Tuning Magnetic Properties of BiFeO3 Thin Films by Controlling Rare

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Tuning Magnetic Properties of BiFeO Thin Films by Controlling Rare-Earth Doping: Experimental and First-Principles Studies Hoa Hong Nguyen, Ngo Thu Huong, Tae-Young Kim, Souraya Goumri-Said, and Mohammed Benali Kanoun J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.5b03834 • Publication Date (Web): 03 Jun 2015 Downloaded from http://pubs.acs.org on June 15, 2015

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The Journal of Physical Chemistry

Tuning Magnetic Properties of BiFeO3 Thin Films

by

Controlling

Rare-Earth

Doping:

Experimental and First-Principles Studies

Nguyen Hoa Hong 1*, Ngo Thu Huong 1,2, Tae-Young Kim 1, Souraya GoumriSaid3, and Mohammed Benali Kanoun4, †

1

Nanomagnetism Laboratory, Department of Physics and Astronomy, Seoul National University, Seoul 151-747, Korea 2

3

Hanoi University of Science, 334 Nguyen Trai, Thanh Xuan, Hanoi, Vietnam

School of Chemistry and Biochemistry and Center for Organic Photonics and Electronics, Georgia Institute of Technology, Atlanta, Georgia 30332-0400, United States.

4

School of Physics, Georgia Institute of Technology, Atlanta, Georgia 30332-0400, United States.

1

ACS Paragon Plus Environment


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ABSTRACT Rare Earth (RE) - doped BiFeO3 (BFO) thin films were grown on LaAlO3 (LAO) substrates by using pulsed laser deposition technique. All of BFO films doped with 10% of RE exhibit a rhombohedral single phase. As for the Pr and Nd doping cases, the ferromagnetic phase is less favored because Fe2+ amount is not dominant. When dopant concentration was increased up to 20%, the RE-doped BFO films have gone through a structural transition from rhombohedral to either pure orthorhombic phase (for Ho, Sm), or a mixed phase of orthorhombic and tetragonal (for Pr, Nd), or pure tetragonal (for Eu). As an important consequence, magnetic properties of RE-doped BFO films have drastically changed. Our results give a guide for how to tailor the ferromagnetism of BFO films by appropriate controlling the type of RE dopant as well as dopant concentration. The experimental findings are completed by performing density functional theory calculations to explore the effect of RE doping in BFO for the considered three phases.

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INTRODUCTION Multiferroics are an interesting group of materials that exhibit both ferroelectricity and ferromagnetism with coupled electric and magnetic order parameters [1-3]. Multiferroism is currently the subject of intensive scientific investigation, since they potentially offer a wide range of interesting applications [2-6]. BiFeO3 (BFO) is known to be the only ABO3-type simple perovskite that shows multiferroic at room-temperature and, thus, is considered to be the most promising candidate for practical applications among multiferroic materials [3, 5, 6]. At room temperature, BiFeO3 exhibits a distorted perovskite structure with rhombohedral polar R3c symmetry. At higher temperatures (≈1100 K), the rhombohedral (R) phase undergoes a first order phase transition to a GdFeO3-like Pbnm structure [7-9] and a (probable) orthorhombic γphase [10]. Basically, BiFeO3 should be G-type antiferromagnetic due to the local spin ordering of Fe3+, that forms a cycloidal spiral spin structure [11]. There are several ways to stress the spiral magnetic ordering by applying a very high magnetic field, or reducing the dimensions of the samples, or by replacing Bi3+ or Fe3+ by other ions of comparable ionic sizes [12]. Dimension reduction seems to be an effective method to enhance the magnetic moment in BFO thin films and in nanoparticles [6, 13]. Some groups have reported about the increase of magnetization in the bulk, thin films, and nanoparticles of BFO, either by substituting on the Bi-site by trivalent rare-earth and divalent ions, or on the Fe-site by transition metal ions. Thakuria and Joy had showed that the magnetic moment of the nanoparticles could be enhanced 3 times by substituting Bi by Ho. However, the reported saturated magnetization is still found only at a quite high field as of 6 T [12, 14]. Partial substitution of Bi by Rare-Earth (RE) ions is known to induce a 3



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ferromagnetic response which has been attributed to suppression of the spiral modulation [15, 16]. In particular, doping of the Bi3+ A site of the perovskite with rare earth cations (RE) has received extensive attention, with a variety of symmetries, and magnetic and electric behaviors reported with increasing values of x in Bi1-xRExFeO3 and/or decreasing ionic radii of the rare earth cation [8, 9, 17–18]. Recent studies of bulk and thin films of (Bi,RE) FeO3 (RE = Nd, Sm) have revealed a formation of a stable antipolar, PbZrO3-like structure in a narrow rare-earth concentration range [8, 18-19] In this respect, ab initio calculations based on the density-functional theory (DFT) have played an important role in the description, understanding, as well as prediction—via identification of suitable material design rules—of magnetic, ferroelectric and magnetoelectric properties of multiferroics, due to its ability to describe the many active degrees of freedom within a comparable level of accuracy [3, 20-23]. Theoretical modeling based DFT approach and effective Hamiltonian scheme have been subject of few recent works on RE doped BFO [19, 23, 24] where they mainly reported the dependence of critical temperature on the RE compositions. In the present study, we attempt to tune the magnetic properties of BiFeO3 by several ways such as: selecting the suitable RE ion for doping; screening the appropriate concentration, targeting to obtain the largest magnetization possible at room temperature, and at a relatively low field. The aim of the present work is to understand the relationship between structural and magnetic properties of RE-doped BFO films by combining ab initio calculations and thin-film growth experiments in order to control magnetism of this family of compounds, with hope to guide correctly the materials strategy for spintronic and magnetic applications.

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EXPERIMENTAL AND COMPUTATIONAL DETAILS RExBi1-xFeO3 ceramic targets (where RE= Ho, Sm, Pr, Nd, and Eu; x = 0; 0.1, and 0.2) were prepared by a sol-gel auto ignition method [14]. RE-doped BiFeO3 (RBFO) thin films have been deposited by Pulsed Laser Deposition (PLD) technique (eximer KrF laser with λ= 248 nm; the repetition rate was 13 Hz and the energy density was 2.1 J/cm2), with a typical thickness as of 200 nm. All the films were grown on (001) LaAlO3 (LAO) substrates. During deposition, the substrate temperature was kept at 700°C and the oxygen partial pressure (PO2) was 1.4×10-3 Torr. After deposition, the sample was kept in the chamber at 500°C with the same oxygen partial pressure as during deposition for 30 min, and then finally cooled down slowly to room temperature [see also Ref. 14 for details]. The structural analysis was carried out by High Resolution X-ray diffraction (HRXRD) with Cu Kα radiation. The M-T and M-H curves were collected by a Quantum Design Superconducting Quantum Interference Device (SQUID) system with magnetic field (H) ranging from 0 up to 0.5 T and temperatures (T) ranging from 350 K down to 5 K. The oxidation states of RE-doped BFO thin films were characterized by X-ray photoelectron spectroscopy (XPS, KRATOS, AXIS-HSi). XPS measurements were performed with an Mg/Al X-ray source. The energy calibrations were made against the C 1s peak and the Shirley background subtraction was used [as in Ref. 14]. The chemical elements’ content was also checked by Energy-dispersive X-ray diffraction spectroscopy (EDX) at room temperature Our calculations were performed using all electron linearized augmented plane wave method with local orbitals basis set based on DFT as implemented in the WIEN2k computer program [25]. Exchange and correlation were treated within the generalized gradient approximation (GGA) of Perdew-Burke-Ernzerhof (PBE) [26]. Onsite Hubbard interaction between the 5f 5



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electrons was treated within the fully rotationally invariant version [27]. Here, U (equal to what is often called Ueff = U − J) is taken as the on-site interaction term as suggested in Ref. 28. The cutoff Rmt*Kmax was set to 7.0 to determine the basis sets. To obtain more information on RE doped-BiFeO3, we were carried out DFT calculations by adopting the following supercell approach: (i) a 2×2×2 supercell that contains 12 FeBiO3 formula cells for the hexagonal R3c structure, and (ii) a 2×2×1 supercell that contains 8 formula cells for the orthorhombic Pbnm and tetragonal P4mm structures. In each supercell, a Bi atom was substituted by one RE impurity, in order to obtain Bi1-xRExFeO3 with x = 0.0833 for hexagonal structure and 0.25 orthorhombic and tetragonal structures. For the integration over the Brillouin zone, a 4×4×1 Monkhorst–Pack kpoint mesh [29] was used for the rhombohedral (hexagonal) cell while a 5×5×3 Monkhorst–Pack k-point grid was adopted for the orthorhombic and tetragonal cells. The convergence of selfconsistent calculations was attained with a total energy convergence tolerance of 0.1 mRy.

RESULTS AND DISCUSSION As we know, the EDX method could not give a very precise evaluation of content of each element in the case of thin films due to the fact that it is just most sensitive to the surface of the film but not as the whole. However, one can see a slight tendency of deviation in resulting concentration if comparing to the starting doping concentration. For example, if we substituted 20% of Bi by Ho, then the resulting Ho:Bi ratio is 22:78, if we substitute 20% of Bi by Sm, then the resulting Sm :Bi ratio is 26.4: 73.6. Therefore, thoroughly in this report we keep naming the compounds as their starting stoichiometry. The XRD data show that for the case of doping of 10%, there is no significant change in structures in comparison with the pristine BiFeO3. All RE0.1Bi0.9O3 films have a rhombohedral structure showing very strong peaks of the BFO phase. 6


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Doping with different RE elements only causes some certain shift of the peak positions in the spectra, showing some change in lattice parameters. The out-of-plane lattice parameter is 3.925, 3.925, 3. 9329, 3.933, and 4.585 Ǻ for Sm, Ho, Nd, Pr and Eu doped BFO, respectively (as discussed partially in Ref. 14]. When the RE doping concentration increases up to 20%, a drastic change in structure of the compound has appeared: As for doping of Ho and Sm , the structure is single phase orthorhombic, while for Pr and Nd, it has become a mixed phase of orthorhombic and tetragonal, and for Eu doping case, it is single phase tetragonal. Some typical spectra are shown in Figure 1 for comparison between structures of 10% and 20% doping cases, showing changing from rhombohedral to orthorhombic for Ho doping case, and changing from rhombohedral to a mix of orthorhombic and tetragonal for Pr doping case. In order to make it easier later for reference, we summarize the structure types for all RExFe1-xO3 thin films in Table 1.

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60

Kα

LAO (003)

LAO (003)

LAO (002) LAO (002)

(002)

(006)

(012) LAO (001)

0 100 50

(a)

(b)

LAO (002)

Kα (104) (113) (006)

50

(012)

Kα

LAO (001)

0 100

80

(c)

LAO (002)

(012)

50

40

Kβ (002)

(001)

20 LAO (001)

100

Intensity (a.u.)

(d) LAO (003)

Kα

(131)

Kα

(214)

Kβ

Kα

(113)

(006)

50

LAO (001)

100

Kα

0

(012)

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0 20

40

60

80

2 theta (degree) Figure 1: X-ray diffraction patterns for 200 nm-thick- films grown on LaAlO3 substrates of (a) Ho0.1Bi0.9FeO3-showing Rhombohedral phase; (b) Ho0.2Bi0.8FeO3-showing Orthorhombic phase; (c) Pr0.1Bi0.9FeO3-showing Rhombohedral phase; and (d) Pr0.2Bi0.9FeO3-showing a mix of Orthorhombic and tetragonal phases.

Table 1: List of structural phases and Fe2+:Fe3+ ratio calculated from XPS data for RExBi1-xFeO3 films (RE= Ho, Sm, Eu, Pr, and Nd).

element

Pr

Nd

Ho

Sm

Eu

x = 0.1

Rhombohedral

Rhombohedral

Rhombohedral

Rhombohedral

Rhombohedral

x = 0.2

Orthorhombic

Orthorhombic

Orthorhombic

Orthorhombic

Tetragonal

concentration

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Tetragonal

Tetragonal Ratio Fe2+: Fe3+

x = 0.1

40 : 60

x = 0.2

51.8 : 48.2

40 : 60 48.5 : 51.5

50 : 50

50 : 50

48.6 : 51.4

45.6 : 54.4

43.7 : 56.3

48.7 : 51.3

Magnetization versus magnetic field taken at 300 K for RE0.1Bi0.9O3 film is shown in Fig. 2 (a). One can see that among all RE doped films, Eu-doped, Sm-doped and Ho-doped BFO films have rather large magnetic moments. One cannot attribute this big ferromagnetic signal to any impurity. If impurities are more than 5% as resolution of the apparatus, one must see from XRD spectra (as we see that no alien peak other than BFO peaks in XRD data), but if they are less than 5%, then the contribution should not have been that big in magnitude. Some other group also got ferromagnetic ordering in Ho-doped BFO, however with much more modest magnitude, and the Ms was obtained at a much greater field (as of 6 T) [12, 14], while we got much a larger Ms but at much lower field (as of 0.2 T). This makes a difference and it is quite meaningful for applications. In comparison to those three dopants mentioned above, the Pr- and Nd-doped BFO films show much weaker magnetism: the curves in Figure 2 (a) show a much smaller magnitude for magnetization, and the curves of M(H) at room temperature are almost linear indicating a paramagnetic phase [14, 30, 31].

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Figure 2: (a) Magnetization versus magnetic field taken at 300 K for (a) 200 nm-thick RE0.1Bi0.9FeO3 films; (b) 200 nm-thick RE0.2Bi0.8FeO3 films (RE= Ho, Sm, Eu, Pr, and Nd), and (c) Zoom for the low field region of the M-H curve of 200 nm-thick Sm0.2Bi0.8FeO3; The insets in (b) shows the zoomed M-H curves of the two typical 10% doping cases (Pr and Sm) in order to compare directly with the 20% doping case.

In order to identify the origin of magnetism of our films, XPS measurements were performed. The peak of Fe 2p could be well observed at 711 eV for Fe3+ and at 709.5 eV for Fe2+ [14]. From the shape of the peaks, it is seen that there are Fe2+ and Fe3+ in all doping cases (Ho, Sm, Nd, Pr, and Eu). The rough fitting analysis to the peaks reveals the oxidation state of Fe in our films which is listed in Table 1. One can see from Table 1 that for the 10% doping case, Fe2+: Fe3+ ratio is about 50%: 50% for the cases of Ho and Sm, roughly so for Eu, but only about 40%: 60% for the Pr and Nd cases [14]. The coexistence of Fe2+ and Fe3+ is thought to be in favor of the ferromagnetic phase in BFO films due to the double exchange between Fe2+ and Fe3+ via the role of oxygen as intermediates [32, 33]. When the amount of Fe3+ is more favored, the ferromagnetism get weaker due to the fact the Fe3+- Fe3+ interaction is in favor of antiferromagnetic ordering [33]. This explains well the corresponding SQUID data shown in Fig. 2 (a), we may understand why in Pr and Nd-doped BFO films, the ferromagnetic phase is not favored [14]. When the doping concentration is increased up to 20%, the Fe2+:Fe3+ ratio has changed as the results of changes in structure (recalling the difference in phase obtained that was discussed earlier concerning Figure 1). In Table 1, one can see that as for Pr and Nd, there is a significant increase of Fe2+ amount, leading to an enforcing of the ferromagnetic ordering. On the contrary, as for Ho, Sm, and Eu doping cases, Fe2+ seems to be decreased, relating to the weakening of ferromagnetic phase. This can be seen clearly from M-H curves for the 20% doping case shown 10


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in Figure 2 (b). As for the Pr case, whose phase is a mix of orthorhombic and tetragonal, one can see that magnetic phase has changed from pure paramagnetic (see in the inset on the left for the 10% doping case) into mixed paramagnetic and ferromagnetic (the M-H curve of Pr0.2Bi0.8FeO3 film shows an accumulation of a part as of paramagnetic and another part as weak ferromagnetic phase). Differently, as for Sm case (similarly for Ho and Eu cases also), it has changed from very strong ferromagnetic into “almost” paramagnetic (compare the curve seen in Fig. 2 (b) for the 20% Sm case with that of the 10% Sm case shown in the inset on the right hand side). It is totally in accord with the structural analysis shown above. The zoom for the low field region of a typical ferromagnetic sample (in this case as of 20%Sm-doped BFO film) is shown in Figure 2(c) to see clearly the loop showing their ferromagnetic behavior. Indeed our samples have coercivity HC ranging from 25-70 Oe (Pr-BFO: ~27 Oe; Nd-BFO~55Oe; Ho-BFO ~52 Oe, Sm-BFO~48 Oe, Eu-BFO~ 70 Oe). In fact doping RE does not change much the coercivity of the host compound (BFO target and film have almost the same magnitude of HC (61 Oe- see in the M -H data show in Figure 3 (a) and (b)).

However, doping RE has enhanced the

magnetization to about 2 orders in comparing to that of the pristine BiFeO3 (refer Figure 3 (a) and (b)). Previously we have tried to dope RE with lower concentration as of 5%. Our results showed that 10% samples have larger magnetic moment. One can see a typical example shown in Figure 3 (c), and more details in Ref. 14. From this work, one could see that starting from doping 20%, samples are not single phase. Therefore, we assume that roughly for the sake of applications, doping 10% seems to be most suitable. Paudel et al have proposed some model to suppose that ferromagnetic moment in BFO may be due to intrinsic defects [34]. But no one has confirmed that defects should be the main source of magnetism in this type of compounds, either undoped BFO or Rare-Earth doped BFO. For films, moreover, it is hard to determine oxygen 11



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vacancies and from the M-H data, “oxygen vacancies” seem to be unlikely the reason for magnetic moment here. One can see that with the same environment, magnetic behaviors of different compounds (pristine, RE-doped BFO with different RE element, or with different RE concentration) are different. And Fe2+:Fe3+ ratio is seen to be quite different. It is more logical to think that this ratio should be the main reason to alter the magnetic properties of RE-BFO compounds. From other aspect, the strains might influence properties of BFO somehow. We also noticed that surface magnetism could also play a role (for films, due to strains of substrates [14]. However in this work we conclude for films grown by PLD on LaAlO3 substrates only, not to generalize for all cases including bulks. It should be too much complicated to include all on one plate.

Figure 3: Magnetization versus magnetic field taken at 300 K for (a) undoped-BiFeO3 ceramic target; (b) 130 nm-thick undoped-BiFeO3 film; and (c) for 200nm thick-Ho0.05Bi0.95FeO3 and Ho0.1Bi0.9FeO3 films

From our theoretical calculation, we calculated the magnetic moments values for orthorhombic, tetragonal and rhombohedral phases, as shown in Figure 4(a-d). In Fig. 4 (a), the rhombohedral phase shows the largest total magnetic moment for all the RE elements doping BFO. The magnetic moment per RE atom shows that the magnetic moment value of the Eu is

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bigger than Pr, Nd, Sm, and Ho values, as shown in Figure 4 (b). The RE doped- BFO leads to a change of magnetic moment of Fe atom. However, the magnetic moment per Fe atoms, as displayed in Figure 4 (c), shows no trend in its variation. In fact, its value varies between 2.6 and 3.6 µB per Fe atom for all considered RE dopants in three phases. This might be resulting from the difference of structure and symmetry. In fact, the RE doping enhanced the magnetic moment of BFO but it also induces the polarization of oxygen atoms situated in nearest site to RE and Fe. The magnetic moment values of O are also presented in Figure 4 (d), where we see that the largest polarization of O is found in the tetragonal phase [35, 36].

Figure 4: (a) Total magnetic moment values for RExBi1-xFeO3 for three phases and the magnetic moment per atom: (b) RE, (c) Fe, and (d) O. The y-axis shows the magnetic moment values given in µB.

In order to understand the mechanism of magnetization and the role of RE doping in BFO for different phases, we have calculated the electronic structures. In the present work, we limit 13



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ourselves to the Sm-doped BFO for three phases calculated within GGA+U approximation, in order to avoid repetition with similar behavior of remaining RE dopants. Figure 5 shows the total and partial DOS (atom and orbital decomposed) of Sm-doped BFO in three phases. The top of the valence band is set to zero, as indicated by a dotted vertical line. Interestingly, the DOS results shown in Figure 5 indicate that the rhombohedral and tetragonal phases are almost similar regarding orbital occupation and different than orthorhombic phase. The total DOS curves of rhombohedral and tetragonal phases suggest that the contributions are larger from the spin-up (down) states below (above) the Fermi. From PDOS, we can see that the highest occupied Sm 4f states for spin up are situated around -6.5 eV below the Fermi energy. Above the Fermi level, we find the unoccupied f-levels and concentrate them into a sharp peak around 1 eV. The fundamental band gap separates the valence band maximum and unoccupied Sm 4f is about 0.44 eV. For minority-spin bands, they are empty and situated in the conduction band minimum with a more localized peak about 5 eV. The valence band maximum for the occupied majority is dominated by the O 2p state. The main O 2p valence band is found between 0 and -7 eV. The main Fe 3d valence band focus in a narrow range near -6 eV lightly mixed with the O 2p bonding state. The peak of the majority-spin of the Bi 6s state is found between -10 and -11 eV. Because of the coupling between the Bi 6s and the O 2p state, the antibonding Bi 6s state is found at the top of the valence band. The conduction band minimum is mainly dominated by the Fe 3d state, but O 2p and Bi 6p states also contribute. For the minority spin, we can see that the states in the energy range from -0.5 to 2.0 eV belong to the Fe 3d and O 2p states. Note that the peak from the O 2p DOS at 1 eV comes from separate bands without 4f contributions and thus these 4f states are localized and not hybridized with the O 2p states. The difference of DOS between the tetragonal phase and the rhombohedral phase is that the minority-spin Fe 3d states 14


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of the tetragonal phase near the Fermi level move towards the high energy direction and become completely unoccupied with two mainly peaks. In the case of the orthorhombic phases, the lowest energy majority- and minority-spin band of the valence band extends in the energy range from -11 to -10 eV, and it is mainly composed of the Bi 6s and O 2p states. The majority- and minority-spin bands occur around the Fermi level, in the energy range from -1.0 to 1.0 eV, mostly due to the Fe 3d states interact with the O 2p states. It is found that Sm 4f states in the majority spin are inserted below the Fermi level around -1.5 eV, whereas in the unoccupied minority-spin Sm 4f states are around 9 eV above the Fermi level. It can also be observed contrarily to the previous phases that there are two distinct peaks in the DOS of the majority spin where the first sharp peak appears at the Fermi level is half occupied and second peak is unoccupied that appear in conduction band around 5.5 eV.

Figure 5: The total and partial spin-polarized DOSs for ferromagnetic Sm-doped BFO. Spin up and spin down correspond to positive and negative values, respectively. The vertical dashed line denotes the Fermi level.

CONCLUSION 15



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Our study on RE- doped BiFeO3 thin films have reported that all the films of 10% of RE doping show a single phase of rhombohedral structure. The saturated magnetization in Ho-, Sm-, and Eu-doped films is much greater than those reported in literature previously, and was observed at a much lower field as of 0.2 T. The magnetic moment of Ho-doped, Sm-doped, and Eu-doped BFO films are roughly largest. The observed ferromagnetism in our RE-doped BFO films were supposed to result from the coexistence of Fe2+ and Fe3+ that favor double exchange via oxygen. When the dopant concentration is increased up to 20%, the BFO films has gone through a structural transition: while Ho-, Sm-, Eu-doped BFO has become orthorhombic, along with the fact that the Fe2+ amount is reduced, Pr and Nd-doped BFO has changed to be mixed phase of orthorhombic and tetragonal, with the consequence is that Fe2+ is increased, leading to a favor of ferromagnetic ordering. This measurement was followed by further analysis using firstprinciples calculations based on density functional theory within GGA+U to elucidate the effect of RE doping BFO on the electronic structures. The calculation of total magnetic moment and contribution of each atom in different phases shows the change in polarization of BFO following the RE doping and the symmetry/structure. Our experimental and theoretical findings reveal that by controlling the type and concentration of Rare-Earth doping, one can tune the magnetic properties of BiFeO3 films in order to fit the requirements of spintronic and magnetic applications.

AUTHOR INFORMATION Corresponding Authors *

Email: [email protected]

†

Email: [email protected] 16


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ACKNOWLEDGMENTS The authors thank Dr. Raghavender for measurements of the undoped BFO, and Prof. Kurisu, Prof. Konishi and J.G. Kim for EDX measurements. N.H.H. and T.Y.K. would like to thank project 3348-20120033 of National Research Foundation of Korea for its financial support. We are grateful for the help of the SNU National Center for Inter-University Research Facilities on XRD measurements and SNU Center for Materials Analysis on XPS measurements. N. T. H thanks Korea Foundation for Advanced Studies for giving her the KFAS fellowship during 20132014, and the BK 21 Plus program of Department of Physics and Astronomy, SNU, for fellowship of 2014-2015. REFERENCES (1) Hill. N. A. Why Are There so Few Magnetic Ferroelectrics?. J. Phys. Chem. B 2000, 104, 6694-9709. (2) Catalan, G.; Scott, J. F. Physics and Applications of Bismuth Ferrite Adv. Mater. 2009, 21, 2463-2485. (3) Hao, X. F.; Stroppa, A.; Barone, P.; Filippetti, A.; Franchini, C.; Picozzi, S. Structural and Ferroelectric Transitions in Magnetic Nickelate PbNiO3. New J. Phys. 2014, 16, 015030. (4) Eerenstein, W.; Mathur, N. D.; Scott, J. F. Multiferroic and Magnetoelectric Materials. Nature 2006, 442, 759-765. (5) Selbach, S. M.; Tybell, T.; Einarsrud, M.-A.; Grande, T. Size-Dependent Properties of Multiferroic BiFeO3 Nanoparticles. Chem. Mater. 2007, 19, 6478-6484. (6) Wang, J.; Neaton, J. B.; Zheng, H.; Nagarajan, V.; Ogale, S. B.; Liu, B.; Viehland, D.; Vaithyanathan, V.; Schlom, D. G.; Waghmare, U. V.; et al.

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Multiferroic Thin Film Heterostructures. Science 2003, 299, 1719-1722.

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(17) Rusakov, D. A.; Abakumov, A. M.; Yamaura, K.; Belik, A. A.; Van Tendeloo, A.; Takayama-Muromachi, E. Structural Evolution of the BiFeO3−LaFeO3 System. Chem. Mater. 2011, 23, 285-292. (18) Karimi, S.; Reaney, I. M.; Han, Y.; Pokorny, J.; Sterianou, I. Crystal Chemistry and Domain Structure of Rare-Earth Doped BiFeO3 Ceramics. J. Mater. Sci. 2009, 44, 51025112. (19) Xu, B.; Wang, D.; Íñiguez, J.; Bellaiche, L. Finite-Temperature Properties of Rare-EarthSubstituted BiFeO3 Multiferroic Solid Solutions. Adv. Funct. Mater. 2014, 25, 552-558. (20) Franchini, C. Hybrid Functionals Applied to Perovskites. J. Phys.: Condens. Matter 2014, 26, 253202. (21) Stroppa, A; Picozzi, S. Hybrid Functional Study of Proper and Improper Multiferroics. Phys. Chem. Chem. Phys. 2010, 12, 5405-5416. (22) Ravindran, P.; Vidya, R.; Kjekshus, A.; Fjellvåg, H.; Eriksson, O. Theoretical Investigation of Magnetoelectric Behavior in BiFeO3. Phys. Rev. B 2006, 74, 224412. (23) Gavriliuk, A. G.; Struzhkin, V. V.; Lyubutin, I. S.; Ovchinnikov, S. G.; Hu, M. Y.; Chow, P. Another Mechanism for the Insulator-Metal Transition Observed in Mott Insulators. Phys. Rev. B 2008, 77, 155112. (24) Lee, J.-H.; Oak, M.-A.; Choi, H. J.; Son, J. Y.; Jang, H. M. Rhombohedral–Orthorhombic Morphotropic Phase Boundary in BiFeO3-Based Multiferroics: First-Principles Prediction. J. Mater. Chem. 2012, 22, 1667-1672. (25) Blaha, P.; Schwarz, K.; Madsen, G. K. H.; Kvasnicka, D.; Luitz J. WIEN2k: An Augmented Plane Wave and Local Orbitals Program for Calculating Crystal Properties. T.U. Wien, Austria, 2001. (26) Perdew, J. P.; Burke, K.; Ernzerhof, M. Generalized Gradient Approximation Made Simple. Phys. Rev. Lett. 1996, 77, 3865-3868.

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Table of content

21



60 LAO (003) LAO (003)

Kα

LAO (002)

(b)

LAO (002)

(002)

(006)

(012) LAO (001)

0 100 50

(a)

(c)

LAO (002)

Kα (104) (113) (006)

50

(012)

Kα

LAO (001)

0 100

80

LAO (002)

Kβ (002)

(012)

50

40 LAO (001)

20 (001)

100

Intensity (a.u.)

(d) LAO (003)

Kα

(131)

Kα

(214)

Kβ

Kα

(113)

(006)

50

LAO (001)

100

Kα

0

(012)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

0 20

40

60

80

2 theta (degree)

FIG 1


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Fig. 2

15

EBFO

(a)

10

SBFO

M (emu/cm3)

5 0 NBFO PBFO

-5

HBFO

-10

-15 -0.50

-0.25

0.00

0.25

0.50

H (T)

10

SBFO

10% PBFO

5

EBFO

0 -5

-10 -0.50 -0.25 0.00

HBFO 0.25

0.50

H (T)

0

10

NBFO

M (emu/cm 3)

M (emu/cm3)

25

M (emu/cm3)

50

-25 PBFO

(c)

10% SBFO

5 0 -5

-10 -0.50

-0.25

0.00

0.25

0.50

H (T)

-50 -2

-1

0

1

2

H (T)

10

Magnetization (emu/cm3)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Bi0.8Sm0.2FeO3

(c) 5

0

-5

-10 -500

-250

0

250

500

Field (Oe)



M-H BiFeO3 Target

M [ emu / g ]

2

(a)

1 0 -1 -2 -1.0

-0.5

0.0 H[T]

0.5

1.0

BFO - 130 nm Thin film

1.5

(b)

3

M [ emu /cm ]

1.0 0.5 0.0 -0.5 -1.0 -1.5

-6

Magnetization (emu/cm3)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

15

-4

-2

0 2 H[T]

4

6

(c)

10 5 0 -5 -10

5% HBFO 200nm 10% HBFO 200nm

-15 -0.4

-0.2

0.0

0.2

0.4

Field (T)


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Tuning Magnetic Properties of BiFeO3 Thin Films by Controlling Rare

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