Introduction

Perovskite oxides, represented by the ABO3 structure, are celebrated for their distinctive crystal and electronic structures1,2,3,4. These unique configurations engender sensitive and controllable spin-lattice-orbit interactions, which imbue these materials with significant potential for scientific research and engineering applications5,6,7,8. However, the intrinsic properties resulting from the inherent structure of the material that often limit the potential for further development. For example, the star material room-temperature multiferroics, BiFeO3. It has three typical structures such as the rhombohedral phase, the pseudo-cubic phase, and the super-tetragonal phase, each of which has its own characteristics and been widely researched. The first is the rhombohedral phase, which belong to the bulk phase BiFeO3 structure with room-temperature ferroelectric and antiferromagnetic properties. However, there are leakage currents exist in films due to the immaturity of the fabrication process in early studies, resulting in observing spontaneous polarization of only 6 μC cm−2 in the rhombohedral phase BiFeO3 films9. To achieve better properties, it has been found that pseudo-cubic phase BiFeO3 films obtained by using pulsed laser deposition onto single crystal substrates (SrTiO3) with small lattice mismatches can exhibit a spontaneous polarization of about 55 μC cm-210. The super-tetragonal phase BiFeO3 film with a large axial ratio (\(c/a\)) has likewise received much attention. One of its distinctive features is that it possesses a large \(c/a\) of  1.24, which makes it possess a large polarization value of ~ 150 μC cm−2 11. Although it has been reported to have long-range ordered ferroelectric polarization, the optimization of magnetic order using effective methods is also of great research importance to further exploit its application potential.

BiFeO3 is a notable room-temperature multiferroic material, classified as a type-I multiferroic12,13. This classification indicates that its ferroelectric and ferromagnetic orderings occur independently. The ferroelectric order parameters are related to the 6s2 lone-pair electrons of Bi3+14, while the ferromagnetic ordering is associated with the spin configuration of 3d orbitals of Fe3+ in the BiFeO315. Within its structural framework, the oxygen anion engages in strong electronic hybridization with B-site cations and simultaneously participates in complex electronic interactions with A-site cations16. Consequently, the substitution or modulation of the oxygen anion can exert a dual influence, affecting both the A–O and B–O hybridization states and impacting the overall material structure and properties17,18. In our previous study, we found that the ferroelectric order parameters can be modulated in [001]-oriented pseudo-cubic epitaxial BiFeO3 thin films17. This modulation was achieved through the strategy of introducing sulfur anions, leading to sulfur-induced rotation of the polarization state. In the anion substitution of other systems (e.g., SrNbO3), the structural and electronic can be modulated through nitrogen doping, showing the mechanism behind a significant metal-to-insulator transition19. These findings, along with an increasing number of studies, highlight the critical role of anion modulation. This modulation significantly influences the crystal structure, electronic structure, and physical properties of perovskite oxide thin films, thereby broadening the scope for the development of materials with innovative and unique properties.

Here, we propose a handy strategy for modulating the magnetic properties of perovskite oxide thin films by the anionic strategy with shear strain, which drives the reconfiguration of the lattice and electronic structures. The structural and ferromagnetic orderings of large \(c/a\) ratio super-tetragonal BiFeO3 (T’-BFO) epitaxial films were successfully modulated. A remarkable finding was sulfurization-induced rotation in the ferromagnetic ordering direction of T’-BFO thin films, which resulted in enhanced magnetic collinearity. In addition, this process was characterized by a shift in the direction of the easy magnetic axis from out-of-plane to in-plane. This anion modulation approach introducing shear strain is broadly applicable to most perovskite oxide. It holds the potential to modulate localized electronic hybridization states, offering new pathways to explore magnetic ordering phenomena in oxide materials.

Result and discussion

The magnetic properties of BiFeO3 originate from the spiral spin configuration in the 3d orbital of Fe ions20, which exhibit a periodic long-range antiferromagnetic order between the neighboring Fe ions along the [111] direction21, resulting in a near-zero macroscopic magnetic moment22. However, the antiferromagnetic order can be rearranged due to the lattice structural distortion. The BiFeO3 thin film characterized by the tetragonal-like structure and large \(c/a\) ratio represents the significant variant within their systems, which breaks the original antiferromagnetic order and exhibits weak ferromagnetism under the influence of substrate clamping strain and structural deformation23. Based on that, we hope to induce structural distortion3 and modify the electronic hybridization of Fe ions with surrounding anions in the thin films by introducing shear strain through sulfurization (Fig. 1a). This approach is intended to trigger a reconfiguration of the electron spins of Fe ions, thereby facilitating the transformation in the magnetic ordering of thin films. In addition, the spin cycloid of the bulk BiFeO36,24 is susceptible to the effects of substrate clamping stress and sulfurization, which may regulate the magnetic axis direction of the thin film.

Fig. 1: Design strategy for modulating the magnetic properties.
figure 1

a Schematic diagram of lattice structure evolution of BiFeO3 in the presence of LAO and sulfurization. b M represents the net magnetization vector in the film, and the green arrows show the direction of magnetic moments of Fe ions in T’- and O-phase BiFeO3 thin films. c, d Design strategies for local polymorphic distortion and magnetic collinearity enhancement to achieve large saturation and remanent magnetizations.

Full size image

As shown in Fig. 1b, the magnetic moment of Fe ions generates two distinct components along the out-of-plane and in-plane directions, resulting in magnetic anisotropy25,26. It means that the different magnetic fields are required to reach saturation magnetization in out-of-plane and in-plane directions, and the direction requiring a lower magnetic field is the easy magnetization axis. Here, the advisable sulfurization method has been applied to the BiFeO3 thin film with a large \(c/a\) ratio tetragonal-like structure27, introducing the shear strain and provoking the redistribution of Fe–O hybridization states. This also presents an opportunity to obtain the large saturation and remanent magnetizations in the sulfurized T’-BFO thin film (O-BFOS) (Fig. 1c). It is worth noting that this enhancement of the magnetic properties is inseparable from the increase of the magnetic collinearity in the O-BFOS thin film according to the Landau phenomenological theory28,29. This is complemented by an increase in the energy barrier required for magnetization, leading to a more dispersed energy distribution in the film after sulfurization (Fig. 1d). Hence, this strategy can maximally preserve the hysteresis while maintaining high magnetization, which is significant for regulating the magnetic properties of BiFeO3 films.

T’-BFO thin films were epitaxially grown on (001)-oriented LaAlO3 (LAO) substrates by radio-frequency (RF) magnetron sputtering. Figure 2a shows the out-of-plane synchrotron X-ray diffraction (XRD) patterns of both the T’-BFO and O-BFOS. The XRD patterns of the T’-BFO film indicated an out-of-plane lattice constant c≈4.66Å and a \(c/a\) ratio \(\approx 1.23\), suggesting the successful fabrication of the super tetragonal-like phase structure of BiFeO3 films30,31,32. The absence of additional diffraction peaks in the XRD pattern, except for the (00 l) diffraction peaks of the film and substrate, confirms that both the T’-BFO and O-BFOS thin films are single-phase epitaxial films. The thickness of T’-BFO and O-BFOS films are 78 nm and 80 nm, respectively (Supplementary Fig. 1). This indicates that the sulfurization process has a negligible influence on the macroscopic thickness of the films. Notably, the out-of-plane diffraction peaks of the T’-BFO films shift to a higher angle following sulfurization, and the phenomenon also observed in sulfurized pseudo-cubic phase BiFeO3 films17. This shift not only demonstrates the successful sulfurization in T’-BFO films but also suggests that the incorporation of S anions induces lattice strains as well.

Fig. 2: Crystal structure characterization.
figure 2

a Out-of-plane XRD pattern. b Φ scans normal to the (101) plane of T’-BFO, O-BFOS, and LAO, respectively. c High-resolution (103) plane synchrotron RSM of T’-BFO and O-BFOS thin films. RSM diffraction features corresponding the MC-phase and the O-phase. d Polarization-dependent SHG patterns of T’-BFO and O-BFOS thin films at p-out and s-out configurations. e, f HADDF-STEM images of MC-phase and O-phase structures and their corresponding in-plane and out-of-plane lattice fringes and GPA analysis maps.

Full size image

To substantiate the effect of sulfurization on the symmetry of the films, Φ scans were performed on the (101) plane of the T’-BFO, O-BFOS film, and LAO substrate. As illustrated in Fig. 2b, all scans exhibit four diffraction peaks following a 360° rotation, with an equal angular spacing of 90°, indicative of a fourfold symmetry. This pattern confirms that T’-BFO films can form a good match with the LAO lattice when grown epitaxially on (001)-oriented single-crystal LAO substrates, and this epitaxial relationship persists even after sulfurization. This finding is highly consistent with the phenomenon observed in STEM (Supplementary Fig. 2). Notably, the Φ scanning diffraction peak of the T’-BFO thin film undergoes a slight splitting compared to the sulfurized films, which is consistent with the monoclinic tetragonal-like phase as reported in the literature (MC-phase)33,34. Furthermore, the full width at half of the maximum (FWHM) of the Φ scanning diffraction peaks of the films before and after sulfurization were further compared to prove the transformation in structure. The results show that there is an increase in the average FWHM of the sulfurized film compared to the MC-phase BiFeO3 thin film, from 2.63 to 3.13 (Supplementary Fig. 3). It is also important to note that all four diffraction peaks of the films after sulfurization are symmetrical, whereas the two peak intensities are not same after splitting in super-tetragonal phase BiFeO3 films. Therefore, the differences in Φ scanning gradually begin to unveil a veil of stress action induced by sulfur anions.

Subsequently, the reciprocal space mapping (RSM) around (103) reflections of T’-BFO and O-BFOS thin films are performed to further investigate their crystal structure (Supplementary Fig. 4). The results show that the sulfurized O-BFOS film experiences compressive strain in the out-of-plane direction, with in-plane relaxation, suggesting a distortion of the lattice structure due to the sulfurization effect. As depicted in Fig. 2c, high-resolution RSM mappings of the (103) reflections of T’-BFO and O-BFOS thin films were compared. It was observed that the T’-BFO thin film displayed three distinguishable diffraction points in the (103) RSM mapping, a consequence of its monoclinic distortion along the [100] direction relative to the tetragonal phase, characteristic of a typical monoclinic MC-phase film. Interestingly, the (103) RSM mapping of the O-BFOS film tends to evolve into an elliptical shape, indicating the diminishing monoclinic distortion along the [100] direction. This leads to the different in-plane lattice constant for a≠b (Fig. 2d), inducing a transformation to the orthorhombic phase (O-phase), indicating that the introduction of S anions induces lattice distortion through chemical strain. In summary, the bulk R-phase BiFeO3 was epitaxially grown by RF magnetron sputtering on a (001)-oriented LAO substrate and the structural distortion of BiFeO3 tends to form a monoclinic MC-phase due to substrate clamping strain. Subsequently, the introduction of chemical strain through sulfurization causes further structural distortion, leading to the formation of the O-phase.

Fast Fourier transform (FFT) patterns of O-BFOS films along the a-axis direction further confirm the films’ transition to the O-phase after sulfurization, as observed in the microscopic morphology. The FFT analysis for regions b and c reveals a superlattice diffraction spot at \(\left(\frac{1}{2}\,\frac{1}{2}\,\frac{1}{2}\right)\), indicated by the red circle in Supplementary Fig. 5, in addition to the FFT for region a, which provides significant evidence for the O-phase transition. The cross-sectional scanning transmission electron microscopy energy dispersive X-ray spectroscopy (STEM-EDX) mapping and X-ray absorption spectroscopy (XAS) analysis confirm the presence of sulfur. The XAS analysis shows that the O-BFOS had a prominent S L-edge, which was not observed in T’-BFO, confirming the success of sulfurization (Supplementary Fig. 6). The STEM-EDX mapping display that the O-BFOS films and the LAO substrate have the clear interface, indicating good epitaxial quality (Supplementary Fig. 7). The elements Bi, Fe, O, and S are found in O-BFOS films, and all elements are evenly distributed throughout the film with no obvious areas of elemental aggregation. In addition, the presence of S elements can be demonstrated under low-resolution EDX mapping, providing further evidence for the successfully sulfurization of T’-BFO films. We observed that the sulfur content in O-BFOS films was approximately 0.48% using STEM-EDX (Supplementary Fig. 8). This outcome provides essential data for a deeper understanding and manipulation of the composition and structure of the thin film, as well as for subsequent research on anion substitution. As is well-known, the properties of BiFeO3-based thin films are intimately related to their lattice structure. The lattice distortion induced by sulfurization alters the hybridization state between the anions and cations in the films, thereby affecting the associated physical properties. Consequently, the modified films are expected to exhibit novel physical properties characteristics.

The \(c/a\) ratio of the BiFeO3 thin film increases significantly under biaxial clamping stress, showing a tetragonal-like phase with a large axial ratio. It has been reported that the remanent polarization of the film can reach about 130 ~ 150 μC cm−2 11,35. The optical second-harmonic generation (SHG) indicates that the films still has robust ferroelectricity after sulfurization (Supplementary Fig. 9)36. We extracted the relative displacement between the Fe and Bi ions (along the [001] direction) using high-resolution STEM images, the result showed that the relative displacement in the O-phase was reduced with 0.06 Å compared to the MC-phase (Supplementary Fig. 10). Based on the empirical linear relationship between polarization strength and relative displacement, the polarization intensity of the films before and after sulfurization was semi-quantitatively obtained to be about 129.9 and 103.9 μC cm−2, respectively. Out-of-plane phase images, along with the local phase and amplitude hysteresis loops of the piezoelectric response also confirm the intrinsic ferroelectricity of T’-BFO and O-BFOS films (Supplementary Figs. 11 and 12). Notably, time-dependent phase images of O-BFOS films, the following writing at + 10 V DC voltage, indicate strong ferroelectric retention even after modification with sulfurization (Supplementary Fig. 13). In addition, the SHG intensity of T’-BFO and O-BFOS films were also measured in s-out and p-out modes, which were obtained by varying the analyzer polarization vertical or parallel to the incident light field, respectively (Fig. 2d). The s-out patterns of the T’-BFO and O-BFOS films show distinct peaks near 45°, 135°, 225°, and 315°, which can be attributed to the fact that the films have 4 mm symmetry. The change of peaks in the p-out indicates that sulfurization causes a tendency for the film to shift from dual rotational symmetry to quadruple symmetry, which also responds to the effect of sulfur anion-induced stress on the structure of the film.

The MC- and O- phases were analyzed based on the large-scale high-angle annular dark field scanning transmission electron microscopy (HAADF-STEM) images along the [100] zone axis (Fig. 2e1, f1). The inverse FFT lattice fringes in both in-plane and out-of-plane directions clearly show the difference between the two structures before and after sulfurization (Fig. 2e2, e3, f2, and f3). In addition, the shrinkage of lattice parameters in the in-plane and out-of-plane directions after sulfurization are also consistent with the trends of RSM and XRD. Specifically, The MC-phase exhibits a pronounced wave shape in the in-plane direction (Fig. 2e2) while the O-phase is characterized by the longer and straighter line (Fig. 2f2), confirming that sulfurization leads to a transition towards the more symmetrical O-phase. Subsequently, the strain states in the in-plane and out-of-plane directions of the MC- and O- phases were quantified using geometric phase analysis (GPA). The GPA mappings show that the O-phase is subjected to higher compressive strain in the in-plane direction compared to the MC-phase (Fig. 2e4, f4) and lower tensile strain in the out-of-plane direction compared with the MC-phase (Fig. 2e5, f5). These findings not only highlight the impact of chemical strain from sulfurization on the lattice structure but also provide important support for the origin of the improved magnetic ordering of the films.

The magnetic hysteresis loops of the films were measured at 300 K by a Magnetic properties measurement system (Fig. 3a, b), indicating that the films before and after sulfurization exhibit significant ferromagnetic properties. It can be observed that both films are magnetically anisotropic before and after sulfurization. Notably, the saturation magnetization of the O-BFOS thin film is larger than that of the T’-BFO thin film both in the out-of-plane and in-plane directions. The zero-field cooling (ZFC) and field cooling (FC) curves show that the T’-BFO and O-BFOS films gradually becomes paramagnetic with the increasing temperature, and the macroscopic magnetic moment approaches zero (Fig. 3c). Their magnetic ordering temperatures can be obtained by derivation of FC curves for T’-BFO and O-BFOS films, which are 573 K and 586 K (Fig. 3d), respectively. The results show that sulfurization can enhance the magnetic transition temperature of the films, which is much higher than the room temperature. Moreover, the result is also in general agreement with the reported transition temperatures calculated using the Monte Carlo method (MC)37. The MC simulations show that the higher magnetic transition temperature of the film increases as the in-plane lattice constant \(a/b\) expands in MC-phase T’-BFO films, which could also be a potential reason for the increase in the magnetic transition temperature of O-BFOS films after sulfurization37,38.

Fig. 3: Characterization of macro magnetic properties.
figure 3

a, b Out-of-plane and in-plane magnetic hysteresis loops of T’-BFO and O-BFOS thin films at 300 K. c ZFC and FC curves of T’-BFO and O-BFOS thin films with a test magnetic field of 1000 Oe and orientated along the out-of-plane direction, the inset shows a local magnification of the out-of-plane hysteresis loops of T’-BFO and O-BFOS thin films at 300 K. d The first-order derivatives of FC curves for T’-BFO and O-BFOS thin films. e The ratio of magnetic properties of O-BFOS to T’-BFO thin film, where the saturation magnetization and remanent magnetization are denoted as Ms and Mr, respectively. The inset shows the magnetic field-dependent SHG of O-BFOS films at 280 K. f The magnetic fields required for T’-BFO and O-BFOS thin films to reach Ms along the in-plane and out-of-plane directions.

Full size image

The changing trend of the magnetic properties after sulfurization is demonstrated in Fig. 3e. The saturation magnetization of the O-BFOS film in the out-of-plane direction is enhanced from 23.6 to 28.3 emu cc−1, which is about 20% higher than that of the T’-BFO film, whereas the enhancement is from 22.3 to 29.5 emu cc−1 in the in-plane direction, which is about 32% higher. However, the enhancement of remanent magnetization of the O-BFOS film even more remarkably compared to saturation magnetization changes. In the out-of-plane direction, the remanent magnetization of the O-BFOS film increases to 5.7 emu cc−1, which is 1.6 times that of T’-BFO. In the in-plane direction, the remanent magnetization increases by 62 times to about 13.3 emu cc−1, demonstrating that sulfurization has a significant effect on the magnetic properties of the film. The large remanent magnetization ensures the advantages of stability, sensitivity, accuracy, and low power consumption in potential application scenarios. We have investigated the magnetoelectric coupling effect in films using the magnetic field-dependent optical SHG technique at 280 K (Supplementary Fig. 14). The findings reveal that the optical magnetoelectric coupling constant of O-BFOS films measure 69.66 (absolute value), which is comparable in magnitude to that of T’-BFO, indicating the robustness of the magnetoelectric coupling effect against sulfurization. Moreover, magnetic field-dependent SHG measurements of O-BFOS films at p-out and s-out configurations show that the SHG intensity changes with the variation of the applied magnetic field strength in both modes, showing an obvious magnetic field dependence (Supplementary Fig. 15). Notably, the magnetoelectric coupling in O-BFOS films persists at near room temperature, highlighting its promising potential for applications in multifunctional devices. Additionally, sulfurization not only reduced the saturation magnetic field of the T’-BFO thin film in the in-plane direction by one-third, from 1.8 T to 1.2 T but also increased the saturation magnetic field in the out-of-plane direction to 1.8 T, ultimately leading to a shift of the magnetic easy axis of the film from out-of-plane to in-plane direction (Fig. 3f and Supplementary Fig. 16).

We provide direct evidence for ferromagnetic behaviors, magnetic anisotropy, and magnetic easy axis flipping in thin films by obtaining XAS of Fe L-edge. As shown in Fig. 4a, the detection of Fe L-edge XAS uses left and right circularly polarized light to verify the spin magnetic moment and magnetic anisotropy of Fe in thin films. Due to the difference in the density of spin-up and spin-down electronic states around Fe ions in thin films, there is a difference in the absorption of left- and right-spin circularly polarized X-rays, resulting in the phenomenon of X-ray magnetic circular dichroism (XMCD)3. The Fe L-edge XAS and XMCD spectra in T’-BFO and O-BFOS thin films measured under out-of-plane and in-plane magnetic field are shown in Fig. 4c, d. The XMCD patterns, observable with the magnetic field aligned along either direction, confirm that the films’ magnetism arises from the electronic spin magnetic moments of Fe, indicating magnetic anisotropy.

Fig. 4: X-ray absorption measurements.
figure 4

a, b Schematic diagram of the XMCD and XMLD measurement. The difference in the absorption of left- and right-spin circularly polarized (LCP and RCP) X-rays, resulting in the XMCD spectra. The polarization direction of the X-rays was fixed to horizontal line polarization, and the X-ray incidence angles were adjusted to 30° and 90° to obtain the XMLD spectra. c, d Out-of-plane and in-plane XMCD of T’-BFO and O-BFOS thin films. e, f The XAS Fe L-edge of T’-BFO and O-BFOS thin films were obtained by changing the incident angle with a fixed linearly polarized X-ray beam. g Fe L-edge XMLD spectra of T’-BFO and O-BFOS thin films.

Full size image

The shift of the magnetic easy axis direction of the sulfurized film from the out-of-plane direction to the in-plane direction was further verified by X-ray magnetic linear dichroism (XMLD)39. As shown in Fig. 4b, the Fe L-edge spectra of T’-BFO and O-BFOS films were obtained by keeping the direction of the linearly polarized light as horizontal linear polarization, changing the angle between the film and the X-rays to 30° and 90°, respectively, and applying a magnetic field along the direction of the X-rays. Furthermore, the in-plane and out-of-plane orbital information were obtained by \({I}_{{\mbox{ab}}}={I}_{90^\circ }\) and \({I}_{{\mbox{c}}}=\left({I}_{90^\circ }-{I}_{30^\circ }{\left(\sin 30^\circ \right)}^{2}\right)/{\left(\cos 30^\circ \right)}^{2}\), respectively. As shown in Fig. 4e, f, the variation indicates that the coordination environments of Fe ions are differ between the in-plane and out-of-plane directions, thereby confirming the presence of magnetic anisotropy in the films both before and after sulfurization40,41. Meanwhile, the XMLD spectra of T’-BFO and O-BFOS films are completely different (Fig. 4g). The XMLD spectra of the T’-BFO film show a trend of first increasing, then decreasing, and then increasing, while the XMLD of the O-BFOS film shows the opposite trend. The magnetic properties of the BiFeO3-based films originate from the electron spin of the Fe 3d orbitals, and the orientation of the magnetic easy axis is closely related to the hybridization between the Fe 3d and the nearby O 2p orbitals42,43. Accordingly, the differences in the directions of three peaks A, B, and C indicate the alteration of the electronic hybridization states of Fe ions for both in-plane and out-of-plane directions in the T’-BFO and O-BFOS films. This observation serves as compelling evidence supporting that sulfurization induces a reorientation of the easy magnetic axis in the O-BFOS films from the out-of-plane to the in-plane direction44.

In general, the lattice distortion or rotation of the BO6 octahedron in the ABO3-type structure will initiate a change in the local electronic structure, which leads to fundamental differences in the macroscopic physical properties of films45,46,47. For the T’-BFO film, the \(c/a\) ratio can reach ~ 1.23, causing the Fe ions to shift from the center of the FeO₆ octahedron due to the LAO substrate clamping stress, which exceeds 4.5%. According to the ligand field theory and additional Jahn-Teller-type level splitting, the local structure in the T’-BFO films became the pyramid-like FeO5 structure (Supplementary Fig. 17). At the same time, due to the large elongation of T’-BFO films, the Fe ions in square pyramidal site symmetry were shifted along the out-of-plane direction to mitigate the strong Coulombic repulsion between the Fe ions and neighboring oxygen ions in the in-plane direction. After sulfurization, the absorption peaks of Fe L-edge of the O-BFOS film exhibit significant t2g and eg orbitals compared to that of the T’-BFO film, indicating that the local structure have been rebuilt and trended to a new balance (Fig. 5a). The result provides necessary evidence for the local structural transitions to FeO6 octahedral site symmetry due to the effect of sulfur-introduced stress. Despite this remarkable transformation in the local coordination geometry, the Fe ions remain predominantly in a trivalent state48,49. The O K-edge absorption spectra consist of three peaks, a, b, and c, which represent the hybridization of O 2p with Fe 3d, Bi 6sp, and Fe 4sp orbitals, respectively42. The change of the a peak in the film before and after sulfurization is similar to the trend of the Fe L-edge XAS. Furthermore, a decreasing trend is observed for the c peak in the O-BFOS film, which further indicates that the local electronic structure surrounding Fe ions is modified under the effect of sulfurization.

Fig. 5: Mechanisms of magnetism transition.
figure 5

a XAS of Fe L-edge and O K-edge of T’-BFO and O-BFOS thin films and the schematic of Fe 3d orbital in different thin films. b Total energy difference with respect to the lowest-energy state with magnetization axis along [100] direction in O-BFOS. The energy with a spin axis of [100] is set to zero. c The charge density difference of T’-BFO and O-BFOS thin films. Each one was calculated as the difference between the charge density of the crystal and the superposition of the charge densities of the isolated atoms. The isosurfaces around oxygen indicate electron gain, while those around iron signify electron loss; hence two colors have been used to distinguish them.

Full size image

In the BiFeO3 system, the origin of magnetism lies in the fine arrangement of the electron spins in the 3d orbitals of the Fe ions. Since the electron spin of neighboring Fe ions in BiFeO3 is mediated by non-magnetic O ions, this interaction triggers Fe–O–Fe superexchange interactions49,50, resulting in the anomalous sensitivity of individual Fe ions hybridized with surrounding anions to slight variations in the crystal structure. However, the lattice structure of films shifts towards the O-phase during sulfurization, which significantly increases the single ion anisotropy of Fe ions, inducing notable changes in the magnetic properties of the films. The key point is that both XAS and STEM results show a shift of the Fe–O hybridization from the square pyramidal site symmetry structure to the octahedral site symmetry, leading to a shift of the spatial-inversion symmetry breaking. Therefore, our analysis suggests that the magnetic enhancement may be related to the following aspects: (i) The change in symmetry breaking will result in a shift of the equilibrium positions of positive and negative charges within the unit cell, as well as alterations in the relevant Dzyaloshinskii–Moriya (DM) interactions51, positively influencing magnetism. In addition, the observed weakening of ferroelectricity in the sulfurized films aligns with the conclusion that ferromagnetism and ferroelectricity are typically mutually exclusive52, which further supports the reliability of the enhancement of the magnetic properties. (ii) The structural changes driven by sulfurization lead to alterations in the electronic overlap states between Fe ions and surrounding anions, as well as a significant increase in the single-ion anisotropy, which may play an important role in enhancing the remanent magnetic moment through its influence on the superexchange interactions. In other words, these changes may be more favorable for the stable existence of a larger magnetic moment at zero magnetic field, thereby leading to a significant alteration in the remanent magnetization. In addition, there are report indicate the observation of diverse changes in the cycloidal states in BiFeO3 films with thicknesses of approximately 90 nm and above, with these changes exhibiting a strong dependence on film thickness53. Due to the potential for sulfurization-induced structural and electronic reconstruction to alter the cycloidal states, which may further impact the uncompensated magnetic moments. Furthermore, considering the complex cycloidal states in highly strained BiFeO3 films, we predict that a possible reason may be associated with changes in the modulation of the cycloidal spirals. This warrants further investigation and attention in future research. In conclusion, the changes in the magnetism of the films are the result of the interplay of various factors influenced by lattice structure changes and localized electronic reconstruction induced by anion strategies, and we also look forward to more systematic studies in the future under specific conditions.

In addition, the Fe ions in T’-BFO films hybridize with the neighboring O ions under strong in-plane compressive stresses to form a pyramid-like of the FeO5 structure. The magnetic easy axis of the Fe ion is in the out-of-plane direction due to its closest distance and maximum hybridization with the top O ion, which makes it easier for the outermost electrons of the Fe ion to spin around the out-of-plane direction and generate magnetic moments. Subsequently, Fe ions are driven back to FeO6 octahedral hybridization with the neighboring O ions by sulfurization. Importantly, the lattice of the O-BFOS film relaxes in the in-plane direction due to S ions, which leads to increased hybridization of Fe ions with in-plane O ions while rotating the outermost electrons around the in-plane direction, thus tilting its easy magnetic axial more towards the in-plane direction.

To investigate the evolution of electronic and magnetic properties upon anion doping, we have performed spin-polarized density functional theory (DFT) calculations with the inclusion of spin-orbit coupling (SOC). Since O-BFOS has an orthorhombic structure, we calculated the total energies of the model by orientating the spins in different directions of [100], [010], and [001]. The energetic profile is shown in Fig. 5b, with the lowest-energy state aligned along the [100] spin direction, which serves as the energy reference set to zero. This indicates that the preferred magnetization axis of O-BFOS is [100], which is consistent with our experiment. As shown in Fig. 5c, the charge density distribution of Fe ions and anions in T’-BFO and O-BFOS films is visualized. The results indicate that compared to the pristine T’-BFO, the doping of S ions leads to an increased electron occupancy around the anions in the O-BFOS films, signifying an enhancement of the hybridized orbitals. Moreover, the change in the orientation of electron occupancy around the anions could be an important piece of evidence for the rotation of the easy magnetic axis.

In conclusion, the effective modulation of magnetic order in super tetragonal phase T’-BFO films are closely related to the anionic strategy, which drives the lattice structure and electronic hybridization towards new equilibrium states, thus influencing the physical characteristics. This process demonstrated the transition of the film structure from a monoclinic tetragonal phase to an orthorhombic phase in the presence of sulfur anions. The origin of the sulfur-induced enhancement in the degree of magnetic ordering and the rotation of the easy magnetic axis was further investigated through the strong correlation observed between changes in lattice structure and local electronic hybridization, as evidenced by techniques such as XMLD, XMCD, and atomic scale imaging. Furthermore, anion substitution engineering can be broadly extended to explore the local structure, magnetic ordering, and ferroelectricity ordering of perovskites. It holds the potential to revolutionize the current landscape of multiferroic and magnetoelectric materials, paving the way for the development of advanced materials with tailored properties for various applications in electronics and spintronics.

Methods

Thin-film synthesis

The epitaxial T’-BFO thin films were grown on single-crystalline (001) LAO substrates by using a radio frequency (RF) magnetron sputtering system (VJC-300, Beijing VNANO Vacuum Technology CO., LTD.). The LAO substrates have the lattice parameter of a = b = c = 3.79Å while the pseudo-cubic bulk BiFeO3 possesses the lattice parameter of a = b = c = 3.96Å. The epitaxial growth of BiFeO3 thin films on LAO substrates results in a significant lattice mismatch, causing substantial in-plane compressive strain of approximately 4.5%. These strain conditions lead to an expansion of the out-of-plane lattice constant, forming tetragonal-like films with large \(c/a\) ratios, denoted as T’-BFO. First, the epitaxial T’-BFO thin films were grown at 0.4 Pa and 500 °C under an O/Ar = 3/1 atmosphere. The sputtering time was 60 min, followed by in situ annealing for 20 min to obtain T’-BFO films. The anion-substituted thin films were prepared through a secondary treatment. A 0.2 M thiourea solution, dissolved in ethanol and water (solvent volume ratio of 4:1), was spin-coated onto the T’-BFO thin films at 3000 rpm, followed by pyrolysis at 350 °C for 5 min and annealing at 600 °C for 20 min to obtain the O-BFOS thin films.

Structural and chemical characterization

The structure, thickness, and orientation of the thin films were investigated via synchrotron XRD, X-ray reflectivity (XRR), and RSM measurements, which were conducted at the Diffuse X-ray Scattering Station of the Beijing Synchrotron Radiation Facility (BSRF-1W1A). Besides, XAS measurements were performed at the photoelectron spectroscopy station of BSRF-4B9B.

Electrical and electron microscopy analyses

The piezoresponse force microscopy (PFM) was performed by Bruker Icon and FSM FM-Nanoview-R. The STEM experiments were performed with a 200 kV JEOL ARM electron microscope equipped with double aberration correctors. The STEM–EDX experiments were conducted at 200 kV with the Super EDX detectors. Geometric phase analysis (GPA) be used to estimate the spatial distribution of strain by combining the Real-space and Fourier-space information of the STEM images.

Magnetic measurements

The magnetic field-dependent magnetization (M-H) loops were measured using a Quantum Design magnetic property measurement system (MPMS XL-7) in the temperature range of 10–700 K with the magnetic field applied along the in-plane and out-of-plane direction of the substrate and were obtained after subtracting the diamagnetic background of the substrate. It should be noted that the ZFC-FC were measured at an out-of-plane magnetic field of 1000 Oe, thus resulting in a lower ZFC-FC curve for the O-BFOS than that of the T’-BFO film. XMCD and XMLD characterizations were performed in total electron yield mode at Beamline BL07U and BL08U at the Shanghai Synchrotron Radiation Facility. Due to the difference in the density of spin-up and spin-down electronic states around Fe ions in thin films, there is a difference in the absorption of left- and right-spin circularly polarized X-rays, resulting in the phenomenon of XMCD. XMCD measurements were collected in total electron yield (TEY) mode at 300 K with the magnetic fields of H = 6000 Oe. For the XMLD, the Fe L-edge spectra of the T’-BFO and O-BFOS thin films were obtained by keeping the direction of the linearly polarized light as horizontal linear polarization, changing the angle between the film and the X-rays to 30° and 90°, respectively, and applying a magnetic field (H = 4000 Oe) along the direction of the X-rays. Furthermore, the in-plane and out-of-plane orbital information were obtained by \({I}_{{\mbox{ab}}}={I}_{90^\circ }\) and \({I}_{{\mbox{c}}}=\left({I}_{90^\circ }-{I}_{30^\circ }{\left(\sin 30^\circ \right)}^{2}\right)/{\left(\cos 30^\circ \right)}^{2}\), respectively. The XMLD spectra can be obtained by making a difference between the Fe L-edge XAS spectra obtained at Ic and Iab, denoted as \({I}_{{\mbox{XMLD}}}={I}_{{\mbox{c}}}-{I}_{{\mbox{ab}}}\).

SHG measurements

The incident laser beam for the SHG measurements was generated by a Spectra-Physics Maitai SP Ti:Sapphire oscillator whose central wavelength is 800 nm and repetition frequency is 82 MHz. The incident power was fixed at 50 mW and focused on the sample surface with a diameter of ~ 100 μm. The p-out and s-out configurations were adopted during measurement, which denote the arrangements where the analyzer polarization is parallel and vertical to the plane of the incident light field, respectively. The polarization direction φ of the incident light field is adjusted by the rotation of the λ/2 waveplate driven by a rotating motor.

DFT calculations

To investigate the evolution of electronic and magnetic properties upon anion doping, we have performed spin-polarized density functional theory (DFT) calculations with the inclusion of spin-orbit coupling (SOC). The T’-BFO is simulated by using a 2 × 2 × 2 40-atom BiFeO3 superlattice with monoclinic distortion. In addition, we used a 40-atom orthorhombic superlattice model with the virtual crystal approximation to approximate the O-BFOS. In these models, the lattice parameters are constrained onto the experimentally measured values.

Specifically, the spin-polarized density functional theory (DFT) calculations of BFO and S-doped BFO are carried out using the Vienna Ab initio Simulation Package (VASP)54,55. The exchange functional, Perdew, Burke, and Ernzerhof56 parameterized in the generalized gradient approximation was used with the energy cutoff of 520 eV for the plane-wave basis set. The ab initio calculation jobs were partially managed by the Quacc code57. It is important to include the Hubbard U correction for Fe-3d electrons. The former DFT studies58,59,60 have shown that the Hubbard U of 6 eV is able to well reproduce the magnetic and electronic properties of T’-BFO or O-BFOS. Therefore, we have applied the effective U of 6 eV in all the calculations using Dudarev’s method61. The spin-orbit coupling (SOC) is included in the DFT calculations to investigate the magnetic ground state. We use 40-atom BFO superlattices in the simulations, where the cell parameters were constrained to the experimental values, and the internal coordinates were fully optimized to meet the energy criterion of 10−6 eV and the force criterion of 5 × 10−3 eV/Å. Because the magnetocrystalline anisotropy energy is a quite small quantity compared to the system’s total energy, we have improved the energy accuracy to 10−7 eV when comparing the total energies between different magnetic structures. The k-mesh uses 4 × 4 × 3. We also tested with k-mesh of 6 × 6 × 6 and cutoff energy of 600 eV and found the magneto crystalline anisotropy energy is nearly not changed, indicating the calculations were well converged.

Reporting summary

Further information on research design is available in the Nature Portfolio Reporting Summary linked to this article.