acknowledgments
This project is supported by NSF through the Nebraska MRSEC, DMR 0213808, by the NRC/NRI supplement to MRSEC, and by CNFD and C-SPIN, part of STARnet, a Semiconductor Research Corporation program sponsored by MARCO and DARPA. The Canadian Light Source is supported by NSERC, NRC, CIHR, Province of Saskatchewan, WEDC and University of Saskatchewan. Fruitful discussion with K D Belashchenko is gratefully acknowledged. References [1] He X, Wang Y, Wu N, Caruso A N, Vescovo E, Belashchenko K D, Dowben P A and Binek Ch 2010 Nature Mater. 9 579 9
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[37] McIntyre N S and Zetaruk D G 1977 Anal. Chem. 49 1521 [38] Brion D 1980 Appl. Surf. Sci. 5 133 [39] Bender M, Ehrlich D, Yakovkin I N, Rohr F, B¨aumer M, Kuhlenheck H, Freund H-J and Staemmler V 1995 J. Phys.: Condens. Matter 7 5289 [40] Takano T, Wilde M, Matsumoto M, Okano T and Fukutani K 2006 J. Surf. Sci. Nanotechnol. 4 534 [41] L¨ubbe M and Moritz W 2009 J. Phys.: Condens. Matter 21 134010 [42] Chambers S A and Droubay T 2001 Phys. Rev. B 64 075410 [43] Wang X-G and Smith J R 2003 Phys. Rev. B 68 201402 (R) [44] Diebold U 2003 Surf. Sci. Rep. 48 53 [45] St¨ohr J, Padmore H and Anders S 1998 Surf. Rev. Lett. 5 1297 [46] Scholl A, Ohldag H, Nolting F, Anders S and St¨ohr J 2002 Study of ferromagnet–antiferromagnet interfaces using polarization dependent PEEM Magnetic Microscopies of Nanostructures ed H Hopster and H P Oepen (Berlin: Springer)
[47] Scholl A et al 2000 Science 287 1014 [48] Kaznatcheev K V, Karunakaran C, Lanke U D, Urquhart S G, Obst M and Hitchcock A P 2007 Nucl. Instrum. Methods Phys. Res. A 582 96 [49] Martin T J and Anderson J C 1966 IEEE Trans. Magn. 2 446 [50] Borca C N, Komesu T and Dowben P A 2002 J. Electron Spectrosc. Relat. Phenom. 122 259 [51] de Groot F M F, Fuggle J C, Thole B T and Sawatzky G A 1990 Phys. Rev. B 42 5459 [52] Brice-Profeta S, Arrio M-A, Tronc E, Letard I, Cartier dit Moulin Ch and Sainctavit Ph 2005 Phys. Scr. T115 626 [53] Goering E J, Gold S, Lafkioti M and Sch¨utz G 2007 J. Magn. Magn. Mater. 310 e249 [54] St¨ohr J and Siegmann H C 2006 Magnetism: From Fundamentals to Nanoscale Dynamics (Springer Series in Solid State Sciences vol 152) (Berlin: Springer) [55] Martin L W and Ramesh R 2012 Acta Mater. 60 2449
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Magnetoelectric Fe 2TeO 6 thin films
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Journal of Physics: Condensed Matter J. Phys.: Condens. Matter 26 (2014) 055012 (9pp)
doi:10.1088/0953-8984/26/5/055012
Magnetoelectric Fe2TeO6 thin films Junlei Wang1 , Juan A Colón Santana2 , Ning Wu1 , Chithra Karunakaran3 , Jian Wang3 , Peter A Dowben1 and Christian Binek1 1
Department of Physics and Astronomy, University of Nebraska, Theodore Jorgensen Hall, 855 North 16th Street, PO Box 880299, Lincoln, NE 68588-0299, USA 2 Department of Electrical Engineering, W Scott Engineering Center, University of Nebraska, North 16th Street, Lincoln, NE 68588-0656, USA 3 Canadian Light Source, University of Saskatchewan, 44 Innovation Boulevard, Saskatoon, SK, S7N 2V3, Canada E-mail: [email protected] Received 22 August 2013, in final form 5 November 2013 Published 17 January 2014
Abstract
We demonstrate that Fe2 TeO6 is a magnetoelectric antiferromagnet with voltage-controllable boundary magnetization. This provides experimental evidence of the theoretical prediction that boundary magnetization is a universal property of magnetoelectric antiferromagnets including boundary magnetization at a surface orthogonal to the polar direction. Highly (110) textured Fe2 TeO6 thin films were grown by pulsed laser deposition, terminating in Te–O at the (110) surface due to a surface reconstruction. Magnetic dc susceptibility measurements of both Fe2 TeO6 powder and thin film samples confirm antiferromagnetic long-range order. Finally, measurements of x-ray magnetic circular dichroism combined with photoemission electron microscopy (XMCD–PEEM) provide a lower bound to the spin and angular magnetic moment of the surface Fe ions. Keywords: boundary polarization, magnetoelectric, XMCD–PEEM, antiferromagnetic domains S Online supplementary data available from stacks.iop.org/JPhysCM/26/055012/mmedia (Some figures may appear in colour only in the online journal)
1. Introduction
magnetization is scarce and microscopic evidence has only been provided for the Cr2 O3 (0001) surface using x-ray magnetic circular dichroism–photoemission electron microscopy (XMCD–PEEM) [5], spin-polarized inverse photoemission [5] and spin-polarized photoemission [1]. In order to bring the concept of voltage-controlled boundary magnetization into a broader experimental context, and demonstrate general applicability, evidence for voltage-controllable boundary magnetization is needed from additional magnetoelectric antiferromagnetic insulators. Fe2 TeO6 is another material that has been demonstrated to be both antiferromagnetic and magnetoelectric [6], and therefore is expected to show boundary magnetization at surfaces and interfaces, based upon general theoretical considerations [3, 4]. Motivated by the prospect of largely dissipationless voltage-controlled interface magnetization in future ultralow power post-CMOS technology, we report on thin film
Voltage-controlled boundary magnetization of dielectric materials, such as in Cr2 O3 based voltage-controlled exchange bias systems [1, 2], has the potential to revolutionize spintronics. Boundary magnetization is a roughness insensitive net magnetization which emerges at the surface or interface of a magnetoelectric antiferromagnet in a single-domain state [1, 3]. The bulk linear magnetoelectric effect is quantified by the magnetoelectric susceptibility tensor, αi j . This tensor determines the magnetization induced in linear response due to the presence of an applied electric field. Symmetry arguments reveal that equilibrium boundary magnetization is a generic property of magnetoelectric antiferromagnets [3, 4], that may be obtained by magnetoelectric field cooling, with simultaneously applied electric and magnetic fields, thus utilizing the magnetoelectric effect to select an antiferromagnetic singledomain state. However, experimental evidence of boundary 0953-8984/14/055012+09$33.00
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Figure 2. Temperature dependence of the magnetic moment (open
circles) of the powder sample and deduced temperature dependence of the parallel component of the magnetic susceptibility (solid circles). Vertical arrow indicates N´eel temperature.
Figure 1. X-ray diffraction pattern (dots) with the calculated profile
(line) of Fe2 TeO6 . The peak positions and difference pattern are shown at the bottom of the chart. Major peaks are labeled.
general trirutile input structure of A2 BO6 compounds [9], provides the lattice parameters indexed for a tetragonal cell, with space group P42 /mnm, and atomic positions, as summarized in table 1. The temperature dependence (5 K < T < 300 K) of the field-induced magnetization was measured for approximately 0.1 cm3 of the Fe2 TeO6 powder sample in an applied magnetic field of 0.6 T, after zero-field cooling using a superconducting quantum interference device, SQUID (Quantum Design, MPMS XL). Figure 2 shows the field-induced magnetic moment (open circles), m, versus temperature, T. A decrease of the moment sets in at T < 230 K, indicating the onset of antiferromagnetic order [10, 11]. The N´eel temperature, TN , is determined from the inflection point in m versus T close to the maximum at 230 K. The antiferromagnetic long-range order establishes with mutually compensating sublattice magnetization. Our data suggest that TN ≈ 210 K, which is in good agreement with the prior literature value [6]. Due to the random orientation of the Fe2 TeO6 crystallites in the powder sample, the magnetization is isotropic and the susceptibility, hχi = ∂ M/∂ H ≈ M/H , is a weighted superposition of the parallel and perpendicular susceptibilities, χk , and χ⊥ , according to hχi = χk /3 + 2χ⊥ /3. Using χ⊥ (T < TN ) ≈ const [12] and χk (T TN ) χ⊥ (T < TN ), we extracted χk for T < TN (the susceptibility χk is plotted as the solid circles in figure 2) from hχ i versus T. One could use the assumption of low anisotropy, which would allow one to continue extrapolation of the parallel component of the susceptibility, χk to above TN using the measured hχ i(T ) for T > TN based on the mean-field approximation χ⊥ (T > TN ) = χk (T > TN ). This assumption is not adopted as we prefer to keep the number of assumptions at a minimum and rather not to deduce the χk values for T > TN . Close to TN , the T-dependence of the parallel magnetoelectric susceptibility αk is given by αk = 8λvµ0 Bη χk . Here η is the antiferromagnetic order parameter, v0 is the volume of the unit cell, µB is the Bohr magneton, and λ is the coupling strength which is determined by the compound’s spin
growth of the magnetoelectric antiferromagnet Fe2 TeO6 , a representative in a class of materials where dissipationless switching of boundary magnetization is feasible through voltage-control. In this work, we discuss the growth of highly (110) textured Fe2 TeO6 thin films utilizing pulsed laser deposition (PLD), and compare some of the magnetic properties in both powder and thin film samples. In addition, we provide direct evidence of voltage-controlled boundary magnetization in Fe2 TeO6 via XMCD–PEEM. Our results strongly support the generality of the concept of boundary magnetization in antiferromagnetic magnetoelectrics and demonstrate that chromia is not an exclusive material for voltage-controlled boundary magnetization. Indeed, we show that the boundary magnetization for a magnetoelectric can occur at a surface orthogonal to the polar direction. 2. Synthesis and characterization of Fe2 TeO6 powder
The starting point for thin film growth is the synthesis of fine Fe2 TeO6 powders, which we used to prepare Fe2 TeO6 targets by successive sintering and annealing. The synthesis of fine Fe2 TeO6 powders was done via a solid state reaction routine [6]. The synthesized powder was characterized by x-ray diffraction confirming the correct chemical structure of the sample. A Rigaku D/Max-B diffractometer using monochromatized Cu Kα radiation was used to record diffraction intensity in the 2θ range from 15◦ to 100◦ in steps of 0.02◦ scanning with speed of 0.2◦ min−1 . Figure 1 shows the x-ray diffraction pattern of Fe2 TeO6 powder and is in good agreement with the x-ray diffraction data reported in the literature [7]4 . In order to derive the structure, Rietveld refinement of the x-ray diffraction pattern was conducted utilizing the Bruker AXS DIFFRAC TOPAS 4.2 software [8]. The Rietveld least squares profile refinement, based on the 4
International Centre for Diffraction Data (ICDD) Powder Diffraction File (PDF) 04-009-3444, and 01-070-7443. 2
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Table 1. Lattice parameters and crystal parameters describing atomic positions within the unit cell obtained by Rietveld refinement of the
x-ray diffraction pattern. Wavelength (nm) a (Å) c (Å) Space group Crystal density (g cm−3 ) Rexp (%) Rwp (%) Goodness of fit Crystal parameters
0.154 06 4.605 98(95) 9.091 58(74) P42 /mnm (#136) 5.773 8.20 10.24 1.25
Atom Fe Te O1 O2
y 0 0 0.303 61(68) 0.305 43(46)
Site 4e 2a 4f 8j
x 0 0 0.303 61(68) 0.305 43(46)
z 0.334 94(12) 0 0 0.336 26(37)
B (Å2 ) 0.368(69) 0.624(67) 0.71(17) 0.58(12)
structure [13–16]. Since the T-dependence of αk is known for bulk Fe2 TeO6 [6] and we measured the T-dependence of χk we are able to deduce η ∝ (TN − T )β with β = 0.37. The value of β, the critical exponent of the order parameter, indicates 3D Heisenberg type criticality. Certainly, additional accurate data in the critical regime are needed to ultimately conclude on the critical behavior. It is tantalizing to correlate the potential similarity in the criticality of chromia and Fe2 TeO6 since the measurements of the critical exponent for Fe2 TeO6 are remarkably similar to the critical behavior of the archetypical magnetoelectric antiferromagnet chromia [17, 18]. 3. Growth of Fe2 TeO6 thin films and structure characterization below the surface
The growth of the Fe2 TeO6 thin films starts with preparation of targets, which are not commercially available at present. The target was prepared by pressing synthesized Fe2 TeO6 powder into a pellet with a hydraulic and cold isostatic press (CIP KJY s150-300). The CIP processed pellets were taken back to the furnace for annealing at 975 K using a ramping speed of 2 K min−1 . Our targets serve for both RF sputtering and pulsed laser deposition. Attempts to grow Fe2 TeO6 thin films by sputtering (ATC 2000-F, AJA International) resulted in only amorphous films. Some improvement towards formation of a polycrystalline material could be achieved through post-growth annealing, as indicated by (110) and (103) peaks in the XRD θ –2θ scan (not shown). The alternative, pulsed laser deposition (PVD MBE/PLD 2000), was employed to grow thin films of Fe2 TeO6 , and led to the growth of superior films. The energy of the KrF excimer laser was set to 130 mJ per pulse with 10 Hz pulse rate in order to prevent the target surface from overheating. During growth the target was kept at a distance of 10 cm relative to the substrate. The Al2 O3 (0001) substrate was maintained at 300 ◦ C with oxygen partial pressure environment, set to 10 mbar to keep the material stable. A deposition time of 1 h was used to achieve relatively thick films (nominally several hundred nanometers) for characterization and further analysis.
Figure 3. X-ray diffraction of the PLD grown sample (thick line)
with the sapphire substrate pattern (thin line). The prominent (110) Bragg peak of the Fe2 TeO6 thin film has been labeled.
Figure 3 shows the XRD pattern of the PLD grown Fe2 TeO6 /Al2 O3 (0001) sample, and in sharp contrast to the sputtered thin films, there is a single, pronounced, and narrow (110) peak present in the wide-angle diffraction pattern. The presence of the single sharp Bragg peak indicates that the thin film is highly (110) textured. The PLD grown thin films were used for all further analysis. 4. Magnetic characterization of (110) Fe2 TeO6 thin film
Magnetic susceptibility data for Fe2 TeO6 are scarce and to the best of our knowledge only available for powder samples [11, 19, 20] and this limits the comparisons available. There is a further complication in that susceptibilities of these antiferromagnetic thin films are challenging to measure due to the unfavorable film-to-substrate signal ratio. The sapphire substrate (0.5 mm × 3.9 mm × 6.8 mm) creates a diamagnetic background of the order of 10−4 emu = 10−7 A m2 in an 3
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Figure 4. (a) Linear combination 1D versus T (squares) of two m versus T data sets measured in orthogonal in-plane magnetic fields of µ0 H = 0.9 T applied along the respective sample edges. (b) Circles (left axis) show the sum of squared residuals S versus Tmin resulting from least squares fits of a Curie–Weiss law to 1D versus T data when systematically varying the fitting interval (see text). Upper lines are 1 versus T linear interpolations of the S versus Tmin data. Their intercept is used as a marker for TN (vertical arrow). Triangles show 1D (right axis). A corresponding linear fit (lower line, see text) is extrapolated towards interception with the T-axis. Onset of deviation of the data from linear behavior is used as a marker for TN .
applied magnetic field of µ0 H = 0.9 T or about two orders of magnitude larger than the signal for a Fe2 TeO6 thin film with a thickness of a few hundred nanometers. In addition, the temperature dependence of the combined substrate and sample holder signal complicates the analysis. The standard approach to address this problem is to measure the magnetic behavior of the substrate separately and then subtract the background signal from the overall sample signal. Note, however, that with this methodology it is not possible to measure the genuine substrate on which the film has been deposited. In cases where the ratio of film-to-substrate signal is 10−2 at most, it is more appropriate to utilize a method which allows subtraction of the background signal of the genuine substrate and holder in the presence of the magnetic film. In order to overcome the problems outlined above, we measured the temperature dependence of the magnetic moment of our sample in two different geometries. The measurements took place with an in-plane applied magnetic field of µ0 H = 0.9 T parallel and orthogonal to the (001) axis (as established by reflection high-energy electron diffraction: see supplementary material available at stacks.iop.org/JPhysCM/26/055012/ mmedia). Hence, the two orthogonal m versus T measurements probe the parallel and perpendicular susceptibilities of the antiferromagnet. The measurements lead to data sets D1,2 , where D1 = m k + bs + bother and D2 = m ⊥ + bs + αbother . Here bs (T ) is the diamagnetic background of the sapphire substrate and bother (T ) is a background signal, which depends on the respective conditions of sample mounting involving different amounts of diamagnetic polymer materials in each configuration. The factor α takes into account that the magnitude of this background contribution varies between the two configurations.
A least squares minimization algorithm was used to derive a constant factor β which allowed us to remove the background when linearly combining the two data sets according to 1D = D2 − β D1 . The algorithm providing β relies on a universal property of antiferromagnets. Above TN , any linear combination of m k and m ⊥ has a temperature dependence which follows the functional form of a Curie–Weiss behavior f (T ) = T C−θ . Here C is a constant and θ is the Curie–Weiss PN temperature. By minimizing S(β, C, θ ) = i=1 (1D(Ti ) − f (Ti ))2 with Tmin = 230 K ≤ Ti ≤ 300 K and Ti+1 − Ti ≈ 5 K, we obtain the parameter β = 0.42, in addition to the auxiliary parameters C and θ . Obviously one can always find a parameter β nulling the background difference bs + αbother − β (bs + bother ). In general, β has the distinct potential to be temperature dependent; however, by systematically increasing Tmin from 230 to 280 K we found that in our case, β remains virtually constant within an error of less than 2.5%. Figure 4(a) shows 1D versus T (squares) calculated from the raw data sets D1,2 using β = 0.42. Indeed, the temperature dependence resembles a linear combination of parallel and perpendicular magnetic responses and, in particular, shows the expected Curie–Weiss type behavior in the high temperature limit. Using the susceptibility data from [20] and, assuming that our film has a magnetic response similar to the powder, we make an order of magnitude estimate of the thickness of our film. It yields a thickness of 200 nm in agreement with the target nominal thickness sought in the PLD growth. Specifically, we utilized the mean-field approximation for the perpendicular susceptibility of antiferromagnets, which implies that the susceptibility remains at the constant value χ⊥ (T < TN ) = χ (TN ). With this and the fact that the parallel susceptibility approaches zero with decreasing temperature 4
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below TN we use the low temperature value of 1D as an estimate for the perpendicular magnetic moment at TN . With the definition of susceptibility as moment per sample volume and applied magnetic field we are able to estimate the sample thickness by comparison with the powder data from [20]. Note that, when applying the approximation from above, the low temperature value of 1D multiplied by 1 − β = 0.58 provides 1D(TN ). Indeed, the data in figure 4(a) satisfy this relation remarkably well, thus providing independent confirmation of the validity of our analysis and the determination of TN outlined below. In order to extract the antiferromagnetic long-range ordering temperature of the film from the SQUID data we follow two complementary approaches. The triangles in figure 4(b) 1 show the 1D versus T dependence, which, in mean-field approximation, is expected to have linear temperature dependence for antiferromagnets for temperatures T ≥ TN [21]. Linear temperature dependence is indeed observed for T > 201 K. For T < 201 K the data deviate increasingly from 1 the linear function, obtained from a linear regression of 1D versus T in the temperature interval 245 K < T < 300 K. The deviation from the Curie–Weiss behavior below TN is clearly evident from a comparison with the extrapolation of the fit towards the intercept with the T-axis, i.e. the lower line in figure 4(b). A virtually identical value of TN is determined by a more sophisticated independent analysis using 1D versus T and fitting the Curie–Weiss law f (T ) in various temperature intervals Tmin ≤ T ≤ 300 K where Tmin has been systematically decreased from 250 K down to 160 K. Figure 4(b) (circles) shows the sum of squared residuals S versus Tmin resulting from the individual least squares fits. The intercepting lines mark a prominent increase of S versus Tmin for temperatures Tmin < 201 K. The Curie–Weiss behavior is strictly speaking an approximation which is asymptotically exact for temperatures T TN . Nevertheless, one expects a very pronounced deviation from the qualitative Curie–Weiss type functional form for antiferromagnets in the temperature region T < TN . Here the parallel susceptibility decreases with decreasing temperature and the perpendicular susceptibility remains virtually constant in accordance with mean-field approximation. From the temperature dependence of S, shown in figure 4(b) (circles), and the T-dependence of 1/1D (triangles in figure 4(b)), we conclude that TN ≈ 201 K or a value slightly higher. This is based on the fact that S can only increase above the noise level if some sizable temperature interval below TN is included in the sequence of fits [22]. Nevertheless, a possible ordering temperature of 201 K, slightly reduced from the powder value of 210 K (see figure 2), is expected in the thin film geometry, as observed elsewhere as well, e.g., in thin films of the magnetoelectric chromia [23].
in this complication. Therefore, angle-resolved x-ray photoemission spectroscopy (ARXPES) [24–29] was employed to determine the surface composition of the PLD grown Fe2 TeO6 thin films. ARXPES studies were performed using an x-ray source with a Mg Kα (1253.6 eV) line from a fixed anode. X-ray photoemission experiments were performed over a wide range of temperature, but the core level binding energies were obtained from data taken at a sample temperature of 600 K to suppress any photovoltaic charging effects in this insulating material. The photoemission spectra, taken in normal emission, provided core level binding energies of 576.4 ± 0.1 eV (figure 5(a)) and 711.6 ± 0.2 eV (figure 5(c)) for the Te (3d5/2 ) and Fe (2p3/2 ) core levels respectively and 531 ± 0.1 eV (figure 5(b)) for the O (1s) core level. The increase of the binding energy, by 3.4 eV for the Te (3d5/2 ) core level with respect to the metallic state (573 eV [30–33]), indicates that the Te underwent oxidation. Specifically, the Te (3d5/2 ) binding energy 576.4 ± 0.1 eV suggests the formation of an oxide TeOx with x = 2 (binding energy 575.9 [34] to 576.4 [30]) or even possibly x = 3 resembling TeO3 (binding energy 576.2 [35] to 576.4 [36]). Likewise, the increase in the Fe (2p3/2 ) binding energy (measured value of 711.6 ± 0.2 eV) from the metallic iron value of 706.8 eV, by about 4.8 eV suggests the formation of an oxidation state akin to Fe2 O3 (binding energy 710.9 eV [37] to 711 eV [38]) leading again to the conclusion that the oxidization states of the metals are as expected for the Fe2 TeO6 compound. Using angle-resolved core level photoemission [24–29], the relative surface concentration of a binary alloy of the form Ax B1−x can also be determined via the expression A R(θ ) = ( IIAB (θ)/σ (θ)/σB )(
p
E kin (A)−C ), p E kin (B)−C
where IA (θ ), IB (θ ) represent
the intensities of the respective element’s core level as a function of emission angle θ, normalized by the photoemission cross sections σA , σB , and the analyzer transmisp sion function E kin (A) − C at the kinetic energy of the core level of interest (see supplementary material available at s tacks.iop.org/JPhysCM/26/055012/mmedia). The oxygen to tellurium core level photoemission intensity ratios, normalized for cross section and transmission function, are roughly 6.2:1, indicating that the chemical stoichiometry of the film is preserved during deposition. As we probe more of the surface region at increased emission angles, the O (1s) to Te (3d3/2 ) ratio decreases slightly (supplementary material available at stacks.iop.org/JPhysCM/26/055012/mmedia) while the Fe (2p3/2 ) signal remains significantly smaller. This suggests a Te–O surface termination and that the tellurium stands slightly out of the plane of the surface compared to oxygen; the latter being consistent with the measured variation in relative intensity with emission angle variation is expected, as tellurium is slightly larger than oxygen, although complications arising from surface oxygen vacancies cannot be excluded. The surface photovoltaic effects are commonly observed in oxides and were also evident in the Fe2 TeO6 film as shown in figures 5(d)–(f). In general, surface photovoltage refers to the non-zero net charge induced at the surface of semiconductor (or insulator) material. In the case of photoemission spectroscopy,uncompensated charges are likely due to the
5. Surface characterization via angle-resolved x-ray photoemission spectroscopy
For many oxides, the chemical composition at the surface differs from that of the bulk [24–27] and Fe2 TeO6 films share 5
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Figure 5. Left panel shows x-ray photoemission spectra obtained for (a) Te (3d), (b) O (1s) and (c) Fe (2p) core levels. The spectra were taken at a temperature of 600 K. Photoelectrons were collected along the surface normal. Right panel shows the core level binding energy corresponding to (d) Te (3d3/2 ), (e) O (1s) and (f) Fe (2p3/2 ) photoemission peaks as a function of temperature. The dashed lines indicate temperatures at which the binding energies of the photoemission peaks begin to shift.
6. XMCD–PEEM investigation of Fe2 TeO6 thin film
creation of a hole after the ejection of the photoelectron from the films. In order to characterize any surface photovoltaic effects, binding energies were also measured as a function of temperature, with an accuracy of ±5 K. Here the binding energies of Te (3d3/2 ) (figure 5(d)), O (1s) (figure 5(e)) and Fe (2p3/2 ) (figure 5(f)) core levels were determined as a function of temperature. The data indicate a surface with a higher conductivity than that of the bulk. This is indicated in figure 5(d) by the shift in the binding energy of the Te (3d3/2 ) core level at a lower temperature (315 K) in comparison to the other core level peaks. This is indeed consistent with a Te–O terminated surface. The Te–O termination of the (110) surface might appear to be at odds with the XRD results, because this is not the expected termination of the bulk structure of this trirutile. However, significant surface reconstructions are well known. Examples with similarities to our system are the insulating, magnetoelectric antiferromagnet chromia [39–43] and the manganites [25–27, 27, 28]. In chromia, the details of the surface reconstructions are still under investigation due to their complicated temperature dependence with reentrant phenomena, and sensitivity to small energy changes. Note that complexity of the Fe2 TeO6 (110) reconstruction can be expected when considering related complex examples such as the well-known non-planar rutile (110) reconstruction [44]. With this in mind, it is no longer surprising that surface reconstructions, with significant deviations from the unrelaxed bulk structure, take place in the insulating, antiferromagnetic magnetoelectric Fe2 TeO6 .
6.1. Evidence for magnetoelectricity
X-ray magnetic circular dichroism combined with photoemission electron microscopy (XMCD–PEEM) [45–47] studies were carried out on our (110) textured Fe2 TeO6 thin film utilizing the spectromicroscopy (SM) beamline at the Canadian Light Source [48]. The SM beamline is capable of producing linearly and circularly polarized photons with energies ranging from 130 to 2500 eV. The photon source is an elliptically polarized APPLE II type undulator (EPU). The spatial resolution of the X-PEEM (Elmitec Gmbh) microscope is better than 30 nm for an ideal flat sample. Left and right circularly polarized light comes in at an angle of 74◦ relative to the surface normal. The X-PEEM measures, as used here, the local difference of intensities, 1I , in electron yield from left and right circularly polarized light. A piezo shutter was used in the data acquisition and the photon flux was reduced by closing the exit slits to 100 µm × 100 µm to keep charging effects at a minimum while enabling focused imaging. Our XMCD–PEEM investigation replicates the strategy of our recent experimental study on chromia [5]. This includes the concept of magnetoelectrically annealing the system to discriminate a single-domain state with virtually uniform boundary magnetization from a multi-domain state with zero net boundary magnetization. The domain pattern can be strongly affected by an electric field applied during cooling when a small but symmetry breaking magnetic field such as the Earth’s magnetic field is present. 6
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Figure 6. (a) Reference XMCD–PEEM image of the PLD grown (110) Fe2 TeO6 film measured at T = 300 K. The inset is a qualitative
depiction of the XMCD experiment indicating the spin polarization with respect to positively circularly polarized incident light. (b) XMCD–PEEM image after cooling to below TN in the absence of an applied electric field. The two black ellipses indicate regions where the spectra in figure 7 are taken. Each ellipse encloses roughly an area of 200 µm2 . (c) XMCD–PEEM image after 13.3 kV mm−1 electric field cooling to below TN in the simultaneous presence of Earth’s magnetic field. The bar to the right of the figures provides identical color codes used in all three images. Zero spin polarization is marked in the center of the color code while the upper and lower ends of the bar correspond to positive and negative spin polarization of the surface.
the degeneracy between the two antiferromagnetic domains is lifted thanks to the transverse magnetoelectric component, as long as the projection of the magnetic field on the surface normal does not vanish. In order to determine the components of the magnetization allowed for a thermodynamic equilibrium surface of a certain orientation, we need to determine how the existence of a polar vector designating the macroscopic orientation of the surface reduces the magnetic point group of the crystal. First, the magnetic point group of Fe2 TeO6 contains a plane m 0 that is normal to the tetragonal axis. This plane is not removed by any surface orthogonal to the c-axis, and antireflection in this plane, which is reflection combined with time reversal, forbids the magnetization component along c. Therefore, for any such surface orientation the equilibrium surface magnetization cannot have a component parallel to the c-axis. Moreover, for the (110) surface the magnetic point group becomes 2m 0 m 0 with the 2-fold axis normal to the surface, which allows only the out-of-plane component of the magnetization. Since the 2-fold axis interchanges the two antiferromagnetic sublattices, the out-of-plane component of the magnetization appears due to relativistic effects [3]. In the geometry of our experiment, the photon helicity has a non-zero projection onto the surface normal, and a spinpolarized signal is expected from the out-of-plane component of the magnetization. However, the spin contrast should also have a significant in-plane contribution, because the oblique incidence of the photon beam further reduces the symmetry. Indeed, in the absence of macroscopic time reversal symmetry, vanishing of the in-plane components of the spin-polarized response of a magnetoelectric surface is only protected by reflections and/or rotations retained by that surface. In the experimental setup used here, the beam lies in the plane containing the tetragonal axis and the surface normal. The symmetry of the combined beam–surface system is reduced from 2m 0 m 0 to m 0 . Here m 0 is the plane of incidence, which allows an axial vector component parallel to the tetragonal axis. The structurally equivalent antiferromagnetic sublattices are not oriented in the same way with respect to the incident x-ray beam giving rise to a microscopic origin of the associated XMCD signal. This may be dominant because both the
Figure 6(a) displays the reference XMCD image at T = 300 K > TN ≈ 201 K. Here virtually no contrast is visible due to the absence of antiferromagnetic long-range order. Figure 6(b) is recorded after cooling to slightly below the N´eel temperature in the absence of an applied electric field. The image shows that long-range order has been established within individual domains forming collectively a mesoscopic multi-domain state. Each domain carries boundary magnetization at the surface with an orientation determined by the registration of the antiferromagnetic order parameter of the respective domain. As a result a very prominent contrast in the XMCD–PEEM signal evolves. Interestingly, this domain contrast is far more pronounced when compared to the related XMCD–PEEM images of the magnetoelectric chromia [5]. Figure 6(c) shows the XMCD–PEEM image recorded after a 13.3 kV mm−1 electric field cooling across the N´eel temperature, obviously in the presence of the Earth’s magnetic field. Such a magnetoelectric field cooling procedure is known to lift the degeneracy of the antiferromagnetic domains [1, 49] and brings the system closer to a single-domain state. The observed XMCD–PEEM image implies that after magnetoelectric annealing, the overwhelming majority of surface spins are uniformly aligned. The XMCD intensity is proportional to the projection of the magnetization on the x-ray polarization direction. The XMCD absorption reflects electronic transitions from the 2p core level to the unoccupied majority-spin (minority-spin) 3d states. These transitions are induced predominantly (but not exclusively) by photons with positive (negative) helicity [50]. When magnetization is aligned parallel (antiparallel) to the photon angular momentum, there is a maximum (minimum) intensity in the absorption yield spectra. The contrast at the L3 edge is used to image domains in the boundary magnetization via XMCD–PEEM as shown in figure 6(b). The magnetic structure of Fe2 TeO6 [10] corresponds to the magnetic point group 4/m 0 m 0 m 0 , for which all the diagonal components of α are non-zero [6]. Therefore, magnetoelectric annealing is possible for almost any geometry and direction of the electric and magnetic fields. In particular, for the (110) surface with the electric field normal to the surface, 7
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Figure 7. (a) Right (circles) and left (line) circular polarized spectra taken from the blue area as shown in figure 6(b). (b) XMCD signal deduced from (a). (c) Right (circles) and left (line) circular polarized spectra taken from the red area as shown in figure 6(b). (d) XMCD signal deduced from (b). Green lower curves are the two-step background fits to the data, respectively. The spin and orbital magnetic moments are denoted accordingly, as derived using standard sum rules [54] from the data.
out-of-plane surface magnetization is related to spin–orbit coupling and the projection of angular momentum, along the surface normal, may be small. In order to distinguish these sources of contrast, it is necessary to analyze the dependence of the spin contrast on the angle of incidence. We can conclude that the XMCD–PEEM measurements provide clear evidence that our Fe2 TeO6 thin film is magnetoelectric. In addition, symmetry analysis reveals that boundary magnetization associated with the (110) surface must be normal to the (110) surface. Although XMCD–PEEM is sensitive to the contrast in perpendicular spin polarization one needs to keep in mind that interpretation of the contrast is based on a combination of two effects inherent to the single-domain magnetoelectric state: the out-of-plane surface magnetization, and the nonequivalence of the antiferromagnetic sublattices due to the oblique incidence of the photon beam. The relative magnitude of these effects was not determined in this work.
Fe2 TeO6 (110) as discussed above. The two absorption peaks at the shoulder of the L2 edge, observed at photon energy of 726.9 and 729.6 eV, are likely the result of the octahedral ligand field of the O2− ions, in accordance with multiplet ligand field calculations [51]. They indicate that the atomic multiplet spectrum of Fe3+ ions has two peaks at the L2 edge and the effect of the octahedral ligand field is to increase the splitting of these two peaks. Once the ligand field is larger than about 1 eV, the degeneracy of the second multiplet peak is lifted and thus the spectrum of Fe3+ at the L2 edge should have three distinguishable peaks. This is in very good agreement with our experiment where three peaks can be resolved. Different from γ -Fe2 O3 [52] and magnetite [53], Fe3+ ions in Fe2 TeO6 actually all reside in the octahedral sites. Figures 7(a) and (c) show the XMCD spectra taken from the different domains: the dark blue and red areas shown in figure 6(b), respectively. In figure 7(a), at the Fe L3 edge, the spectrum taken under left circular polarized light has a higher intensity than that taken under right circular polarized light. This intensity contrast is opposite at the Fe L2 edge. Such an XMCD contrast is characteristic of a ferromagnetic material but obtained here from the surface of an antiferromagnet. On the other hand, as shown in figure 7(c), the right circular polarized spectrum has a higher intensity at the Fe L3 edge and a lower intensity at the Fe L2 edge. Therefore, figures 7(a) and (c) exhibit XMCD spectra taken from opposite magnetic domains. A two-step background signal (green curves figures 7(a) and (c)) was subtracted, in the standard way [54], and subtracted from the difference of left and right circular polarized spectra respectively. Using regular sum rules [54], we deduce a lower bound for the local spin magnetic moment of 0.52 µB per
6.2. Lower bound to the spin and angular magnetic moment of the surface Fe ions
The XMCD spectra, recorded for the two different magnetic domains, are shown in figure 7. The two regions where the spectra have been taken are indicated by black ellipses in figure 6(b). Each ellipse encloses roughly an area of 200 µm2 . The Fe L3 edge and L2 edge are located at 711.8 eV and 724 eV, respectively. Both L3 and L2 edges in our measurement are at higher photon energy compared with spectra of Fe2+ ions (L3 at 710 eV and L2 at 723 eV) [51], indicating that the iron ions in Fe2 TeO6 should be in a higher oxidation state. In fact the Fe3+ configuration suggested by the L3 and L2 edge energies is consistent with the x-ray photoemission spectra for 8
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Fe atom in the blue area (figure 6(b)). This value is likely underestimated due to the possible misalignment of the Fe 3d spin moments and the photon polarization direction. The orbital magnetic moment in the blue area is 0.17 µB . In the red area (figure 6(b)), the spin magnetic moment and orbital magnetic moment contributions are estimated to be about 0.88 µB and 0.33 µB respectively. These large values rule out the possibility that the Fe moment measured is the result of the bulk magnetoelectric effect: the magnetoelectric effect contribution to the XMCD contrast would be very small as the bulk magnetization from the magnetoelectric effect is also small. Since XMCD– PEEM is not completely surface sensitive (it is a signal that is only predominantly from the surface region, not strictly a top surface layer measurement) and because of misalignment of the light angular momentum and the surface magnetization direction(s), our estimates of apparent spin magnetic moment must represent a lower bound and not the true spin magnetic moment and orbital magnetic moment contributions.
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7. Summary and conclusions
For the first time, highly (110) textured thin films of the ME antiferromagnet Fe2 TeO6 are fabricated with the help of pulsed laser deposition methodology. The raw material is synthesized from a solid state reaction and used to sinter targets for thin film deposition. Magnetometry of the Fe2 TeO6 powder suggests 3D Heisenberg universality. Magnetometry on the film sample confirms in-plane magnetic anisotropy and the presence of long-range antiferromagnetic order below the N´eel temperature of approximately 201 K. Photoemission supports correct chemical stoichiometry and bulk electrically insulating properties with an increased conductivity at the reconstructed Te–O terminated (110) surface of the Fe2 TeO6 film. Further investigations of Fe2 TeO6 films are necessary to clarify details of the reconstruction. XMCD–PEEM investigations give evidence of magnetoelectricity and voltage-controlled boundary magnetization in thin film geometry. The finding shows that boundary magnetization is not a unique functional property of chromia but can be found in other ME antiferromagnets supporting the theoretical expectation. This result puts voltage-controlled boundary magnetization in a broader context allocating it a key role in future voltage-controlled spintronic devices [55]. Acknowledgments
This project is supported by NSF through the Nebraska MRSEC, DMR 0213808, by the NRC/NRI supplement to MRSEC, and by CNFD and C-SPIN, part of STARnet, a Semiconductor Research Corporation program sponsored by MARCO and DARPA. The Canadian Light Source is supported by NSERC, NRC, CIHR, Province of Saskatchewan, WEDC and University of Saskatchewan. Fruitful discussion with K D Belashchenko is gratefully acknowledged. References [1] He X, Wang Y, Wu N, Caruso A N, Vescovo E, Belashchenko K D, Dowben P A and Binek Ch 2010 Nature Mater. 9 579 9
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