Large and Robust Exchange Bias in Co/CoO film: Implication for Flexible Spintronics

Abstract

Developing a robust and flexible exchange bias (EB) system is crucial for wearable spintronic devices. In this work, large and robust EB effect is demonstrated in flexible Co/CoO thin films grown by magnetron sputtering. The obtained EB field is as large as 6172 Oe at low temperatures on a flexible substrate. Based on the atomistic spin models, the unexpectedly large EB field is explained by the decrease of ferromagnetic grain size and the increase of ferromagnetic /antiferromagnetic contact area. Furthermore, the strain‐dependent coercivity and EB effect are observed in flexible Co/CoO films. Finally, the flexible Co/CoO thin film shows no degradation (even a slight increase) of performances in all figures of merit after 500 bending cycles, indicating excellent mechanical durability. This study sheds light on Co/CoO as a robust EB material for flexible spintronics.

Introduction

Ferromagnetic (FM)/antiferromagnetic (AFM) film system is an essential structure in spintronic devices. The interface exchange coupling of FM/AFM, known as the exchange bias (EB) effect,^[1] plays a key role in practical applications, such as fixing moment of spin valve (or magnetic tunnel junction),^[2] providing bias fields for spin-orbit torque.^[3] Rigid FM/AFM films are typically favorited in spintronic devices. However, with the growing interest in flexible spintronic devices, rigid FM/AFM film systems face challenges due to the difficulty in meeting flexibility and mechanical durability.^[4–9] Developing flexible and robust FM/AFM film systems has become an important topic in the spintronic community. For the fabrication of flexible films, “film transferring” and “direct deposition on flexible substrate” have been developed in recent years. Some flexible FM/AFM films were successfully realized via these approaches, such as NiMn/Fe₄N, BiFeO₃/La_0.7Sr_0.3MnO₃, Co/CoO, etc.^[10–16] The reported FM/AFM films not only exhibit decent mechanical durability but also have tunable magnetic properties. For instance, large strain can be induced in flexible FM/AFM film through large curvature bending,^[12,15] whereas the induced strain in the rigid counterparts is usually small. In this case, exchange bias effect and other magnetic properties can be tailored by the strain. These flexible features are very appealing for both fundamental research and technology applications.

On the other hand, strong exchange bias coupling is also of importance in terms of application, which is related to the stability and energy efficiency of spintronic devices.^[17] To date, the reported exchange bias fields on the flexible films are typically small, in the range of a few hundred Oersted.^{[10,13,15,18]} Continuous efforts are being devoted to achieve strong exchange bias coupling in flexible FM/AFM film systems. Among the various candidate materials, CoO has a simple crystal and magnetic structure with a transition temperature close to room temperature.^[19–21] As such, CoO is ideal AFM material for studying exchange bias.^[13,22–26] In the CoO-based FM/AFM films, enhanced exchange bias coupling was recently realized via employing a porous structure or strong magnetocrystalline anisotropy FM layer.^[11,27] Thus, the strong exchange bias and strain-tuning effect in polycrystalline FM/AFM systems are critical for flexible spintronics.

In this study, we have prepared flexible Co/CoO thin films by magnetron sputtering, in which polyethylene naphthalate two formic acid glycol ester (PEN) has been used as the flexible substrate. The magnetic properties and flexibility of Co/CoO thin films from 2 to 300 K were investigated. Interestingly, our sample shows a large exchange bias, which is up to 6172 Oe at 2 K. The unexpected strong exchange bias effect is attributed to increase the FM/AFM contact area. In addition, the mechanical strain is demonstrated to effectively tune the exchange bias in polycrystalline Co/CoO thin films. Our results demonstrate that the classical Co/CoO polycrystalline thin films show an unexpected large exchange bias field with excellent flexibility, which paves the way for delicate design of future spintronic devices.

Results and Discussion

First, we study the microstructure and morphology of Co/CoO thin films using transmission electron microscopy (TEM) characterization. Energy dispersive spectroscopy (EDS) elemental mapping is also employed to investigate atomic and chemical configuration of the film. Figure 1 shows the typical cross-section and morphology of a Co/CoO thin film prepared on the Si substrate. A clear interface between the film and substrate is observed. The thickness of the thin film is estimated ≈10 nm (Figure 1a). The EDS mapping results indicate that the film contains Co and O elements. which indicates that neither Co nor O distribution is uniform, as shown in Figure 1b,c. The Co content close to the bottom is significantly higher than that near the surface. In contrast, the O element is mainly distributed in the surface region with an average depth of ≈4 nm, suggesting the existence of cobalt oxide. It is worth noting that the O penetrates deep into the surface of Si substrate in some regions, where Co also shows a reduced distribution. The O distribution at the bottom is attributed to native oxide of silicon substrate, as shown in Figure S1 (Supporting Information). As the oxidation was performed at room temperature, the formed cobalt oxide is assumed to be CoO because the formation of Co₃O₄ requires a heat treatment in an oxygen ambience.^[28,29] Figure 1d–f displays the atomic-resolution high-angle annular dark-field (HAADF) images of the film. The interfaces between the Co layer and CoO layer are highly jagged. It is thus difficult to define the exact thickness of Co layer and CoO layer. Clearly, both Co and CoO are polycrystalline. The grain size of Co varies from grain to grain, with an average lateral dimension of 6 nm. The corresponding fast Fourier transform (FFT) images indicate that the Co crystalline is a hexagonal close packing structure, while CoO is the cubic structure. The TEM results demonstrate the successful growth of Co/CoO thin films with a large FM/AFM contact area.

[IMAGE OMITTED. SEE PDF]

Using a superconducting quantum interference device (SQUID) magnetometer, the magnetic hysteresis loops of Co/CoO films were measured after +1.5 T field cooling (FC) down to 2 K. We studied the magnetic properties of Co/CoO films on both Si and PEN substrates, as shown in Figure S2 (Supporting Information). The magnetic hysteresis loops of Co/CoO films on both substrates are almost the same at different temperatures, suggesting that the magnetic properties of Co/CoO film are not sensitive to substrate material. Therefore, we tentatively assume that the two substrates result in no significant difference in microstructure and morphology of Co/CoO thin films. Figure 2 shows the magnetic properties of a representative Co/CoO film on the PEN substrate (sample A). The magnetic hysteresis loop of in-plane and out-of-plane directions at 300 K are plotted in Figure 2a. For the in-plane measurement geometry, the loop exhibits symmetrical rectangle and large remanence. In the case of out-of-plane, the remanence is approximately zero and the saturation field is much larger than that of in-plane geometry. This result reveals that the Co/CoO film is a ferromagnet with an in-plane easy axis. In this study, we mainly focus on in-plane magnetic properties. Clearly, at low temperatures, the center of hysteresis loops is shifted toward the negative magnetic field after FC (Figure 2b). In addition to the very large negative EB effect, the coercivity of flexible Co/CoO film is also greatly enhanced. The value of EB field H_EB is even comparable with the coercivity field H_C at low temperatures (e.g., H_EB = –4081 Oe, H_C = –4680 Oe, at 5 K), where the H_C and H_EB are determined by H_C = (H_C+ – H_C−)/2 and H_EB = (H_C− + H_C+)/2 respectively. H_C- (H_C+) denotes the coercivity field of the descending (ascending) branch. Another feature is a small kink near zero fields in the descending branch, suggesting a two-stage magnetization reversal. At higher temperatures, both H_C and H_EB become smaller (Figure 2b, 90 K). When temperature increases to 200 K, the magnetic hysteresis loop is symmetric shape and H_C also decreases significantly.

[IMAGE OMITTED. SEE PDF]

Figure 2c displays the absolute value of H_C versus temperature, which is calculated using H_C± of the magnetic hysteresis loops at different temperatures. The temperature dependence of H_C± is plotted in Figure S3 (Supporting Information). The value of H_C+ decreases monotonically with increasing temperature, while the H_C- increases first and then decreases. At some temperature points between 150 and 200 K, the right and left coercive fields are equal, indicating that the EB effect vanishes. The temperature at which the EB effect disappears is called as the blocking temperature T_B. It is well known that one of the characteristics associated with EB effect is the enhancement of coercivity.^[30] Indeed, at higher temperatures, the H_C varies slightly with temperature. When the Co/CoO system is cooled below the blocking temperature, the H_C largely increases, and the linear temperature dependence of H_C is consistent with other Co/CoO films.^[13,31]

To better understand the EB of our Co/CoO thin film, we summarize the dependence of H _EB with temperatures from 2 to 300 K in Figure 2c. The exact T_B is calculated using the relationship H_EB (T) = H_EB0 (1−T/T_B)ⁿ.^[32,33] The data fitting yields H _EB0 = 4485 Oe, T_B = 195 K, n = 3.27 for sample A. Therefore, the blocking temperature of sample A is determined to be 195 K, which is lower than the Néel temperature T_N of bulk CoO (≈293 K). Studies on CoO suggest that the T_N of CoO is related to its thickness.^[30,34,35] When the thickness of AFM CoO is less than 20 nm, the T_N decreases with thickness. In our case, a lower T_B is reasonable due to the small size of CoO. Linear (or exponent) with n≈1 dependence and T_B<2/3 T_N have been reported in different CoO-based exchange bias systems.^{[25,32,33,36]} The results can be explained by the cubic AFM anisotropy or the thermal instabilities of AFM grains.^[32,33] In addition, other mechanisms can also result in a larger exponent, such as coupling FM layer to the independent AFM grains, FM cluster with wide size distribution.^[25,36] For our flexible Co/CoO thin films, the Co/CoO layer is actually composed of numerous Co-CoO granules (Figure 1). These Co/CoO granules show distributions in shape and size. Furthermore, the granules contain either single grain or multiple grains, which thus have different anisotropy energies. The anisotropy energy of Co/CoO granules should be sufficiently large compared to the thermal energy to stabilize the magnetization and contribute to the exchange bias field. As the temperature decreases, more Co/CoO granules are thermally stable and contribute to the exchange bias field. Therefore, the H _EB rapidly increases with the decrease in temperature. The kink in loops is also in line with this mechanism. First, the small Co grains reverse at small external fields due to the weak anisotropy, resulting in a kink near zero field, as shown in Figure 2b. Next, the large Co granules reverse at large fields. The small steps also indicate that the magnetic moment of small Co granules accounts for a small proportion of the total magnetic moment of the Co/CoO thin film.

Significantly, our flexible Co/CoO thin films exhibit a large exchange bias field. As shown in Figure 2c, the maximum value of H_EB is up to approximately 6172 Oe at 2 K for Sample B (Sample A and Sample B are different patches on the same flexible substrates). Disregarding the highly jagged interface between Co and CoO layer, we consider a Co/CoO film with a clear and flat interface, and roughly estimate the interfacial exchange energy per unit area ∆E. According to the Meiklejohn–Bean model, the ∆E is expressed as ∆E = µ₀H_EB × M_s × t_Co, where M_s and t_Co are the saturation magnetization and thickness of Co layer, respectively. From the TEM images, the average size of Co is t_Co ≈ 6 nm. For our flexible Co/CoO films, the saturation magnetization is determined to be M_s ≈209.7 emu g⁻¹ (1.4 × 10⁶ A m⁻¹, 2 K). Thus the ∆E ≈ 5.35 mJ m⁻² (5.35 erg cm⁻², 2 K). Note that the calculated ∆E value is overestimated due to the simplified FM/AFM coupling area. When replacing the flat interface with hemispherical interface, this value is modified to be ∆E ≈1.70 mJ m⁻². This value is comparable with the ∆E for Co/CoO film system in previous reports.^[30] In Figure 2d, we summarize the reported exchange bias fields on different flexible substrates. Although the exchange bias fields are obtained at different temperatures, our flexible Co/CoO film shows relatively large values, which suggests that Co/CoO thin film is a promising candidate for flexible spintronic devices.

Various Co/CoO films typically show relatively small exchange bias field (<3000 Oe),^{[13,32,37–45]} while our samples show unusually large H_EB. To better understand the enhanced EB, we performed a simulation with atomistic spin models subjected to Landau–Lifshitz–Gilbert–Brown equation: 1 $$$\begin{equation}\frac{{\partial {{\bf{m}}}_{\rm{i}}}}{{\partial t}} = - {\gamma }_0{{\bf{m}}}_{\rm{i}} \times {\bf{B}}_{\rm{i}}^{{\rm{eff}}} + \alpha {{\bf{m}}}_{\rm{i}} \times \frac{{\partial {{\bf{m}}}_{\rm{i}}}}{{\partial t}}\end{equation}$$$in which m_i is the local magnetic moment of the i^th site, γ₀ = 1.76 × 10⁻⁴ (fs⁻¹ T⁻¹) is the gyromagnetic ratio, α is the Gilbert damping factor which is set to 0.1 here. B_i^eff is the effective magnetic fields acting on the ith magnetic moment, and B_i^eff = B_i^xc + B_i^ani + B_i^therm, including the exchange field B_i^xc = Σ_jJ_ijm_j, the magnetic anisotropy field B_i^ani, and stochastic field B_i^therm for mimicking the thermal fluctuations. The exchange coupling constant J_ij is set to −1 between ferromagnetic sites, representing the magnetic coupling in Co. J_ij is set to +1 between antiferromagnetic sites, representing the magnetic coupling in CoO.

Three types of magnetic structures are considered as shown in Figure 3a. The vertexes stand for the local magnetic moments in Co/CoO, and the edges denote the magnetic couplings between them. The regions filled with the red lines are the Co region, in which the couplings J_ij are ferromagnetic. The regions filled with the blue lines are the CoO region, in which the couplings J_ij are antiferromagnetic. The couplings between the Co regions and CoO regions are assumed to be ferromagnetic. The area of the Co regions of the three types are set to be the same in Figure 3a, while the lengths of boundaries are in the order of Type 1 > Type 2 > Type 3 due to the differences in geometry shapes. Figure 3b shows Calculated M-H loops for the three types at 2 K. The obtained exchange bias fields are 4429, 4165, and 3578 Oe, with the order Type 1 > Type 2 > Type 3, confirming that the exchange bias field is directly related to the size of nanoparticles. Furthermore, Figure 3c shows the calculated temperature-dependent M-H loops of the three types. The exchange bias field are 3540 and 3013 Oe at 5 and 50 K, respectively, which agree well with our experimental results. The shrinking of EB from 5 to 50 K comes from the thermal suppressing of interfacial coupling between Co and CoO, and the vanishing of EB at 500 K is due to the disorder of magnetism in CoO region. Our Co/CoO thin film has small grain size and large FM/AFM contact area, as shown in Figure 1. Therefore, the calculation results well explain the unexpected large EB field of our flexible Co/CoO thin film.

[IMAGE OMITTED. SEE PDF]

In addition to the magnetic properties of unstrained Co/CoO films, we also investigate the strain-tuning of exchange bias effect. To apply mechanical strain, the Co/CoO thin film was bent with a certain curvature. The applied strain is estimated by simplified equation, i.e., ζ∼ t/(2r),^[46,47] where ζ is the value of strain, t is the sample thickness, and r is the bending radius. In this study, the applied tensile strain (bending out) is estimated to be 3.1% for a 4-mm radius. The temperature dependence of H_EB (or H_C) can provide more insight into the strain-tuning effect, as shown in Figure 4. Figure 4a displays the coercivity after bending. The inset of Fig. 4a compares the hysteresis loops at 30 K with different bent states (see Figure S4, Supporting Information for more magnetic hysteresis loops). The flexible Co/CoO film show evident strain-tuning effect. Below 200 K, the coercivity of bending state exhibit significantly smaller values compared to the unbent state. The strain-tuning effect of H_EB is not as pronounced as that of H_C. For the bending state, the temperature dependence of H_EB maintains a similar evolution to the unbent state, except that the H_EB is slightly smaller at certain temperatures. Such a strain-tuning behavior is similar to the previous study on epitaxial Co/CoO/mica films, in which the H_EB and H_C showed a clear change with the applied strains along the CoO (100) direction.^[13] This strain-tuning effect of H_EB is believed to result from the change of AFM anisotropy, which is induced by the symmetry broken of Co²⁺ orbitals and variation of orbital energy of Co in CoO (100). Meanwhile, the change of AFM also influences the anisotropy of FM layer and leads to the change of H_C. Although Co and CoO are polycrystalline in our samples, the same mechanism should also play an important role. The exact mechanism, however, requires more detailed studies.

[IMAGE OMITTED. SEE PDF]

Finally, the magnetic properties and mechanical durability of flexible films are evaluated before and after 500 bending cycles. Figure 5 shows the representative magnetic hysteresis loops and summarizes H_C (or H_EB). In Figure 5a, after bending, the H_C- even increases slightly at 5 K, while the H_C+ shows a very small decrease. Meanwhile, the kink is more obviously after bending. Both H_C- and H_C+ increases slightly after bending at 200 K. The temperature dependence of coercivity H_C after bending is same as that before bending, except that H_C slightly increases at low-temperature range after bending (Figure 5c). Similar result is also observed in the temperature dependence of H_EB (Figure 5d). For the small increase in H_C and H_EB, the possible reason might be related to the increased number of small Co/CoO granules after bending. In other words, the increased coupling area of FM/AFM contributes to the overall H_C and H_EB. Our flexible Co/CoO thin film shows no degradation (even a slight increase) of performances in all figures of merit after 500 bending cycles, which proves that our flexible Co/CoO film exhibits decent mechanical durability.

[IMAGE OMITTED. SEE PDF]

Conclusion

In summary, we have demonstrated unexpected large exchange bias field with excellent flexibility in classical Co/CoO thin films grown by magnetron sputtering. The unexpectedly large EB field is well explained by the decrease in ferromagnetic grain size and the increase of ferromagnetic/antiferromagnetic contact area. Furthermore, our flexible Co/CoO shows strain-tuning effect and excellent mechanical durability after 500-time bending. This work provides an effective approach for the integration of large exchange bias in flexible spintronics.

Experimental Section

Film Acquisition

Co/CoO thin film were prepared on 125 µm thick flexible PEN substrate using DC magnetron sputtering in a high vacuum chamber. At the same time, reference samples were grown on rigid Si substrate for structural and morphology characterization. The base vacuum was lower than 10⁻⁴ Pa and the work pressure was 0.5 Pa. To achieve oblique deposition, the sputtering target was oriented at a 30° angle to the normal axis of the thin film plane. Before growth, the sputtering power was set to 80 W and the deposition rate was calculated to be 0.16 nm s⁻¹. The Co thin film was grown at a pure argon gas atmosphere. After the Co deposition, CoO layer was obtained by in situ exposure to pure oxygen at a pressure of 0.1 Pa for 3 min. This typical step produced a dense CoO layer that prevented further oxidization in the air and provided a good FM/AFM interface quality. All deposition procedures were performed at room temperature.

Characterizations

The structural properties and thicknesses of sample were studied using transmission electron microscopy (TEM) from Jem3200fs. Focus ion beam (FIB, Helios 5 UX) was used to observe the cross-section of the material. The elemental distribution was measured by using EDS (X-Max N 100TLE) mapping. Magnetic properties were measured using SQUID magnetometer (Quantum Design MPMS 3). The detailed measurement procedures were as follows. First, Co/CoO thin films were cooled down from 300 to 2 K in the presence of an external field of 1.5 T. Then, the thin films were warmed up to a setting temperature and magnetic hysteresis loops were measured. The maximum magnetic field was 1.5 T for magnetic hysteresis loops measurement. To explore the dependence of magnetic properties on the mechanical strain, the flexible Co/CoO thin films were bent with different curvatures, and the direction of magnetic field was parallel to the thin film plane during the bent measurements. Mechanical durability was evaluated over 500 bending cycles.

Acknowledgements

Y.S. and W.T. contributed equally to this work. This work was supported by the National Natural Science Foundation of China (grant nos. 52273298 and 62174044), the Shenzhen Science and Technology Foundation (grant nos. JCYJ20180507182246321, JCYJ20220818100204010, and JCYJ20200109105825504), and the Guangdong Basic and Applied Basic Research Foundation (grant nos. 2022A1515010649 and 2022A1515010940).

Note: The cited years in refs. [9] and [17] were corrected on 11 January 2023, after initial publication online.

Conflict of Interest

The authors declare no conflict of interest.

Data Availability Statement

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

References

W. H. Meiklejohn, C. P. Bean, Phys. Rev. 1956, 102, 1413.

S. Gider, B. U. Runge, A. C. Marley, S. S. P. Parkin, Science 1998, 281, 797.

S. Fukami, C. Zhang, S. DuttaGupta, A. Kurenkov, H. Ohno, Nat. Mater. 2016, 15, 535.

M. Melzer, D. Makarov, A. Calvimontes, D. Karnaushenko, S. Baunack, R. Kaltofen, Y. Mei, O. G. Schmidt, Nano Lett. 2011, 11, 2522.

S. Zhao, Z. Zhou, C. Li, B. Peng, Z. Hu, M. Liu, ACS Nano 2018, 12, 7167.

L. M. Loong, W. Lee, X. Qiu, P. Yang, H. Kawai, M. Saeys, J. H. Ahn, H. Yang, Adv. Mater. 2016, 28, 4983.

J. Y. Chen, Y. C. Lau, J. M. Coey, M. Li, J. P. Wang, Sci. Rep. 2017, 7, [eLocator: 42001].

H. Matsumoto, S. Ota, A. Ando, D. Chiba, Appl. Phys. Lett. 2019, 114, [eLocator: 132401].

X. Liao, S. Ulusoy, R. Huang, E. Wetterskog, K. Gunnarsson, Y. Wang, H. Liang, Y.‐J. Zeng, G. Salazar‐Alvarez, P. Svedlindh, Sci. China Mater. 2022, 65, 193.

C. Jin, Y. Zhu, W. Han, Q. Liu, S. Hu, Y. Ji, Z. Xu, S. Hu, M. Ye, L. Chen, Appl. Phys. Lett. 2020, 117, [eLocator: 252902].

J. G. Ovejero, V. Godinho, B. Lacroix, M. A. García, A. Hernando, A. Fernández, Mater. Des. 2019, 171, [eLocator: 107691].

X. Shi, W. Mi, Q. Zhang, X. Zhang, Appl. Phys. Lett. 2020, 117, [eLocator: 132401].

T. D. Ha, M. Yen, Y. H. Lai, C. Y. Kuo, C. T. Chen, A. Tanaka, L. Z. Tsai, Y. F. Zhao, C. G. Duan, S. F. Lee, C. F. Chang, J. Y. Juang, Y. H. Chu, Nanoscale 2020, 12, 3284.

Y. Ling, Y. Hu, H. Wang, B. Niu, J. Chen, R. Liu, Y. Yuan, G. Wang, D. Wu, M. Xu, Z. Han, J. Du, Q. Xu, ACS Appl. Mater. Interfaces 2021, 13, [eLocator: 24285].

Z. Zhang, E. Liu, W. Zhang, P. K. J. Wong, Z. Xu, F. Hu, X. Li, J. Tang, A. T. S. Wee, F. Xu, ACS Appl. Mater. Interfaces 2019, 11, 8258.

W. C. Zheng, D. X. Zheng, Y. C. Wang, D. Li, C. Jin, H. L. Bai, J. Magn. Magn. Mater. 2019, 481, 227.

X. Liao, P. Svedlindh, G. Salazar‐Alvarez, Y.‐J. Zeng, R. Huang, S. Yang, Y. Wang, J. Phys. D: Appl. Phys. 2021, 54, [eLocator: 045001].

H. Li, C. Wang, D. Li, P. Homm, M. Menghini, J. P. Locquet, C. Van Haesendonck, M. J. Van Bael, S. Ruan, Y. J. Zeng, J. Phys.‐Condes. Matter 2019, 31, [eLocator: 155301].

W. L. Roth, Phys. Rev. 1958, 110, 1333.

D. Herrmann‐Ronzaud, P. Burlet, J. Rossat‐Mignod, J. Phys. C: Solid State Phys. 1978, 11, 2123.

T. Gredig, I. N. Krivorotov, E. D. Dahlberg, Phys. Rev. B 2006, 74, [eLocator: 094431].

K. Takano, R. H. Kodama, A. E. Berkowitz, W. Cao, G. Thomas, Phys. Rev. Lett. 1997, 79, 1130.

S. Laureti, S. Y. Suck, H. Haas, E. Prestat, O. Bourgeois, D. Givord, Phys. Rev. Lett. 2012, 108, [eLocator: 077205].

K. Gandha, R. P. Chaudhary, J. Mohapatra, A. R. Koymen, J. P. Liu, Phys. Lett. A 2017, 381, 2092.

M. Perzanowski, M. Marszalek, A. Zarzycki, M. Krupinski, A. Dziedzic, Y. Zabila, ACS Appl. Mater. Interfaces 2016, 8, [eLocator: 28159].

H. L. Liu, S. Brems, Y. J. Zeng, K. Temst, A. Vantomme, C. Van Haesendonck, J. Phys.‐Condes. Matter 2016, 28, [eLocator: 196002].

T. D. Ha, J. W. Chen, M. Yen, Y. H. Lai, B. Y. Wang, Y. Y. Chin, W. B. Wu, H. J. Lin, J. Y. Juang, Y. H. Chu, ACS Appl. Mater. Interfaces 2020, 12, [eLocator: 46874].

L. Lukashuk, K. Fttinger, E. Kolar, C. Rameshan, D. Teschner, M. Hvecker, A. Knop‐Gericke, N. Yigit, H. Li, E. McDermott, J. Catal. 2016, 344, 1.

K. Przybylski, W. W. Smeltzer, J. Electrochem. Soc. 1981, 128, 897.

J. Nogués, I. K. Schuller, J. Magn. Magn. Mater. 1999, 192, 203.

T. J. Moran, J. M. Gallego, I. K. Schuller, J. Appl. Phys. 1995, 78, 1887.

M. Gruyters, D. Riegel, Phys. Rev. B 2000, 63, [eLocator: 052401].

A. P. Malozemoff, J. Appl. Phys. 1988, 63, 3874.

M. Molina‐Ruiz, A. F. Lopeandía, F. Pi, D. Givord, O. Bourgeois, J. Rodríguez‐Viejo, Phys. Rev. B 2011, 83, [eLocator: 140407(R)].

A. D. Lamirand, M. M. Soares, A. Y. Ramos, H. C. N. Tolentino, M. De Santis, J. C. Cezar, A. de Siervo, M. Jamet, Phys. Rev. B 2013, 88, [eLocator: 140401(R)].

S. Sahoo, S. Polisetty, Y. Wang, T. Mukherjee, X. He, S. S. Jaswal, C. Binek, J. Phys.‐Condes. Matter 2012, 24, [eLocator: 096002].

J. K. Jochum, T. Saerbeck, V. Lazenka, V. Joly, L. Shan, H.‐G. Boyen, K. Temst, A. Vantomme, M. J. Van Bael, J. Magn. Magn. Mater. 2019, 483, 76.

Z. Hussain, V. Raghavendra Reddy, J. Appl. Phys. 2017, 122, [eLocator: 103903].

J. B. Yi, J. Ding, J. Magn. Magn. Mater. 2006, 303, [eLocator: e160].

A. N. Dobrynin, D. Givord, Phys. Rev. B 2012, 85, [eLocator: 014413].

B. Kocaman, K. Yıldız Aktaş, A. C. Basaran, J. Magn. Magn. Mater. 2021, 530, [eLocator: 167926].

S. Singh, P. Kumar, A. Gupta, D. Kumar, J. Magn. Magn. Mater. 2020, 513, [eLocator: 167186].

E. Menendez, J. Demeter, J. Van Eyken, P. Nawrocki, E. Jedryka, M. Wojcik, J. F. Lopez‐Barbera, J. Nogues, A. Vantomme, K. Temst, ACS Appl. Mater. Interfaces 2013, 5, 4320.

C. Ge, X. Wan, E. Pellegrin, Z. Hu, S. Manuel Valvidares, A. Barla, W. I. Liang, Y. H. Chu, W. Zou, Y. Du, Nanoscale 2013, 5, [eLocator: 10236].

M. Öztürk, E. Sınır, E. Demirci, M. Erkovan, O. Öztürk, N. Akdoğan, J. Appl. Phys. 2012, 112, [eLocator: 093911].

G. Dai, Q. Zhan, Y. Liu, H. Yang, X. Zhang, B. Chen, R.‐W. Li, Appl. Phys. Lett. 2012, 100, [eLocator: 122407].

Z. Tang, B. Wang, H. Yang, X. Xu, Y. Liu, D. Sun, L. Xia, Q. Zhan, B. Chen, M. Tang, Y. Zhou, J. Wang, R.‐W. Li, Appl. Phys. Lett. 2014, 105, [eLocator: 103504].