Interface‐Coupled BiFeO₃/BiMnO₃ Superlattices with Magnetic Transition Temperature up to 410 K

Enhanced magnetic transition temperatures are demonstrated in 50 nm thick (001)-oriented (BiFeO3)m/(BiMnO3)m (BFO/BMO) superlattices (SL). Highly strained ultra-short SLs showed a magnetic transition, TC2, up to ≈410 K as a result of ferrimagnetic interaction between BFO/BMO. All SL showed a new ferromagnetic transtion, TC1, of ≈150 K coming from Fe4+-Fe4+ interaction explained by electron transfer/leakage from Fe3+ to Mn3+. In the field of multiferroics, the coexistence of magnetism and ferroelectricity in perovskite oxides (ABO3) has drawn increasing interest due to their potential for multifunctional devices, for example, magnetoelectric random access memory, tunable multifunctional spintronic devices including four-state memory devices, and spin filters.1-3 Room temperature operating multiferroics are currently of intense scientific interest for these devices.1-3 An important requirement for multiferroics is a strong coupling between the ferroic properties. Two of the most widely studied multiferroics are BiMnO3 (BMO) and BiFeO3 (BFO). BMO is one of the rare multiferroic materials that can possess both ferromagnetism (FM) and ferroelectricity (FE). It has a paramagnetic–ferromagnetic transition temperature (TC) of 105 K and a magnetic saturation moment (MS) of 3.6 Bohr magneton/formula unit (μB/f.u.). Unfortunately, its electrical polarization is very small (0.1 μC cm−2 at ≈90 K) or it is even paraelectric.4 In BFO, strong FE behavior (Ferroelectric transition temperature, TC,FE ≈ 1103 K) coexists together with antiferromagnetism (AFM). BFO has an incommensurate cycloidally modulated G-type AFM structure with a large period of 62 nm (Néel temperature, TN ≈ 650 K).5, 6 Despite having ferroelectric and magnetic ordering at high temperatures, BFO has drawbacks for practical applications. Most notably the weak FM-like magnetic behavior results from destruction of its incommensurate magnetic structure leading to AFM.6 With strong ferromagnetic behavior in BMO and strong ferroelectric behavior in BFO, various studies have suggested combining BMO and BFO in order to achieve room-temperature (RT) multiferroicity.7-10 One approach is to modify the local spin configuration of Fe by chemical doping at the B-site.8, 9 By 50% doping, this simple concept could result in ordered double perovskite ferrites of Bi3+2Fe(3+n)+TM(3−n)+O2−6 (TM = transition metal: Cr, Mn, Co, Ni, and Cu) giving antiferromagnetically coupled high spin state of the TM and Fe. However, this concept has not been widely realized due to the lack of cation ordering as a result of negligible differences in ionic valences and ionic radii.11-15 On the other hand, disordered double perovskites (notably BiFe0.5Mn0.5O3 (BFMO) and Bi2CoMnO6 (BCMO)) show a strongly enhanced TC albeit with lower magnetic moment than expected for the ordered structure.16, 17 Pálová et al. extended the idea of ordered double perovskites by theoretically studying BFMO with an ordered nanoscale checkerboard structure (NCB–BFO/BMO, see Figure 1a), i.e., columnar B-site ordering.10 Their calculations predict that an ordered NCB–BFO/BMO would be ferroelectric and ferrimagnetic (FiM) with MS of 3.8 μB/Fe–Mn pair and TC of 406 K, as a result of magnetic ordering arising from the superexchange coupling between the neighboring AFM–Fe and FM–Mn.10 Unfortunately, spontaneous columnar B-site ordering is not expected and high quality growth of artificial 1:1 BFO/BMO superlattices along the 〈110〉 direction is very hard. Therefore, it is worth considering (001)-oriented BMO/BFO superlattices (SL–BFO/BMO) instead (Figure 1b). Similar to NCB–BFO/BMO, interfacial Fe–Mn ferromagnetic ordering and a concomitant increase in TC can be expected in SL–BFO/BMO. In this structure, small local variations in magnetic behavior will also be expected since the number of Fe–Mn nearest neighbors is halved for a 1:1 SL–BFO/BMO compared to NCB–BFO/BMO.10 Various studies on short period perovskite superlattices have shown that the magnetic behavior can be enhanced by the interaction between the B-site cations. For example, Ueda et al. reported ferromagnetism at 375 and 230 K in (111)-oriented 1:1 LaCrO3/LaFeO3 and LaMnO3/LaFeO3 superlattices, respectively, while the parent bulk materials are AFM.18, 19 Moreover, interfacial FM has been observed at 95 and 75 K in (001)-oriented CaMnO3/CaRuO3 and LaNiO3/CaMnO3 superlattices, respectively even though CaMnO3 is AFM and both CaRuO3 and LaNiO3 are paramagnetic.20-22 Additional changes in the B-B′ magnetic interactions may result from internal charge transfer from B to B′.20-24 Internal charge transfer has been observed in, e.g., (001)-oriented LaNiO3/LaMnO3 superlattices from Mn3+ to Ni3+, similar to rock salt ordered La2NiMnO6.23, 24 As a result, interface-induced ferromagnetism was found in the paramagnetic LaNiO3 layer. Moreover, the lattice strain from the substrate as well as the superlattice interfaces can affect the magnetic behavior as a result of changes in orbital ordering or bonding distortion of the BO6 octahedra.25-29 Here, in SL–BFOm/BMOm (m = 1, 2, 4, and 8 unit cells (u.c.)), tuning Fe–Mn magnetic interactions through interfacial strain and charge engineering is demonstrated. Detailed analysis of both structural and magnetic features of superlattices using X-ray diffraction (XRD) and X-ray absorption spectroscopy (XAS) supports our conclusions. We show the presence of three magnetic transitions in SL–BFO/BMO. For all SL films, TC1 is observed at around 145 ± 15 K, which is significantly higher than the TC of bulk BMO (≈100 K).4 m = 1, 2 superlattices even showed ferromagnetic behavior at room temperature with a TC2 of ≈410 K for m = 1 and ≈320 K for m = 2, whereas m = 4, 8 showed a TC,BMO of 96 K which is close to bulk BMO. Recently, a spin glass was observed for BFO/BMO multilayers with m = 1.30 These samples were suggested to be relaxed and had a single out-of-plane film lattice parameter of ≈3.93 Å which is considerably smaller than cpc = 3.970 ± 0.005 Å (pc is pseudocubic), as measured for our m = 1 films, as shown later. This difference is likely the result of the use of different growth conditions since the creation of highly epitaxial in-plane strained BFO and BMO is difficult. We have found that it is necessary to carefully optimize the growth conditions to create the desired highly strained, heteroepitaxial films with strong ferrimagnetic coupling between BFO and BMO. To verify both the epitaxy as well as the phase purity of the superlattices of this work, we measured out-of-plane XRD scans (see Figure 2a). The scans show that all superlattices grew epitaxially on SrTiO3 (STO), showing pseudocubic (001)pc reflections.31 No superlattice peaks from the multilayer were observed; however, this does not indicate the lack of a high quality SL–BFO/BMO structure since any minor interface roughening reduces the intensity of these peaks dramatically due to the small difference between scattering factors of BMO and BFO. A minor contribution of monoclinic BMO (<1%) was observed in all films. The monoclinic phase could be expected to form as a result of strain relaxation.31 For the superlattices with m = 4, 8, a minor contribution of Bi2O3 was present. Note that the impurity phases were not observed by transmission electron microscopy (TEM). TEM images across the SL layers showed a single uniform structure (see inset of Figure 2a for a representative TEM image of a 1/1 SL–BFO/BMO taken around the substrate–film interface). To investigate the out-of-plane lattice parameters in more detail, symmetric reciprocal space maps (RSMs) around STO (004) were conducted (see Figure 2b). For the short period superlattice (m = 1), a single film peak (cpc = 3.970 ± 0.005 Å) corresponding to the SL–BFO/BMO was observed. The single peak indicates that the two lattices are strain coupled to give a single out-of-plane lattice parameter. The large c parameter is consistent with in-plane compression of the film lattices by the smaller STO lattice (apc–bulk BMO = 3.935 Å, apc–bulk BFO = 3.962 Å and abulk STO = 3.905 Å), and consequent out-of-plane elastic lattice extension of both the BMO and BFO lattices. As indicated by the dashed lines, with increasing m a clear separation of the out-of-plane film peaks for BMO and BFO occurs. A distinct increase of the BFO out-of-plane lattice parameter was found (to cpc = 3.992 ± 0.005 Å) when m increased to 8. On the other hand, for BMO only a small decrease of the out-of-plane lattice parameter occurred with increasing m (from cpc = 3.970 ± 0.005 Å for m = 1 to cpc = 3.950 ± 0.005 Å for m = 8). The small decrease in cpc for BMO with increasing m is consistent with the BMO bulk lattice being closer to STO and hence being subject to less in-plane compression. This was explored further from asymmetric RSM scans around the STO (103) reflection which reveals both the in-plane and out-of-plane strain states in the films (Figure 2c). In Figure 2c, the same (H00) positions were observed for the film and the substrate which confirmed coherent growth of the superlattices and full in-plane strain of both the BMO and BFO films in the superlattice structure. Interestingly, the BMO (H0L) peak hardly relaxed with m, while some peak broadening of the BFO (H0L) peak was observed indicating some in-plane relaxation. This means that although small changes in the BMO structure were observed, the change in tetragonal distortion (c/a) of the BMO layer was small (c/a = 1.018 for m = 1 to c/a = 1.013 for m = 8). For BFO a shift to lower reciprocal L value was also observed with increasing m, consistent with the downward shift of STO (004) reflection of Figure 2b. The results show that minor in-plane relaxation of the BFO occurs for m = 4, 8 and the out-of-plane lattice parameter increases toward bulk values, i.e., cpc of BFO increases with m. The appearance of in-plane relaxation with increasing m indicates that the BMO/BFO interface supports BFO to maintain the strain state.32 To compare the magnetic properties of SL–BFO/BMO with bulk BMO, we measured the field-cooled (FC, 500 Oe) temperature dependence of the magnetic moment (M–T) for all superlattices (Figure 3a). To determine the transition temperature, we undertook first derivatives of the M–T curves (dM/dT). For m = 4 and 8, a magnetic transition is observed at 96 K (the value for bulk BMO, TC,BMO). For all the films, a clear magnetic transition is observed at around 145 ± 15 K (TC1), which is significantly higher than the TC of bulk BMO. Interestingly, for m = 1, 2, a second magnetic transition is observed above room temperature (TC2), at 410 K for m = 1, and at 320 K for m = 2. To determine the origin of TC1 and TC2, we deconvoluted the M–T curves into two different transitions by using the saturation magnetic moment from the M–H loops measured at various temperatures (see Figure S1 in the Supporting Information). From the deconvoluted M–T curves, we can deduce that the high TC2 for m = 1, 2 originates from the change of the Mn environment in BMO. The environment of Mn in BMO for m = 1, 2 is different to m = 4, 8 since BMO is directly coupled to BFO, giving a single, strained BFO–BMO lattice (Figure 2) similar to the NCB–BFO/BMO lattice (Figure 1a). TC1 is observed for all the films. A plot of TC1, TC2, and TC,BMO versus m for the various superlattices is given in the inset of Figure 3a. Figure 3b shows the relationship between TC1 and the out-of-plane lattice parameter of BFO. To get a better understanding of the magnetic behavior of these superlattices, we also measured the magnetic hysteresis loops (M–H) of m = 1 and 8 superlattices at 10 K (Figure 3c). The MS (at 5 T) for the m = 1 superlattice was found to be 90 emu cc−1 (≈0.59 μB/f.u.) whereas the MS for the m = 8 was slightly lower, 80 emu cc−1 (≈0.52 μB/f.u.). The measured MS values are lower than that of bare BMO films grown on STO (1.87 μB/f.u.) and the estimated values of antiferromagnetically ordered Fe3+–O–Mn3+ (1 μB/f.u.) or ferromagnetically ordered Fe3+–O–Mn3+ (4.5 μB/f.u.).14, 19, 33 The lower MS value of our superlattices is likely to be caused by complex effects resulting from a mixture of AFM and FM ordering which can arise from internal charge transfer leading to mixed ionic valences for Mn and Fe.19 Also note that the in-plane BMO/BFO interface is present along the step edges. However, the effect of in-plane Fe–Mn interactions on the MS of our superlattices can be expected to be within the error of the measured values as the fraction is very small compared to the fraction of out-of-plane Fe–Mn interactions («1 %). The superlattices show an increase in coercivity (HC) (Figure 3c) compared to bare BMO films grown on STO (HC = 470 Oe).33 While for the m = 8 superlattice the enhancement of HC was doubled (HC = 900 Oe), for the m = 1 superlattice a much more dramatic enhancement was found (HC = 5.6 kOe). To clarify the origin of the increase of HC, we deconvoluted the experimental M–H loops into two different transitions (loop1 and loop2, see Figure S2 in the Supporting information). A squarer and narrower loop (loop1) is simulated by subtracting a trial function of the form M = MS tanh{α(H ± HC)} from the experimental data. The residual M–H loop (loop2) is of a reasonable form and has a much higher coercivity. For the m = 8 sample, the HCs are 400 Oe (loop1) and 4.5 kOe (loop2), and the HC of loop1 is close to that of plain BMO films.33 For the m = 1 sample, the M–H loop is decomposed into two M–H curves with HCs of 1.2 (loop1) and 9.2 (loop2) kOe. Therefore, it seems reasonable to assume that in both cases the loop1 relates to BMO. The enhanced HC for m = 1 can be understood by the general thickness dependence of HC which tends to decrease with increase of magnetic layer thickness.34, 35 The broader loop2 presumably originates from a ferromagnetic BFO phase with TC1 of ≈145 K for which the magnetic hardness can be understood in terms of an exchange interaction between FM– and a residual AFM–BFO matrix.36, 37 XAS near the Mn L2,3 and Fe L2,3 absorption edges was carried out to determine the Mn and Fe oxidation states (see Figure 4a,b). A clear change in the fingerprint of the Mn spectra versus m is present, indicating a variation of the oxidation state of our Mn spectra with of simple the of the m = 8 superlattice is close to that of indicating a of By m, the peak around and well the peak of our m = 1 superlattice with of simple the peak of Mn for the m = 1 superlattice is similar to the peak of that of all Mn are by the Fe a small change in versus m is present as The of the m = 8 superlattice shows a peak at with a clear at lower This strongly a Fe3+ with a small fraction of By m, the observed at lower is into the strong peak at our spectra with the this change in indicates an increase of when The observed for Fe and Mn states versus m are in with the XAS for the m = 1 superlattices, we clear mixed states of the B-site and while for m = 8 superlattices we have a of Fe3+ and also some and are present. This that charge transfer occurs from Fe to Mn near the the of both BFO and BMO, charge transfer from BFO to BMO is since is than and An for the presence of and would be of from the into the BMO layer. is to be on the order of in order to interfacial FM in and is suggested to in SL of and Note that a has only been observed when CaRuO3 and LaNiO3 are BFO is an BFO films are not highly to the presence of in the structure, small charge In charge transfer has been suggested to at the interface of BFO and to interface For all our superlattices, we found two magnetic transitions and clear at temperatures above the TC of BMO. For the very short period superlattices, m = 1, 2, was even found above room the change in tetragonal distortion of BMO was and above 1, the variation in magnetic behavior is not expected to result from a change in Therefore, to the origin of the magnetic states and the enhancement in we the various superexchange interactions between and that may be present in our m = 1 superlattices, the into of to these most of the interactions are AFM Also the interaction can be expected to be AFM as found in The presence of these AFM interactions may only decrease the MS of the Therefore, to the magnetic behavior of the BMO/BFO superlattices we on the ferromagnetic and and of the interactions present in BMO are likely to be present in the superlattices interactions give to an AFM However, due to the presence of BMO has magnetic structure, leading to FM For the m = 4, 8 superlattices, we can BMO to be present. For these superlattices, a clear magnetic transition is present around K, which is close to the TC of bulk BMO. Therefore, this transition likely originates from the interactions and is For the m = superlattices, we would the interaction to be This also clear transition is found around ferromagnetic interactions are observed in leading to a TC of K and a MS of Interestingly, the TC of is close to the TC1 of our Moreover, shows an enhancement of TC with increasing Interestingly, this is similar to the behavior of TC1 in our superlattices the unit of BFO increases with m as a result of the increase in out-of-plane lattice 2 and 3c). These strongly that TC1 may from the FM Therefore, a closer the found magnetic moment for of the deconvoluted M–T curves with TC1 show that magnetic moment increases when m (see Figure in the Supporting information). This is in with the increase of for lower m. The higher magnetic moment found for m = 2 compared to m = 1 can be explained by the of the interactions for m = 2. MS is at 10 K and 5 from the deconvoluted M–H loop2 shown in Figure S2 in the Supporting was found for m = 1, when the magnetic moment would from we assume of the Fe to be in the single layers of BFO, we may a magnetic moment of 1.2 for m = However, a magnetic moment of 1.2 would that Fe3+ and are within the layer. As Fe3+ and in-plane is the presence of AFM interactions would the magnetic the observed MS of for m = The between and our superlattices, and the of in the XAS spectra that TC1 originates from ferromagnetic interactions in The origin of TC1 is in Figure As in the ferromagnetic interactions between AFM BFO and FM BMO are to result in a TC of K for NCB–BFO/BMO.10 The high TC2 of up to 410 K in our very short period SL–BFO/BMO is very close to this SL–BFO/BMO films have the same magnetic ordering as ferrimagnetic ordering of AFM–Fe and (Figure The of TC2 for m = 4, 8 superlattices is consistent with the increasing out-of-plane lattice parameter in BFO structure with increasing m, and the separation of BFO and BMO into two distinct lattices 2 and 3c) than a single strain BFO–BMO Here, the of the BO6 at the BMO/BFO interfaces may be not for magnetic Similar effects of strongly magnetic properties to structural distortion have been observed in superlattices of and of TC were found in these strained superlattices, K for and 650 K for whereas in relaxed bulk and plain films of and lower of only and K, respectively, were that we for the origin of the high TC2 phase in our films is double exchange FM coupling of in BMO as a clear mixed of and Mn3+ was observed in the m = 1, 2 superlattices (Figure to on electron e.g., FM ordering with TC of K has been in for up to TC increases as the of However, the TC of is much lower than TC2 of SL–BFO/BMO. this could be due to a difference in BO6 which has been shown to affect the TC of In this double exchange FM coupling of result in behavior TC2 caused by electron as for with = However, of m, all our superlattices are highly which indicates that the is not the Therefore, we strongly the ferrimagnetic interaction between AFM–BFO and to be the origin of In in BFO/BMO superlattices of very high quality we observed above room temperature magnetic transition temperatures for very short period A TC of 410 K was observed for m = 1, and a TC of 320 K was observed for m = 2. This TC TC2, well to a nanoscale checkerboard BFO/BMO which was shown by to have a TC of 406 K as a result of FM BFO A lower TC1, was observed in all the superlattices (m = and can be explained by ferromagnetic interactions in BFO. superlattices of ferroelectric BFO and BMO with interface-induced high temperature magnetism are an important for practical magnetoelectric superlattices (m = 1, 2, 4, 8 of nm thickness were grown on (001)-oriented STO by using a = BFO and BMO were at a of and a of The substrate to was at to a growth which was to maintain coherent growth the substrate temperature was and the was at The films were in by the To determine the growth of BMO and BFO, the of single phase BFO and BMO films were by X-ray The growth for BFO and for BMO, were to the of the It was found that the BMO not which could be expected since the growth is by the highly as well as the large lattice between pseudocubic BMO and It is that the interface roughening of our superlattices increased with film thickness as confirmed by were carried out by XRD high and high transmission electron microscopy operating at 400 versus temperature and field were using a and a To investigate the state of Fe and Mn in the films, XAS was on at the The spectra were measured by electron at K at kOe. 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