Volume 7 Number 8 Materials August 2020 Pages 1923–2172 Horizons rsc.li/materials-horizons ISSN 2051-6347 COMMUNICATION Jordi Sort, Ignasi Fina et al. Local manipulation of metamagnetism by strain nanopatterning Materials Horizons This article is licensed under a Creative Commons Attribution-NonCommercial 3.0 Unported Licence. View Article Online COMMUNICATION View Journal | View Issue Local manipulation of metamagnetism by strain nanopatterning† Cite this: Mater. Horiz., 2020, Open Access Article. Published on 14 May 2020. Downloaded on 11/4/2020 3:23:56 PM. 7, 2056 Michael Foerster,a Enric Menéndez,b Emerson Coy, c Alberto Quintana,b Received 10th April 2020, Carles Gómez-Olivella, d Daniel Esqué de los Ojos,b Oriol Vallcorba, a Accepted 13th May 2020 Carlos Frontera, e Lucia Aballe,a Josep Nogués,fg Jordi Sort *bg and DOI: 10.1039/d0mh00601g Ignasi Fina *e rsc.li/materials-horizons Among metamagnetic materials, FeRh alloys are technologically appealing due to their uncommon antiferromagnetic-to-ferromagnetic New concepts metamagnetic transition which occurs at a temperature T* just above Control of the transition temperature (T*) in metamagnetic materials, such as room temperature. Here, a controlled increase of T* (DT* B 20 8C) is the transition from antiferromagnetic to ferromagnetic states in FeRh, has been so far realized by means of different approaches: voltage-driven piezo- induced in pre-selected regions of FeRh films via mechanical strain strain, application of strong magnetic fields (spin–flip or spin-flop transitions) nanopatterning. Compressive stresses generated at the vicinity of pre- or chemical doping, amongst others. These methods typically affect the entire defined nanoindentation imprints cause a local reduction of the FeRh FeRh films, causing global changes in their magnetic properties. Here, a crystallographic unit cell parameter, which leads to an increase of T* in straightforward alternative concept, which allows for a large enhancement of these confined micro-/nanometric areas. This enhances the stability of the stability of the AFM phase both locally and globally, is presented. Nanoindentation patterning (i.e., localized generation of compressive the antiferromagnetic phase in these localized regions. Remarkably, mechanical stress in pre-selected micro- and nanoscale actuated regions) is generation of periodic arrays of nanopatterned features also allows shown to induce a significant shift of T* in these pre-defined areas of the FeRh modifying the overall magnetic and electric transport properties across films. This method results in an array of confined regions which exhibit an large areas of the FeRh films. This approach is highly appealing for the enhanced antiferromagnetic stability with respect to the adjacent non-stressed design of new memory architectures or other AFM-spintronic devices. film. By properly tuning the heating conditions (and not far from room temperature) it is possible to obtain an array of antiferromagnetic dots embedded in a ferromagnetic matrix. The concept can be extrapolated to Introduction other materials displaying functional properties that are dependent on the crystallographic unit cell volume. Therefore, this method constitutes an Nanopatterning of magnetic structures is an ubiquitous approach interesting approach towards functional nanostructuration of physical properties, eventually breaking the light diffraction limits, and allowing for widely used in applications involving magnetic materials,1–4 the implementation of novel micro-/nano-device architectures. particularly in storage media. Nanopatterning is also interesting for the basic understanding of effects such as exchange bias, a ALBA Synchrotron Light Facility, Cerdanyola del Vallès, Carrer de la Llum 2-26, current-induced magnetic switching and other spintronic Barcelona E-08290, Spain phenomena.5–8 From the broad range of methods to perform b Departament de Fı́sica, Universitat Autònoma de Barcelona, Bellaterra, nanopatterning, mechanical nanoimprinting, which uses local Barcelona E-08193, Spain. E-mail: jordi.sort@uab.cat c NanoBioMedical Centre, Adam Mickiewicz University, Wszechnicy Piastowskiej 3, stress applied by means of a tip, is a convenient technique 61-614, Poznan, Poland because, in spite of its simplicity, it allows generating structures d Departament Fı́sica Aplicada i Òptica, Universitat de Barcelona, Martı́ i Franquès 1, eventually smaller than those of standard optical methods 08028 Barcelona, Spain which cannot easily go below the light diffraction limit. Via e Institut de Ciència de Materials de Barcelona (ICMAB-CSIC), Campus UAB, mechanical nanopatterning, properties such as crystallinity,9,10 Bellaterra E-08193, Catalonia, Spain. E-mail: ifina@icmab.es f Catalan Institute of Nanoscience and Nanotechnology (ICN2), CSIC and BIST, nanoparticle arrangement,11 ferroelectricity,12,13 electric transport13 Campus UAB, Bellaterra, E-08193 Barcelona, Spain or metal-insulator transitions,14 amongst others, can be controlled g ICREA, Pg. Lluı́s Companys 23, E-08010 Barcelona, Spain at the nanoscale. Remarkably, mechanical nanopatterning of non- † Electronic supplementary information (ESI) available: Supplementary materials ferromagnetic alloys can also be used to obtain nanomagnetic includes materials and methods and Fig. S1–S8. Supplementary materials Video S1 motifs (i.e., magnetic nanolithography).15–17 corresponds to direct PEEM image at Fe L3 absorption edge at ca. 707 eV while increasing temperature to the video was recorded on the same region shown in FeRh in its a phase shows an antiferromagnetic (AFM)- Fig. 2. The temperature ranges from 85 1C and 115 1C. Topographic contribution to-ferromagnetic (FM) phase transition slightly above room has not been subtracted. See DOI: 10.1039/d0mh00601g temperature.18 This particular functionality makes FeRh interesting 2056 | Mater. Horiz., 2020, 7, 2056--2062 This journal is © The Royal Society of Chemistry 2020 View Article Online Communication Materials Horizons for magnetic and spintronic applications, such as heat-assisted The as-grown FeRh film exhibits the conventional magnetic recording,19–23 tunneling24 and AFM-based memory devices,25–27 properties of this alloy,22 where the temperature dependence of magnetocalorics and barocalorics.28,29 This AFM–FM transition the magnetization, M(T), displays an abrupt transition from near has been widely studied as well from a fundamental point of zero magnetization at room temperature, in the AFM phase, to This article is licensed under a Creative Commons Attribution-NonCommercial 3.0 Unported Licence. view.30–33 The temperature at which the AFM order disappears about 1000 emu cm 3 at high temperature, in the FM phase (metamagnetic transition temperature, T*) can be tailored by the (Fig. 1(b)). As expected, the system shows thermal hysteresis,22 application of external magnetic fields,34 changing composition/ where the transition temperature (T*) during heating is 90 1C, doping,35,36 electrochemical processes,37 hydrostatic pressure38 whereas it is 80 1C upon cooling. or biaxial strain imposed by the substrate.39–46 The two latter To locally adjust the magnetic properties of the FeRh films, effects are related to the change of the FeRh unit cell across we apply nanoindentations using different loads (see Materials the transition, being the unit cell parameter smaller in the low and methods in ESI†). As it can be seen in Fig. 1(c), the Open Access Article. Published on 14 May 2020. Downloaded on 11/4/2020 3:23:56 PM. temperature AFM phase than in the high temperature FM indentations produce triangular-shaped depressions on the phase.47 Thus, when the FeRh unit cell is reduced by external film and do not induce any important changes on the overall stimuli, the AFM phase is stabilized and, consequently, T* increases. sample morphology except for the typical pile-up effect near the In the present work, we use mechanical nanopatterning in indentation imprint. The compression in the FeRh film produced FeRh films to induce, locally, a shift of T* towards higher by the mechanical indentation is expected to stabilize the AFM temperatures. This is different from previous studies in which phase due to the smaller unit cell volume, resulting in a T* increase. strain is induced in the entire FeRh film, often through application of voltage across an adjacent piezoelectric substrate.39–46 The local Local distribution of the AFM phase near the indented regions increase of T* is evidenced using results from photo-emission In order to gain insight into the changes produced in the electron microscopy (PEEM). Finite element modelling reveals that magnetic properties of the material near the indented region, the T* increase is not solely induced by the imposed nanoindenta- we use X-ray Magnetic Circular Dichroism (XMCD) in combination tion stresses but also mediated by the mechanical properties of the with PEEM50 (XMCD-PEEM, Fig. 2(a); see Materials and methods substrate (MgO), which are transmitted to FeRh by epitaxial strain. in ESI†). In these experiments, black and white regions account for Finally, by using Kerr magnetometry and electric transport char- FM domains and grey regions can account for either AFM acterization, we show that arrays of mechanically patterned struc- domains or domains with net magnetization perpendicular to tures can be used to selectively tune the average T* of large areas the incident light. within in the FeRh films. First, we focus on an indented region where the applied load force is 300 mN. This force is strong enough to ensure that the Results tip reached the MgO substrate (see Fig. S3, ESI†). Shown in Fig. 2(b)–(d) are the XMCD-PEEM images collected while heating Sample structural and magnetic properties and mechanical at 75 1C, 95 1C, and 115 1C, respectively, where the dashed patterning triangle encloses the indentation region (i.e., a 6.5 mm triangle An epitaxial 55 nm FeRh film is grown on a MgO(001) substrate for a 300 mN indentation). At low temperatures, in regions both by RF sputtering (see Materials and methods in ESI†). The ‘‘far’’ from (black square) and ‘‘near’’ (red square) the indenta- structural characterization (see Fig. S1 and S2, ESI†),48,49 tion, the magnetic contrast is negligible (grey contrast), indicat- reveals that the as-deposited film grows in the a phase and it ing the absence of FM domains in both regions. The absence of has a slight tetragonal distortion [a = 2.98(1) Å and c = 2.99(1) Å FM domains is in agreement with the predominant AFM state of lattice parameters] compared to its bulk value.47 The FeRh film the sample at this temperature (Fig. 1(b)). Only starting at the grows epitaxially on the MgO substrate with a [100]FeRh(001)// vertexes of the triangular indentation and along the [110], [1 10], [110]MgO(001) epitaxial relation. Namely, the FeRh pseudo- [ 110], [ 1110] FeRh-directions, needles of black/white domains, cube grows with a 451 rotation with respect to the MgO cubic which account for FM regions, are found at 75 1C (Fig. 2(b)), structure, as sketched in Fig. 1(a).19 schematically depicted in green in the inset of Fig. 2(b). This indicates that near these particular directions there is unit cell expansion, probably resulting from cracks in the film, favouring locally the FM phase, which results in a very confined reduction of T*. At intermediate temperatures (e.g., 95 1C, Fig. 2(c)), where the AFM-to-FM transition starts, black/white domains corres- ponding to FM regions are observed in areas far (black square) from the indentation. Conversely, near the indentation (red Fig. 1 (a) Sketch of the FeRh film grown onto MgO(001). The FeRh square) the contrast is virtually inexistent. These regions with pseudo-cubic structure grows at 451 rotated respect to the MgO one. lower contrast indicate the predominance of the AFM phase. (b) Magnetization, M, versus temperature measured while increasing and Therefore, in these particular regions, the transition to the FM decreasing the temperature and applying 500 Oe after zero field cooling the sample from above 200 1C. (c) Representative atomic force micro- state has not yet started at this temperature. The gradual scopy scans of a FeRh region with an indentation imprint obtained evolution of the magnetic contrast (domain pattern) with applying 10 mN. The field of view is 5  5 mm2. increasing temperature is shown in Video S1, ESI.† Indentations This journal is © The Royal Society of Chemistry 2020 Mater. Horiz., 2020, 7, 2056--2062 | 2057 View Article Online Materials Horizons Communication This article is licensed under a Creative Commons Attribution-NonCommercial 3.0 Unported Licence. Open Access Article. Published on 14 May 2020. Downloaded on 11/4/2020 3:23:56 PM. Fig. 2 (a) Schematic representation of the experimental set-up for XMCD-PEEM experiments. The black and white areas account for domains with net magnetization along and opposite to the direction of the incident light, respectively. Grey regions correspond to either domains without net magnetization (in our case AFM domains), or domains with net magnetization perpendicular to the incident light. CCW and CW correspond to anticlockwise and clockwise X-ray light, respectively. (b)–(d) XMCD-PEEM magnetic domain images of the FeRh film region with a 300 mN indentation measured at 75 1C, 95 1C and 105 1C, respectively. Insets show schematic representations of the image symmetry. (e) and (f) Sketch of the AFM regions after indentation in the FeRh/MgO system, as deduced from the XMCD-PEEM characterization, at the indicated temperatures. (g) Magnetic contrast of the XMCD images (SDXMCD) of the regions enclosed in red (near indentation) and black (far from the indentation) in (b)–(d) and Fig. S8, ESI.† with 5 mN, 10 mN and 20 mN display similar features, although The 4-fold shape of the grey region of Fig. 2(c) along the the AFM areas at high temperature are smaller (see Fig. S4, [100], [010], [ 100] and [0 10] FeRh-directions is in agreement ESI†). Importantly, in all cases, the grey regions have 4-fold with the maximum compression of the {110} planes reported symmetry and they are mainly confined along the [100], [010], for MgO,51,52 as sketched in Fig. 2(e). Notably, according to the [ 100] and [0 10] FeRh-directions, schematically depicted in FeRh mechanical parameters,53 the maximum deformation of a yellow in the inset of Fig. 2(c). At high temperature (Fig. 2(d)), stand-alone FeRh film is expected to be along the {110} FeRh where the FM state should be predominant throughout the planes, as demonstrated by finite-element calculations (see sample, large black/white domains corresponding to FM regions Fig. S5, ESI,† based on low-load indentation data54 of Fig. S6, are found both far and near the indentation. Only very close to ESI†). This is in contrast with our results, which show that the the indentation the contrast remains grey. In this case, one can maximum deformation in FeRh is along the [100], [010], [ 100] infer that the grey region is a mixture of 4-fold with 3-fold and [0 10]. This indicates that the mechanical properties of symmetries along the directions normal to the sides of the the FeRh itself do not play a crucial role in the overall triangular indentation, dictated by the indenter shape (dashed behaviour; instead, it is the deformation of the MgO substrate yellow line). The 3-fold symmetry is schematically depicted in and its propagation to the FeRh film what governs the shift yellow in the inset of Fig. 2(d). in T*. Given the epitaxial growth of the FeRh layer on the MgO, 2058 | Mater. Horiz., 2020, 7, 2056--2062 This journal is © The Royal Society of Chemistry 2020 View Article Online Communication Materials Horizons the larger strain along the {110} MgO planes is transmitted to residual strain produces a clear 20 1C increase of the transition the {100} planes of the FeRh, originating the observed 4-fold temperature. symmetry (Fig. 2(e)). Nevertheless, the mechanical properties of the FeRh and the film interaction with the Berkovich Large-area effects for device functionalization This article is licensed under a Creative Commons Attribution-NonCommercial 3.0 Unported Licence. indentation tip are relevant very close to the indentation as In order to understand the effects of indentation in larger areas evidenced by the 3-fold symmetry of the remaining ‘‘grey (50  50 mm2) of the film, arrays of indentations (Fig. 3(a)) are domains’’ at high temperatures sketched in Fig. 2(f), in this studied by magneto-optic Kerr effect (MOKE) and electrical case in agreement with the finite element simulations of the resistance. The temperature dependence of the Kerr rotation pure FeRh (Fig. S7, ESI†). In fact, very close to the indentation, (proportional to M), yK(T), in the as-obtained film (black line in T* is even higher than 115 1C (note the remaining grey contrast Fig. 3(b)) shows a similar temperature dependence as the M(T) in Fig. 2(d)). Hence, the enhanced T* originates from a complex measured by vibrating sample magnetometer (Fig. 1(b)), Open Access Article. Published on 14 May 2020. Downloaded on 11/4/2020 3:23:56 PM. interplay of the mechanical properties of the film and the although the transition obtained by MOKE is more abrupt, substrate. The AFM–FM transition in the non-indented regions probably due to the improved crystallinity of the sample surface takes place isotropically (see Fig. S4 and S8, ESI†), thus ruling (MOKE probes mainly the film top surface; Efirst tens of nm in out any contribution from the intrinsic crystalline directions in depth) in this 50  50 mm2 region. As seen in Fig. 3(b), in the the anisotropic AFM–FM transition of indented regions (Fig. 2). array of indentations (blue line), the transition occurs at a The evolution of the magnetic contrast of the XMCD images higher temperature and it is less abrupt. Note that the abruptness (SDXMCD; see Methods in ESI†) collected at several temperatures of the transition is reduced in the arrays due to the strain (see Fig. S8, ESI,† for the whole set of images) is plotted as a gradients surrounding each indentation imprint. function of temperature in Fig. 2(g). The figure reveals that Since the AFM-to-FM phase transition in FeRh films is the SDXMCD value, which is indicative of the magnetization, accompanied by an important decrease (increase) of the electric increases at around T* = 90 1C for the region far from the resistance upon heating (cooling),34 we next study the transport indentation (black line), in agreement with the M(T) of the properties of the indented FeRh (arrays with different maximum as-grown films (Fig. 1(b)). On the other hand, the SDXMCD for the indentation loads of 0 mN, 10 mN, 20 mN and 50 mN) using the region near the indentation increases at a higher temperature, contact configuration shown in Fig. 3(c) (see Methods in ESI†). T* = 110 1C (see Video S1, ESI†). Thus, the indentation-induced Similar to the magnetic measurements, the transition can be Fig. 3 (a) Atomic force microscopy image of a 50  50 mm2 region with 10  10 indentations obtained applying 10 mN. (b) MOKE signals normalized to the values at 150 1C, measured at several temperatures, on regions of FeRh without (black squares) and with an array of 50 mN indentations (blue circles). (c) Optical microscopy image including a 50  50 mm2 array of 10  10 indentations (10 mN) and the gold pads used to perform electric transport measurements. I+/ and V+/ indicate the pads where the current is injected and the voltage drop measured, respectively. (d) Temperature dependence of the resistance increment with respect to its value at 127 1C, DR, as function of temperature for a non-indented region and indented regions of 10, 20 and 50 mN. (e) T* evaluated from transport measurements in d at halfway through of the transition (approx. for DR = 2 O) for increasing and decreasing temperatures as a function of the force applied in the indentations. The lines are guides to the eye. This journal is © The Royal Society of Chemistry 2020 Mater. Horiz., 2020, 7, 2056--2062 | 2059 View Article Online Materials Horizons Communication also clearly observed in the resistance dependence on tempera- and C. F. carried out the X-ray characterization experiments. ture (Fig. 3(d)). Interestingly, the normalized resistance curves I. F. wrote the manuscript with input from M. F., E. M., J. N. (DR) at different indentation loads (Fig. 3(d)) reveal a progressive and J. S. increase of T*. This difference is more clearly seen in Fig. 3(e), This article is licensed under a Creative Commons Attribution-NonCommercial 3.0 Unported Licence. where the T* for the cooling and heating branches is plotted as a function of the indentation load. Note that the increase of T* Data and materials availability with indentation load is only about 10 1C (for the 50 mN region), Data and materials are available from the corresponding author whereas the XMCD-PEEM and MOKE characterization revealed upon reasonable request. changes in T* of at least 20 1C. This can be understood by the higher surface sensitivity of XMCD-PEEM and MOKE with respect to transport measurements. This suggests that the Conflicts of interest Open Access Article. Published on 14 May 2020. Downloaded on 11/4/2020 3:23:56 PM. temperature shift is larger at the surface than in the bulk of the There are no conflicts to declare. film, coherent with previous findings of different behaviour of the bulk and the surface of FeRh films.55 Thus, although the FM domains are large (micron size, Fig. 2(d)), they are expected to be Acknowledgements inhomogeneous along the thickness, as shown by Gatel et al.55 Dr Patxi López-Barberá is acknowledged for his assistance with for equivalent FeRh thin films. This inhomogeneity can be easily the MOKE experiments. Vicente Garcı́a-Juez from Real Casa de solved by reducing the film thickness. Note also that the width of la Moneda – Fábrica Nacional de Moneda y Timbre is acknowl- the transition increases in the MOKE experiments, whereas in the edged for his scientific advice. Dr Florencio Sánchez is acknowl- transport experiments it is roughly constant, with a slight edged for the growth of the samples. Dr Bernat Bozzo is decrease. The higher surface sensitivity of MOKE experiments acknowledged for the electric transport characterization. ALBA results in their higher sensitivity to strain gradients produced synchrotron is acknowledged for the provision of beamtime at by indentation, which induce an increase in the width of the the MSPD (proposal number 2017092412) and CIRCE (proposal transition, whereas the transition width variations in transport numbers 2017092462 and 2018022818) beamlines. Financial experiments are not significant and probably masked by small support from the Spanish Ministry of Economy and Competitive- lateral inhomogeneity in the film composition.56 ness, through the ‘‘Severo Ochoa’’ Programme for Centres of Excellence in R&D (SEV-2015-0496 and SEV-2017-0706) and the MAT2017-85232-R, RTI2018-095303-B-C53. MAT2014-56063-C2-1- Conclusions R, MAT2017-86357-C3-1-R (and associated FEDER) and MAT2015- We demonstrate the possibility to locally modify the AFM-to-FM 73839-JIN projects, the Generalitat de Catalunya (2014 SGR 734 transition temperature in FeRh films by means of nano- and 2017 SGR 292), AGAUR (2018 LLAV 00032 and 2019 LLAV indentation. We show that nanoindentation stresses induce 00050) and the European Union’s Horizon 2020 research and compressive strains (mediated by the underlying compression innovation programme under the Marie Skłodowska-Curie grant of the MgO substrate), which cause a shift of T* by at least agreement no. 665919 is acknowledged. This work was supported 20 1C. In addition, we show that the developed methodology by the European Research Council under the SPIN-PORICS 2014- can be used to globally increase T*, across large areas, as Consolidator Grant, Agreement No. 648454 and the MAGIC- evidenced by Kerr magnetometry and transport measurements. SWITCH 2019-Proof of Concept Grant, Agreement No. 875018. Interestingly, this approach can be potentially extrapolated to ICN2 is funded by the CERCA Programme/Generalitat de other metamagnetic materials or, in general, to any system in Catalunya. I. F. acknowledges his RyC contract RYC-2017-22531. which a magnetic order phase transition is accompanied by a E. C. acknowledges the partial financial support from the structural transformation. The T* increase by nanoindentation is, National Science Centre of Poland (NCN) by the PRELUDIUM at first instance, an irreversible process, however, annealing or project UMO-2015/17/N/ST5/01988 and the SONATA Project No. other procedures could, at least partially, recover the initial state. UMO-2016/23/D/ST3/02121. We acknowledge support of the publica- The presented results are relevant for memory devices based tion fee by the CSIC Open Access Publication Support Initiative on the lateral co-existence of contiguous antiferromagnetic- through its Unit of Information Resources for Research (URICI). ferromagnetic regions, or for devices based on AFM materials, such as those where magnetic information must be locally Notes and references stored in an AFM phase. 1 J. I. Martı́n, J. Nogues, K. Liu, J. L. Vicent and I. K. Schuller, J. Magn. Magn. Mater., 2003, 256, 449–501. Author contributions 2 A. Kosiorek, W. Kandulski, H. Glaczynska and M. Giersig, Small, 2005, 1, 439–444. I. F and J. S. planned the study. M. F., A. Q., L. A. and I. F. 3 J. R. Jeong, S. Kim, S. H. Kim, J. Bland, S. C. Shin and carried out the XMCD-PEEM experiments. A. Q., E. M. and I. F. S. M. Yang, Small, 2007, 3, 1529–1533. carried out the MOKE experiments under assistance of J. N. E. 4 J. Fassbender and J. McCord, J. Magn. Magn. Mater., 2008, C. and E. M. carried out the nanoindentations. O. V., C. G.-O. 320, 579–596. 2060 | Mater. Horiz., 2020, 7, 2056--2062 This journal is © The Royal Society of Chemistry 2020 View Article Online Communication Materials Horizons 5 P. Grünberg, J. Magn. Magn. Mater., 2001, 226, 1688–1693. 29 Q. Hu, J. Li, C. Wang, Z. Zhou, Q. Cao, T. Zhou, D. Wang and 6 S. Wolf, D. Awschalom, R. Buhrman, J. Daughton, S. Von Y. Du, Appl. Phys. Lett., 2017, 110, 222408. Molnar, M. Roukes, A. Y. Chtchelkanova and D. Treger, 30 V. Uhlı́ř, J. A. Arregi and E. E. Fullerton, Nat. Commun., 2016, Science, 2001, 294, 1488–1495. 7, 13113. This article is licensed under a Creative Commons Attribution-NonCommercial 3.0 Unported Licence. 7 J. Nogués, J. Sort, V. Langlais, V. Skumryev, S. Suriñach, 31 T. Seifert, U. Martens, S. Günther, M. A. W. Schoen, F. Radu, J. Muñoz and M. Baró, Phys. Rep., 2005, 422, 65–117. X. Z. Chen, I. Lucas, R. Ramos, M. H. Aguirre, P. A. 8 S. S. Parkin, M. Hayashi and L. Thomas, Science, 2008, 320, Algarabel, A. Anadón, H. S. Körner, J. Walowski, C. Back, 190–194. M. R. Ibarra, L. Morellón, E. Saitoh, M. Wolf, C. Song, 9 J.-J. Kim, Y. Choi, S. Suresh and A. Argon, Science, 2002, 295, K. Uchida, M. Münzenberg, I. Radu and T. Kampfrath, Spin, 654–657. 2017, 7, 1740010. 10 J. Fornell, E. Rossinyol, S. Surinach, M. D. Baró, W. H. Li 32 S. Bennett, A. Herklotz, C. Cress, A. Ievlev, C. Rouleau, Open Access Article. Published on 14 May 2020. Downloaded on 11/4/2020 3:23:56 PM. and J. Sort, Scr. Mater., 2010, 62, 13–16. I. Mazin and V. Lauter, Mater. Res. Lett., 2018, 6, 106–112. 11 J. A. Lefever, T. D. Jacobs, Q. Tam, J. L. Hor, Y.-R. Huang, 33 Y. Wang, M. M. Decker, T. N. Meier, X. Chen, C. Song, D. Lee and R. W. Carpick, Nano Lett., 2016, 16, 2455–2462. T. Grünbaum, W. Zhao, J. Zhang, L. Chen and C. H. Back, 12 H. Lu, C.-W. Bark, D. E. De Los Ojos, J. Alcala, C.-B. Eom, Nat. Commun., 2020, 11, 275. G. Catalan and A. Gruverman, Science, 2012, 336, 59–61. 34 L. Lewis, C. Marrows and S. Langridge, J. Phys. D: Appl. 13 A. Gómez, J. M. Vila-Fungueiriño, R. Moalla, G. Saint-Girons, Phys., 2016, 49, 323002. J. Gázquez, M. Varela, R. Bachelet, M. Gich, F. Rivadulla and 35 N. Baranov and E. Barabanova, J. Alloys Compd., 1995, 219, A. Carretero-Genevrier, Small, 2017, 13, 1701614. 139–148. 14 Y. Kim, S. J. Kelly, A. Morozovska, E. K. Rahani, E. Strelcov, 36 W. Lu, J. Fan and B. Yan, Sci. China: Phys., Mech. Astron., E. Eliseev, S. Jesse, M. D. Biegalski, N. Balke and N. Benedek, 2011, 54, 1223–1226. Nano Lett., 2013, 13, 4068–4074. 37 M. Jiang, X. Chen, X. Zhou, B. Cui, Y. Yan, H. Wu, F. Pan 15 J. Sort, A. Concustell, E. Menéndez, S. Suriñach, K. V. Rao, and C. Song, Appl. Phys. Lett., 2016, 108, 202404. S. C. Deevi, M. D. Baró and J. Nogués, Adv. Mater., 2006, 18, 38 R. Wayne, Phys. Rev., 1968, 170, 523. 1717–1720. 39 R. Cherifi, V. Ivanovskaya, L. Phillips, A. Zobelli, I. Infante, 16 J. Sort, A. Concustell, E. Menéndez, S. Suriñach, M. Baró, E. Jacquet, V. Garcia, S. Fusil, P. Briddon and N. Guiblin, J. Farran and J. Nogués, Appl. Phys. Lett., 2006, 89, 032509. Nat. Mater., 2014, 13, 345. 17 J. Sort, L. F. Bonavina, A. Varea, C. Souza, W. J. Botta, 40 Y. Lee, Z. Liu, J. Heron, J. Clarkson, J. Hong, C. Ko, C. S. Kiminami, C. Bolfarini, S. Surinach, M. D. Baró and M. Biegalski, U. Aschauer, S.-L. Hsu and M. Nowakowski, J. Nogués, Small, 2010, 6, 1543–1549. Nat. Commun., 2015, 6, 5959. 18 J. Kouvel and C. Hartelius, J. Appl. Phys., 1962, 33, 1343–1344. 41 Z. Liu, L. Li, Z. Gai, J. Clarkson, S. Hsu, A. Wong, L. Fan, 19 J.-U. Thiele, S. Maat and E. E. Fullerton, Appl. Phys. Lett., M.-W. Lin, C. Rouleau and T. Ward, Phys. Rev. Lett., 2016, 2003, 82, 2859–2861. 116, 97203. 20 J.-U. Thiele, S. Maat, J. L. Robertson and E. E. Fullerton, 42 M. Loving, R. Barua, C. Le Graët, C. Kinane, D. Heiman, IEEE Trans. Magn., 2004, 40, 2537–2542. S. Langridge, C. Marrows and L. Lewis, J. Phys. D: Appl. 21 G. Ju, J. Hohlfeld, B. Bergman, R. J. van de Veerdonk, Phys., 2017, 51, 024003. O. N. Mryasov, J.-Y. Kim, X. Wu, D. Weller and 43 J. Clarkson, I. Fina, Z. Liu, Y. Lee, J. Kim, C. Frontera, B. Koopmans, Phys. Rev. Lett., 2004, 93, 197403. K. Cordero, S. Wisotzki, F. Sanchez and J. Sort, Sci. Rep., 22 S. Maat, J.-U. Thiele and E. E. Fullerton, Phys. Rev. B: 2017, 7, 15460. Condens. Matter Mater. Phys., 2005, 72, 214432. 44 I. Fina, A. Quintana, J. Padilla-Pantoja, X. Marti, F. Macia, 23 E. E. Fullerton, S. Maat and J.-U. Thiele, US Pat., 7372116 B2, F. Sanchez, M. Foerster, L. Aballe, J. Fontcuberta and J. Sort, 2008. ACS Appl. Mater. Interfaces, 2017, 9, 15577–15582. 24 X. Chen, J. Feng, Z. Wang, J. Zhang, X. Zhong, C. Song, L. Jin, 45 I. Fina, A. Quintana, X. Marti, F. Sanchez, M. Foerster, B. Zhang, F. Li and M. Jiang, Nat. Commun., 2017, 8, 449. L. Aballe, J. Sort and J. Fontcuberta, Appl. Phys. Lett., 2018, 25 X. Marti, I. Fina, C. Frontera, J. Liu, P. Wadley, Q. He, 113, 152901. R. J. Paull, J. D. Clarkson, J. Kudrnovsky, I. Turek, J. Kunes, 46 I. Fina and J. Fontcuberta, J. Phys. D: Appl. Phys., 2019, D. Yi, J.-H. Chu, C. T. Nelson, L. You, E. Arenholz, 53, 023002. S. Salahuddin, J. Fontcuberta, T. Jungwirth and R. Ramesh, 47 A. Zakharov, A. Kadomtseva, R. Levitin and E. Ponyatovskii, Nat. Mater., 2014, 13, 367–374. Sov. Phys. JETP, 1964, 19, 1348–1353. 26 T. Moriyama, N. Matsuzaki, K.-J. Kim, I. Suzuki, 48 F. Fauth, I. Peral, C. Popescu and M. Knapp, Powder Diffr., T. Taniyama and T. Ono, Appl. Phys. Lett., 2015, 107, 122403. 2013, 28, S360–S370. 27 I. Fina, N. Dix, E. Menéndez, A. Crespi, M. Foerster, L. Aballe, 49 O. Vallcorba and J. Rius, J. Appl. Crystallogr., 2019, 52, 478. F. Sánchez and J. Fontcuberta, ACS Appl. Mater. Interfaces, 50 L. Aballe, M. Foerster, E. Pellegrin, J. Nicolas and S. Ferrer, 2020, 12, 15389–15395. J. Synchrotron Radiat., 2015, 22, 745–752. 28 Y. Liu, L. C. Phillips, R. Mattana, M. Bibes, A. Barthélémy 51 J. Amodeo, S. Merkel, C. Tromas, P. Carrez, S. Korte-Kerzel, and B. Dkhil, Nat. Commun., 2016, 7, 11614. P. Cordier and J. Chevalier, Crystals, 2018, 8, 240. This journal is © The Royal Society of Chemistry 2020 Mater. Horiz., 2020, 7, 2056--2062 | 2061 View Article Online Materials Horizons Communication 52 T. Nazarova and M. Nazarov, Philos. Mag. A, 1996, 74, 1311–1318. 55 C. Gatel, B. Warot-Fonrose, N. Biziere, L. Rodrı́guez, 53 U. Aschauer, R. Braddell, S. A. Brechbühl, P. M. Derlet and D. Reyes, R. Cours, M. Castiella and M.-J. Casanove, Nat. N. A. Spaldin, Phys. Rev. B: Condens. Matter Mater. Phys., Commun., 2017, 8, 15703. 2016, 94, 014109. 56 V. Saidl, M. Brajer, L. Horak, H. Reichlova, K. Vyborny, This article is licensed under a Creative Commons Attribution-NonCommercial 3.0 Unported Licence. 54 E. Coy, L. Yate, Z. Kabacińska, M. Jancelewicz, S. Jurga and M. Veis, T. Janda, F. Trojanek, M. Marysko, I. Fina, X. Marti, I. Iatsunskyi, Mater. Des., 2016, 111, 584–591. T. Jungwirth and P. Nemec, New J. Phys., 2016, 18, 083017. Open Access Article. Published on 14 May 2020. Downloaded on 11/4/2020 3:23:56 PM. 2062 | Mater. Horiz., 2020, 7, 2056--2062 This journal is © The Royal Society of Chemistry 2020