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Evidence of ferromagnetic ordering at low temperatures in epitaxial TbMnO3 thin films grown by means of DC magnetron sputtering
Journal Pre-proofs Evidence of ferromagnetic ordering at low temperatures in epitaxial TbMnO3 thin films grown by means of DC magnetron sputtering J.L. Izquierdo, A. Astudillo, J. Martínez, G. Bolaños, O. Morán PII: DOI: Reference:
S0304-8853(19)31041-8 https://doi.org/10.1016/j.jmmm.2019.166141 MAGMA 166141
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Journal of Magnetism and Magnetic Materials
Received Date: Revised Date: Accepted Date:
20 March 2019 14 October 2019 11 November 2019
Please cite this article as: J.L. Izquierdo, A. Astudillo, J. Martínez, G. Bolaños, O. Morán, Evidence of ferromagnetic ordering at low temperatures in epitaxial TbMnO3 thin films grown by means of DC magnetron sputtering, Journal of Magnetism and Magnetic Materials (2019), doi: https://doi.org/10.1016/j.jmmm.2019.166141
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Evidence of ferromagnetic ordering at low temperatures in epitaxial TbMnO3 thin films grown by means of DC magnetron sputtering J. L. Izquierdo1, A. Astudillo2, J. Martínez2, G. Bolaños2, O. Morán3 1
Institución Universitaria Pascual Bravo, Centro de Investigación y Desarrollo en Materialografía, Calle 73 No. 73A-226, C.P. 050001 Medellín, Colombia. 2 Universidad del Cauca, Departamento de Física, Grupo de bajas temperaturas, A.A. 1226 Popayán, Colombia. 3 Universidad Nacional de Colombia, Campus Medellín, Facultad de Ciencias, Departamento de Física, Advanced Oxides Group, A.A. 568 Medellín, Colombia Abstract. Thin films of the archetypal multiferroic TbMnO3 were epitaxially deposited on (001)oriented SrTiO3 substrates by means of DC magnetron sputtering. A substrate temperature of 1053 K and an oxygen pressure of 300 Pa were determined to be appropriate growing conditions for obtaining high-quality, epitaxial TbMnO3 films using this physical growing technique. The structural and magnetic properties of ~100 nm thin TbMnO3 films were investigated. X-ray diffraction patterns showed that the films are single-phase and (00ℓ)-oriented. Rocking curves of the (002) peak confirm the good crystalline quality of the films. X-ray photoelectron spectroscopy analysis confirmed that the nominal valence of the Mn ions is 3+. The films grown on the SrTiO3 substrates displayed substantial epitaxial strain, which decidedly changed the magnetic ground state of the bulk-like counterpart. In particular, the films displayed ferromagnetic-like interactions below a temperature close to the magnetic ordering temperature of the Mn3+ spins (TN42 K). Well-defined hysteresis loops with coercive fields as high as 0.1 T were observed at 5 K. The variation of the coercive field with the temperature clearly indicated the presence of low-temperature ferromagnetism in the TMO films. The origin of the anomalous ferromagnetism in the films is discussed in terms of the strain-induced distortion generated by the lattice mismatch between the film and substrate. The reported results can give insight into accurate strategies for controlling single-phase spin-driven multiferroic states. Corresponding author: E-mail: [email protected] Phone: 57-4-4309327 Fax: 57-4- 2604489
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1. Introduction Multiferroic materials have attracted considerable attention over the last decade after observations of the coexistence of ferromagnetic and ferroelectric polarizations, as well as the coupling effect between them [1]. Apart from their fascinating physical properties, these materials could also have technological applications [2]. Among the many multiferroic materials, orthorhombic TbMnO3 (TMO) is one of the few in which ferroelectricity arises from the magnetic spiral order (spin-driven ferroelectrics) [3]. Bulk-like TMO is a multiferroic material that exhibits a complex magnetoelectric phase diagram, including antiferromagnetism below ~40 K. [4]. However, scaling TMO crystals into thin film form may generate epitaxial strain that would drastically affect their multiferroic properties [5]. In this regard, the possibility of tuning the exchange interactions and inducing ferromagnetic behavior via strain could lead to single-phase ferrimagnetic ferroelectrics [6]. Such sophisticated devices are highly interesting for a variety of applications and are not commonly reported in the literature. The fact that epitaxial strain may alter the magnetic properties of rare-earth-based manganites when they are grown on crystalline substrates via physical or chemical methods has been demonstrated in recent papers [7]. Generally speaking, the origin of the ferromagnetism in epitaxially strained TMO films seems to be related to a change in the electronic structure of the material [8]. From the analysis of the results of neutron diffraction experiments, it has been established that in highly strained films, the magnetic order changes from the bulk-like incommensurate bc-cycloidal structure to the commensurate magnetic order [9]. This change in the magnetic order can lead to a complex coexistence of magnetic order parameters [10]. In the case of TMO thin films, a pronounced increase of the ferroelectric polarization and a change in its direction from along the c- to the a-axis was observed, which seemed to be concomitant with the
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magnetic order change in strained TMO films [9]. Competing magnetic interactions at low temperatures tuned by small lattice distortions could be the mechanism responsible for the switchable electric polarization in TMO thin films [9]. Hence it is possible that the complex ferromagnetic properties reported for TMO thin films could be the result of strain-induced deformation of the unit cell and the subsequent unbalancing of magnetic interactions [1]. As is evident from the previous paragraph, the fabrication of high-quality thin films is a fundamental step toward the possibility of controlling the different multiferroic phases in novel oxides by means of the epitaxial strain, dimensionality, and size effect, without varying the chemical composition. In most of the previous investigations reporting on the preparation, characterization, and multiferroic function of TMO thin films, pulsed laser deposition (PLD) was employed as a physical method for depositing the thin layers on crystalline substrates [11–13]. The reported results are diverse. Nevertheless, reports on the growth of epitaxial TMO thin films by other methods are scarce [14]. Although the PLD technique definitely has several advantages, such as high deposition rate, short test period, and low substrate temperature requirements, the deposited films can be affected by small molten particles or target fragments, which are generated during the laser-induced explosion [15]. The presence of these particles can severely diminish the quality of the films. Sputter deposition, although it equally presents several difficulties, such as the differentiation between the nature of the compound target (whether metal or insulator), is a versatile and inexpensive physical vapor deposition method for preparing thin films of alloys and complex materials for research and industrial purposes [16]. Generally speaking, sputter deposition can be considered to be a natural method for the deposition of complex ceramic materials [16]. The large particle energy, the atomic transfer of the target materials, and the possibility of using
3
reactive gas components (reactive sputtering) at a large partial pressure open up a broad parameter range for growing high-quality epitaxial thin films and superlattices of complex oxides, often at a significantly lower cost than other techniques. Sputter is compatible with high oxygen-pressure, and calibration and rate control are readily obtained. Films grown via sputter techniques are highly reproducible, and a homogeneous composition can be obtained over a rather large area (typically 6.4x10-3 m) [17]. With each wave of interest in a particular class of complex oxide materials, a route to high-quality oxide films through sputtering has usually been found, and the resulting films have been important for scientific progress. In the present investigation, we report on the successful growth of c-axis-oriented epitaxial TMO films on (001)oriented SrTiO3 (STO) substrates by means of DC magnetron sputtering. Specifically, it is shown that the magnetic ground state of TMO thin films is clearly modified by epitaxial strain. 2. Experimental details The preparation of the TMO targets was described in a previous paper [18]. The singlephase TMO targets were coupled to a sputter cannon and pre-sputtered before film deposition in order to remove possible contaminated surface layers. (001)-oriented STO was chosen as substrate material. STO possesses a cubic perovskite structure, and the lattice constant is 3.905 Å. TMO films (~100 nm) were deposited in an oxygen atmosphere at a pressure of 300 Pa, a power of 30 W, and a substrate temperature of 1053 K. After deposition, the substrate temperature was lowered to 723 K, and the chamber was then flooded with oxygen at a pressure of ~8x104 Pa. The films were annealed under these conditions for 1 h. X-ray diffraction patterns of the TMO films were registered with a four-circle Rigaku diffractometer using CuKα radiation (λ=0.15418 nm). The elemental composition of the films and the depth profile with high
4
resolution spectra were studied by means of X-ray photoelectron spectroscopy (XPS) using a multi-technique scanning XPS microprobe (PHI Versa Probe II). The source of X-rays used was a monochromatic AlK lamp (300 W) at a photon energy of 2.38x10-16 J. The photoelectron detection was carried out by means of a concentric hemispherical analyzer (CHA) (PHI 10360 Model) operating with an analytical aperture 0.8 m in diameter. The binding energy resolution of the instrument is 1.12x10-19 J. Magnetization (M) as a function of temperature and magnetic field (H) was measured using a superconducting
quantum
interference
device
(SQUID,
Quantum
Design)
magnetometer. The applied magnetic field was always kept parallel to the plane of the film. The field-cooled (FC) and zero-field-cooled (ZFC) magnetization measurements were carried out on heating from 5 to 350 K, after the sample was first cooled from 300 K down to the lower temperature with and without the magnetic field. The isothermal M(H) curves were recorded at 40, 20, and 5 K after FC (7 T) from 300 K. 3. Results and discussion
Figure 1(a) shows the XRD pattern of one typical ~100 nm thin TMO film grown on a (001)-oriented STO substrate by means of DC magnetron sputtering. The film is pure caxis oriented and free from secondary phases. The displacement of the peaks of the substrate and the film is due to the mismatch between the lattice parameters of the film and those of the substrate. From the XRD pattern, the out-of-plane lattice parameter was found to be 0.745 nm, which is slightly larger than the c0.740 nm for TMO bulk single crystals [19]. This finding suggests that the films that were grown were subjected to an in-plane compressive strain, which resulted in the tetragonally distorted orthorhombic phase of TMO. The achieved result agrees with previous reports on TMO thin films grown on STO substrates [20]. Figure 1(b) shows the rocking curve around the (002) reflection for the TMO film. The full width at half maximum is 0.4°. Here it is 5
important to note that whereas TMO possesses an orthorhombic structure with lattice parameters a05.293 Å, b05.838 Å, and c07.403 Å [19], the STO substrates have a cubic crystalline structure with lattice parameter a=3.905 Å. Therefore, a large lattice misfit should be taken into consideration in the epitaxial growth of TMO on STO substrates (-4.1 % along the a axis and 5.7 % along the b axis). Hence it is to be expected that the orhorhombic unit cell of TMO be grown on a 3.9052 Åx3.9052 Å square STO lattice, which can better accommodate the in-plane film lattice. This growth model has previously been suggested for TMO and DyMnO3 films deposited on STO substrates [11, 21]. In this model, the strained TMO films should share the pseudocubic lattice parameters of the substrate and therefore should be compressed in the plane directions from about 3.94 Å of the bulk to 3.905 Å [3]. In Ref [3], a deviation of 90° of the pseudocubic angle of the TMO films was reported. This deviation, which represents the orthorhombic distortion, was determined to be smaller than that of the bulk counterpart. As a result, these TMO films grew with a highly-compressed orthorhombic bo and a slightly expanded ao axis. Therefore, the enlargement of the out-of-plane co lattice parameter in the TMO films is consistent with an overall compressive in-plane strain. The oxidation states of the Tb 3d, Mn 2p, and O 1s in the TMO films was verified via XPS measurements. The peak positions were referenced to carbon at 4.61x10-17 J. Figure 2 shows a wide scan spectrum of a representative TMO film. Peaks of Tb, Mn, O, and C can clearly be observed. The narrow scan spectra for Tb 3d is shown in Figure 3(a). Two characteristic peaks can be seen at binding energies of 1.98x10-16 J (Tb 3d5/2) and 2.04x10-16 J (Tb 3d3/2), in agreement with values previously reported [22]. In turn, the characteristic peaks of Mn3+ 2p3/2 (1.02x10-16 J) and 2p1/2 (1.04x10-16 J) are shown in Fig. 3(b). The binding energy values of the Mn3+ agree with those already reported for
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the compound [3,22]. The experimental peaks were successfully fit, as indicated by the solid lines in Fig. 4(a) and (b). In particular, the absence of a secondary peaks (shake-up peaks) in the Mn 2p XPS edge suggests that Mn exists in the 3 oxidation state in the TMO films grown on STO substrates by means of DC magnetron sputtering. Here it is useful to point out that Mn is normally characterized by a 3 oxidation state in Mnbased perovskite oxides [22]. The Mn 3p level splitting can be seen directly in Fig. 3(b) and equals 1.89x10-18 J. The experimental reference values reported for Mn 2p ions is exactly 1.89x10-18 J [23]. Thus the absence of shake peaks in the Mn 2p spectrum and the magnitude of the Mn 2p splitting discount the existence of mixed +2/+3 valence in the films. The existence of mixed +2/+3 valence in Mn-based compound is commonly associated with the presence of oxygen vacancies in the sample [24]. Therefore, these TMO films grown via DC magnetron sputtering do not seem to be affected by this kind of defect. Figure 3(c) shows the O 1s XPS spectrum of the TMO film. A pronounced peak and a clear shoulder are located at 8.49x10-17 and 8.53x10-17 J, respectively. Such splitting indicates that the oxygen in the TMO films exists in two kinds of chemical states. The binding energy at 8.49x10-17 J eV is mainly due to the contribution of the crystal lattice oxygen (OL) [22]. The OL signal probably stems from the contribution of Tb–O and Mn–O in the TMO lattice. In turn, the shoulder at 8.53x10-17 J can arise from chemisorbed oxygen species (OH) on the surface (e.g. chemisorbed water) [22]. The temperature-dependent FC magnetization (M) data of the polycrystalline TMO target are shown in Fig. 4. The data were recorded in a field of 0.02 T. A pronounced increase in the value of the magnetization can be seen at T<45 K, which is close to the TN bulk value of TMO (T~42 K). Below TN, the Mn3+ sublattice starts to develop an incommensurate sinusoidal antiferromagnetic (AFM) alignment [25]. Superexchange interaction due to the spin reorientation of Mn3+ ions is considered to be the mechanism
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behind this phenomenon [26]. The inverse magnetic susceptibility, -1, versus the temperature for the polycrystalline TMO target is plotted in the inset of Fig. 4. It can be seen that the experimental data follow the Curie-Weiss law =C/(T-) within the temperature range 42–350 K. Thus the Curie-Weiss law is fulfilled down to 42 K, below which the Mn3+ spins of TMO start to develop an incommensurate sinusoidal AFM structure [25]. The values of the Curie constant C and the Curie-Weiss temperature were found to be 620 Am2K/kgT and -34.6 K, respectively. The AFM ordering at low temperatures is confirmed by the negative value of . The values found for the parameters C and are in accordance with those commonly reported in the literature on the compound [27]. On the other hand, the incommensurate-commensurate transition Tlock, occurring at ~28 K, was not observed in the M(T) (or -1(T)) measurements. The transition at Tlock is accompanied by magnetoelastically-induced lattice modulation and the appearance of ferroelectricity [2]. Probably the dominant paramagnetic signal of the Tb spins, which possess a larger magnetic moment than the Mn spins, causes this transition not to be observed [8]. The ZFC and FC in-plane DC susceptibility as a function of the temperature for a 100 nm thin TMO film is presented in Fig. 5(a). The data were recorded in a 0.1 T Oe magnetic field with a cooling field of 7 T for the FC mode. The diamagnetic contribution from the substrate was subtracted from the raw magnetization data. The ZFC and the FC curves show a distinct anomaly at 43 K (inset of Fig. 5(a)), which is very close to the anomalies in the C(T) and M(T) dependences observed in TMO single crystals and ascribed to the transition from the paramagnetic phase to the sinusoidal antiferromagnetic structure of the manganese spins [1]. Since the Tb3+ moments are located in the ab plane of a single crystal [28], it is possible that the peak appearing at
8
~10 K (inset of Fig. 5(a)) is associated with the Tb3+ spin ordering in the strained films (the field is applied parallel to the film plane). Moreover, the magnetic measurements did not reveal any feature related to the stabilization of the spiral antiferromagnetic ordering, concomitant with the onset of ferroelectricity, which occurs in bulk at ∼27 K [1]. A split between FC and ZFC measurements can also be observed below the Mn3 spin ordering temperature. This effect suggests that ferromagnetic interactions are present in the films [3]. Splitting between the FC and ZFC modes reported for TMO in bulk form has been attributed to extrinsic effects related to uncompensated moments at the grain boundaries and clearly depends on the magnetic field strength [4]. The existence of the low-temperature FM phase in the TMO films is verified by the clearly hysteretic behavior observed at low temperatures in the 6 T field range after cooling from 300 K in a 7 T magnetic field (Fig. 5(b)). A coercive field of about 0.1 T at 5 K (inset of Fig. 5(b)) confirms the presence of a FM order at low temperatures in the analyzed films. This behavior drastically departs from the antiferromagnetic order observed for bulk TMO. The results obtained for the TMO films are in good agreement with
polarized
neutron
reflectometry
measurement
data
[29].
Interestingly,
orthorhombic YbMnO3 and YMnO3 thin films exhibit similar effects, which speaks to a general mechanism in manganite thin films [30, 31]. As stated above, the weak ferromagnetism present in the TMO films has no counterpart in the bulk. In particular, the magnetic measurements showed evidence of a transition that is very close to the magnetic ordering temperature (TN42 K) of the Mn3+ spins. Nevertheless, the measurements did not show clear evidence of stabilization of the spiral antiferromagnetic ordering and the concomitant onset of ferroelectricity, as in the bulk counterpart. Hence it is apparent that the observed ferromagnetism in the thin TMO films should be explained by another scenario. Generally speaking, the lattice misfit
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between the films and the substrates generates strain effects in the thin films. A modified microstructure is the direct consequence of this effect. Following this line of argument, it is probable that the origin of the ferromagnetism in epitaxially strained TMO films is related to a change of the electronic structure in the material [32]. For bulk RMnO3 compounds, it is well stablished that each Mn3+ ion has four unpaired d electrons [1]. Here, the low-energy t2g orbitals are occupied by three electrons and the high-energy eg orbitals by only one. Moreover, the two degenerate eg states on each spin site split into two nondegenerate orbitals as a consequence of the Jahn-Teller distortion. Only one of these two nondegenerate orbitals is occupied. In this case, the occupied eg orbitals on two adjacent spin sites are almost orthogonal to each other in the ab plane [33]. On the other hand, it has been suggested that the orbital order in epitaxiallystrained TMO films is very different [12]. First, the two eg orbitals on each spin site are almost degenerate. This implies that both can be considered to be half-occupied, and the electron configuration of each Mn3+ (d4) ion could roughly be described as (t2g)3(d3z2−r2)1/2(dx2−y2)1/2. Such electron configuration leads to slightly different Mn-O bond lengths in the different directions of the MnO6 octahedron (2, 2.02, and 2.08 Å in the x, y, and z directions, respectively.) The O-Mn-O bond angles are close to 90° or 180°. It is interesting to analyze the interaction between the eg orbitals of two adjacent spin sites along the c direction for the FM arrangement. In this case, it has been determined that the majority-spin eg (d3z2−r2) orbitals of both sites with the same energy level can couple with each other [33]. This will lead to the formation of an empty higher-energy level and an occupied lower-energy level. Thus the total energy will be reduced by the quantity t (>0), which represents the hopping integral between these two d3z2−r2 orbitals. For the AFM configuration, it has also been established that the majority-spin eg orbital of one Mn site couples with the minority-spin eg orbital of the
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other Mn site [34]. Hence the reduction in the value of the total energy can be expressed as U/2-(t2 + U2/ 4)1/2. In this equation, U (>0) corresponds to the energy difference between eg orbitals with opposite spin components. From these results, it is evident that the FM state has a lower energy. Interestingly, it has been verified that the in-plane exchange interaction Jab is governed by the same mechanism [35]. Recalling that symmetric exchange interactions between spins of Mn3+ ions characterize a ferromagnetic ground state of materials with perovskite-type structures [33], it is apparent that ferromagnetic interactions are favored in epitaxially-strained TMO films. In addition, it has also been established that nearest-neighbor Mn-Mn distances in strained TMO films decrease and Mn-O-Mn angles become larger as compared to those of the bulk counterpart [32]. This will cause as consequence the enhancement of the coupling of eg orbitals (increase of t) and hence an increase in the magnitude of the spin exchange interaction parameters, which is consistent with the increase in the symmetric spin exchange parameters. On the other hand, some authors have suggested that strain effects alone cannot account for the difference between the magnetic properties of the bulk material and thin films [31]. The morphology was then considered to be an additional factor influencing the magnetic response of epitaxial TMO films. In a recent and interesting investigation, the influence of a ferroelectric STO substrate on the magnetic properties of GdMnO3 and YbMnO3 thin films fabricated via RF-magnetron sputtering was reported [36]. There, a clear difference in the values of the DC susceptibility of GdMnO3 films grown on ferroelectric STO and LiNbO3 substrates was observed. A difference in the value of the splitting temperature between the ZFC and FC curves was also evident for the films grown on the two substrates. The influence of the polarization vector of the substrate on atomic positions at the interface and the consequent change in the magnitude and direction of the magnetic moments in the
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GdMnO3 films was suggested as a possible mechanism behind the observed differences in their magnetic behavior. According to the authors, this mechanism is supported by the fact that the permittivity values of STO are one order of magnitude higher than those of LiNbO3. Although a rather homogeneous magnetization profile of orthorhombic TMO films, with some increase in the magnetic moment at the bottom interface, has been verified by means of spin-polarized neutron reflectometry [29], it is evident that a more complex analysis of the full structure of the films is needed to explain the observed ferromagnetism in this material when grown in thin film form. 4. Summary and conclusions
Epitaxial TMO films were successfully grown on (001)-STO substrates by means of DC magnetron sputtering. XRD diffraction analysis showed that films were single-phase and (00ℓ)-oriented. The XPS spectra suggested that Mn exists in the 3 oxidation state in the studied TMO films. The films displayed ferromagnetism below ~40 K. No evidence of a second low-temperature magnetic transition, similar to that associated with the onset of ferroelectricity in bulk TMO, was found. It was verified that epitaxial strain can lead to very different magnetic properties of TMO. The appearance of ferromagnetism in epitaxially-strained TMO films seems to be associated with a modified microstructure of the films with respect to bulk TMO. The results show that strain engineering is an effective method for modulating the multiferroic properties of TMO, which could be important for the development of practical applications based on these thin films. Acknowledgments
This investigation was supported by Universidad Nacional de Colombia, Sede Medellín. J.L. Izquierdo acknowledges the financial support of the “Dirección de Investigaciones” of the Institución Universitaria Pascual Bravo.
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Figure captions Figure 1. (a) 2- scan of a ~100 nm thin TMO films grown on a (001)-oriented STO substrate by means of DC magnetron sputtering. (b) rocking curve around the (002) reflection for the TMO film. Figure 2. Wide XPS scan spectrum of a TMO film grown on an STO substrate. Figure 3. Narrow XPS scan spectra around the Tb 3d (a), Mn 2p, and O 1S (c) edge for a TMO film grown on an STO substrate. Figure 4. Magnetization versus temperature for a polycrystalline TMO sample. The data were recorded under warming conditions in a field of 0.02 T. Inset: temperature variation of inverse magnetic susceptibility with the Curie-Weiss fit (solid line). Figure 5. (a) ZFC and FC temperature variation of the DC susceptibility for a 100 nm thin TMO film. The data were recorded upon warming with a 0.1 T measuring field applied parallel to the film’s plane. Inset: first derivative dDC/dT plotted on an enlarged scale (b) variation of the magnetization with the magnetic field for the TMO film at different temperatures. Inset: M(H) dependence for the TMO film in the low-field range. References [1] T. Kimura, T. Goto, H. Shintani, K. Ishizaka, T. Arima, Y. Tokura, Magnetic control of ferroelectric polarization, Nature 426 (2003) 55. [2] W. Eerenstein, N.D. Mathur, J.F. Scott, Multiferroic and magnetoelectric materials, Nature (London) 442 (2006) 759. [3] D. Rubi, C.D. Graaf, C.J.M. Daumont, D. Mannix, R. Broer, B. Noheda, Ferromagnetism and increased ionicity in epitaxially grown TbMnO3 films, Phys. Rev. B 79 (2009) 014416. [4] M. Staruch, D. Violette, M. Jain, Structural and magnetic properties of multiferroic bulk TbMnO3, Mater. Chem. Phys. 139 (2013) 897. [5] Y. Cui, Y. Tian, A. Shan, C. Chen, R. Wang, Magnetic anisotropy and anomalous transitions in TbMnO3 thin films, Appl. Phys. Lett. 101 (2012) 122406. [6] N.A. Hill, Why Are There so Few Magnetic Ferroelectrics, J. Phys. Chem. B 104 (2000) 6694.
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[7] F. Jiménez-Villacorta, J.A. Gallastegui, I. Fina, X. Marti, J. Fontcuberta, Straindriven transition from E-type to A-type magnetic order in YMnO3 epitaxial films. Phys. Rev. B 86 (2012) 024420. [8] Y. S. Hou, J. H. Yang, X. G. Gong, H. J. Xiang, Prediction of a multiferroic state with large electric polarization in tensile-strained TbMnO3, Phys. Rev. B 88 060406(R) (2013). [9] K. Shimamoto, S. Mukherjee, S. Manz, J.S. White, M. Trassin, M. Kenzelmann, L. Chapon, T. Lippert, M. Fiebig, C.W. Schneider, C. Niedermayer, Tuning the multiferroic mechanisms of TbMnO3 by epitaxial strain Scientific Reports 7 (2017) 44753. [10] H. Wadati, J. Okamoto, M. Garganourakis, V. Scagnoli, U. Staub, Y. Yamasaki, H. Nakao, Y. Murakami, M. Mochizuki, M. Nakamura, M. Kawasaki, Y. Tokura, Origin of the Large Polarization in Multiferroic YMnO3 Thin Films Revealed by Soft- and HardX-Ray Diffraction. Phys. Rev. Lett. 108 (2012) 047203. [11] C.J.M. Daumont, D. Mannix, S. Venkatesan, G. Catalan, D. Rubi, B.J. Kooi, J.T. M. D. Hosson, B. Noheda, Epitaxial TbMnO3 thin films on SrTiO3 substrates: a structural study, J. Phys.: Condens. Matter. 21 (2009) 182001. [12] X. Marti, V. Skumryev, C. Ferrater, M. V. Garcıa-Cuenca, M. Varela, F. Sanchez, and J. Fontcuberta, Emergence of ferromagnetism in antiferromagnetic TbMnO3 by epitaxial strain, Appl. Phys. Lett. 96 (2010) 222505. [13] F. Perez-Osuna, J.M. Siqueiros, A. Durán, M.P. Cruz, L. Salamanca-Riba, J. Heiras, Magnetic properties of Al doped TbMnO3 thin films grown by pulsed laser deposition, J. Appl. Phys. 112 (2012) 033914. [14] A. Sazanovicha, Y. Nikolaenko, W. Paszkowicz, V. Mikhaylov, K. Dyakonov, Y. Medvedev, V. Nizhankovskii, V. Dyakonov, H. Szymczak, Magnetic and Ferroelectric Ordering in the TbMnO3 Film, Acta Physica Polonica A 125 (2014) 128. [15] Andrew H. Simon, in Handbook of Thin Film Deposition (Third Edition), Elsevier (2012) 55. [16] R. Wördenweber, Woodhead Publishing Series in Electronic and Optical Materials (2011) 3. [17] R. Wördenweber, Comprehensive Semiconductor Science and Technology 4 (2011) 177. [18] A. Astudillo, J. Izquierdo, G. Bolanos, O. Moran, Emergence of Ferromagnetism in TbMnO3 Bulk by Al-Doping, IEEE Trans. on Magn. 49 (2013) 4590. [19] J.A. Alonso, M.J. Martínez-Lope, M.T. Casais, M.T. Fernández-Díaz, Evolution of the Jahn−Teller Distortion of MnO6 Octahedra in RMnO3 Perovskites (R=Pr, Nd, Dy, Tb, Ho, Er, Y): A Neutron Diffraction Study, Inorg. Chem. 39 (2000) 917.
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[20] S. Venkatesan, C. Daumont, B. J. Kooi, B. Noheda, and J. T. M. D. Hosson, Nanoscale domain evolution in thin films of multiferroic TbMnO3, Phys. Rev. B 80 (2009) 214111. [21] C. Lu, S. Dong, Z. Xia, H. Luo, Z. Yan, H. Wang, Z. Tian, S. Yuan, T. Wu, J. Liu, Polarization enhancement and ferroelectric switching enabled by interacting magnetic structures in DyMnO3 thin films, Scientific reports 3 (2013) 3374. [22] R. D. Kumar, M. Subramanian, M. Tanemura, R. Jayavel, Synthesis, annealing effect and magnetic behavior of TbMnO3 nanoparticles, J. Nanopart. Res. 16 (2014) 2501. [23] J.F. Moulder, W.F. Stickle, P.E. Sobol, K.D. Bomben, Handbook of X-Ray Photoelectron spectroscopy, Ulvac-Phi-Inc Japan, 1995. [24] V. R. Galakhov, M. Demeter, S. Bartkowski, M. Neumann, N. A. Ovechkina, E. Z. Kurmaev, N. I. Lobachevskaya, Y. M. Mukowskii, J. Mitchell, and D. L. Ederer, Mn 3s exchange splitting in mixed-valence manganites, Phys. Rev. B 65 (2002) 113102. [25] J. Blasco, C. Ritter, J. García, J.M. de Teresa, J. Pérez-Cacho, M.R. Ibarra, Structural and magnetic study of Tb1-xCaxMnO3 perovskites, Phys. Rev. B 62 (2000) 5609 [26] R. Das, A. Jaiswal, S. Adyanthaya, P. Poddar, Effect of particle size and annealing on spin and phonon behavior in TbMnO3, J. Appl. Phys. 109 (2011) 064309. [27] V. Dyakonov, Low Temp. Phase transitions in TbMnO3 manganites, Low Temp. Phys. 38 (2012) 216. [28] R. Kajimoto, H. Yoshizawa, H. Shintani, T. Kimura, and Y. Tokura, Magnetic structure of TbMnO3 by neutron diffraction, Phys. Rev. B 70 (2004) 012401. [29] B.J. Kirby, D. Kan, A. Luykx, M. Murakami, D. Kun-daliya, I. Takeuchi, Anomalous ferromagnetism in TbMnO3 thin films, J. Appl. Phys. 105 (2009) 07D917. [30] X. Marti, V. Skumryev, V. Laukhin, F. Sánchez, M.V. García-Cuenca, C. Ferrater, M. Varela, J. Fontcuberta, Dielectric anomaly and magnetic response of epitaxial orthorhombic YMnO3 thin films, J. Mater. Res. 22 (2007) 2096. [31] D. Rubi, S. Venkatesan, B. J. Kooi, J. Th. M. De Hosson, T. T. M. Palstra, B. Noheda, Magnetic and dielectric properties of YbMnO3 perovskite thin films, Phys. Rev. B 78 (2008) 020408(R). [32] Y.S. Hou, J.H. Yang, X.G. Gong, H.J. Xiang, Prediction of a multiferroic state with large electric polarization in tensile-strained TbMnO3, Phys. Rev. B 88 (2013) 060406(R). [33] Y. Tokura, Y. Tomioka, Colossal magnetoresistive manganites, J. Magn. Magn. Mater. 200 (1999) 1. 15
[34] Y. Tokura, Critical features of colossal magnetoresistive manganites, Rep. Prog. Phys. 69 (2006) 797. [35] M. Kenzelmann, A. B. Harris, S. Jonas, C. Broholm, J. Schefer, S. B. Kim, C. L. Zhang, S.-W. Cheong, O. P. Vajk, J.W. Lynn, Magnetic Inversion Symmetry Breaking and Ferroelectricity in TbMnO3, Phys. Rev. Lett. 95 (2005) 087206.
10
1
20
40 2-
(a)
60
40 FWHM=0.4°
0
23
Tb -3d3/2
C 1S
0,5
Tb -3d5/2
O 1S
5
Intensity [x10 c/s]
1,0
Mn -2p3/2 Mn -2p1/2
Figure 1. Morán
1,5
500
(b)
80 TMO (002)
Intensity [c/s]
3
TMO (004)
10
STO (002)
5
TMO (002)
10
STO (001)
Intensity [c/s]
[36] R. Eremina, Z. Seidov, I. Ibrahimov, M. Najafzade, M. Aljanov, D. Mamedova, T. Gavrilova, I. Gilmutdinov, V. Chichkov, N. Andreev, Magnetization of manganite thin films on ferroelectric substrates, J. Magn. Mag. Mater. 440 (2017) 179.
1000 B.E. [eV]
16
24
25
Fig. 2. Morán
+3
Mn 2p1/2
655
1260 1240 B.E. [eV]
[x10 T/f.u.]
Curie-Weiss fit
1
5
-1
-2
M [x10 B/f.u.]
Fig. 2. Morán
2
1
0
42 K
0
200 T [K]
FC
0
H=0.02 T
0
100
200 T [K]
300
Figure 4. Morán
17
(c)
Mn 2p3/2
Intensity [a.u.]
Intensity [a.u.] 1280
+3
(b)
+3
Tb 3d3/2
Intensity [a.u.]
+3
(a) Tb 3d5/2
650 645 B.E. [eV]
640
OL
OH
534
532 530 B.E. [eV]
52
FC ZFC
M [B/Mn]
d(M/H)/dT
M [B/Mn]
d(M/H)/dT
0.1
0.0
2
10 T [K] 20
30 40 50 60 T [K]
20
T [K]
40
1 0
-1
H=0.1 T
0
(b)
40 K 20 K 5K
M [B/Mn]
(a)
-2
60
0,2 0,0
-0,2 -0,2
0,0
0,2
H [T]
-6
-3
0 3 H [T]
6
Fig. 5. Morán
Highlights
Title: Evidence of ferromagnetic ordering at low temperatures in epitaxial TbMnO3 thin films grown by means of DC magnetron sputtering
Epitaxial TbMnO3 thin films are grown via dc magnetron sputtering.
The TbMnO3 films are epitaxial and (00ℓ)-oriented.
The nominal valence of the Mn ions is 3+
The films show well-defined transition at T<42 K
Weak ferromagnetism is observed in undoped TbMnO3 films
Hysteresis loops with coercive fields of 1000 Oe are observed at 5 K
18
Conflict of interest statement, in good faith, the authors of the manuscript Evidence of ferromagnetic ordering at low temperatures in epitaxial TbMnO3 thin films grown by means of DC magnetron sputtering declare that there is no conflict of interest arising from conflicting financial or other interests.
Yours sincerely
Dr. Oswaldo Morán, corresponding author Physics Department, National University of Colombia, Medellín Campus A. A. 568, Medellín, Colombia e-mail: [email protected] Phone: 57-4-4309327 Fax: 57-4- 2604489
19
S0304-8853(19)31041-8 https://doi.org/10.1016/j.jmmm.2019.166141 MAGMA 166141
To appear in:
Journal of Magnetism and Magnetic Materials
Received Date: Revised Date: Accepted Date:
20 March 2019 14 October 2019 11 November 2019
Please cite this article as: J.L. Izquierdo, A. Astudillo, J. Martínez, G. Bolaños, O. Morán, Evidence of ferromagnetic ordering at low temperatures in epitaxial TbMnO3 thin films grown by means of DC magnetron sputtering, Journal of Magnetism and Magnetic Materials (2019), doi: https://doi.org/10.1016/j.jmmm.2019.166141
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© 2019 Published by Elsevier B.V.
Evidence of ferromagnetic ordering at low temperatures in epitaxial TbMnO3 thin films grown by means of DC magnetron sputtering J. L. Izquierdo1, A. Astudillo2, J. Martínez2, G. Bolaños2, O. Morán3 1
Institución Universitaria Pascual Bravo, Centro de Investigación y Desarrollo en Materialografía, Calle 73 No. 73A-226, C.P. 050001 Medellín, Colombia. 2 Universidad del Cauca, Departamento de Física, Grupo de bajas temperaturas, A.A. 1226 Popayán, Colombia. 3 Universidad Nacional de Colombia, Campus Medellín, Facultad de Ciencias, Departamento de Física, Advanced Oxides Group, A.A. 568 Medellín, Colombia Abstract. Thin films of the archetypal multiferroic TbMnO3 were epitaxially deposited on (001)oriented SrTiO3 substrates by means of DC magnetron sputtering. A substrate temperature of 1053 K and an oxygen pressure of 300 Pa were determined to be appropriate growing conditions for obtaining high-quality, epitaxial TbMnO3 films using this physical growing technique. The structural and magnetic properties of ~100 nm thin TbMnO3 films were investigated. X-ray diffraction patterns showed that the films are single-phase and (00ℓ)-oriented. Rocking curves of the (002) peak confirm the good crystalline quality of the films. X-ray photoelectron spectroscopy analysis confirmed that the nominal valence of the Mn ions is 3+. The films grown on the SrTiO3 substrates displayed substantial epitaxial strain, which decidedly changed the magnetic ground state of the bulk-like counterpart. In particular, the films displayed ferromagnetic-like interactions below a temperature close to the magnetic ordering temperature of the Mn3+ spins (TN42 K). Well-defined hysteresis loops with coercive fields as high as 0.1 T were observed at 5 K. The variation of the coercive field with the temperature clearly indicated the presence of low-temperature ferromagnetism in the TMO films. The origin of the anomalous ferromagnetism in the films is discussed in terms of the strain-induced distortion generated by the lattice mismatch between the film and substrate. The reported results can give insight into accurate strategies for controlling single-phase spin-driven multiferroic states. Corresponding author: E-mail: [email protected] Phone: 57-4-4309327 Fax: 57-4- 2604489
1
1. Introduction Multiferroic materials have attracted considerable attention over the last decade after observations of the coexistence of ferromagnetic and ferroelectric polarizations, as well as the coupling effect between them [1]. Apart from their fascinating physical properties, these materials could also have technological applications [2]. Among the many multiferroic materials, orthorhombic TbMnO3 (TMO) is one of the few in which ferroelectricity arises from the magnetic spiral order (spin-driven ferroelectrics) [3]. Bulk-like TMO is a multiferroic material that exhibits a complex magnetoelectric phase diagram, including antiferromagnetism below ~40 K. [4]. However, scaling TMO crystals into thin film form may generate epitaxial strain that would drastically affect their multiferroic properties [5]. In this regard, the possibility of tuning the exchange interactions and inducing ferromagnetic behavior via strain could lead to single-phase ferrimagnetic ferroelectrics [6]. Such sophisticated devices are highly interesting for a variety of applications and are not commonly reported in the literature. The fact that epitaxial strain may alter the magnetic properties of rare-earth-based manganites when they are grown on crystalline substrates via physical or chemical methods has been demonstrated in recent papers [7]. Generally speaking, the origin of the ferromagnetism in epitaxially strained TMO films seems to be related to a change in the electronic structure of the material [8]. From the analysis of the results of neutron diffraction experiments, it has been established that in highly strained films, the magnetic order changes from the bulk-like incommensurate bc-cycloidal structure to the commensurate magnetic order [9]. This change in the magnetic order can lead to a complex coexistence of magnetic order parameters [10]. In the case of TMO thin films, a pronounced increase of the ferroelectric polarization and a change in its direction from along the c- to the a-axis was observed, which seemed to be concomitant with the
2
magnetic order change in strained TMO films [9]. Competing magnetic interactions at low temperatures tuned by small lattice distortions could be the mechanism responsible for the switchable electric polarization in TMO thin films [9]. Hence it is possible that the complex ferromagnetic properties reported for TMO thin films could be the result of strain-induced deformation of the unit cell and the subsequent unbalancing of magnetic interactions [1]. As is evident from the previous paragraph, the fabrication of high-quality thin films is a fundamental step toward the possibility of controlling the different multiferroic phases in novel oxides by means of the epitaxial strain, dimensionality, and size effect, without varying the chemical composition. In most of the previous investigations reporting on the preparation, characterization, and multiferroic function of TMO thin films, pulsed laser deposition (PLD) was employed as a physical method for depositing the thin layers on crystalline substrates [11–13]. The reported results are diverse. Nevertheless, reports on the growth of epitaxial TMO thin films by other methods are scarce [14]. Although the PLD technique definitely has several advantages, such as high deposition rate, short test period, and low substrate temperature requirements, the deposited films can be affected by small molten particles or target fragments, which are generated during the laser-induced explosion [15]. The presence of these particles can severely diminish the quality of the films. Sputter deposition, although it equally presents several difficulties, such as the differentiation between the nature of the compound target (whether metal or insulator), is a versatile and inexpensive physical vapor deposition method for preparing thin films of alloys and complex materials for research and industrial purposes [16]. Generally speaking, sputter deposition can be considered to be a natural method for the deposition of complex ceramic materials [16]. The large particle energy, the atomic transfer of the target materials, and the possibility of using
3
reactive gas components (reactive sputtering) at a large partial pressure open up a broad parameter range for growing high-quality epitaxial thin films and superlattices of complex oxides, often at a significantly lower cost than other techniques. Sputter is compatible with high oxygen-pressure, and calibration and rate control are readily obtained. Films grown via sputter techniques are highly reproducible, and a homogeneous composition can be obtained over a rather large area (typically 6.4x10-3 m) [17]. With each wave of interest in a particular class of complex oxide materials, a route to high-quality oxide films through sputtering has usually been found, and the resulting films have been important for scientific progress. In the present investigation, we report on the successful growth of c-axis-oriented epitaxial TMO films on (001)oriented SrTiO3 (STO) substrates by means of DC magnetron sputtering. Specifically, it is shown that the magnetic ground state of TMO thin films is clearly modified by epitaxial strain. 2. Experimental details The preparation of the TMO targets was described in a previous paper [18]. The singlephase TMO targets were coupled to a sputter cannon and pre-sputtered before film deposition in order to remove possible contaminated surface layers. (001)-oriented STO was chosen as substrate material. STO possesses a cubic perovskite structure, and the lattice constant is 3.905 Å. TMO films (~100 nm) were deposited in an oxygen atmosphere at a pressure of 300 Pa, a power of 30 W, and a substrate temperature of 1053 K. After deposition, the substrate temperature was lowered to 723 K, and the chamber was then flooded with oxygen at a pressure of ~8x104 Pa. The films were annealed under these conditions for 1 h. X-ray diffraction patterns of the TMO films were registered with a four-circle Rigaku diffractometer using CuKα radiation (λ=0.15418 nm). The elemental composition of the films and the depth profile with high
4
resolution spectra were studied by means of X-ray photoelectron spectroscopy (XPS) using a multi-technique scanning XPS microprobe (PHI Versa Probe II). The source of X-rays used was a monochromatic AlK lamp (300 W) at a photon energy of 2.38x10-16 J. The photoelectron detection was carried out by means of a concentric hemispherical analyzer (CHA) (PHI 10360 Model) operating with an analytical aperture 0.8 m in diameter. The binding energy resolution of the instrument is 1.12x10-19 J. Magnetization (M) as a function of temperature and magnetic field (H) was measured using a superconducting
quantum
interference
device
(SQUID,
Quantum
Design)
magnetometer. The applied magnetic field was always kept parallel to the plane of the film. The field-cooled (FC) and zero-field-cooled (ZFC) magnetization measurements were carried out on heating from 5 to 350 K, after the sample was first cooled from 300 K down to the lower temperature with and without the magnetic field. The isothermal M(H) curves were recorded at 40, 20, and 5 K after FC (7 T) from 300 K. 3. Results and discussion
Figure 1(a) shows the XRD pattern of one typical ~100 nm thin TMO film grown on a (001)-oriented STO substrate by means of DC magnetron sputtering. The film is pure caxis oriented and free from secondary phases. The displacement of the peaks of the substrate and the film is due to the mismatch between the lattice parameters of the film and those of the substrate. From the XRD pattern, the out-of-plane lattice parameter was found to be 0.745 nm, which is slightly larger than the c0.740 nm for TMO bulk single crystals [19]. This finding suggests that the films that were grown were subjected to an in-plane compressive strain, which resulted in the tetragonally distorted orthorhombic phase of TMO. The achieved result agrees with previous reports on TMO thin films grown on STO substrates [20]. Figure 1(b) shows the rocking curve around the (002) reflection for the TMO film. The full width at half maximum is 0.4°. Here it is 5
important to note that whereas TMO possesses an orthorhombic structure with lattice parameters a05.293 Å, b05.838 Å, and c07.403 Å [19], the STO substrates have a cubic crystalline structure with lattice parameter a=3.905 Å. Therefore, a large lattice misfit should be taken into consideration in the epitaxial growth of TMO on STO substrates (-4.1 % along the a axis and 5.7 % along the b axis). Hence it is to be expected that the orhorhombic unit cell of TMO be grown on a 3.9052 Åx3.9052 Å square STO lattice, which can better accommodate the in-plane film lattice. This growth model has previously been suggested for TMO and DyMnO3 films deposited on STO substrates [11, 21]. In this model, the strained TMO films should share the pseudocubic lattice parameters of the substrate and therefore should be compressed in the plane directions from about 3.94 Å of the bulk to 3.905 Å [3]. In Ref [3], a deviation of 90° of the pseudocubic angle of the TMO films was reported. This deviation, which represents the orthorhombic distortion, was determined to be smaller than that of the bulk counterpart. As a result, these TMO films grew with a highly-compressed orthorhombic bo and a slightly expanded ao axis. Therefore, the enlargement of the out-of-plane co lattice parameter in the TMO films is consistent with an overall compressive in-plane strain. The oxidation states of the Tb 3d, Mn 2p, and O 1s in the TMO films was verified via XPS measurements. The peak positions were referenced to carbon at 4.61x10-17 J. Figure 2 shows a wide scan spectrum of a representative TMO film. Peaks of Tb, Mn, O, and C can clearly be observed. The narrow scan spectra for Tb 3d is shown in Figure 3(a). Two characteristic peaks can be seen at binding energies of 1.98x10-16 J (Tb 3d5/2) and 2.04x10-16 J (Tb 3d3/2), in agreement with values previously reported [22]. In turn, the characteristic peaks of Mn3+ 2p3/2 (1.02x10-16 J) and 2p1/2 (1.04x10-16 J) are shown in Fig. 3(b). The binding energy values of the Mn3+ agree with those already reported for
6
the compound [3,22]. The experimental peaks were successfully fit, as indicated by the solid lines in Fig. 4(a) and (b). In particular, the absence of a secondary peaks (shake-up peaks) in the Mn 2p XPS edge suggests that Mn exists in the 3 oxidation state in the TMO films grown on STO substrates by means of DC magnetron sputtering. Here it is useful to point out that Mn is normally characterized by a 3 oxidation state in Mnbased perovskite oxides [22]. The Mn 3p level splitting can be seen directly in Fig. 3(b) and equals 1.89x10-18 J. The experimental reference values reported for Mn 2p ions is exactly 1.89x10-18 J [23]. Thus the absence of shake peaks in the Mn 2p spectrum and the magnitude of the Mn 2p splitting discount the existence of mixed +2/+3 valence in the films. The existence of mixed +2/+3 valence in Mn-based compound is commonly associated with the presence of oxygen vacancies in the sample [24]. Therefore, these TMO films grown via DC magnetron sputtering do not seem to be affected by this kind of defect. Figure 3(c) shows the O 1s XPS spectrum of the TMO film. A pronounced peak and a clear shoulder are located at 8.49x10-17 and 8.53x10-17 J, respectively. Such splitting indicates that the oxygen in the TMO films exists in two kinds of chemical states. The binding energy at 8.49x10-17 J eV is mainly due to the contribution of the crystal lattice oxygen (OL) [22]. The OL signal probably stems from the contribution of Tb–O and Mn–O in the TMO lattice. In turn, the shoulder at 8.53x10-17 J can arise from chemisorbed oxygen species (OH) on the surface (e.g. chemisorbed water) [22]. The temperature-dependent FC magnetization (M) data of the polycrystalline TMO target are shown in Fig. 4. The data were recorded in a field of 0.02 T. A pronounced increase in the value of the magnetization can be seen at T<45 K, which is close to the TN bulk value of TMO (T~42 K). Below TN, the Mn3+ sublattice starts to develop an incommensurate sinusoidal antiferromagnetic (AFM) alignment [25]. Superexchange interaction due to the spin reorientation of Mn3+ ions is considered to be the mechanism
7
behind this phenomenon [26]. The inverse magnetic susceptibility, -1, versus the temperature for the polycrystalline TMO target is plotted in the inset of Fig. 4. It can be seen that the experimental data follow the Curie-Weiss law =C/(T-) within the temperature range 42–350 K. Thus the Curie-Weiss law is fulfilled down to 42 K, below which the Mn3+ spins of TMO start to develop an incommensurate sinusoidal AFM structure [25]. The values of the Curie constant C and the Curie-Weiss temperature were found to be 620 Am2K/kgT and -34.6 K, respectively. The AFM ordering at low temperatures is confirmed by the negative value of . The values found for the parameters C and are in accordance with those commonly reported in the literature on the compound [27]. On the other hand, the incommensurate-commensurate transition Tlock, occurring at ~28 K, was not observed in the M(T) (or -1(T)) measurements. The transition at Tlock is accompanied by magnetoelastically-induced lattice modulation and the appearance of ferroelectricity [2]. Probably the dominant paramagnetic signal of the Tb spins, which possess a larger magnetic moment than the Mn spins, causes this transition not to be observed [8]. The ZFC and FC in-plane DC susceptibility as a function of the temperature for a 100 nm thin TMO film is presented in Fig. 5(a). The data were recorded in a 0.1 T Oe magnetic field with a cooling field of 7 T for the FC mode. The diamagnetic contribution from the substrate was subtracted from the raw magnetization data. The ZFC and the FC curves show a distinct anomaly at 43 K (inset of Fig. 5(a)), which is very close to the anomalies in the C(T) and M(T) dependences observed in TMO single crystals and ascribed to the transition from the paramagnetic phase to the sinusoidal antiferromagnetic structure of the manganese spins [1]. Since the Tb3+ moments are located in the ab plane of a single crystal [28], it is possible that the peak appearing at
8
~10 K (inset of Fig. 5(a)) is associated with the Tb3+ spin ordering in the strained films (the field is applied parallel to the film plane). Moreover, the magnetic measurements did not reveal any feature related to the stabilization of the spiral antiferromagnetic ordering, concomitant with the onset of ferroelectricity, which occurs in bulk at ∼27 K [1]. A split between FC and ZFC measurements can also be observed below the Mn3 spin ordering temperature. This effect suggests that ferromagnetic interactions are present in the films [3]. Splitting between the FC and ZFC modes reported for TMO in bulk form has been attributed to extrinsic effects related to uncompensated moments at the grain boundaries and clearly depends on the magnetic field strength [4]. The existence of the low-temperature FM phase in the TMO films is verified by the clearly hysteretic behavior observed at low temperatures in the 6 T field range after cooling from 300 K in a 7 T magnetic field (Fig. 5(b)). A coercive field of about 0.1 T at 5 K (inset of Fig. 5(b)) confirms the presence of a FM order at low temperatures in the analyzed films. This behavior drastically departs from the antiferromagnetic order observed for bulk TMO. The results obtained for the TMO films are in good agreement with
polarized
neutron
reflectometry
measurement
data
[29].
Interestingly,
orthorhombic YbMnO3 and YMnO3 thin films exhibit similar effects, which speaks to a general mechanism in manganite thin films [30, 31]. As stated above, the weak ferromagnetism present in the TMO films has no counterpart in the bulk. In particular, the magnetic measurements showed evidence of a transition that is very close to the magnetic ordering temperature (TN42 K) of the Mn3+ spins. Nevertheless, the measurements did not show clear evidence of stabilization of the spiral antiferromagnetic ordering and the concomitant onset of ferroelectricity, as in the bulk counterpart. Hence it is apparent that the observed ferromagnetism in the thin TMO films should be explained by another scenario. Generally speaking, the lattice misfit
9
between the films and the substrates generates strain effects in the thin films. A modified microstructure is the direct consequence of this effect. Following this line of argument, it is probable that the origin of the ferromagnetism in epitaxially strained TMO films is related to a change of the electronic structure in the material [32]. For bulk RMnO3 compounds, it is well stablished that each Mn3+ ion has four unpaired d electrons [1]. Here, the low-energy t2g orbitals are occupied by three electrons and the high-energy eg orbitals by only one. Moreover, the two degenerate eg states on each spin site split into two nondegenerate orbitals as a consequence of the Jahn-Teller distortion. Only one of these two nondegenerate orbitals is occupied. In this case, the occupied eg orbitals on two adjacent spin sites are almost orthogonal to each other in the ab plane [33]. On the other hand, it has been suggested that the orbital order in epitaxiallystrained TMO films is very different [12]. First, the two eg orbitals on each spin site are almost degenerate. This implies that both can be considered to be half-occupied, and the electron configuration of each Mn3+ (d4) ion could roughly be described as (t2g)3(d3z2−r2)1/2(dx2−y2)1/2. Such electron configuration leads to slightly different Mn-O bond lengths in the different directions of the MnO6 octahedron (2, 2.02, and 2.08 Å in the x, y, and z directions, respectively.) The O-Mn-O bond angles are close to 90° or 180°. It is interesting to analyze the interaction between the eg orbitals of two adjacent spin sites along the c direction for the FM arrangement. In this case, it has been determined that the majority-spin eg (d3z2−r2) orbitals of both sites with the same energy level can couple with each other [33]. This will lead to the formation of an empty higher-energy level and an occupied lower-energy level. Thus the total energy will be reduced by the quantity t (>0), which represents the hopping integral between these two d3z2−r2 orbitals. For the AFM configuration, it has also been established that the majority-spin eg orbital of one Mn site couples with the minority-spin eg orbital of the
10
other Mn site [34]. Hence the reduction in the value of the total energy can be expressed as U/2-(t2 + U2/ 4)1/2. In this equation, U (>0) corresponds to the energy difference between eg orbitals with opposite spin components. From these results, it is evident that the FM state has a lower energy. Interestingly, it has been verified that the in-plane exchange interaction Jab is governed by the same mechanism [35]. Recalling that symmetric exchange interactions between spins of Mn3+ ions characterize a ferromagnetic ground state of materials with perovskite-type structures [33], it is apparent that ferromagnetic interactions are favored in epitaxially-strained TMO films. In addition, it has also been established that nearest-neighbor Mn-Mn distances in strained TMO films decrease and Mn-O-Mn angles become larger as compared to those of the bulk counterpart [32]. This will cause as consequence the enhancement of the coupling of eg orbitals (increase of t) and hence an increase in the magnitude of the spin exchange interaction parameters, which is consistent with the increase in the symmetric spin exchange parameters. On the other hand, some authors have suggested that strain effects alone cannot account for the difference between the magnetic properties of the bulk material and thin films [31]. The morphology was then considered to be an additional factor influencing the magnetic response of epitaxial TMO films. In a recent and interesting investigation, the influence of a ferroelectric STO substrate on the magnetic properties of GdMnO3 and YbMnO3 thin films fabricated via RF-magnetron sputtering was reported [36]. There, a clear difference in the values of the DC susceptibility of GdMnO3 films grown on ferroelectric STO and LiNbO3 substrates was observed. A difference in the value of the splitting temperature between the ZFC and FC curves was also evident for the films grown on the two substrates. The influence of the polarization vector of the substrate on atomic positions at the interface and the consequent change in the magnitude and direction of the magnetic moments in the
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GdMnO3 films was suggested as a possible mechanism behind the observed differences in their magnetic behavior. According to the authors, this mechanism is supported by the fact that the permittivity values of STO are one order of magnitude higher than those of LiNbO3. Although a rather homogeneous magnetization profile of orthorhombic TMO films, with some increase in the magnetic moment at the bottom interface, has been verified by means of spin-polarized neutron reflectometry [29], it is evident that a more complex analysis of the full structure of the films is needed to explain the observed ferromagnetism in this material when grown in thin film form. 4. Summary and conclusions
Epitaxial TMO films were successfully grown on (001)-STO substrates by means of DC magnetron sputtering. XRD diffraction analysis showed that films were single-phase and (00ℓ)-oriented. The XPS spectra suggested that Mn exists in the 3 oxidation state in the studied TMO films. The films displayed ferromagnetism below ~40 K. No evidence of a second low-temperature magnetic transition, similar to that associated with the onset of ferroelectricity in bulk TMO, was found. It was verified that epitaxial strain can lead to very different magnetic properties of TMO. The appearance of ferromagnetism in epitaxially-strained TMO films seems to be associated with a modified microstructure of the films with respect to bulk TMO. The results show that strain engineering is an effective method for modulating the multiferroic properties of TMO, which could be important for the development of practical applications based on these thin films. Acknowledgments
This investigation was supported by Universidad Nacional de Colombia, Sede Medellín. J.L. Izquierdo acknowledges the financial support of the “Dirección de Investigaciones” of the Institución Universitaria Pascual Bravo.
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Figure captions Figure 1. (a) 2- scan of a ~100 nm thin TMO films grown on a (001)-oriented STO substrate by means of DC magnetron sputtering. (b) rocking curve around the (002) reflection for the TMO film. Figure 2. Wide XPS scan spectrum of a TMO film grown on an STO substrate. Figure 3. Narrow XPS scan spectra around the Tb 3d (a), Mn 2p, and O 1S (c) edge for a TMO film grown on an STO substrate. Figure 4. Magnetization versus temperature for a polycrystalline TMO sample. The data were recorded under warming conditions in a field of 0.02 T. Inset: temperature variation of inverse magnetic susceptibility with the Curie-Weiss fit (solid line). Figure 5. (a) ZFC and FC temperature variation of the DC susceptibility for a 100 nm thin TMO film. The data were recorded upon warming with a 0.1 T measuring field applied parallel to the film’s plane. Inset: first derivative dDC/dT plotted on an enlarged scale (b) variation of the magnetization with the magnetic field for the TMO film at different temperatures. Inset: M(H) dependence for the TMO film in the low-field range. References [1] T. Kimura, T. Goto, H. Shintani, K. Ishizaka, T. Arima, Y. Tokura, Magnetic control of ferroelectric polarization, Nature 426 (2003) 55. [2] W. Eerenstein, N.D. Mathur, J.F. Scott, Multiferroic and magnetoelectric materials, Nature (London) 442 (2006) 759. [3] D. Rubi, C.D. Graaf, C.J.M. Daumont, D. Mannix, R. Broer, B. Noheda, Ferromagnetism and increased ionicity in epitaxially grown TbMnO3 films, Phys. Rev. B 79 (2009) 014416. [4] M. Staruch, D. Violette, M. Jain, Structural and magnetic properties of multiferroic bulk TbMnO3, Mater. Chem. Phys. 139 (2013) 897. [5] Y. Cui, Y. Tian, A. Shan, C. Chen, R. Wang, Magnetic anisotropy and anomalous transitions in TbMnO3 thin films, Appl. Phys. Lett. 101 (2012) 122406. [6] N.A. Hill, Why Are There so Few Magnetic Ferroelectrics, J. Phys. Chem. B 104 (2000) 6694.
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[7] F. Jiménez-Villacorta, J.A. Gallastegui, I. Fina, X. Marti, J. Fontcuberta, Straindriven transition from E-type to A-type magnetic order in YMnO3 epitaxial films. Phys. Rev. B 86 (2012) 024420. [8] Y. S. Hou, J. H. Yang, X. G. Gong, H. J. Xiang, Prediction of a multiferroic state with large electric polarization in tensile-strained TbMnO3, Phys. Rev. B 88 060406(R) (2013). [9] K. Shimamoto, S. Mukherjee, S. Manz, J.S. White, M. Trassin, M. Kenzelmann, L. Chapon, T. Lippert, M. Fiebig, C.W. Schneider, C. Niedermayer, Tuning the multiferroic mechanisms of TbMnO3 by epitaxial strain Scientific Reports 7 (2017) 44753. [10] H. Wadati, J. Okamoto, M. Garganourakis, V. Scagnoli, U. Staub, Y. Yamasaki, H. Nakao, Y. Murakami, M. Mochizuki, M. Nakamura, M. Kawasaki, Y. Tokura, Origin of the Large Polarization in Multiferroic YMnO3 Thin Films Revealed by Soft- and HardX-Ray Diffraction. Phys. Rev. Lett. 108 (2012) 047203. [11] C.J.M. Daumont, D. Mannix, S. Venkatesan, G. Catalan, D. Rubi, B.J. Kooi, J.T. M. D. Hosson, B. Noheda, Epitaxial TbMnO3 thin films on SrTiO3 substrates: a structural study, J. Phys.: Condens. Matter. 21 (2009) 182001. [12] X. Marti, V. Skumryev, C. Ferrater, M. V. Garcıa-Cuenca, M. Varela, F. Sanchez, and J. Fontcuberta, Emergence of ferromagnetism in antiferromagnetic TbMnO3 by epitaxial strain, Appl. Phys. Lett. 96 (2010) 222505. [13] F. Perez-Osuna, J.M. Siqueiros, A. Durán, M.P. Cruz, L. Salamanca-Riba, J. Heiras, Magnetic properties of Al doped TbMnO3 thin films grown by pulsed laser deposition, J. Appl. Phys. 112 (2012) 033914. [14] A. Sazanovicha, Y. Nikolaenko, W. Paszkowicz, V. Mikhaylov, K. Dyakonov, Y. Medvedev, V. Nizhankovskii, V. Dyakonov, H. Szymczak, Magnetic and Ferroelectric Ordering in the TbMnO3 Film, Acta Physica Polonica A 125 (2014) 128. [15] Andrew H. Simon, in Handbook of Thin Film Deposition (Third Edition), Elsevier (2012) 55. [16] R. Wördenweber, Woodhead Publishing Series in Electronic and Optical Materials (2011) 3. [17] R. Wördenweber, Comprehensive Semiconductor Science and Technology 4 (2011) 177. [18] A. Astudillo, J. Izquierdo, G. Bolanos, O. Moran, Emergence of Ferromagnetism in TbMnO3 Bulk by Al-Doping, IEEE Trans. on Magn. 49 (2013) 4590. [19] J.A. Alonso, M.J. Martínez-Lope, M.T. Casais, M.T. Fernández-Díaz, Evolution of the Jahn−Teller Distortion of MnO6 Octahedra in RMnO3 Perovskites (R=Pr, Nd, Dy, Tb, Ho, Er, Y): A Neutron Diffraction Study, Inorg. Chem. 39 (2000) 917.
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[20] S. Venkatesan, C. Daumont, B. J. Kooi, B. Noheda, and J. T. M. D. Hosson, Nanoscale domain evolution in thin films of multiferroic TbMnO3, Phys. Rev. B 80 (2009) 214111. [21] C. Lu, S. Dong, Z. Xia, H. Luo, Z. Yan, H. Wang, Z. Tian, S. Yuan, T. Wu, J. Liu, Polarization enhancement and ferroelectric switching enabled by interacting magnetic structures in DyMnO3 thin films, Scientific reports 3 (2013) 3374. [22] R. D. Kumar, M. Subramanian, M. Tanemura, R. Jayavel, Synthesis, annealing effect and magnetic behavior of TbMnO3 nanoparticles, J. Nanopart. Res. 16 (2014) 2501. [23] J.F. Moulder, W.F. Stickle, P.E. Sobol, K.D. Bomben, Handbook of X-Ray Photoelectron spectroscopy, Ulvac-Phi-Inc Japan, 1995. [24] V. R. Galakhov, M. Demeter, S. Bartkowski, M. Neumann, N. A. Ovechkina, E. Z. Kurmaev, N. I. Lobachevskaya, Y. M. Mukowskii, J. Mitchell, and D. L. Ederer, Mn 3s exchange splitting in mixed-valence manganites, Phys. Rev. B 65 (2002) 113102. [25] J. Blasco, C. Ritter, J. García, J.M. de Teresa, J. Pérez-Cacho, M.R. Ibarra, Structural and magnetic study of Tb1-xCaxMnO3 perovskites, Phys. Rev. B 62 (2000) 5609 [26] R. Das, A. Jaiswal, S. Adyanthaya, P. Poddar, Effect of particle size and annealing on spin and phonon behavior in TbMnO3, J. Appl. Phys. 109 (2011) 064309. [27] V. Dyakonov, Low Temp. Phase transitions in TbMnO3 manganites, Low Temp. Phys. 38 (2012) 216. [28] R. Kajimoto, H. Yoshizawa, H. Shintani, T. Kimura, and Y. Tokura, Magnetic structure of TbMnO3 by neutron diffraction, Phys. Rev. B 70 (2004) 012401. [29] B.J. Kirby, D. Kan, A. Luykx, M. Murakami, D. Kun-daliya, I. Takeuchi, Anomalous ferromagnetism in TbMnO3 thin films, J. Appl. Phys. 105 (2009) 07D917. [30] X. Marti, V. Skumryev, V. Laukhin, F. Sánchez, M.V. García-Cuenca, C. Ferrater, M. Varela, J. Fontcuberta, Dielectric anomaly and magnetic response of epitaxial orthorhombic YMnO3 thin films, J. Mater. Res. 22 (2007) 2096. [31] D. Rubi, S. Venkatesan, B. J. Kooi, J. Th. M. De Hosson, T. T. M. Palstra, B. Noheda, Magnetic and dielectric properties of YbMnO3 perovskite thin films, Phys. Rev. B 78 (2008) 020408(R). [32] Y.S. Hou, J.H. Yang, X.G. Gong, H.J. Xiang, Prediction of a multiferroic state with large electric polarization in tensile-strained TbMnO3, Phys. Rev. B 88 (2013) 060406(R). [33] Y. Tokura, Y. Tomioka, Colossal magnetoresistive manganites, J. Magn. Magn. Mater. 200 (1999) 1. 15
[34] Y. Tokura, Critical features of colossal magnetoresistive manganites, Rep. Prog. Phys. 69 (2006) 797. [35] M. Kenzelmann, A. B. Harris, S. Jonas, C. Broholm, J. Schefer, S. B. Kim, C. L. Zhang, S.-W. Cheong, O. P. Vajk, J.W. Lynn, Magnetic Inversion Symmetry Breaking and Ferroelectricity in TbMnO3, Phys. Rev. Lett. 95 (2005) 087206.
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1
20
40 2-
(a)
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40 FWHM=0.4°
0
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Tb -3d3/2
C 1S
0,5
Tb -3d5/2
O 1S
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Intensity [x10 c/s]
1,0
Mn -2p3/2 Mn -2p1/2
Figure 1. Morán
1,5
500
(b)
80 TMO (002)
Intensity [c/s]
3
TMO (004)
10
STO (002)
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TMO (002)
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STO (001)
Intensity [c/s]
[36] R. Eremina, Z. Seidov, I. Ibrahimov, M. Najafzade, M. Aljanov, D. Mamedova, T. Gavrilova, I. Gilmutdinov, V. Chichkov, N. Andreev, Magnetization of manganite thin films on ferroelectric substrates, J. Magn. Mag. Mater. 440 (2017) 179.
1000 B.E. [eV]
16
24
25
Fig. 2. Morán
+3
Mn 2p1/2
655
1260 1240 B.E. [eV]
[x10 T/f.u.]
Curie-Weiss fit
1
5
-1
-2
M [x10 B/f.u.]
Fig. 2. Morán
2
1
0
42 K
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200 T [K]
FC
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0
100
200 T [K]
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Figure 4. Morán
17
(c)
Mn 2p3/2
Intensity [a.u.]
Intensity [a.u.] 1280
+3
(b)
+3
Tb 3d3/2
Intensity [a.u.]
+3
(a) Tb 3d5/2
650 645 B.E. [eV]
640
OL
OH
534
532 530 B.E. [eV]
52
FC ZFC
M [B/Mn]
d(M/H)/dT
M [B/Mn]
d(M/H)/dT
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10 T [K] 20
30 40 50 60 T [K]
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40
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-2
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0,2 0,0
-0,2 -0,2
0,0
0,2
H [T]
-6
-3
0 3 H [T]
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Fig. 5. Morán
Highlights
Title: Evidence of ferromagnetic ordering at low temperatures in epitaxial TbMnO3 thin films grown by means of DC magnetron sputtering
Epitaxial TbMnO3 thin films are grown via dc magnetron sputtering.
The TbMnO3 films are epitaxial and (00ℓ)-oriented.
The nominal valence of the Mn ions is 3+
The films show well-defined transition at T<42 K
Weak ferromagnetism is observed in undoped TbMnO3 films
Hysteresis loops with coercive fields of 1000 Oe are observed at 5 K
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Conflict of interest statement, in good faith, the authors of the manuscript Evidence of ferromagnetic ordering at low temperatures in epitaxial TbMnO3 thin films grown by means of DC magnetron sputtering declare that there is no conflict of interest arising from conflicting financial or other interests.
Yours sincerely
Dr. Oswaldo Morán, corresponding author Physics Department, National University of Colombia, Medellín Campus A. A. 568, Medellín, Colombia e-mail: [email protected] Phone: 57-4-4309327 Fax: 57-4- 2604489
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