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There have been reports for different systems that there is an asymmetry in the saturation magnetisation in exchange bias systems [1, 2]. This report contains the first comprehensive study of these effects in the CoFe/IrMn system which is used in the read head of almost all HDDs [3].&lt;br/&gt;&lt;br/&gt;The magnetic behaviour of exchange bias systems has two quasi-independent contributions [3]. The first contribution derives from the ferromagnetic (F) layer which in device applications is in the range 2 to 5 nm. This thickness produces the maximum loop shift when the structure is field cooled. However, it has been shown that there can be a further significant contribution to the magnetisation from spins lying at the interface which can have their origin either from the F layer or the antiferromagnetic (AF) layer. It has been shown that ordering of interface spins increases the loop shift by increasing the coupling between the F and AF layers and that doping the interface in the CoFe/IrMn system with atomic layers of the high-moment Mn can increase the spin ordering significantly [3, 4].&lt;br/&gt;&lt;br/&gt;In this work the asymmetry in the saturation moment (&lt;i&gt;m&lt;/i&gt;&lt;sub&gt;s&lt;/sub&gt;) and its temperature (T) dependence to &amp;lt;4.2 K have been measured for sputtered polycrystalline multilayer Si/Ta (2)/Ru (2)/IrMn (10)/Mn (&lt;i&gt;d&lt;/i&gt;&lt;sub&gt;Mn&lt;/sub&gt;)/CoFe (5)/Ru (5) (thicknesses in nm) thin films where 0≤&lt;i&gt; d&lt;/i&gt;&lt;sub&gt;Mn&lt;/sub&gt; ≤0.2 nm in a LakeShore 8600 Series VSM. The York Model of Exchange Bias was used to separate and interpret the effect of the F/AF interface [3]. Low-temperature measurements were undertaken to ensure that the bulk of the AF grains remained magnetically frozen. Thus there was no reorientation of the magnetic moments in the bulk of the AF grains in subsequent measurements.&lt;br/&gt;&lt;br/&gt;The magnetisation of an exchange bias system undergoes an athermal training effect whereby the first and second magnetisation loops exhibit different loop shifts after the initial field cooling, even when the temperature of measurements is such that the moment of the bulk of the AF grain does not vary. This effect on the coercivity, which is defined as half loop width, has been shown to originate in the interfacial spins [5].&lt;br/&gt;&lt;br/&gt;Figure 1 shows the measured hysteresis loop, the structure of which is shown in the insert, at 5.5 and 300 K. For the system measured we established by separate measurements (not shown) that the loop shift for the system shown was thermally stable at T = 300 K by holding the magnetisation at reverse saturation for a period of 30 mins after which time the hysteresis loop show at 300 K reproduced. In figure 1 the difference between the loops measured was due to changes in the orientation and ordering of interfacial spin clusters. The increase in the measured loop shift comes from ordering of spins at the interface. Furthermore, the first loop training effect is very small at T = 300 K but becomes significant at T = 5.5 K. This shows that the training effect is caused by interfacial spins, as is the coercivity.&lt;br/&gt;&lt;br/&gt;There is an asymmetry in the magnetisation. This data was obtained for the same sample without it being removed from the magnetometer. A significant increase in &lt;i&gt;m&lt;/i&gt;&lt;sub&gt;s&lt;/sub&gt; occurs at 5.5 K, which is greater at negative saturation than is the case in a positive field. We designate this difference in &lt;i&gt;m&lt;/i&gt;&lt;sub&gt;s&lt;/sub&gt; in positive (&lt;i&gt;m&lt;/i&gt;&lt;sub&gt;s&lt;/sub&gt;&lt;sup&gt;+&lt;/sup&gt;) and negative (&lt;i&gt;m&lt;/i&gt;&lt;sub&gt;s&lt;/sub&gt;&lt;sup&gt;–&lt;/sup&gt;) fields as&lt;br/&gt;&lt;br/&gt;Δ&lt;i&gt;m&lt;/i&gt;&lt;sub&gt;s&lt;/sub&gt; = |&lt;i&gt;m&lt;/i&gt;&lt;sub&gt;s&lt;/sub&gt;&lt;sup&gt;+&lt;/sup&gt;|–|&lt;i&gt;m&lt;/i&gt;&lt;sub&gt;s&lt;/sub&gt;&lt;sup&gt;–&lt;/sup&gt;|. (1)&lt;br/&gt;&lt;br/&gt;Figure 2 shows the temperature dependence of Δ&lt;i&gt;m&lt;/i&gt;&lt;sub&gt;s&lt;/sub&gt; for samples where there is a variation in the thickness of the Mn interface doping layer (&lt;i&gt;d&lt;/i&gt;&lt;sub&gt;Mn&lt;/sub&gt;). It has been found that the optimum loop shift for this structure is affected by &lt;i&gt;d&lt;/i&gt;&lt;sub&gt;Mn&lt;/sub&gt; and has been found to be a layer of subatomic thickness of the order of 1/6 of an atom. The use of layers of an atomic thickness or greater were found to result a reduction in Δ&lt;i&gt;m&lt;/i&gt;&lt;sub&gt;s&lt;/sub&gt; [4]. Figure 2 shows Δ&lt;i&gt;m&lt;/i&gt;&lt;sub&gt;s&lt;/sub&gt; for the first (i) and second (ii) loops measured immediately after the field annealing process. There is a change in the interfacial spin order which gives rise to the interface training effect but also effects the change in Δ&lt;i&gt;m&lt;/i&gt;&lt;sub&gt;s&lt;/sub&gt;.&lt;br/&gt;&lt;br/&gt;For all films figure 2 shows that the increase in Δ&lt;i&gt;m&lt;/i&gt;&lt;sub&gt;s&lt;/sub&gt; is small at T &amp;gt; 100 K. We attribute this to the interface spins not being ferromagnetically ordered and hence not responding to the modest fields used to measure the loops. Below 100 K the change in Δ&lt;i&gt;m&lt;/i&gt;&lt;sub&gt;s&lt;/sub&gt; with T becomes more rapid as it appears these spins become F ordered, probably in a spin cluster structure similar to a spin glass [6]. It appears that at &amp;lt;6 K Δ&lt;i&gt;m&lt;/i&gt;&lt;sub&gt;s&lt;/sub&gt; goes through a peak for all samples which is consistent with a spin cluster structure.&lt;br/&gt;&lt;br/&gt;It would be expected that &lt;i&gt;d&lt;/i&gt;&lt;sub&gt;Mn&lt;/sub&gt; would not be optimum for a layer of subatomic thickness. It is most likely that the Mn layer is making a compositional adjustment at the interface. In studies of a similar system by Baltz &lt;i&gt;et al&lt;/i&gt;. using 3-D atom probe tomography on an exchange bias system containing Co and Mn significant diffusion was found to occur [7]. CoMn alloys have complex properties existing in both F, AF and superparamagnetic phases dependent on composition and temperature but no high resolution phase diagram has been determined [8]. In the production of an exchange bias system the field annealing requires up to 600 K which promotes interfacial diffusion [3, 7]. We conclude that the role of the Mn interface layer is to replace Mn deficiencies in the bulk of the AF.

**References:**

[1] U. Nowak et al., Phys. Rev. B, Vol. 66(1), p. 014430 (2002)&lt;br/&gt;&lt;br/&gt;[2] J. Keller et al., Phys. Rev. B, Vol. 66(1), p. 014431 (2002)&lt;br/&gt;&lt;br/&gt;[3] K. O&#39;Grady et al., J. Magn. Magn. Mater., Vol. 322(8), p. 883-899 (2010)&lt;br/&gt;&lt;br/&gt;[4] R. Carpenter et al., IEEE Trans. Magn., Vol. 48(11), p. 4351–4354 (2014)&lt;br/&gt;&lt;br/&gt;[5] R. Carpenter et al., J. Appl. Phys., Vol. 115(17), p. 17D715 (2014)&lt;br/&gt;&lt;br/&gt;[6] J. L. Tholence and R. Tournier, Phys. B+C, Vol. 86-88, p. 873-874 (1977)&lt;br/&gt;&lt;br/&gt;[7] F. Letellier, et al., J. Appl. Phys., Vol. 116, p. 203906 (2014)&lt;br/&gt;&lt;br/&gt;[8] A. Z. Men’shikov, et al., Sov. Phys. JETP, Vol. 62, p. 734 (1985)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/c4880f4f4f141c9f2e76b5ec7dfed9c9.jpg)
Figure 1: The first (i) and second (ii) hysteresis loops at 5.5 and 300 K for a film with structure shown inset.

![](https://assets.underline.io/uploads/markdown_image/1/image/b484480c5ac00e6406cf86b341a6f8a8.jpg)
Figure 2: The temperature dependence of the magnetisation offset (Δ&lt;i&gt;m&lt;/i&gt;&lt;sub&gt;s&lt;/sub&gt;) measured at ± 5 kOe shown for the first (i) and second (ii) magnetisation reversal for the thin films shown inset where 0≤&lt;i&gt; d&lt;/i&gt;&lt;sub&gt;Mn&lt;/sub&gt; ≤0.2 nm.

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 Magnetisation Asymmetry in Exchange Bias Systems

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There have been reports for different systems that there is an asymmetry in the saturation magnetisation in exchange bias systems [1, 2]. This report contains the first comprehensive study of these effects in the CoFe/IrMn system which is used in the read head of almost all HDDs [3]. The magnetic behaviour of exchange bias systems has two quasi-independent contributions [3]. The first contribution derives from the ferromagnetic (F) layer which in device applications is in the range 2 to 5 nm. This thickness produces the maximum loop shift when the structure is field cooled. However, it has been shown that there can be a further significant contribution to the magnetisation from spins lying at the interface which can have their origin either from the F layer or the antiferromagnetic (AF) layer. It has been shown that ordering of interface spins increases the loop shift by increasing the coupling between the F and AF layers and that doping the interface in the CoFe/IrMn system with atomic layers of the high-moment Mn can increase the spin ordering significantly [3, 4]. In this work the asymmetry in the saturation moment (ms) and its temperature (T) dependence to &lt;4.2 K have been measured for sputtered polycrystalline multilayer Si/Ta (2)/Ru (2)/IrMn (10)/Mn (dMn)/CoFe (5)/Ru (5) (thicknesses in nm) thin films where 0≤ dMn ≤0.2 nm in a LakeShore 8600 Series VSM. The York Model of Exchange Bias was used to separate and interpret the effect of the F/AF interface [3]. Low-temperature measurements were undertaken to ensure that the bulk of the AF grains remained magnetically frozen. Thus there was no reorientation of the magnetic moments in the bulk of the AF grains in subsequent measurements. The magnetisation of an exchange bias system undergoes an athermal training effect whereby the first and second magnetisation loops exhibit different loop shifts after the initial field cooling, even when the temperature of measurements is such that the moment of the bulk of the AF grain does not vary. This effect on the coercivity, which is defined as half loop width, has been shown to originate in the interfacial spins [5]. Figure 1 shows the measured hysteresis loop, the structure of which is shown in the insert, at 5.5 and 300 K. For the system measured we established by separate measurements (not shown) that the loop shift for the system shown was thermally stable at T = 300 K by holding the magnetisation at reverse saturation for a period of 30 mins after which time the hysteresis loop show at 300 K reproduced. In figure 1 the difference between the loops measured was due to changes in the orientation and ordering of interfacial spin clusters. The increase in the measured loop shift comes from ordering of spins at the interface. Furthermore, the first loop training effect is very small at T = 300 K but becomes significant at T = 5.5 K. This shows that the training effect is caused by interfacial spins, as is the coercivity. There is an asymmetry in the magnetisation. This data was obtained for the same sample without it being removed from the magnetometer. A significant increase in ms occurs at 5.5 K, which is greater at negative saturation than is the case in a positive field. We designate this difference in ms in positive (ms+) and negative (ms–) fields as Δms = |ms+|–|ms–|. (1) Figure 2 shows the temperature dependence of Δms for samples where there is a variation in the thickness of the Mn interface doping layer (dMn). It has been found that the optimum loop shift for this structure is affected by dMn and has been found to be a layer of subatomic thickness of the order of 1/6 of an atom. The use of layers of an atomic thickness or greater were found to result a reduction in Δms [4]. Figure 2 shows Δms for the first (i) and second (ii) loops measured immediately after the field annealing process. There is a change in the interfacial spin order which gives rise to the interface training effect but also effects the change in Δms. For all films figure 2 shows that the increase in Δms is small at T &gt; 100 K. We attribute this to the interface spins not being ferromagnetically ordered and hence not responding to the modest fields used to measure the loops. Below 100 K the change in Δms with T becomes more rapid as it appears these spins become F ordered, probably in a spin cluster structure similar to a spin glass [6]. It appears that at &lt;6 K Δms goes through a peak for all samples which is consistent with a spin cluster structure. It would be expected that dMn would not be optimum for a layer of subatomic thickness. It is most likely that the Mn layer is making a compositional adjustment at the interface. In studies of a similar system by Baltz et al. using 3-D atom probe tomography on an exchange bias system containing Co and Mn significant diffusion was found to occur [7]. CoMn alloys have complex properties existing in both F, AF and superparamagnetic phases dependent on composition and temperature but no high resolution phase diagram has been determined [8]. In the production of an exchange bias system the field annealing requires up to 600 K which promotes interfacial diffusion [3, 7]. We conclude that the role of the Mn interface layer is to replace Mn deficiencies in the bulk of the AF.

**References:**

[1] U. Nowak et al., Phys. Rev. B, Vol. 66(1), p. 014430 (2002) [2] J. Keller et al., Phys. Rev. B, Vol. 66(1), p. 014431 (2002) [3] K. O'Grady et al., J. Magn. Magn. Mater., Vol. 322(8), p. 883-899 (2010) [4] R. Carpenter et al., IEEE Trans. Magn., Vol. 48(11), p. 4351–4354 (2014) [5] R. Carpenter et al., J. Appl. Phys., Vol. 115(17), p. 17D715 (2014) [6] J. L. Tholence and R. Tournier, Phys. B+C, Vol. 86-88, p. 873-874 (1977) [7] F. Letellier, et al., J. Appl. Phys., Vol. 116, p. 203906 (2014) [8] A. Z. Men’shikov, et al., Sov. Phys. JETP, Vol. 62, p. 734 (1985)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/c4880f4f4f141c9f2e76b5ec7dfed9c9.jpg)
Figure 1: The first (i) and second (ii) hysteresis loops at 5.5 and 300 K for a film with structure shown inset.

![](https://assets.underline.io/uploads/markdown_image/1/image/b484480c5ac00e6406cf86b341a6f8a8.jpg)
Figure 2: The temperature dependence of the magnetisation offset (Δms) measured at ± 5 kOe shown for the first (i) and second (ii) magnetisation reversal for the thin films shown inset where 0≤ dMn ≤0.2 nm.

technical paper

The thermoelectric generation based on the anomalous Nernst effect (ANE) has become one of the promising approaches to realize efficient energy harvesting from waste heat and heat flux sensor due to its several advantages over conventional devices based on the Seebeck effect (SE) 1,2. However, till date, the observed ANE thermopower is very small and a systematic material investigation to realize large anomalous Nernst coefficient is necessary. Recent theoretical and experimental investigations identified a Weyl semimetal as a potential material having giant ANE which can be achieved by a little tuning of Fermi level (EF) near Weyl points3,4,5. However, systematic tuning of EF and investigation of thermoelectric properties based on many films with different compositions are time-consuming. Furthermore, in this process, one losses many intermediate compositions due to the limitation of number of films that can be prepared. In this regard, we introduce an efficient approach to tune EF for Weyl semimetals and demonstrate the approach through a layer-by-layer combinatorial deposition of a Co2MnAl1-xSix (CMAS) thin film with x varied from 0 to 1 on a single substrate. The fabrication procedure is represented schematically in Fig.1 (a). The structural characterization reveals the formation of single-phase CMAS alloys throughout the composition range. Hard X-ray photoemission spectroscopy (HAXPES) directly confirmed a continuous shifting of EF from Co2MnAl to Co2MnSi with a maximum shift of 0.4 eV (see Fig.1. (b)), which is consistent with the theoretical prediction. We also measured the anomalous Ettingshausen effect (AEE), the reciprocal of ANE, for the all x range using a single strip along the composition gradient (see Fig.1. (c)) by means of the lock-in thermography technique6,7. Only one LIT image tells us that large AEE is obtained between x = 0.06 to 0.12, which well agrees with the composition dependence of the ANE signals (see Fig.1. (d)), measured by making many parallelly aligned Hall bars elongated to the orthogonal direction to the composition gradient (see Fig.1. (c)). The findings clearly demonstrate that the AEE measurement using the composition spread film is an effective approach to investigate the composition dependence of ANE and find the highest performance without fabricating many films and performing systematic measurements. These demonstrations will enable the optimization of the composition in Weyl semimetals and exploration of new materials to obtain giant ANE/AEE, which is essential to realize applications based on the transverse thermoelectric effects.

**References:**

[1] Y. Sakuraba et al., Appl. Phys. Express 6, 033003 (2013) [2] W. Zhou et al., Appl. Phys. Express 13, 043001 (2020) [3] K. Sumida et al., Commun. Mater. 1, 89 (2020) [4] J. Noky et al., NPJ Comput. Mater. 6, 77 (2020) [5] Y. Sakuraba et al., Phys. Rev. B 101, 134407 (2020) [6] S. Daimon et al., Nat. Commun. 7, 13754 (2016) [7] T. Seki et al., Appl. Phys. Lett. 112, 152403 (2018)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/58d6615be315220821aa1a7a138fb225.jpg)
Fig.1 (a) Schematic representation of composition-spread CMAS combinatorial film deposited with wedge shaped layer by layer deposition (b) Valance band HAXPES spectra for CMAS composition-spread film at Co2MnAl (CMA) and Co2MnSi (CMS) side (c) Schematic measurement configuration used for ANE and AEE (d) AAEE/jc and SANE as a function of x. Where AAEE is the amplitude of temperature modulation due to AEE, jc is the charge current density and SANE is anomalous Nernst coefficient. The inset shows observed AAEE image in LIT.

 Efficient tuning of electronic, transport, and thermoelectric properties of Weyl semimetal Co2MnAl1-xSix composition-spread thin film

Lately, Magnetic Random-Access Memory (MRAM) has stimulated significant interests compared to traditional semiconductor memory devices due to its non-volatility combined with expeditious read and write endurance. In the preceding studies, conventional methods that had been investigated widely include spin-orbit torque (SOT)1, spin-transfer torque (STT)2 and current-driven switching. However, magnetization switching in MRAM utilizing such techniques fall short to observe the desired switching and inevitably supplemented by thermal issues, leading to inefficient functionality. Recently, considerable endeavour towards a new paradigm of straintronic devices (STR)3 for efficient magnetization switching is explored. A magnetoelectric STR constitutes of ferromagnetic and piezoelectric order parameters, in which an applied electric field across the piezoelectric order parameter instigate mechanical strain at the heterostructure interface giving rise to the magnetization switching due to the Villari effect. Though both out-of-plane and in-plane magnetization switching are examined in the aforementioned studies to explore the STR devices, magnetization switching in the in-plane direction is obscured on account of thermal instability in smaller element size and multi-domain formation in larger element size. Also, the self-biased STR-MRAM device that extinguishes the external bias magnetic field is still undiscovered. In this work, we investigated the size-dependent in-plane magnetization switching of FeGaB/PMN-PT (ferromagnetic/piezoelectric) heterostructure (Fig. 1(a)) using micromagnetic finite element method (FEM) coupled with electrodynamic and elastic analysis. Three elliptical nanodiscs of FeGaB magnetoelastic material having cross-sectional areas 100×90.9 nm2, 100×66.67 nm2, and 100×50 nm2 are investigated by varying the thickness from 2 nm to 28 nm on a 10×10×0.5 mm3 single-crystal PMN-PT (011) piezoelectric substrate. Initially, minimum thickness (dmin) at pre-stress condition for three elliptical nanodiscs are investigated (Fig. 1(b)) using in-plane induced energy well (△E) scaled by kbT (where kb is Boltzmann constant and T is the temperature in Kelvin). △E/kbT≤70 is an assumed condition for thermal instability to ensure dmin for STR-MRAM devices. To explore the maximum thickness related to self-biasing (dsb), pre-stress equilibrium exchange and demagnetization energies associated with the three elliptical nanodiscs are analyzed against thickness. The larger size magnets make an effort to minimize the energy by aligning the magnetization with the elliptical nanodisc boundary and gives rise to vortex state. On contrary, smaller size magnets minimize the energy by aligning the magnetization along the easy axis (0° or 180°) of the elliptical FeGaB nanomagnets without applying the external magnetic bias field (Fig. 2). This result is pivotal as such self-biasing would give primacy to future STR-MRAM devices by omitting the necessity for an extrinsic magnetic bias field. Although the above analysis is performed to obtain optimum size where STR-MRAM device is stable and self-biased, the maximum thickness (dmax) for in-plane 180° magnetization switching is limited by the critical stress requirement. It is observed that dsb is not necessarily the same as dmax for in-plane 180° magnetization switching. When compressive strain equal or greater than the critical compressive strain is applied from the pre-stress state, where magnetization is along the easy axis, in-plane 90° magnetization switching (along 90° or 270°) is observed. This happens as exchange and stress anisotropy energies are almost negligible and the demagnetization energy rises as the magnetization is switched from the easy to the hard axis. Magnetic precession is used to obtain in-plane 180o switching. When the magnetization component along the x-axis is negative during the precession, removal of external strain causes another 90o switching because the ferromagnet tries to reduce its total energy by gravitating to its nearest easy axis. In-plane fastest 180o magnetization switching time and associated external strain are investigated also. For fixed thickness, compressive strain requirement rises with increasing cross-sectional area. As the thickness is increased, tradeoff between fastest switching time and associated switching energy is observed for the same cross-sectional area. Whereas faster switching time and lesser switching energy are required for larger cross-sectional nanodiscs having the same thickness. This work can be potentially useful in the future STR-MRAM devices that eliminates the necessity of extrinsic magnetic field bias and completes in-plane 180o magnetization switching.

**References:**

1C. Zhang, S. Fukami, H. Sato, F. Matsukura, H. Ohno, Appl. Phys. Lett., 107, 012401 (2015). 2C.F. Pai, L. Liu, Y. Li, H. W. Tseng, D. C. Ralph, R. A. Buhrman, Appl. Phys. Lett., 101, 122404 (2012). 3K. Roy, S. Bandyopadhyay, J. Atulasimha, Appl. Phys. Lett., 99, 063108 (2011).

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/f80b0892f94f3338bf712e9f149fd36a.jpg)
FIG 1: (a) Proposed structure consists of FeGaB/PMN-PT (ferromagnetic/piezoelectric) heterostructure and corresponding magnetization components. (b) △E/kbT vs thickness curve to examine the minimum thickness (dmin) requirement for thermal stability.

![](https://assets.underline.io/uploads/markdown_image/1/image/58c037c03f736d7b2641a003318c63ad.jpg)
FIG 2: (a) Simulated in-plane 90o magnetization switching from initial random magnetization state after applying critical stress. (b) Simulated in-plane 180o magnetization switching after magnetization switches to in-plane easy axis for 100×90.9×9.5 nm3 nanodisc.

 Micromagnetic study of strain-induced magnetization switching in FeGaB nanomagnets for self-biased MRAM applications

In this study, we report room temperature negative magneto capacitance due to 60% Fe doping and positive magneto capacitance due to 60% Ti doping at Mn site in La0.7Pb0.3MnO3. We have studied the structural, mixed valence studies using synchrotron X-ray absorption spectroscopy (XAS), magnetic and magneto capacitance properties of La0.7Pb0.3Mn0.4Fe0.6O3-δ and La0.7Pb0.3Mn0.4Ti0.6O3-δ. La0.7Pb0.3Mn0.4Fe0.6O3-δ and La0.7Pb0.3Mn0.4Ti0.6O3-δ were synthesized by solid state method. The x-ray diffraction (XRD) patterns show (i) successful substitution of Fe and Ti at Mn site and (ii) change in lattice parameters of La0.7Pb0.3Mn0.4Fe0.6O3-δ and La0.7Pb0.3Mn0.4Ti0.6O3-δ. L edge absorption studies using synchrotron XAS show the presence of mixed valence states in these compositions. La0.7Pb0.3Mn0.4Fe0.6O3-δ show maximum ~64 % negative magneto capacitance at 500 gauss(G) magnetic field, ~60 % at 5kG magnetic field and -56 % at 10 kG magnetic field at 300 K. La0.7Pb0.3Mn0.4Fe0.6O3-δ show ~51% negative magneto capacitance at 15 kG magnetic field while La0.7Pb0.3Mn0.4Ti0.6O3-δ show ~ 12 % positive magneto capacitance at 15 kG magnetic field. Magneto capacitance values in these compositions at room temperature in accordance with the Maxwell Wagner effect, which was explained by Catalan et al. Interaction between Mn+3 and Mn+4 via oxygen atom in the lattice originates the local magnetic region in the middle of non magnetic regions which originated negative magneto capacitance while Ti doped samples show positive magneto capacitance due to the presence quantum interference and coulombic interaction effects. Presence of unpaired electrons in the d orbitals of Fe, Mn and Ti dictates the magneto capacitance properties of these materials.

**References:**

1. G. Catalan Appl Phys Lett 88, 102902 (2006) 2. P. Chen , D. Y. Xing and Y. W. Du Phys. Rev. B 64, 104402 (2001) 3. S. S. Manoharan, B. Singh, R.K. Sahu, et al. J. Phys.: Condens. Matter 20, 235205 (2008) 4. B. Singh Mater. Let. 156, 76 (2015)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/d8756a16e22f94e3a35f045e93a84cc9.png)
Fig. 1 Magneto capacitance vs frequency plots at different applied Magnetic field for La0.7Pb0.3Mn0.4Fe0.6O3-δ polycrystalline sample.

![](https://assets.underline.io/uploads/markdown_image/1/image/60a94b74ab5d40cb913da542d67977db.png)
Fig. 2 Magneto capacitance vs frequency plots at different applied Magnetic field for La0.7Pb0.3Mn0.4Ti0.6O3-δ polycrystalline sample.

 Huge room temperature negative magneto capacitance in La0.7Pb0.3Mn0.4Fe0.6O3-δ and positive magneto capacitance in La0.7Pb0.3Mn0.4Ti0.6O3-δ

Coupling lattice and spin degrees of freedom without the use of external magnetic fields allows for energy-efficient spintronic devices. In this context, hybrid materials composed of an oxide layer undergoing a first-order, structural phase transition (SPT), such as V2O3, and a ferromagnetic layer, such as Ni, constitute a very promising paradigm for the development of such applications. In the case of a Ni/V2O3 heterostructure, the structural domains inherent to the first-order phase transition in the V2O3 cause a differential strain pattern in the Ni layer. By virtue of the inverse magnetorestrictive effect, this strain pattern is converted into a magnetic anisotropy distribution in the Ni. This offers a promising alternative to voltage-controlled magnetism or other switching mechanisms without a magnetic field [1,2]. Here we show by direct imaging of the thermal evolution of the Ni spin structure that the magnetic domains can be tuned in both size and orientation upon crossing the structural phase transition (SPT) of the proximal V2O3 layer [3]. We find a drastic temperature-driven reorientation of the Ni magnetic domains across the SPT which is responsible for an abrupt increase in the coercive field (up to 500%). By synchrotron-based X-ray microscopy we show a reconfiguration of the ferromagnetic domain pattern in the Ni layer across the V2O3 SPT together with changes in magnitude and direction of the magnetic anisotropy. These observations are in good agreement with both angular dependent ferromagnetic resonance measurements and micromagnetic simulations. Our simulations successfully replicate the change in the coercivity as well as the reorientation of the magnetization of the Ni domains after the implementation of a simple model of strain-dependent magnetic anisotropies. Furthermore, simulations also exhibit a phase coexistence of different magnetic domains across the SPT in the V2O3. Direct observations of the lateral correlation length of the Ni domains show an increase of almost an order of magnitude at the SPT compared to room temperature as well as a broad spatial distribution of the local transition temperatures. This corroborates the phase coexistence of Ni anisotropies due to the V2O3 structural domains coexistence across the SPT. Our data reveal that the reorientation of Ni domains is controlled by the reconfiguration of the structural domains of the oxide layer across the SPT [4,5], due to strain induced proximity. Our findings reveal a novel pathway to control magnetic domains without a magnetic field through proximity to a material undergoing a first-order structural phase transition, which allows for engineering of coercive fields for novel data storage architectures and novel device concepts based on “straintronics”. Acknowledgements: This work was supported by the Spanish MINECO projects MAT2015-68772-P and PGC2018-097789-B-I00 and European Union FEDER funds. A.F.R. and M.G.M acknowledge financial support from the EU CALIPSO Transnational Access programme. J.G.R. acknowledges support from Colciencias under grant 120471250659. The work by I.V., C.W. and I.K.S. at UCSD were supported by the Office of Basic Energy Science, U.S. Department of Energy, BES-DMS funded by the Department of Energy’s Office of Basic Energy Science, DMR under grant DE FG02 87ER-45332.

**References:**

[1] de la Venta, J., Wang, S., Ramirez, J. G., et al., Appl. Phys. Lett. 102, 1 (2013); [2] de la Venta, J., Wang, S., Saerbeck, T., et al., Appl. Phys. Lett. 104, 062410 (2014). [3] Valmianski, I., Fraile Rodríguez. A., Rodriguez-Alvarez, J., et al., (in review). [4] McLeod, A. S., Van Heumen, E., Ramirez, J. G., et al. Nat. Phys. 13, 80 (2017). [5] Gilbert, D. A., Ramírez, J. G., Saerbeck, T., et al., Sci. Rep. 7 13471 (2017).

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/609df7dde4c22e1de9b043b18a9dd651.jpg)
Fig. 1. Top row panels show representative XMCD-PEEM images recorded at temperatures below (150 K), across (170 K) and above (200 K) the SPT of V2O3. The insets show their Fourier transforms. The initial state (blue contrast) at low temperature (LT) splits into a pattern of small domains (yellow) which progressively rotate and grow across the V2O3 SPT as the high-temperature (HT) phase is reached. Bottom row panels show histograms of the absolute value of the average rotation angle of the Ni magnetization, Df, at selected temperatures (calculated from the images in the top row). Sharp Df distributions are found for the LT (150 K) and HT (200 K) phases in contrast to broader distributions in the middle of the V2O3 SPT (170 K).

![](https://assets.underline.io/uploads/markdown_image/1/image/0691162ae0a39dfe6b44c065dd312054.jpg)
Fig. 2. Coercivity as a function of temperature of Ni(10 nm)/V2O3(100 nm) across the V2O3 SPT recorded while cooling the sample from 300 K to 120 K (only data from 200 K are shown). A sharp rise in the coercive field, Hc, of about a factor of four, is found as the temperature goes below the V2O3 SPT. The garnet circles represent the coercivity values obtained from the experimental loops while the green triangles correspond to those obtained from micromagnetic simulations.

 Nanoscale manipulation of magnetic domains by interfacial strain-induced proximity

Systematic studies of magnetic properties observed in epitaxial W/Co/Pt layered structures grown by molecular beam epitaxy are discussed. In this system contributions to interfacial Dzyaloshinskii-Moriya interaction (DMI) from the asymmetric interfaces between ferromagnet (Co) and non-magnetic heavy metals with high spin-orbit coupling (W or Pt) are of additive nature. DMI can stabilize the chiral magnetic texture, vortex and skyrmions, being promising in future data storage technology. Presented progressive investigations encompass layered structures containing: a single Co layer (magnetic anisotropy), two Co layers (interlayer coupling) and multilayers (dipolar interactions, complex domain structure). Layered structures containing a single Co layer exhibit spin reorientation transition (SRT) with increasing thickness, dCo,from perpendicular (perpendicular magnetic anisotropy, PMA) to in-plane aligned magnetisation. The range of dCo with PMA and the critical thickness dSRT depend on the type and sequence of adjacent covers. In the Pt/W/Co/Pt system dCo range with PMA shifts towards higher values (from 0.6-0.9 nm to 2.1-2.4 nm) with increasing bottom W layer thickness, dW, from 2.4 nm to 8.0 nm. This behaviour differs from the similar system with Mo component leyer, studied earlier [1]. Hysteresis loops measured by polar magnetooptical Kerr effect (PMOKE) magnetometry exhibit shapes typical of fine domain structure with perpendicular magnetisation. In the W/Co(1.0 nm)/Pt(dPt)/W(dW)/Co(1.0 nm)/Pt interlayer antiferromagnetic coupling of perpendicular magnetization with the strength as high as 0.6 kOe is very distinct for the component spacers of equal thickness, dPt = dW, in the range between 0.7 and 1.1 nm. For dW ≠ dPt (keeping the same total thickness dW + dPt) the coupling strength decreases. Particularly, the single spacer systems (Pt or W; dPt or dW kept in the corresponding range from 1.4 nm to 2.0 nm) do not exhibit such feature, indicating a crucial role of the complex spacer structure affecting both coupling and PMA (Figure 1). The multilayer structures W/Co(0.6 &lt; dCo &lt; 1.0 nm)/Pt (10-20 repetitions) are magnetised in perpendicular direction to the film plane and display at the remanent state typical labyrinth domain structure with tendency towards bubble domains formation with diameter of around 75 nm (Figure 2). DMI strength determined form (aligned) domain structure reaches the value of D = 2.49 mJ/m2 [2]. Such high value of D is a result of high quality crystalline structure characteristic for epitaxial systems and additive contribution to DMI from asymmetric well-defined interfaces. The above discussed structures are compared to those with different component layer thickness and opposite sequence (i.e. Pt/Co/W). This work is supported by IEEE educational seed funding 2020 (beneficiary:-Sukanta Kumar Jena), Foundation for Polish Science (FNP) under the European Regional Development Fund – Program REINTEGRATION 2017 OPIE 14-20 and by the National Science Centre in Poland, projects no: 2016/23/G/ST3/04196 and 2020/37/B/ST5/02299.

**References:**

[1] A Wawro et al. J. Phys. D: Appl. Phys. 50 215004 (2017). [2] S.K. Jena et al. to be submitted (2021).

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/f62fbee3caf0455305aedeea9ea695c0.jpg)
“Fig.1” Hysteresis PMOKE loop from W/Co/Pt/W/Co/Pt layered structures composed of two Co (1nm) layers separated by the Pt/W spacer with total thickness of 1.4 nm showing distinct antiferromagnetic coupling (red). Corresponding structures with a single spacer: W (black) or Pt (green) 1.4 nm thick do not exhibit such coupling.

![](https://assets.underline.io/uploads/markdown_image/1/image/5285c84da6791eb5744546eceba7f522.jpg)
"Fig.2" Magnetic force microscopy image from as grown W/Co/Pt multilayer with basic trilayer repetition number equal to 20.

 Magnetic anisotropy, interlayer coupling and Dzyaloshinskii-Moriya interaction in epitaxial W/Co/Pt multilayers

The interfacial Dzyaloshinskii-Moriya Interaction has had profound consequences on the possible magnetic configurations that can be stabilized in multi-layer thin films, most notably chiral Neel skyrmions and domain walls.[1] More recently, there have been theoretical calculations and experimental work confirming the existence of an appreciable anisotropic DMI in epitaxial thin films with further reduced symmetry – namely the C2v (or mm2) point group.[2] However, so far the systems under consideration have been based on cubic crystals where the magnetocrystalline anisotropy is nearly isotropic. Here, we identify a new epitaxial relationship with C2v symmetry based on HCP Co(10.0)/FCC Pt(110). Although it is quite common to grow FCC Co(111)/Pt(111), for the case of a Pt(110) underlayer, the lattice mismatch is actually smaller for the HCP Co(10.0) orientation [as opposed to FCC Co(110)] where the c-axis lies along a single direction in the plane of the film (figure 1). Such an orientation is expected to offer some degree of anisotropic DMI based on the heavy metal Pt and reduced symmetry as well as the added feature of significant uniaxial in-plane magnetocrystalline anisotropy. Samples were prepared by DC magnetron sputtering on HF-etched Si(110) single crystal substrates. Following the work of Yang et al., we use a Ag (10nm) buffer layer, which is known to grow with (110) orientation on Si(110) at room temperature, thereby avoiding further unwanted interface diffusion.[3] This is followed by deposition of the Pt (10nm) /Co (10-50nm) bilayers, which are the focus of this work. All depositions were done in an Ar atmosphere with the working pressure fixed at 2.5 mTorr. The films were characterized structurally by standard θ/2θ x-ray diffraction (XRD) measurements and phi-scans to characterize the epitaxial relationship. The XRD scan shows primarily Ag(220), Pt(220), and Co(10.0) peaks with no identifiable Co(220) peak suggesting the film is predominantly HCP. Phi scans (figure 2) on Pt(111), Ag(111), and Co(10.1) confirm the epitaxial relationship shown in figure 1. Moreover, reciprocal space q-scans on the HCP Co spots were leveraged to characterize directional peak broadening due to stacking faults. As it is known that stacking faults cause broadening only of certain peaks (h-k ≠ 3m) along the c* direction (with breadth also determined by the parity of l), it is possible to extract the density of growth and deformation faults in the Co layer, which are found to be approximately 10% and 3% respectively in 50nm Co films. Details on how these stacking fault densities are calculated will presented in detail. Alternating gradient field magnetometry was used to measure M-H loops along different crystallographic directions in the plane of the film. A clear in-plane easy axis is identified along Si[001] (parallel to Co[0001]). The hard axis saturation field of 1.5 kOe suggests a notably smaller magnetocrystalline anisotropy than would be expected for bulk cobalt, likely due to the aforementioned stacking faults. Although the relatively thick Co layer in these films makes any interface induced DMI unlikely, the epitaxial relationship is expected to be preserved in thinner samples along with the reported uniaxial anisotropy. Together with the expected interfacial anisotropic DMI, this in-plane anisotropy makes the proposed material system an interesting playground for exploring new chiral magnetic ground states.

**References:**

[1] Thiaville, A., Rohart, S., Jué, É., Cros, V. and Fert, A., 2012. Dynamics of Dzyaloshinskii domain walls in ultrathin magnetic films. EPL (Europhysics Letters), 100(5), p.57002. [2] Hoffmann, M., Zimmermann, B., Müller, G.P., Schürhoff, D., Kiselev, N.S., Melcher, C. and Blügel, S., 2017. Antiskyrmions stabilized at interfaces by anisotropic Dzyaloshinskii-Moriya interactions. Nature communications, 8(1), pp.1-9. [3] Yang, W., Lambeth, D.N. and Laughlin, D.E., 1999. Unicrystal Co–alloy media on Si (110). Journal of applied physics, 85(8), pp.4723-4725.

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/cf10786ba9d1e164cc2eadffc57f6d46.jpg)
Fig 1. Schematic of the lattice matching for Ag(110)/Pt(110) with bottom left) FCC Co(110) or bottom right) HCP Co(10.0). The HCP (10.0) orientation has a better match along the Pt [001] direction.

![](https://assets.underline.io/uploads/markdown_image/1/image/78b7907db3cc501cd82fd177111f404e.jpg)
Fig 2. Phi scans for Pt (111), Ag (111), and Co (10.1) confirming the epitaxial relationship shown in figure 1.

 Structural and Magnetic Characterization of Epitaxial Co(10.0)/Pt(110) Multi-layers for Future Anisotropic DMI Systems Based on C2v Symmetry

Skyrmions typically appear in systems with perpendicular magnetic anisotropy (PMA) and Dzyaloshinskii-Moriya interactions (DMI). Bloch skyrmion states can be stabilized in magnetic dots with no need of DMI [1]. However, PMA is typically considered to be necessary. Nanoparticle geometry is an additional important factor for the skyrmion stabilization due to the influence of magnetostatic energy. Therefore, the sample shape (e.g. spherical) can result in stabilization of the skyrmions [2]. Recently, stabilization of the half-hedgehog skyrmions in soft magnetic (Permalloy) hemispherical dots by the MFM tip was demonstrated [3]. Here we conduct micromagnetic simulations in 3D soft magnetic dots with planar (cylindrical) or curved (spherical cap) geometry with neither PMA nor DMI. The geometries of both magnetic elements correspond to the base radius 60-120 nm and the aspect ratio thickness/ radius ranged from 0.4 to 2 and magnetic parameters of permalloy. In a wide range of the dot sizes, we observed the co-existence of multiple magnetic states. Particularly, in the spherical caps we stabilized for the same geometry, the in-plane, the out-of-plane (flower-like), the 3D vortex-like (Bloch skyrmion), the 3D Néel skyrmion –like (see Fig.1) and the Bloch point-like states. The Bloch skyrmion is characterized by a vortex configuration in the basal plane while the Neel-like skyrmon – by the radial vortex configuration and is similar to the half hedgehog structure. The existence of magnetic skyrmions without DMI or PMA in cylindrical nanodots was demonstrated by analytical calculations using the Belavin-Polyakov ansatz [4]. Both Néel and Bloch (vortex-like) skyrmions structures have been proven to be metastable or stable states in some range of dot sizes. The Néel skyrmion always has larger energy than the Bloch skyrmion. Direct micromagnetic simulations confirm the co-existence of both types of skyrmions, although in a smaller range of geometrical parameters than predicted by analytical calculations. In spherical caps, the curved geometry provides an additional factor for the Néel skyrmion stabilization which was obtained for a wide range of geometrical parameters. Typically, the Néel skyrmion has the largest energy of all structures, however its energy decreases as a function of the cap basal diameter and aspect ratio. Consequently, for some geometrical parameters we found a Néel skyrmion having smaller energy than that of the Bloch skyrmion. The spherical geometry also leads to a coupling between the skyrmion polarization and radial magnetization components, so that only the skyrmions with positive Polarity (P) and outward chirality (Q) or negative polarity and inward chirality are possible, see Fig. 1 Finally, we present the state diagram for skyrmions in soft magnetic caps and nanopillars. We have also calculated 2D topological charges and the gyrovector values, both depending on the geometrical parameters.

**References:**

[1] K. Guslienko, K. Y. IEEE Magn. Lett. Vol. 6, p,4000104 (2015) [2] V.P. Kravchuk et al Phys. Rev. B, Vol. 85, p.144433 (2012). [3] E.Berganza et al, Nanoscale, Vol.12, p.18646 (2020) [4] A. A. Belavin and A. M. Polyakov, Pis’ma Zh. Eksp. Teor. Fiz. Vol. 22, p. 503 (1975) .

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/c0723cd91b5db4d72e227b49f23c8980.jpg)
Figure.1 (a-b) Chiral (Neel-like) skyrmion magnetic configurations in permalloy caps with diameter 60 nm and height 30 nm colored by the axial-component mz. Left and right figures are cross sections in the XZ plane and in the XY basis plane respectively. P and Q stand for the polarity and the chirality of the skyrmion respectively.

 Skyrmions in 3D soft magnetic nanodots with no Dzyaloshinskii-Moriya interactions

Exchange-spring nanomagnets [1–3] are composed of the interaction of the magnetic exchange coupling between magnetically hard and soft phases. Such frameworks are promising for cutting edge perpetual magnetic applications, as they have a high energy product - the mix of permanent magnetic field and magnetization - contrasted with the conventional, single-phase magnetic materials. Hence, for the last three decades, these exchange spring nanomagnets have taken the keen attention of the scientific community. However, the necessary condition that both the hard and soft phases are controlled at the nanometer scale, to ensure effective exchange coupling, has posed significant preparation challenges. In pursuit of preparing multi-nanolayered-structured nanomagnetic material, sputtering, electron beam lithography/ evaporation (EBL/EBE) have inspired a wealth of experimental results. However, to date, there is a very limited experimental understanding of preparing in situ multilayered nanomagnetic material using electrodeposition technique and the elucidation of their magnetic behavior. Towards this end, using electrodeposition technique with a modified and stable acidic bath chemistry at room temperature, we have prepared in situ multilayered nanomagnetic cobalt phosphorus (CoP) thin film and investigated its characteristics structurally and magnetically combined with micromagnetic simulation utilizing MuMax3. Despite the availability of various techniques, the electrodeposition technique [4–7] has been chosen for the preparation of the nanomagnetic cobalt phosphorous thin films at room temperature using acidic bath chemistry due to its relatively higher deposition rate and low cost. Here, the crossover from the exchange spring structure to exchange coupled structures by the manipulation of thickness in nanometer level in a single amorphous system of cobalt phosphorus nanomagnetic thin film has been demonstrated. We have characterized the developed smooth, stress-free, and shiny electroplated samples structurally by XRD (shown in FIG. 1), SEM, and magnetically by the hysteresis loop tracer (ShB Instruments, USA). The SEM micrographs confirm the fine grain surface morphology of the cobalt phosphorous nanomagnetic thin films. The XRD analysis further ensures the amorphousness of the thin film, as in the XRD pattern, the three dominating peaks only appears from the substrate, i.e., the largest peak near 2θ = 69.22° is due to the Si-wafer and the rest of the two peaks near 2θ = 38° and 83° are from the top conducting seed layer of gold (Au). The elemental analysis by energy dispersive X-Ray (EDX) yields the stoichiometry of the electrodeposited cobalt phosphorous. An optimum exchange coupling can be obtained by tuning the spin configuration with respect to the multi-nano layer thicknesses. The static in-plane magnetic measurement by the hysteresis loop tracer shown in FIG. 2 reveals different hysteresis loops of electrodeposited CoP thin film. FIG. 2(a) confirms the exchange spring amorphous cobalt phosphorus films with a relatively higher coercivity of ~197 Oe for 0.68 µm CoP thin film while the FIG. 2(b) affirms the flat hysteresis loop with a coercivity of 162 Oe for the 11.54 µm thick CoP thin film which signifies the exchange coupled structure. Utilizing the finite difference method-based freely available GPU-accelerated micromagnetic simulator, MuMax3 [8], we have also verified the experimentally obtained results further by doing a micromagnetic simulation. Considering the typical magnetic material parameters (i.e., Saturation Magnetization (MS) as obtained from the static measurement, Exchange Stiffness Constant (AEx), Anisotropy Constant (Ku1)), MuMax3 solves the well-known Landau-Lifshitz-Gilbert (LLG) equations, and so forth gives the static micromagnetic evolution of the nanomagnetic thin films. To proceed with the simulation, we have discretized the overall amorphous cobalt phosphorous thin film into the cells of dimension 5 x 5 x 5 nm3. To include all the exchange interactions in the static magnetization of nanoscale thin film, this cell size in each direction is considered lower than the exchange length (lex = 5.1 nm) of the system, as well as to precisely reproduce the nanomagnetic impression of the thin film under study. The dimensions of the whole atlas we used for the simulation are (l x w x t = 8000 x 6000 x 5400 nm3), where l, w, and t are the length, width, and thickness of the atlas, respectively. Finally, we report the micromagnetic spin configuration at every coordinate point within the nanomagnetic thin film at every instant of field determined by the simulator, which confirms the formation of a multidomain state from a single domain state (shown in FIG. 2(e))

**References:**

[1] E. F. Kneller and R. Hawig, IEEE Trans. Magn. 27, 3560 (1991). [2] E. E. Fullerton, J. S. Jiang, M. Grimsditch, Phys. Rev. B 58, 12193 (1998). [3] H. Zeng, J. Li, J. P. Liu, Z. L. Wang, Nature 420, 395 (2002). [4] P. McCloskey, B. Jamieson, T. O’Donnell, J. Magn. Magn. Mater. 320, 2509 (2008). [5] P. McCloskey, B. Jamieson, T. O’Donnell, J. Magn. Magn. Mater. 322, 1536 (2010). [6] P. McCloskey, T. O’Donnell, B. Jamieson, in Handb. Nanoelectrochem. Electrochem. Synth. Methods, Prop. Charact. Tech.,2016, pp. 1–27. [7] S. Roy, A. Connell, M. Ludwig, J. Magn. Magn. Mater. 290–291, 1524 (2005). [8] A. Vansteenkiste, J. Leliaert, M. Dvornik, AIP Adv. 4, 107133 (2014).

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/ab0568662cd1fd8eb2bea9c14dc8237e.jpg)
FIG. 1. XRD of 0.68 µm, 5.06 µm thick electrodeposited thin films having only the substrate peaks (i.e., Au and Si) confirms the amorphousness of the samples.

![](https://assets.underline.io/uploads/markdown_image/1/image/af142b7a2e9ace085c438741e63feae8.png)
FIG. 2. Hysteresis loop of the electrodeposited CoP thin films measured in XY plane, and here, dir-1+ 90° = dir-2. The thickness of the films is (a) 0.68 µm, (b) 11.54 µm. (c)-(d) The differential of the magnetization of the (a) 0.68 µm, (b) 11.54 µm thick CoP thin films, respectively. (e) Simulated Hysteresis Loop and different micromagnetic spin configurations at different coordinates of the hysteresis loop.

 Exchange Spring Co-Rich CoP Nanomagnetic Thin Films

Controlled manipulation of spin texture in magnetic materials by external means, e.g., magnetic and electric fields, temperature etc., is a prerequisite for the cutting-edge spintronic applications. Magnetic materials with perpendicular magnetic anisotropy (PMA) are of special interest because of their wide-ranging functionalities starting from applications in magnetoresistive random access memory (MRAM) devices to modulation of topological spin textures. Recently, there has been a growing interest in rare earth (RE)-transition metal (TM) alloy thin films due to their tuneable spin configurations as a function of temperature and RE content. For example, GdFeCo shows in-plane spin configuration at room temperature, but PMA occurs near the compensation temperature (TCO).[1] Such spin reorientation is generally attributed to competition between PMA and shape anisotropy. In order to understand the behaviour of magnetic anisotropy around the spin reorientation temperature, several indirect methods were employed to determine the temperature dependence of anisotropy fields, e.g., longitudinal and polar MOKE, FMR etc.[1] Here, we examined the spin reorientation in single layer ferrimagnetic alloy thin films of FeGd by dc magnetometry and radio frequency transverse susceptibility measurement, which is a direct method to determine the effective anisotropy fields. The FeGd system belongs to the RE-TM family and is particularly interesting as TCO of this system can be tuned by varying the chemical composition. We prepared the FexGd100-x(90 nm)/Ta(8 nm) films onto a Si substrate coated with a 10 nm Ta layer using RF and DC sputtering; we focus below on a nominal composition of Fe75.7Gd24.3. From the temperature dependence of the in-plane magnetization M(T) curve, TCO for the Fe75.7Gd24.3 film was estimated to be ~ 70 K (main panel of Fig. 1(a)). Inset of Fig. 1(a) depicts the in-plane M(H) loops for this film measured at T = 10 and 300 K. While the in-plane saturation magnetization has drastically dropped between 300 and 10 K, the coercivity is enhanced significantly. In addition, the ratio of remanent magnetization to the saturation magnetization is an order of magnitude higher at 10 K than at 300 K. These observations indicate the occurrence of spin reorientation between 10 and 300 K. In order to probe the spin reorientation, we performed magnetic field dependent transverse susceptibility (TS) measurements for both in-plane and out-of-plane configurations at different temperatures between T = 300 and 10 K. Our TS measurement set up exploits a self-resonant ultra-sensitive tunnel diode oscillator circuit with a resonant frequency of 13 MHz and is a highly sensitive technique to precisely determine the effective anisotropy fields.[2] The inset of Fig. 1(b) shows the magnetic field dependence of the relative change in TS at 40 K for the Fe75.7Gd24.3 for in-plane configuration. As the dc magnetic field is swept from positive to negative saturation, the first (second) peak to positive (negative) anisotropy field (HK). As shown in the main panel of Fig. 1(b), HK remains the same for in-plane and out-of-plane configurations but begins bifurcating below 250 K. At lower temperatures, the out-of-plane HK shows a broad maximum between 200-150 K and then decreases smoothly. On the other hand, the in-plane HK continues to decrease down to 100 K below which it shows a small plateau around TCO and then decreases gradually. These observations indicate that the spin configuration in the system changes from a combination of in-plane and out-of-plane spins with equal contributions above 250 K, to a combination of in-plane and out-of-plane spins with a major contribution by the out-of-plane spins below 250 K. The origin of such temperature-driven spin reorientation might be related to phase-segregation. With the help of scanning transmission electron microscopy equipped with energy dispersive x-ray spectroscopy (STEM-EDS) and x-ray reflectivity (XRR) measurements, a recent study shows that modulation of chemical composition occurs in FeGd films with thickness more than 20 nm where Fe is partially diffused to the adjacent Ta layer.[3] Since the thickness of our film is ~ 90 nm, it is possible that there are different Fe-deficient (Gd-rich) regions which are separated by regions with nominal chemical composition. In such a hybrid phase-segregated spin system, different regions have distinct spin configurations and hence, different orientations of magnetic anisotropy. We have also performed out-of-plane M(H) measurements on the Fe75.7Gd24.3 film that showed a double-step reversal behaviour below 250 K and hence, confirms the chemical phase-segregation in our system. The spin reorientation in our system is thus determined by the competition between local magnetic anisotropy energies of adjacent phase-segregated regions. Our TS technique determines the average behaviour of the local anisotropy field and hence effectively probes the spin reorientations in these micron-scale phase-segregated magnetic systems. Results for other Fe/Gd compositions will also be discussed.

**References:**

[1] W. He, H. -L. Liu, H.-Y. Wu, J. -W. Cai, and Z. -H. Cheng, Appl. Phys. Lett. 106, 042401 (2015). [2] V. Kalappattil, R. Geng, R. Das, M. Pham, H. Luong, T. Nguyen, A. Popescu, L. M. Woods, M. Kläui, H. Srikanth and M. H. Phan, Mater. Horiz. 7, 1413-1420 (2020). [3] E. Kirk, C. Bull, S. Finizio , H. Sepehri-Amin, S. Wintz , A. K. Suszka, N. S. Bingham, P. Warnicke, K. Hono, P. W. Nutter, J. Raabe, G. Hrkac, T. Thomson and L. J. Heyderman, Phys. Rev. Mater. 4, 074403 (2020).

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/ad612b336fa97c19a15298613ec09cf9.png)
(a) Main panel: in-plane M(T) of Fe75.7Gd24.3 measured in 1 T field and inset: in-plane M(H) loops at 10 and 300 K, (b) main panel: in-plane and out-of-plane HK(T), inset: field dependent TS curve at 40 K for in-plane configuration.

 Competing Magnetic Anisotropies and Spin Reorientation in Phase-Segregated Single Layer Ferrimagnetic FeGd Films

Antiferromagnets (AFMs), due to their unique set of properties, are promising candidates for future spintronic materials. However, their low ordering temperature makes a wide group of AFM materials not applicable. Hence, tuning ordering temperature and spin configuration due to magnetic proximity effect in AFM/AFM systems has been recently identified as an important field of research [1]. In our studies we investigated how the proximity of antiferromagnetic layer with a higher ordering temperature (Néel temperature, TN) influences the TN of FeO in FeO/CoO bilayer. The layers were grown by molecular beam epitaxy on MgO(001) substrate coated with a 5 nm-thick homoepitaxial MgO buffer. CoO layer with a thickness of 2 nm was prepared by reactive deposition of Co in a molecular oxygen atmosphere under a pressure of 1×10-6 mbar. FeO layer was deposited by reactive deposition of 1 nm of Fe under oxygen pressure of 5×10-8 mbar, which corresponds to about 1.7 nm of stoichiometric wüstite (FeO). To enable Mössbauer spectroscopy (MS) studies we used 57Fe isotope during FeO deposition. The sample was capped with 3 nm of MgO. To compare magnetic properties of wüstite in FeO/CoO bilayer with a single FeO layer a MgO/FeO/MgO sample was prepared. The magnetic properties of FeO were studied with Conversion Electron Mössbauer Spectroscopy (CEMS). CEMS measurements were obtained in temperature controlled cryostat under UHV conditions with use of a channeltron detector connected to a standard Mössbauer spectrometer and a 57Co(Rh) source irradiating the sample at normal incidence. Figure 1 shows CEMS spectra collected at 240 K for MgO/57FeO/MgO (Fig. 1(a)) and MgO/57FeO/CoO/MgO (Fig. 1(b)). The temperature of 240 K is far above the TN of FeO (198 K) [2]. The spectrum of FeO (Fig. 1(a)) is characteristic for bulk wüstite in paramagnetic state [3]. It was fitted by two quadrupole sites, which describe the octahedrally coordinated Fe2+ (Q1) atoms and Fe2+ atoms neighbouring defects (Q2). Interestingly, proximity of CoO causes a serious modification of the spectrum, it becomes magnetically split. Similar spectrum was recorded for MgO/FeO/MgO under the TN (not shown). Systematic measurement of CEMS spectra recorded as a function of temperature for FeO/CoO revealed that the magnetic character of spectrum is preserved up to 260 K. Thus, proximity of CoO layer induces enhancement of the TN of FeO up to 260 K. To determine TN of CoO in FeO/CoO bilayer we performed X-ray Magnetic Linear Dichroism (XMLD) measurements [4]. X-ray absorption spectra (XAS) were collected using a linearly polarized x-rays for two grazing incident angles (0° and 60°) as a function of temperature. XMLD effect observed at the Co L3 edge can be attributed to the AFM ordering in CoO. Thus, RL3 difference (ΔRL3) described as ΔRL3 = RL30 - RL360, where RL30 and RL360 are defined as the ratio of the intensities of the XAS peaks at 777 eV and 779.6 eV for grazing incident angles of 0° and 60°, respectively is a measure of AFM ordering in CoO. The AFM contribution to ΔRL3(T) decreases with an increase of temperature and vanishes above TN. The inset in Figure 1 shows ΔRL3(T) dependence obtained for CoO in MgO/FeO/CoO/MgO multilayers. We noted positive ΔRL3 values, which indicates in-plane alignment of CoO AFM spins in FeO/CoO/MgO [5,6]. Dependence of ΔRL3 on temperature is negligible for T ≥ 290 K. Thus, the ordering temperature of CoO layer in FeO/CoO/MgO stack is comparable to the TN of bulk CoO (293 K). Enhancement of ordering temperature of FeO should result in increase in exchange interaction between ferromagnetic (FM) layer and FeO in FM/FeO/CoO multilayer. To elucidate how AFM proximity effect influence exchange interaction with FM a dedicated sample was prepared for which Fe(1 nm)/FeO bilayer was grown on wedge shaped CoO. Figure 2 presents the product of exchange bias field (defined as HEB = (HC1 - HC2)/2, where C1 and C2 represent respectively the coercive fields of ascending and descending branches of hysteresis loop) and CoO layer thickness dCoO as a function of temperature obtained for different CoO thicknesses. The values of HEB were determined from the magnetic hysteresis loops measured with longitudinal magnetooptic Kerr effect (LMOKE). While for Fe/FeO/CoO(0.7 nm) we noted vanishing of HEB * dCoO at around 220 K (so called blocking temperature, TB), for dCoO &gt; 2 nm TB was established at around 260 K, which is much greater than a blocking temperature of Fe/FeO bilayer (160 K). Our study demonstrates the feasibility of tailoring the magnetic properties of AFM layer with a proximity of other AFM. This work was supported by the “Antiferromagnetic proximity effect and development of epitaxial bimetallic antiferromagnets – two routes towards next-generation spintronics” project, which is carried out within the Homing programme of the Foundation for Polish Science co-financed by the European Union under the European Regional Development Fund.

**References:**

 [1]. S. Fukami et al., Journal of Applied Physics 128 070401 (2020) [2] K. Koike, T. Furukawa, Physical Review Letters 77 3921 (1996) [3] A. Koziol-Rachwal et al., Applied Physics Letters 108 041606 (2016) [4] https://synchrotron.uj.edu.pl/en_GB/linie-badawcze/peem-xas [5] J. Zhu et al., Physical Review B 90 054403 (2014) [6] Q. Li et al., Scientific Reports 6 22355 (2016)

**Images:**

![](https://assets.underline.io/uploads/markdown_image/1/image/e8177b448bb39166aa1dc962682247fd.jpg)
Fig. 1 CEMS measurements of FeO (black points) in MgO/FeO/MgO and MgO/FeO/CoO at 240 K. Inset represents ΔRL3(T) dependence for CoO in FeO/CoO/MgO(001).

![](https://assets.underline.io/uploads/markdown_image/1/image/b63fa16fb0795eb0b8aa4752326fde37.jpg)
Fig. 2 HEB*dCoO(T) obtained from LMOKE measurements for MgO/Fe/FeO/CoO/MgO(001) for different CoO thicknesses. Lines are guides for the eye.

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