Strong Enhancement of Room Temperature Ferrimagnetism in Insulating CoFe₂O₄ by Site‐Specific Hydrogen Intercalation

ABSTRACT

Strong enhancement of magnetization is highly desirable for increasing device efficiency, improving temperature stability, enhancing sensitivity, and enabling device miniaturization. Here, we demonstrate a 2.6−4.7 times higher room temperature ferrimagnetic moment (730−770 emu cc⁻¹) compared to 170−260 emu cc⁻¹ in as‐grown CoFe₂O₄ films, achieved by site‐specifically intercalating hydrogen ions near the Fe³⁺ in tetrahedral sites. While the spin moments of the Fe^{3 +} at octahedral and tetrahedral sites cancel each other out in as‐grown films, X‐ray spectroscopy of hydrogenated films reveals that both Fe³⁺ and Co²⁺ at octahedral sites can contribute to the enhanced total magnetic moment. This robust enhancement in insulating CoFe₂O₄, exhibiting high reproducibility over ≈10 cycles and long‐term stability over ≈10 d, suggests that hydrogenic‐ferrimagnetic coupling holds great promise for ultralow power electronics operating at high frequencies.

We enhanced room temperature ferrimagnetic moment (730–770 emu cc⁻¹) by 2.6–4.7 times in hydrogenated CoFe₂O₄ films via site‐specific intercalation of hydrogen ions near tetrahedral Fe³⁺ sites and associated electron transfer. The simultaneous achievement of high reproducibility and long‐term stability suggests great promise for ultralow power spintronic devices.

1. Introduction

Ferrimagnetism, origin of permanent magnetization in lodestone (mineral magnetite, Fe₃O₄) [1, 2], appears when a net magnetization remains due to the non‐compensating magnetic moments of the two sublattices. Most ferrimagnets are electrical insulators so their insulating nature suppresses Joule heating caused by eddy currents, ensuring low energy loss and high permeability even under a rapidly oscillating magnetic field [1, 3]. Ferrimagnetic insulators have attracted great attention as core components in high frequency electronics, including antennas, transformers, and microwave devices [2, 3, 4, 5, 6, 7]. In addition, many ferrimagnetic insulators as oxides exhibit superior corrosion resistance compared to metallic ferromagnets [1], as exemplified by spinel ferrites (CoFe₂O₄, NiFe₂O₄, CuFe₂O₄), garnets (Y₃Fe₅O₁₂, Gd₃Fe₅O₁₂, Dy₃Fe₅O₁₂, Ho₃Fe₅O₁₂), and hexagonal ferrite (BaFe₁₂O₁₉).

Developing strategies to increase magnetization is essential for improving device efficiency, temperature stability, and sensitivity, as well as for facilitating miniaturization. Recently, the intercalation of ions via liquid electrolyte gating has been demonstrated as a simple yet powerful method to enhance the magnetization of the ferrimagnetic insulator CoFe₂O₄ [8, 9, 10]. This enhancement of ≈1.3 times by ionic liquid gating [8, 9, 10]. was comparable to 1.4−2.8 times enhancement obtained by cation substitution [11, 12, 13], ≈2.2 times by strain engineering [14, 15], and ≈1.1 times by interface modulation [16]. The enhanced magnetization in ionically gated CoFe₂O₄ was quite surprising since, conversely, most materials, including La_0.67Sr_0.33MnO₃ and LaCoO₃, typically exhibit deterioration of their intrinsic physical properties with ion intercalation [17, 18]. However, various conflicting experimental findings have led to ongoing debates in this rapidly expanding field. Ionic liquid gating for oxide films includes multiple simultaneous effects: oxygen vacancy formation [19, 20], hydrogen intercalation typically arising from the water in the liquid electrolyte [21, 22], and electric‐field‐driven modulation of the carrier charge density due to the generation of an electric double layer at the film surface [23]. To unambiguously unravel the ionic effects responsible for the magnetic enhancement in CoFe₂O₄ films, it is crucial to employ a method that introduces hydrogen or oxygen exclusively. Key questions remain regarding how these individual ions affect the crystal structure and electronic band structure, and how they modulate magnetic properties within an insulating ground state.

Here, we demonstrate that hydrogen ions play a pivotal role in extremely enhancing the magnetization of CoFe₂O₄. Through Pt‐catalyst‐assisted hydrogenation of the CoFe₂O₄ films, we achieved a 2.6−4.7‐times‐enhanced room temperature ferrimagnetic moment (730−770 emu cc⁻¹) in the hydrogenated CoFe₂O₄ films, representing the highest value among ever‐reported. To elucidate the mechanism, we quantified the hydrogen ion density via time‐of‐flight secondary ion mass spectrometry (TOF‐SIMS), identified the specific hydrogen intercalation sites within the CoFe₂O₄ lattice using X‐ray magnetic circular dichroism (XMCD), and monitored the associated electron transfer via X‐ray photoemission spectroscopy (XPS). For convenience, we will use AG, AH, and AO films as the abbreviations of the as‐grown film, the film annealed in hydrogen, and the film annealed in oxygen, respectively.

2. Strong Enhancement of Magnetization in CoFe₂O₄ via Pt‐Catalyst‐Assisted Annealing in Hydrogen

We deposited (001)‐oriented CoFe₂O₄ films epitaxially on (001)‐oriented SrTiO₃ substrates by employing pulsed laser epitaxy (Figure 1a). The AG film at room temperature exhibited the magnetic field dependence of magnetization with the saturation magnetic moment M _S of ≈210 emu cc⁻¹ at 2 T, and the coercive field was ≈0.15 T (black solid line in Figure 1b). These values are consistent with those typically observed in ferrimagnetic CoFe₂O₄ epitaxial films with 285−330 emu cc⁻¹ and ≈0.25 T [8, 9, 10, 24, 25]. Furthermore, the resistance of the AG films exceeded our measurement limit (>10⁷ Ω), confirming the high‐quality ferrimagnetic insulating nature of our AG‐CoFe₂O₄ epitaxial film at room temperature.

FIGURE 1.

Reversible modulation of the ferrimagnetic signals in an insulating CoFe₂O₄ film by annealing in hydrogen. (a) We intercalate hydrogen ions into CoFe₂O₄ epitaxial films through the catalytic reaction H₂(g) → 2H⁺ + 2e⁻ occurring at the Pt/H₂/CoFe₂O₄ triple phase boundary. (b) The magnetic field dependence of the magnetization at room temperature shows that the annealed films in hydrogen (AH, red solid line) have much higher ferrimagnetic signals than the as‐grown films (AG, black solid line). The magnetic property is recovered to the AG state by annealing the films in oxygen (AO, blue dashed line). (c) The enhancement of saturation magnetization in our work is outstanding, compared with previous trials for CoFe₂O₄, including ionic liquid gating [8−10], and hydrogen reduction [24, 25]. M_s (AH) and M_s (AG) denote the saturation magnetization of AH and AG films at 2 T, respectively. (d) Our M _s‐values in AH‐CoFe₂O₄ are much higher than those of other ferrimagnetic oxides (NiFe₂O₄ [29], and CuFe₂O₄ [30].). (e) Reversibility of the hydrogen control of ferrimagnetic signals in CoFe₂O₄ over at least 10 cycles and (f) long‐term stability of the hydrogenated state for at least 10 d at room temperature in ambient air demonstrates the potential of our ionic‐magnetic coupling for ionic devices.

We distributed 10‐nm‐thick platinum dots with 100‐µm‐diameter on the film surface and annealed the CoFe₂O₄ films in hydrogen. This allowed us to utilize the ‘hydrogen spillover’, a well‐established method for effectively intercalating only hydrogen ions into oxides via the catalytic reaction of H₂(g) → 2H⁺ + 2e⁻ through the triple‐phase Pt/H₂/oxide boundary (Figure 1a) [26, 27, 28]. After we annealed the AG‐CoFe₂O₄ films in hydrogen at 300°C for 12 hours, the M _S‐value of AH‐CoFe₂O₄ increased significantly by a factor of four, reaching ≈780 emu cc⁻¹ (red solid line in Figure 1b). Compared with 0.12 ± 0.04 T for AG films, the coercive fields of AH films were effectively maintained at 0.09 ± 0.02 T, indicating a negligible formation of defects in AH films. Through systematic testing of annealing conditions, we found that the M _S‐values increased abruptly at 250−300°C for 12 h (Figure S1), and more than ≈2 h of annealing was required to reach the maximum M_S ‐value at 300°C (Figure S2). Unless otherwise stated in the paper, we annealed the films at 300°C for 12 h. We checked that the annealing SrTiO₃ substrate in hydrogen produced a negligible magnetic signal (Figure S3), indicating that the magnetic enhancement in AH films originated entirely from the hydrogen control of ferrimagnetism in CoFe₂O₄.

Notably, our hydrogen control with the assistance of Pt catalysts was the most effective strategy to enhance the ferrimagnetic moment in CoFe₂O₄. When we tested nine AH films at room temperature, we consistently achieved higher M_S ‐values of 730−770 emu cc⁻¹, representing a 2.6−4.7 times enhancement compared to the 170−260 emu cc⁻¹ observed in the AG films (Figure 1c and Figure S4 for raw data). This enhancement significantly exceeds the ≈1.2 times increase in ionic‐liquid‐gated films (navy rhombus [8], square [9], and pentagon [10]) and 1.1−2.2 times increase in the annealed films in hydrogen without Pt catalyst (magenta circles [24] and triangles [25]). Furthermore, we achieved higher M_S ‐values of 4.6−6.3 μ _B per formula unit (f.u.) in our AH‐CoFe₂O₄ films compared to conventional ferrimagnetic insulators, including 1.1 μ _B fu⁻¹ in NiFe₂O₄ [29] and 0.4 μ _B fu⁻¹ in CuFe₂O₄ (Figure 1d) [30].

Our hydrogen control of ferrimagnetism in CoFe₂O₄ films holds great potential for ionic device applications. After we annealed AH films in oxygen, the magnetic field dependence of magnetization in AO films returned to that of AG films (blue dashed line in Figure 1b). This reversibility was maintained for at least 10 cycles with a substantial modulation window of 200−780 emu cc⁻¹ (Figure 1e, Figure S5 for raw data). The AH films showed long‐term stability of magnetization when we stored the samples at room temperature in ambient air (Figure 1f, Figure S6 for raw data). After 10 d, the AH films still showed significantly larger magnetic moments than the AG films exhibited. Notably, the AH films showed negligible lattice volume expansion (Figure S7) with very flat surface (Figure S8), preventing the device delamination during the operation. Considering their insulating ground state, our tunable magnetic properties at room temperature in CoFe₂O₄ epitaxial films made them exceptionally attractive for high frequency operations, as the absence of significant eddy current minimized energy loss. These outstanding performance characteristics are expected to be beneficial for the development of ultralow power ionic‐magnetic coupled applications for sensor, energy storage, and neuromorphic devices for long‐term usage in the context of carbon neutrality [31, 32].

3. Hydrogen Intercalation in CoFe₂O₄

A significantly higher concentration of hydrogen ions was detected in the AH‐CoFe₂O₄ compared to the AG and AO films. The depth profiles of hydrogen ions measured via TOF‐SIMS showed a much higher intensity in the AH‐CoFe₂O₄ films than that in AG films (Figure 2a). The intensity in the AH film was much higher near the film surface than near the film/substrate interface, indicating a successful incorporation of hydrogen ions into the CoFe₂O₄ films from the surface. In the AO films, the hydrogen ion intensity reverted to the level observed in the AG films.

FIGURE 2.

Evidence of hydrogen intercalation in CoFe₂O₄. (a) The time‐of‐flight secondary ion mass spectrometry shows that the intensity of hydrogen ions increases in the AH films. The titanium ion signal from the SrTiO₃ (dotted line) is used as a marker to identify the interface between the CoFe₂O₄ film and the SrTiO₃ substrate. (b) The X‐ray photoemission spectroscopy (XPS) spectra near O 1s exhibiting higher intensity of the O−H peak (pink area). In both (a) and (b), the intensities for the AO films revert toward those of the AG films.

To obtain more insights on reversible intercalation of hydrogen ions into CoFe₂O₄ films, we measured the XPS O 1s spectra in 525−535 eV binding energy and deconvoluted them into O−C (532.8 eV), O−H (531.8 eV), and Co/Fe−O (529.9 eV) peaks. (Figure 2b). The spectral intensity of the O−H peak was significantly higher in the AH‐CoFe₂O₄ films than that in AG films. This observation was quite similar to the observation in hydrogenated La_0.67Sr_0.33MnO₃ [17] and VO₂ [19]. This result is consistent with the Grotthuss mechanism when hydrogen ions are absorbed and hop through forming −OH bonds in solids [33]. The O−H XPS intensity of the AO films was reduced toward a level of the AG films.

Our comparative experiments revealed that hydrogen‐driven control was crucial in the strong enhancement of magnetization in AH‐CoFe₂O₄. When we annealed the AG‐CoFe₂O₄ films in vacuum, we did not observe a significant change in magnetic properties (Figure S9), indicating that hydrogen should be intercalated into CoFe₂O₄ for the observed magnetic enhancement. Moreover, the magnetic properties of AG‐CoFe₂O₄ films did not exhibit any changes when the films were annealed in hydrogen without platinum dots (Figure S10). This indicates that there was no catalytic H₂(g) → 2H⁺ + 2e⁻ reaction (Figure 1a) since the triple‐phase Pt/H₂/CoFe₂O₄ boundary was absent [26, 27, 28].

This enhancement of magnetic properties in hydrogenated CoFe₂O₄ is highly surprising, as most materials showed deterioration of physical properties with hydrogen intercalation [17, 18, 19, 21, 28]. For instance, the ferromagnetic properties were weakened in hydrogenated La_0.67Sr_0.33MnO₃ since the double exchange interaction was suppressed by a lattice‐expansion‐induced and an electron‐doping‐induced modulation [17]. The smaller magnetization with hydrogen intercalation was also observed in tensile‐strained LaCoO₃ films, where the subsequent filling of correlated electrons in the Co 3d‐orbitals destabilizes the CoO₆ octahedra, triggering structural transition into CoO₄ tetrahedra and the successive release of oxygen [18]. Notably, the robust positive hydrogenic‐magnetic coupling in CoFe₂O₄ indicates that its mechanism is fundamentally different from previous observations. Using XMCD and XPS, we will elucidate how the intercalated hydrogen ions and accompanying electrons affect the magnetic enhancement of CoFe₂O₄.

4. Mechanism for Colossal Magnetization in Hydrogenated CoFe₂O₄

To understand the strong enhancement of magnetization in AH‐CoFe₂O₄, we first consider the ferrimagnetic origin of AG‐CoFe₂O₄. The CoFe₂O₄, ${\mathrm{Co}}_{\mathrm{Oct}}^{2+}[{\mathrm{Fe}}_{\mathrm{Oct}}^{3+}{\mathrm{Fe}}_{\mathrm{Tet}}^{3+}]{\mathrm{O}}_4$, has an inverse spinel structure, which all divalent Co²⁺ ions occupy octahedral sites (${\mathrm{Co}}_{\mathrm{Oct}}^{2+}$), while the trivalent Fe³⁺ ions are equally distributed in the octahedral sites (${\mathrm{Fe}}_{\mathrm{Oct}}^{3+}$) and the tetrahedral sites (${\mathrm{Fe}}_{\mathrm{Tet}}^{3+}$) (Figure 1a) [1, 2]. The critical factor for ferrimagnetism in CoFe₂O₄ is the superexchange interaction of the spin moments between adjacent Fe³⁺ ions, mediated by a non‐magnetic oxygen anion, located between the magnetic Fe³⁺ ions. The spin moments of all ${\mathrm{Fe}}_{\mathrm{Oct}}^{3+}$ ions are aligned parallel to one another but are oppositely directed to the ${\mathrm{Fe}}_{\mathrm{Tet}}^{3+}$ ions. Thus, there is a cancellation of the spin moments of all Fe³⁺ ions and they do not contribute to the net magnetization of AG‐CoFe₂O₄. Consequently, net magnetization is solely derived from the ${\mathrm{Co}}_{\mathrm{Oct}}^{2+}$ ions, which their moments are aligned in the same direction.

We found that the hydrogen ions were predominantly intercalated near the ${\mathrm{Fe}}_{\mathrm{Tet}}^{3+}$ sites. The XMCD spectra (= I _right − I _left), calculated by measuring the X‐ray absorption spectra with the primary beam of right (I _right) and left (I _left) circularly polarized photons, are one of the powerful and widely‐used methods for investigating cation site occupancy in CoFe₂O₄ [8, 34]. Figure 3a,b shows the XMCD spectra of Fe‐L edges and Co‐L edges in AG (black thin line), AH (red thick line), and AO films (blue squares), respectively (Figure S11 for the corresponding X‐ray absorption spectra). By comparing the experimental XMCD spectra with theoretical spectra of ${\mathrm{Fe}}_{\mathrm{Tet}}^{3+}$ (orange line, strong peak at the photon energies of 708.7 eV), ${\mathrm{Fe}}_{\mathrm{Oct}}^{3+}$ (purple line, 709.5 eV), and ${\mathrm{Fe}}_{\mathrm{Oct}}^{2+}$ (magenta line, 707.6 eV), we found that the XMCD signal of Fe‐L edge in AG‐CoFe₂O₄ was mainly composed of ${\mathrm{Fe}}_{\mathrm{Tet}}^{3+}$, ${\mathrm{Fe}}_{\mathrm{Oct}}^{3+}$, and ${\mathrm{Fe}}_{\mathrm{Oct}}^{2+}$. The comparison experimental XMCD spectra of Co‐L edges with theoretical spectra of ${\mathrm{Co}}_{\mathrm{Tet}}^{2+}$ (green line, strong peak at 778.4 eV), ${\mathrm{Co}}_{\mathrm{Oct}}^{2+}$ (navy line, 778.4 eV), and ${\mathrm{Co}}_{\mathrm{Oct}}^{3+}$ (cyan line, 779.0 eV) indicated that the XMCD signal of Co‐L edge in AG‐CoFe₂O₄ was mainly composed of ${\mathrm{Co}}_{\mathrm{Oct}}^{2+}$. The XMCD signals of Fe‐L and Co‐L edges in AG‐CoFe₂O₄ were quite similar to previous reports [8, 34]. Notably, upon hydrogenation, the XMCD signals showed different behaviors for tetrahedral and octahedral sites. The ${\mathrm{Fe}}_{\mathrm{Tet}}^{3+}$ peak in AH films was completely suppressed, suggesting that hydrogen ions were intercalated into tetrahedral sites. This interpretation aligns with previous reports showing that hydrogen ions preferentially occupy the tetrahedral sites of oxides, in which the hydrogen ions hop by generating −OH bonds [33, 35, 36]. For brownmillerite SrCoO_2.5 and BaInO_2.5, whose crystal structure has unique stacks of oxygen octahedra and tetrahedra alternatively along the [001] direction, the equatorial oxygen sites at the tetrahedral layers were energetically more favorable (by ≈0.37 eV in SrCoO_2.5) than equatorial and interlayer oxygens at the octahedral layers [33, 36]. For the AO films, the XMCD signals were recovered to the level of AG films, consistent with the reversible hydrogen control of ferrimagnetism (Figure 1b) in CoFe₂O₄.

FIGURE 3.

Hydrogen absorption near Fe³⁺ ions in the tetrahedral sites (${\mathrm{Fe}}_{\mathrm{Tet}}^{3+}$) in CoFe₂O₄, evidenced by X‐ray magnetic circular dichroism (XMCD). (a) The ${\mathrm{Fe}}_{\mathrm{Tet}}^{3+}$ peak is suppressed in AH films (red thick line) and is recovered in AO films (blue square) to the level of AG films (black thin line), suggesting that hydrogen ions were intercalated near ${\mathrm{Fe}}_{\mathrm{Tet}}^{3+}$ ions. (b) The ${\mathrm{Co}}_{\mathrm{Oct}}^{3+}$ peak appears in AH films, indicating a change of valence states of ${\mathrm{Co}}_{\mathrm{Oct}}^{2+}$ → ${\mathrm{Co}}_{\mathrm{Oct}}^{3+}$, suggesting electron transfer between these ions. Thin lines represent theoretical XMCD spectra of Fe and Co ions [34].

From the XMCD peak suppression at 776.6 eV and evolution at 793.5 eV of the Co‐L edge (Figure 3), we also observed a change of valence states of ${\mathrm{Co}}_{\mathrm{Oct}}^{2+}$ → ${\mathrm{Co}}_{\mathrm{Oct}}^{3+}$, suggesting electron transfer between these ions. To obtain more insights into the valence change of not only ${\mathrm{Co}}_{\mathrm{Oct}}^{2+}$ but also ${\mathrm{Fe}}_{\mathrm{Tet}}^{3+}$ and ${\mathrm{Fe}}_{\mathrm{Oct}}^{3+}$ by the intercalation of hydrogen ions near the ${\mathrm{Fe}}_{\mathrm{Tet}}^{3+}$, we measured the XPS spectra near Fe 2p_3/2 and Fe 2p_1/2 peaks within a binding energy range of 700−740 eV (Figure 4a) and near the Co 2p_3/2 and Co 2p_1/2 peaks in 770−810 eV binding energy (Figure 4b). We fitted the XPS spectra of Fe‐L edge by ${\mathrm{Fe}}_{\mathrm{Oct}}^{3+}$ 2p_3/2 at 710.3eV, ${\mathrm{Fe}}_{\mathrm{Oct}}^{3+}$ 2p_1/2 at 723.5eV, ${\mathrm{Fe}}_{\mathrm{Tet}}^{3+}$ 2p_3/2 at 712.4 eV, and ${\mathrm{Fe}}_{\mathrm{Tet}}^{3+}$ 2p_1/2 at 725.8 eV (Figure 4a) [29, 33]. We also fitted the XPS spectra of Co‐L edge by ${\mathrm{Co}}_{\mathrm{Oct}}^{2+}$ 2p_3/2 at 780.1 eV, ${\mathrm{Co}}_{\mathrm{Oct}}^{2+}$ 2p_1/2 at 795.5 eV, ${\mathrm{Co}}_{\mathrm{Tet}}^{2+}$ 2p_3/2 at 782.5 eV, and ${\mathrm{Co}}_{\mathrm{Tet}}^{2+}$ 2p_1/2 at 797.4 eV (Figure 4b) [37, 38, 39, 40]. By calculating the relative proportions of the ${\mathrm{Fe}}_{\mathrm{Oct}}^{3+}$ and ${\mathrm{Fe}}_{\mathrm{Tet}}^{3+}$ in the Fe‐2p_3/2 signal and the ${\mathrm{Co}}_{\mathrm{Oct}}^{2+}$ and ${\mathrm{Co}}_{\mathrm{Tet}}^{2+}$ in the Co‐2p_3/2 signal, we experimentally confirmed the similar occupation of the ${\mathrm{Fe}}_{\mathrm{Oct}}^{3+}$, ${\mathrm{Fe}}_{\mathrm{Tet}}^{3+}$, ${\mathrm{Co}}_{\mathrm{Oct}}^{2+}$, and ${\mathrm{Co}}_{\mathrm{Tet}}^{2+}$ in the AG film, indicating $[{({\mathrm{Co}}_{\mathrm{Oct}}^{2+})}_{0.8}{({\mathrm{Co}}_{\mathrm{Tet}}^{2+})}_{0.2}][{\mathrm{Fe}}_{\mathrm{Oct}}^{3+}{\mathrm{Fe}}_{\mathrm{Tet}}^{3+}]{\mathrm{O}}_4$. Although a few ${\mathrm{Co}}_{\mathrm{Tet}}^{2+}$ ions could exist in the crystal lattice of the CoFe₂O₄ films despite their inverse spinel structure [8, 41, 42], the XPS spectra and peak assignments in AG films were quite consistent with previous reports [37, 38]. In the AH film, the ${\mathrm{Fe}}_{\mathrm{Tet}}^{3+}$ signal was suppressed concomitantly with the appearance of the ${\mathrm{Fe}}_{\mathrm{Tet}}^{2+}$ while the ${\mathrm{Fe}}_{\mathrm{Oct}}^{3+}$ signal did not show a significant change. Notably, we observed the suppression of the ${\mathrm{Co}}_{\mathrm{Oct}}^{2+}$ and the significant appearance of ${\mathrm{Co}}_{\mathrm{Oct}}^{3+}$. Therefore, the XPS analysis conclusively revealed that the intercalation of hydrogen ions near ${\mathrm{Fe}}_{\mathrm{Tet}}^{3+}$ sites was accompanied by valence exchange of ${\mathrm{Fe}}_{\mathrm{Tet}}^{3+}$ → ${\mathrm{Fe}}_{\mathrm{Tet}}^{2+}$ and ${\mathrm{Co}}_{\mathrm{Oct}}^{2+}$ → ${\mathrm{Co}}_{\mathrm{Oct}}^{3+}$. In AO films, three peaks were recovered to the level of AG films, consistent with the reversible transition of ferrimagnetism (Figure 1b) and XMCD spectra (Figure 3) in CoFe₂O₄.

FIGURE 4.

Electron transfer between ${\mathrm{Fe}}_{\mathrm{Tet}}^{3+}$ and ${\mathrm{Co}}_{\mathrm{Oct}}^{2+}$, evidenced by X‐ray photoemission spectra near (a) Fe 2p_3/2 and Fe 2p_1/2 and (b) Co 2p_3/2 and Co 2p_1/2. In AH film, we observe the suppression of ${\mathrm{Fe}}_{\mathrm{Tet}}^{3+}$ without a significant change of ${\mathrm{Fe}}_{\mathrm{Oct}}^{3+}$. We observe the suppression of the ${\mathrm{Co}}_{\mathrm{Oct}}^{2+}$ and the significant appearance of ${\mathrm{Co}}_{\mathrm{Oct}}^{3+}$.

We propose a possible mechanism for the strong enhancement of magnetization in hydrogenated CoFe₂O₄ by discussing how site‐specific hydrogen intercalation and/or electron doping can affect the ferrimagnetism, as shown in Figure 5. The hydrogen ions, which were intercalated preferentially into tetrahedral sites in CoFe₂O₄, disrupted the superexchange interaction between ${\mathrm{Fe}}_{\mathrm{Tet}}^{3+}$ and ${\mathrm{Fe}}_{\mathrm{Oct}}^{3+}$. Consequently, the spin moments of Fe³⁺ ions no longer completely cancel each other out, thereby contributing to the enhancement of the total moment. In addition, the charge transfer from ${\mathrm{Co}}_{\mathrm{Oct}}^{2+}$ to ${\mathrm{Fe}}_{\mathrm{Tet}}^{3+}$ increased the magnetic moment of ${\mathrm{Co}}_{\mathrm{Oct}}^{2+}$ from S = 3/2 to S = 2. Therefore, by comprehensively considering two effects, net magnetization increases by a factor of 2.6−4.7.

FIGURE 5.

Proposed mechanism for the robust enhancement of room temperature ferrimagnetism in insulating CoFe₂O₄, based on site‐specific hydrogen intercalation and/or electron transfer. Preferential intercalation of hydrogen ions near ${\mathrm{Fe}}_{\mathrm{Tet}}^{3+}$ in CoFe₂O₄ disturbs the cancellation of the spin moments of ${\mathrm{Fe}}_{\mathrm{Tet}}^{3+}$ and ${\mathrm{Fe}}_{\mathrm{Oct}}^{3+}$. Furthermore, the charge transfer from ${\mathrm{Co}}_{\mathrm{Oct}}^{2+}$ to ${\mathrm{Fe}}_{\mathrm{Tet}}^{3+}$ increases the magnetic moment of Co ions from S = 3/2 to S = 2. Both effects can comprehensively contribute to the 2.6−4.7 times enhancement of net magnetization.

5. Conclusions

We found that the hydrogenated CoFe₂O₄ films exhibited a remarkably high ferrimagnetic moment (730−770 emu cc⁻¹), representing 2.6−4.7 times increase compared to that of 170−260 emu cc⁻¹ measured in as‐grown films. Notably, these total magnetic moments are the highest values reported to date. This enhancement is highly significant, as it contrasts sharply with the deterioration of physical properties typically observed upon hydrogenation in most transition metal oxides. We propose that the hydrogen ions specifically intercalate near the Fe³⁺ ions in the tetrahedral sites of CoFe₂O₄. While the spin moments of Fe³⁺ in tetrahedral and octahedral sites cancelled one another in as‐grown films, X‐ray spectroscopy for hydrogenated films revealed that both Fe³⁺ and Co²⁺ ions contribute to the enhancement of the total magnetic moment. We expect that this proposed mechanism can also work if there is the hydrogen‐induced magnetic enhancement in other spinel ferrites (MFe₂O₄, M = divalent cation such as Zn²⁺, Fe²⁺, Ni²⁺, Cu²⁺, and Mn²⁺) [43]. Excellent reproducibility and long‐term stability for ≈10 d, achieved while maintaining an isostructural spinel transition, suggest that hydrogenic‐ferrimagnetic coupling in CoFe₂O₄ holds great promise for ultralow power spintronic devices as well as high frequency electronics, including antennas, transformers, and microwave devices.

6. Experimental Section

6.1. Epitaxial Growth of CoFe₂O₄ Films

30‐nm‐thick CoFe₂O₄ films were epitaxially grown on (001)‐oriented SrTiO₃ via pulsed laser deposition. To evaporate CoFe₂O₄, we focused an ultraviolet laser beam (a wavelength of 266 nm) onto the target pellet using an Nd:YAG (model Q‐SMART 850; Quantel laser). The deposition was carried out with a laser intensity of 1.5 J cm⁻², and a repetition frequency of 10 Hz, while keeping a target‐to‐substrate distance of 6 cm. During deposition, a lamp heater maintained the substrate at 600°C. Concurrently, the chamber was filled with oxygen, and the oxygen partial pressure was stabilized at 75 mTorr with a constant flow of 20 sccm, controlled by a mass flow controller.

6.2. Intercalation and Deintercalation of Hydrogen Ions

Platinum electrodes were deposited onto the surface of CoFe₂O₄ films via DC magnetron sputtering. To generate plasma, we applied 30 W power in the Ar environment with 2 mTorr partial pressure. We patterned the platinum dots (10‐nm‐thickness, 100‐µm‐diameter, and 65 µm‐spacing between dots) using a metal mask. The intercalation of hydrogen ions was carried out by annealing the films at 300°C for 12 h in a gas mixture of H₂(5%) + Ar(95%). For desorbing hydrogen ions, the hydrogenated films were subsequently annealed at 300°C for 12 h in pure oxygen gas (O₂, 99.9999%). During both annealing processes, the chamber was maintained at a constant pressure of ≈400 Torr.

6.3. Characterization of Physical Properties

The magnetization of CoFe₂O₄ films was evaluated by employing a magnetic property measurement system (model MPMS3, Quantum Design, Inc.). All measurements were performed at room temperature to suppress the thermal desorption of intercalated hydrogen during the measurement. To further prevent unintended hydrogen release, the films containing hydrogen ions were consistently stored at −18°C in a freezer when not used for analysis.

6.4. Detection of Hydrogen Ions

The time‐of‐flight secondary ion mass spectrometry (model TOF.SIMS 5–100; ionTOF) was employed to confirm the successful incorporation of hydrogen ions. To ensure the precise detection of hydrogen and to minimize the noise from other gaseous elements, the chamber was maintained at a base pressure of ≈2 × 10⁻⁹ Torr. Depth profiling was performed using a dual‐beam configuration: a 0.5 keV cesium ion beam was utilized to sputter a 100 × 100 µm² area, while a 30 keV bismuth ion beam was used to analyze the hydrogen ion signal over a region of 50 × 50 µm². Concurrently, the titanium ion signal was monitored to accurately identify the exact interface between the CoFe₂O₄ film and SrTiO₃ substrate.

6.5. Investigation of Structural Properties

The structural integrity and crystalline quality of the films were evaluated by employing a four‐circle high‐resolution X‐ray diffractometer (model Empyrean; PANalytical). A Cu Kα radiation source with a wavelength of 1.54 Å was used. For the θ−2θ scans and reciprocal space mapping, 1/4‐ and 1/16‐ inch divergence slits were employed to optimize the beam resolution, respectively. The morphology and surface roughness of CoFe₂O₄ films were evaluated by employing atomic force microscopy (model XE7, Park Systems). Topography mapping was performed in tapping mode over a randomly chosen 5 × 5 µm² areas of the samples with a scan rate of 0.5 Hz.

6.6. Establishment of Electronic Band Structure

To elucidate the electronic band structure, X‐ray absorption spectroscopy in total electron yield mode was performed at the 2A Beamline of the Pohang Accelerator Laboratory. This laboratory was funded in part by the Ministry of Science and ICT and POSTECH. Considering the highly insulating nature of the CoFe₂O₄ films, the measurements were intentionally performed at room temperature to mitigate charging effects. Furthermore, the chemical oxidation states of the Co and Fe ions were determined using an X‐ray photoemission spectrometer (model ESCALAB 250Xi; Thermo Scientific). The primary beam was generated from a monochromatic Al Kα source. The chamber was maintained at a base pressure of 10⁻⁸ Torr, with the reference binding energy anchored to the C‐1s peak at 284.5 eV.

Author Contributions

J.P. and Y.H. performed the experiments and finalized the manuscript. S.L. supervised the experiment and writing. All authors reviewed the paper.

Conflicts of Interest

The authors declare no conflicts of interest.

Supporting information

Supporting File: adma73974‐sup‐0001‐SuppMat.docx.

Data Availability Statement

The data supporting the findings of this study are available from the corresponding author upon reasonable request.