Perpendicular Local Magnetization Under Voltage Control in Ni Films on Ferroelectric BaTiO3 Substrates

High-resolution magnetoelectric imaging is used to demonstrate electrical control of the perpendicular local magnetization associated with 125 nm-wide magnetic stripe domains in 100-nm-thick Ni films. This magnetoelectric coupling is achieved in zero magnetic field using strain from ferroelectric BaTiO3 substrates to control perpendicular anisotropy imposed by the growth stress. These findings may be exploited for perpendicular recording in nanopatterned hybrid media.


DOI: 10.1002/adma.201404799
that is large enough to compete successfully with strong OOP demagnetizing fi elds, while arranging for this anisotropy to be signifi cantly modifi ed by an electric fi eld. Here we create and destroy an OOP anisotropy in large contiguous areas of Ni fi lms, by repeatedly overcoming the growth stress with electrically controlled strain from ferroelectric domain switching in BaTiO 3 (BTO) substrates. This OOP anisotropy produces an OOP component of magnetization, which is manifested via the formation of magnetic stripe domains, and turned on and off via the applied voltage. To confi rm our interpretation, we show that the same magnetic switching is achieved when quantitatively similar changes of strain are produced by thermally driving phase transitions in the BTO substrates.
Stripe domains [26][27][28][29] arise in magnetic thin fi lms when a uniaxial OOP anisotropy competes with the IP shape anisotropy, yielding a canted local magnetization whose OOP component alternates in sign along an IP direction, minimizing stray-fi eld energy. The IP component of magnetization lies perpendicular to this IP direction, along the same direction in adjacent stripes. In polycrystalline fi lms of negative-magnetostriction Ni, which show stripe domains above a critical thickness, [ 26,28 ] the uniaxial OOP anisotropy arises due to isotropic IP tensile stress [ 30 ] generated at grain boundaries during Volmer-Weber growth. [ 31 ] This builtin stress does not come from the substrate, as stripe domains appear in free-standing polycrystalline fi lms, [ 29 ] and substrate choice does not infl uence stripe width (Supporting Information Note 1). Therefore stripe domains should be ubiquitous in our as-grown samples, even though our BTO substrates are twinned.
In our as-prepared Ni//BTO samples, stripe domains were as expected nominally ubiquitous in all fi lms, and magnetoelectric effects were similar whether substrates were poled after (Samples 1-4) or before (Samples 5-6) fi lm growth. We show that stripe domains are erased by cooling BTO through structural phase transitions, or by electrically moving 90° ferroelectric walls to convert a to c domains. Our high-resolution imaging of electrically driven changes reveals that ferromagnetic and Electrical control of magnetism has been demonstrated in multiferroic compounds [1][2][3][4][5] and ferromagnetic semiconductors, [6][7][8][9] but electrical switching of a substantial net magnetization at room temperature has not been demonstrated in these materials. This scientifi c and technological goal has instead been achieved using ferromagnetic fi lms in which electrically driven magnetic changes arise due to strain [10][11][12][13][14][15][16] or exchange bias [ 17,18 ] from ferroic substrates, or changes of interfacial orbital occupation. [ 19 ] However, in these works and other magnetoelectric studies, it is typically necessary to employ a variable magnetic fi eld, either to repeatedly drive electrically nucleated magnetization reversal, [ 12 ] or to interrogate changes of anisotropy, [ 16,19,20 ] exchange bias, [ 21 ] or coercivity. [ 12,14,22 ] A variable magnetic fi eld permits magnetoelectric phenomena to be investigated comprehensively, but electrical control of magnetization with no magnetic fi eld represents a higher scientifi c goal. Purely electrical switching of in-plane (IP) magnetization has been demonstrated with thin-fi lm ferromagnets, [11][12][13]23 ] but electrical switching of out-of-plane (OOP) magnetization in zero applied magnetic fi eld has only been demonstrated at a subset of small features, in high-quality epitaxial nanocomposites [ 24 ] and polycrystalline heterostructures. [ 25 ] Indeed, it is challenging to achieve an OOP anisotropy ferroelectric domains are correlated, and that stripe domains undergo non-volatile repeatable erasure.
For demagnetized Ni//BTO in non-saturating magnetic fi elds, thermally driven structural phase transitions in the substrate yield sharp hysteretic jumps of OOP magnetization that complement expected [ 10,14,32 ] jumps of IP magnetization ( Figure 1 a). On cooling, these jumps represent an increase of IP magnetization at the expense of OOP magnetization, consistent with two observations. First, there is a dramatic increase of easy-axis loop squareness that evidences the development of a strong IP magnetic anisotropy (Figure 1 b) of biaxial character (Supporting Information Note 2). Second, the hard-axis saturation fi eld increases to a value that corresponds closely to saturation magnetization µ 0 M s = 0.54 T, and thus the demagnetizing fi eld (inset, Figure 1 b), evidencing complete suppression of weak OOP anisotropy. After returning to room temperature in non-saturating magnetic fi elds, IP and OOP magnetizations are slightly modifi ed due to the stochastic nature of magnetic domain-wall depinning (Figure 1 a), but easy and hard-axis loops are recovered (Figure 1 b and inset). Therefore strains on cooling to R-phase BTO appear to be reversible.
Magnetic imaging was performed using photoemission electron microscopy (PEEM) and magnetic force microscopy (MFM). PEEM data were obtained with X-ray magnetic circular dichroism (XMCD) contrast, so images depend on the projection of local magnetization onto the grazing-incidence X-ray beam direction. Therefore the IP and OOP components of magnetization are imaged simultaneously to provide 3D information, and the alternating OOP component of magnetization in stripe domains [26][27][28][29] is apparent for any IP sample orientation.
XMCD-PEEM ( Figure 2 a) revealed stripe domains with a stripe width of ≈125 nm, comparable with our 100 nm fi lm thickness as expected. [ 28 ] As also expected, these stripe domains were nominally ubiquitous in all of our many PEEM/MFM images of as-grown Samples 1-4, fabricated with electrically virgin BTO substrates that primarily comprised a domains with IP polarizations at 90° to one another ( a 1a 2 domains [ 13 ] ). Long thin features with IP magnetizations were observed in Sample 1, and could represent an initial decoration [ 13 ] of c domains with OOP polarizations (Supporting Information Note 3), but using MFM this could not be confi rmed as they were not reproduced after prepoling BTO substrates to produce a signifi cant population of c domains (Samples 5-6). In view of this irreproducibility, the long-thin features are considered no further, but the example in Figure 2 a is helpful because the antiparallel tailto-tail domains set the XMCD asymmetry scale for Figure 2 c.
The Figure 2 a stripe orientation lies approximately perpendicular to the IP projection of the beam, and there is waviness due to inhomogeneous IP stress. Therefore the IP component of magnetization along the local stripe direction [26][27][28][29] has typically no signifi cant projection onto beam direction. On rotating the sample 90° (Figure 2 b), the same fi eld of view reveals that the locally inhomogeneous IP component of magnetization is split into a few large domains that coexist with the alternating OOP component of magnetization, which remains visible.
On cooling to the R-phase of the substrate, stripe domains are annihilated in favor of a uniform IP magnetization within each domain, now clearly visible with slightly modifi ed boundaries (Figure 2 c). This annihilation is due to suppression of the OOP anisotropy responsible for stripe domains, consistent with our macroscopic OOP data (inset, Figure 1 b). The resulting IP domains are antiparallel, evidencing the development of a uniaxial IP anisotropy in our fi eld of view, which thus contains only one of the two twin species responsible for the biaxial anisotropy (Supporting Information Note 2).
Subsequent warming to the O-phase of the substrate nucleates stripes at IP domain walls (Figure 2 d). Returning to the room temperature T-phase recovers stripe domains that are less wavy than before (Figure 2 e). The low-temperature excursion has rendered the IP fi lm stress more homogenous [ 35 ] (image quality in Figure 2 e was compromised by thermal drift, but is restored in a subsequent MFM image, Supporting Information Note 4). A high-temperature excursion to the cubic C-phase of BTO is uninteresting because it irreversibly erases stripe domains (Supporting Information Note 5).
To investigate the electrically driven magnetic changes microscopically, we used PEEM with XMCD contrast to image the Ni fi lm near a zig-zag edge (see Experimental Section), and we used PEEM with X-ray linear dichroism (XLD) contrast [ 36 ] to image nearby ferroelectric domains in the exposed BTO substrate ( Figure 3 ). The defi nition of an edge was necessary because ferroelectric domains cannot be imaged through our 100-nm-thick fi lms using either PEEM, birefringence, or piezoforce microscopy (PFM); and because thinner fi lms show no stripe domains (Supporting Information Note 6).
For this magnetoelectric imaging, a room-temperature electrical cycle (Supporting Information Note 7) produced our Adv. Mater. 2015, 27, 1460-1465 www.advmat.de www.MaterialsViews.com   . Therefore we infer that the electrical annihilation of stripe domains in favor of IP domains is due to the conversion of underlying a to c domains. A widespread MFM investigation of Samples 1-4 was suffi cient to reveal many regions where stripe domains could be electrically switched off and on in a non-volatile manner (e.g., Figure 4 ). Unlike our thermally driven magnetic changes (Figure 2 ), these magnetoelectric effects are not ubiquitous because 90° switching in T-phase BTO is somewhat inhibited by stress from other domains. [ 34 ] Therefore, macroscopic magnetoelectric effects (discussed in Supporting Information Note 9) are localized and small. We next use an essentially qualitative analysis, then a quantitative analysis, to show that thermally/electrically annihilated stripe domains lie on T-phase a domains. We will see that thermal/electrical annihilations both arise due to local IP strains that are nominally uniaxial, compressive, and similar in magnitude (≈1%). The resulting uniaxial stress eliminates the OOP anisotropy created by isotropic stress provided its magnitude is equal or larger (Supporting Information Note 10).
We fi rst use a qualitative argument to deduce that thermally annihilated stripe domains lie on T-phase a domains. On cooling BTO from T to R, lattice parameter a T 297K increases 0.1%, and lattice parameter c T 297K reduces 0.97%. For an × a c template of a domains, cooling to 127 K produces a nominally uniaxial ≈1% contraction, annihilating the OOP stress anisotropy (Supporting Information Note 10) (the opposing ≈0.17% thermal contraction [ 37 ] from Ni may be ignored). Therefore stripe domains that undergo thermal annihilation lie on T-phase a domains whose twinning results in biaxial anisotropy (Supporting Information Note 2).
An essentially qualitative argument shows that electrically annihilated stripe domains also lie on T-phase a domains, as 90° switching interconverts a domains (a c We now present a quantitative analysis to explain our experimental observations of stripe-domain annihilation. The tensile isotropic IP growth stress 0 p σ > may be evaluated from uniaxial OOP anisotropy [ 29 ] K 3 2 z s p λ σ = − (saturation magnetostriction s λ = −32.9 × 10 −6 for room-temperature Ni). For our 125 nm-wide stripes, the micromagnetic theory of stripe domains implies K z 297K = 22 kJ m −3 (Supporting Information Note 11). Alternatively, variable-temperature IP and OOP measurements of macroscopic magnetization along virgin curves imply K z 297K = 24 kJ m −3 (Supporting Information Note 12). These values of K z 297K are similar, and in the 10-60 kJ m −3 range for Ni fi lms. [ 26,27 ] The corresponding value of K 2 3 p zs 1 σ λ = − − ≈0.5 GPa is large, but typical of the growth stress in polycrystalline metallic fi lms, [ 31 ] and similar to the value required to create stripe domains in a Ni-rich fi lm. [ 38 ] Even if the predicted ≈1% contractions are not fully developed in the twinned substrate or fully transmitted to the fi lm, this 0.5 GPa growth stress is indeed small enough to be overcome by our electrically/thermally generated stresses, as the electrically driven ≈1% strain corresponds to a stress of 1.3 GPa, and the thermally driven ≈1% strain corresponds to an even larger stress as the 133 GPa Young's modulus of Ni increases on cooling. (Reversible piezoelectric strain at 200 V is two orders of magnitude smaller and may be ignored, Supporting Information Note 13.) In summary, we have demonstrated systematic electrical and thermal control of OOP magnetization, using growth stress to create the requisite OOP anisotropy, and using stress associated with discontinuous changes of lattice parameter to annihilate this anisotropy. Our magnetoelectric imaging is attractive as images of both ferromagnetic and ferroelectric domains are rare; [ 13,18 ] concomitant changes have only been seen in lowresolution studies; [ 13 ] and published PEEM data have hitherto exploited XMCD contrast without XLD contrast [ 17,18,39 ] except for one single image. [ 36 ] Our high-resolution magnetoelectric measurements are suitable for nanostructures that are too small for magneto-optical Kerr effect and birefringence. [ 13 ] Moreover, our approach is preferable to scanning force microscopy with both [ 24 ] MFM and PFM, because the magnetic information is more direct and complete, and because the two imaging modes may be trivially swapped at any temperature.
Adv. Mater. 2015, 27, 1460-1465 www.advmat.de www.MaterialsViews.com In future, one may use nanopatterned bilayers on inactive substrates for non-volatile voltage control of magnetization at a precise array of locations. If this were achieved by exploiting fi lms patterned down to, for example, the stripe width, then one may electrically switch OOP magnetization in bits that each comprise a single magnetic domain, notably for low-power write in high-density perpendicular data storage.

Experimental Section
Sample Preparation and Structural Charaterization: All six samples comprised 100-nm-thick Ni fi lms grown on 4 mm × 4 mm × 0.5 mm singlecrystal BTO (001) substrates using room-temperature e-beam assisted evaporation with a base pressure of 1.5 × 10 −10 mbar. All fi lms were capped with 4 nm of Cu to prevent oxidation. The Ni deposition rate was ≈0.3 nm min −1 , as determined using a quartz microbalance. Films were either co-deposited (Samples 1 and 2; Samples 5 and 6) or deposited separately (Samples 3 and 4). For Samples 1-4, we used unpoled BTO substrates. For Samples 5 and 6, which were prepared for both the initial decoration study and the magnetoelectric imaging, we poled substrates prior to growth by applying 300 V (6 kV cm −1 ) between a sputter-deposited back electrode of Pt and a sacrifi cial top electrode of 8-nm-thick evaporated Au. The poling process was monitored through this layer of Au by optical birefringence. After etching away the Au, small areas of the BTO surface were protected during Ni deposition using a marker pen. Lift-off in isopropyl alcohol subsequently exposed these small areas, which were bounded by fi lm edges that were sometimes zig-zag. Using a Philips powder diffractometer with Cu K α radiation to study Sample 1, we found a lateral Ni grain size of over 100 nm, and no evidence for preferred orientations.
Magnetic Measurements: Magnetization measurements were performed using a Princeton Measurements Corporation vibrating sample magnetometer (VSM), with electrical access to the sample. [ 40 ] Measurements of macroscopic IP magnetization were performed after initially rotating the sample to maximize remanence.
Magnetic Force Microscopy: MFM measurements were performed in a Digital Instruments Dimension 3100, at lift heights of 40-60 nm, using low-moment ASYMFMLM Asylum Research tips of stiffness 2 N m −1 coated with 15 nm of CoCr. All MFM images were obtained at magnetic remanence after having applied an OOP magnetic fi eld, except for Figure 4b, Supporting Information Note 3, which was obtained while an IP magnetic fi eld was applied using a tapered soft magnet connected to a larger permanent magnet. All MFM image analysis was performed using WSxM software. [ 41 ] Photoemission Electron Microscopy: PEEM in zero magnetic fi eld was performed with the X-ray beam at a grazing-incidence angle of 16°, using an Elmitec SPELEEM-III microscope on beamline I06 at Diamond Light Source. The probe depth was ≈7 nm, and the lateral resolution was typically ≈50 nm. Images of ferromagnetic (ferroelectric) domains without topographical information were obtained by plotting average values of XMCD (XLD) asymmetry. These values represent the projection onto the incident-beam direction of the local surface magnetization (polarization), whose IP and OOP components are both detected. The comparison between intensities obtained on and off resonance avoids the infl uence of any inhomogeneous illumination. Images for each X-ray energy and beam polarization were acquired during 5 s exposure times. Each XMCD-PEEM image that we present was constructed via an averaging process based on 20 such images, but XMCD-PEEM and XLD-PEEM images that we present together were each constructed via an averaging process based on 40 such images.
Voltage V was applied during VSM, PEEM, and MFM measurements between the grounded Ni fi lm, and a sputter-deposited Pt electrode under the BTO substrate.

Supporting Information
Supporting Information is available from the Wiley Online Library or from the author.