Atomic-scale determination of spontaneous magneticreversal in oxide heterostructuresM. Saghayezhiana,1, Summayya Kouserb,c,d,1, Zhen Wanga,e, Hangwen Guoa, Rongying Jina, Jiandi Zhanga, Yimei Zhue,Sokrates T. Pantelidesb,c,d, and E. W. Plummera,2

aDepartment of Physics and Astronomy, Louisiana State University, Baton Rouge, LA 70803; bDepartment of Physics and Astronomy, Vanderbilt University,Nashville, TN 37235; cDepartment of Electrical Engineering and Computer Science, Vanderbilt University, Nashville, TN 37235; dMaterials Science andTechnology Division, Oak Ridge National Laboratory, Oak Ridge, TN 37830; and eCondensed Matter Physics and Materials Science Department, BrookhavenNational Laboratory, Upton, NY 11973

Contributed by E. W. Plummer, April 4, 2019 (sent for review November 15, 2018; reviewed by Jian Liu, Xiaoqing Pan, and David Vanderbilt)

Interfaces between transition metal oxides are known to exhibitemerging electronic and magnetic properties. Here we report in-triguing magnetic phenomena for La2/3Sr1/3MnO3 films on an SrTiO3

(001) substrate (LSMO/STO), where the interface governs the mac-roscopic properties of the entire monolithic thin film. The interface ischaracterized on the atomic level utilizing scanning transmissionelectron microscopy and electron energy loss spectroscopy (STEM-EELS), and density functional theory (DFT) is employed to elucidatethe physics. STEM-EELS reveals mixed interfacial stoichiometry, sub-tle lattice distortions, and oxidation-state changes. Magnetic mea-surements combined with DFT calculations demonstrate that aunique form of antiferromagnetic exchange coupling appears atthe interface, generating a novel exchange spring-type interactionthat results in a remarkable spontaneous magnetic reversal of theentire ferromagnetic film, and an inverted magnetic hysteresis, per-sisting above room temperature. Formal oxidation states derivedfrom electron spectroscopy data expose the fact that interfacial ox-idation states are not consistent with nominal charge counting. Thepresent work demonstrates the necessity of atomically resolved elec-tron microscopy and spectroscopy for interface studies. Theory dem-onstrates that interfacial nonstoichiometry is an essential ingredient,responsible for the observed physical properties. The DFT-calculatedelectrostatic potential is flat in both the LSMO and STO sides (nointernal electric field) for both Sr-rich and stoichiometric interfaces,while the DFT-calculated charge density reveals no charge transfer/accumulation at the interface, indicating that oxidation-state changesdo not necessarily reflect charge transfer and that the concept ofpolar mismatch is not applicable in metal−insulator polar−nonpolarinterfaces.

magnetism | thin films | electron microscopy | oxide interfaces | densityfunctional theory

Transition metal oxides (TMOs) exhibit a wide range ofelectrical, magnetic, and optical properties largely because

they can be easily alloyed and the 10 slots in the transition metalatom d orbitals allow for very diverse spin arrangements. Addingstrong electron−lattice coupling and easily formed oxygenvacancies to this mix, the result is a large number of coupleddegrees of freedom. Therefore, it is not surprising that TMOheterostructures exhibit highly unusual behavior, induced by in-terfaces between different oxides where symmetry discontinuitiesoccur, leading to properties that are absent in bulk (1–5). A well-known example is the LaAlO3/SrTiO3 (001) heterostructure,where a 2D electron gas with high mobility forms at the interfacebetween two insulating oxides (6), including the appearance ofsuperconductivity (7) and background ferromagnetic (FM) or-dering (8). In many cases, the symmetry discontinuity at the in-terface strongly modifies the transition metal−oxygen octahedranetwork, inducing changes in local structure (bond geometry)and local stoichiometry. These changes have been shown to re-sult in exotic magnetic properties such as antiferromagnets (AFM)built from FM layers (9), interface-driven magnetic phases absent

in bulk (5), antiparallel spin alignment at manganite−ruthenateinterfaces (10), and fine control over magnetic anisotropy by in-terfacing manganites and iridates (11). Advances in epitaxial syn-thesis provide a fine control over interfaces that is necessary torealize novel magnetic systems hitherto unseen in TMOs, such asspontaneous magnetic reversal (SMR) and exchange spring (12–14).La2/3Sr1/3MnO3 is a material with interesting magnetic prop-

erties, such as high Curie temperature (15) and nearly perfectspin polarization (16), used for the realization of magnetictunnel junctions (17), magnetoelectric devices (18), and spininjection into cuprate superconductors (19) as well as a plat-form for studying spin-dependent transport in organic materials(20). In this material, like other oxides, the interface-inducedmagnetic properties are usually discussed in terms of chargetransfer (21–24), while the atomic-scale role of interface struc-ture and intermixture are not explicitly investigated (22, 25). TheLa2/3Sr1/3MnO3/SrTiO3 (001) (LSMO/STO) interface is composedof a stacking sequence SrO/TiO2−La2/3Sr1/3/MnO2/La2/3Sr1/3 andso on (26). By viewing formal oxidation states as physical chargeson the atomic layers, it has been concluded that interface inter-mixing in LSMO/STO is driven by polar mismatch, thus degradingmagnetic and electrical properties near the interface (27–30).However, the Thomas−Fermi screening length in LSMO is just

Significance

Transition metal oxide interfaces have shown extraordinarypromise in the quest to design materials with custom elec-tronic, magnetic, and optical properties. In rare cases, inter-faces exhibit electronic and magnetic properties that areradically different from those of the components. An exampleis the emergent two-dimensional electron gas between twoinsulators. In this work, we show that the interface of a non-magnetic oxide substrate and a ferromagnetic metallic thinfilm possesses a local antiferromagnetic coupling which con-trols the reversal of the entire film’s ferromagnetic ordering.Electron microscopy and quantum calculations elucidate theatomic-scale origin of the observed phenomena, and demon-strate that local nonstoichiometry and structure are the keyfactors in interfacial phenomena.

Author contributions: M.S., H.G., J.Z., and E.W.P. designed research; M.S., S.K., Z.W., H.G.,R.J., J.Z., Y.Z., and S.T.P. performed research; M.S., S.K., Z.W., H.G., R.J., J.Z., Y.Z., S.T.P.,and E.W.P. analyzed data; and M.S., S.K., Z.W., H.G., R.J., J.Z., Y.Z., S.T.P., and E.W.P. wrotethe paper.

Reviewers: J.L., University of Tennessee; X.P., University of California, Irvine; and D.V.,Rutgers, The State University of New Jersey.

The authors declare no conflict of interest.

Published under the PNAS license.1M.S. and S.K. contributed equally to this work.2To whom correspondence should be addressed. Email: wplummer@phys.lsu.edu.

This article contains supporting information online at www.pnas.org/lookup/suppl/doi:10.1073/pnas.1819570116/-/DCSupplemental.

Published online May 8, 2019.

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0.31 nm (smaller than a unit cell) (31). The short screening lengthis not consistent with a buildup of an electrostatic potential andthus disagrees with the notion of electrostatic potential divergence(28). Contradictory past studies question the validity of divergenceof energy at polar-mismatched metal−insulator interfaces (27).On the other hand, lack of atomic-scale studies limits the trans-ferability of experimental findings and does not provide usefulinput for or direct comparison with first-principles calculations.For example, Lee et al. (32) and Chen et al. (33) reported “hiddenmagnetic properties” in LSMO/STO heterostructures. The formerwork proposes mixed Mn3+/Mn4+ oxidation state, and possibleelectronic reconstruction at the interface as the reason behind thephenomena, while the latter relates unexpected magnetic proper-ties to the presence of an interfacial AFM-coupled pinned magneticlayer. In the absence of atomically resolved structure, stoichiometry,and oxidation state information, the data could only be discussed interms of phenomenological models and cannot be evaluated the-oretically by density functional theory (DFT) calculations.In this paper, we report the intricate interplay between mag-

netic properties, structure, and chemical composition at the in-terface of LSMO/STO (001) heterostructures, revealingfascinating interface-induced magnetic behavior such as SMRand inverted hysteresis (IH), that are persistent above roomtemperature (TC = 337 K). Such behavior resembles exchangespring interactions which, in the past, have never been observedin monolithic films (2, 14). Atomically resolved electron mi-croscopy and spectroscopy reveal two Sr-rich interface atomiclayers in LSMO, and significant changes in the interfacialstructure (octahedral tilt and elongation of lattice constant) andin formal oxidation state at the interface. The formal oxidationstate of the Mn at the interface is not consistent with what isexpected from the stoichiometry (La/Sr). The above data createa platform for DFT calculations that can be used to explainquantitatively the magnetic interactions and suggest future di-rections. DFT calculations show that presence of strain Sr-richinterface and an octahedral tilt gradient are both essential toobtain the interfacial magnetic orderings that can account for theobserved SMR. Furthermore, DFT results attribute the un-expected observed out-of-plane (OOP) lattice expansion to thepresence of O vacancies, an all too common occurrence in ox-ides. The DFT-calculated electrostatic potential is flat in boththe LSMO and STO sides (no internal electric field), while theDFT-calculated charge density reveals no charge transfer at theinterface. These results, including calculations of a stoichio-metric interface, indicate that models based on interfacial polarmismatch (27) are not applicable to metal−insulator hetero-structures and that oxidation state changes reflect only localorbital rehybridization. Our study exemplifies a methodologywhere combining atomically monitored growth, macroscopicmeasurements, atomically resolved electron microscopy andspectroscopy, and DFT calculations successfully realize and un-ravel the intriguing physical properties. Given the important roleof atomic-scale structure and composition in magnetic proper-ties, we begin with the structural characterization and then themagnetic properties. Following that, we will present our first-principles calculations and underlying physics.

Atomic-Scale Structural Data and Formal Oxidation StatesEpitaxial thin films of LSMO/STO (001) were grown usingpulsed laser deposition (see SI Appendix for details). The struc-ture and composition across the interface were measured usingboth high-angle annular dark field (HAADF) with intensity pro-portional to atomic number (Z contrast) and annular bright field(ABF) imaging, which is sensitive to light elements such as oxygen.The HAADF and ABF scanning transmission electron microscopy(STEM) images from a 50-unit cell thin film are presented in Fig.1 A and B, respectively. The data show that, in this metallic thinfilm, there are no ferroelectric-like displacements in either of the

two materials, as opposed to ultrathin insulating LSMO wherestrong ferroelectric-like displacements have been observed (34). TheABF−STEM image displays a zig-zag arrangement of Mn−O−Mnand Ti−O−Ti chains, due to tilts of the octahedra. The oxygenoctahedral tilt (OOT) angle, namely the rotation of the octahedronabout the [1−10]c direction (SI Appendix, Fig. S4), is shown in Fig.1C as a function of distance from the interface. Near the interface,the first unit cell of STO shows a small tilt that vanishes, recoveringan a0a0a0 bulk tilt pattern. To preserve the continuity of thestructure across the interface (tilt−nontilt interface), the first andsecond unit cells of LSMO (yellow) show a suppressed OOT angle,with an average value of 4.5° ±0.5° and 5.6° ±0.5°, respectively. Afterthe second unit cell, LSMO exhibits an OOT value of 6.1° ±0.5°,close to the bulk value 6.8° (35). The ABF images clearly indicatethat the LSMO system exhibits an a−a−c0 tilt and rotation, agreeingwith X-ray diffraction (XRD) measurements (SI Appendix, Figs.S5 and S6).Fig. 1D shows the OOP lattice constant as a function of distance

from the interface. The STO lattice constant exhibits the bulkvalue, while LSMO has an enlarged spacing near the interface. Thefirst and second unit cells near the interface, shaded in yellow, showan elongation along the c axis, 3.93± 0.05 Å and 3.92± 0.05 Å,respectively, compared with the average XRD value of3.848± 0.003 Å. To ensure that the difference in the OOP latticeconstants near the interface is statistically meaningful, themeasurements were performed in several different places alongthe interface, all exhibiting the same value. Away from the in-terface, the OOP lattice parameter recovers the average XRDvalue (36–38). XRD reciprocal space mapping (SI Appendix, Fig.S5) shows that the thin film is fully strained and the in-planelattice constant retains the same value as STO (3.905 Å), which islarger than the unstrained value of bulk LSMO (3.87 Å). Therefore,under tensile strain, it is surprising that the OOP lattice constant atthe interface is larger than in the film.The stoichiometric variation of each element across the in-

terface was measured using atomically resolved electron energyloss spectroscopy (EELS) and is shown in Fig. 1E for Sr and Laconcentrations of each atomic layer. In the first and secondatomic layers of LSMO (Fig. 1A), the concentration of the A site(La/Sr) is La0.40Sr0.60 and La0.55Sr0.45, respectively. From thethird atomic layer (second unit cell) and beyond, the concen-tration of LaSr recovers its bulk value, i.e., La0.67Sr0.33. In otherwords, two interfacial LSMO layers are Sr-rich, while the in-terfacial STO layers remain stoichiometric. In view of the factthat increased Sr concentration in bulk LSMO leads to a smallerc-axis lattice constant (39), one would again expect a smallerlattice constant in the interfacial region, which is contrary towhat we observed (Fig. 1D).Fig. 1F shows the L2,3 ratio of Mn atoms as extracted from the

L2 and L3 EELS peak intensities in the EELS spectra, which isroutinely correlated with the Mn formal oxidation state (SI Ap-pendix, Fig. S7) (40). The measured L2,3 intensity ratio in the firstunit cell shows a dip, which corresponds to a reduction of the Mnformal oxidation state from 3.2 to 3. However, according to formaloxidation states, if La, Sr, and O are kept at their usual values (+3,+2, and −2), the increase in Sr concentration in the first unit cell(La0.475Sr0.525MnO3) should increase the Mn oxidation state to3.5. Changes in oxidation states at the interface are typicallyinterpreted as reflecting charge transfer to or from the interface(41). These issues will be taken up later in DFT Calculations andExplanation of the Structural and Electronic Data where we presenta qualitative discussion based on the results of DFT calculations.

Magnetic CharacterizationThe magnetic properties of the thin films are measured using asuperconducting quantum interference device. The LSMO/STOsample is first heated to 380 K, well above the Curie temperature,

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then cooled to 2 K in the presence of a 1,000-Oe in-plane mag-netic field, which aligns all of the magnetic moments in the di-rection of the external magnetic field. The remanent magnetizationM(T) is then measured under zero-field warming. Fig. 2A showsthat, as expected, the magnetization is positive at 2 K, the samedirection as the cooling field. Upon warming, the magnetizationdecreases slowly, then rapidly at ∼60 K, exhibiting a sharp SMR at∼75 K. SMR occurs in LSMO films with different thicknesses (SIAppendix, Fig. S8) but disappears when a small magnetic field isapplied during warming (SI Appendix, Fig. S9). The observed SMRis very similar to exchange spring systems, where a soft FM ma-terial exchange couples to a hard FM or AFM material, occurringbecause of interfacial coupling in multimaterial magnetic hetero-structures (2, 13, 42–44). Here we observe that this phenomenoncan occur in a monolithic thin film, obviating the need for multi-material magnetic heterostructuring (2).

Using M(T) and AC susceptibility, we can explore the origin ofthe observed SMR with concepts developed in multimaterialheterostructures (2, 13, 42–44). Given the observed changes inthe structure and the oxidation state of Mn near the interface,the unexpected magnetic interaction is undoubtedly a productof the interface. Since the first unit cell has a different structureand stoichiometry, its unidirectional anisotropy should be dif-ferent from the rest. One way to drive the observed SMR isthrough AFM magnetic coupling between the bulk of the LSMOfilm, treated as a soft magnet, and the interfacial layers, treatedas a hard magnet. After the magnetic field is turned off at 2 K, allspins are initially aligned by the external magnetic field throughcooling as illustrated in Fig. 2A, Inset, marked as the 0th state. Apossible scenario for the SMR is as follows: With increasingtemperature, the interfacial AFM coupling prevails and begins toovercome the magnetic anisotropy barrier, resulting in spin flip.

Fig. 1. (A and B) HAADF and ABF STEM images of LSMO/STO along the [1−10] direction. The red arrow marks the interface. Inset shows the octahedral tiltaway from the interface. (C) Octahedral tilt angle as function of distance from the interface. The yellow shaded region shows the zone suppression of oc-tahedral tilt near the interface. (D) OOP lattice constant as a function of distance from the interface. (E) Layer-by-layer La and Sr concentration as a functionof distance from the interface. La and Sr reach their stoichiometric concentration, i.e., La0.67Sr0.33 after the third atomic column from the interface. (F) L2,3ratio and Mn nominal oxidation state as a function of distance from the interface.

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A layer-by-layer process corresponding to this scenario isdepicted in Fig. 2A, Inset. The first interfacial layer or layers withAFM coupling to the rest of the LSMO film retain their AFMspin alignment as temperature increases, while the rest of thefilm gradually flips. Fig. 2B shows the AC magnetic susceptibilityof the thin film. Three maxima are shown where the magneticsusceptibility rises. The emergence of a maximum in AC sus-ceptibility below Curie temperature is called Hopkinson maxi-mum (45). This effect is observed often in soft magnets such asferrites and in materials that undergo a magnetization reversal(45, 46). This feature is explained in the framework of the Stoner−Wohlfarth model, which is related to domain wall motion (47).In Fig. 2B, three Hopkinson maxima are seen. From high tem-peratures, the first and second maxima (marked by arrows) arerelated to a magnetic transition of the interface layer and the restof the thin film, respectively. This result further confirms thepresence of two magnetically active layers, further supporting theproposed model. The peak around ∼65 K cannot be due to amagnetic transition. It is more likely caused by domain wallmotions in response to the SMR. It is noteworthy that, in theabsence of SMR, the AC susceptibility should follow a steadydecline at low temperature, as is illustrated by the dashed line inFig. 2B (48). A signature of soft-phonon (49) second-order cubic-to-tetragonal structural phase transition of STO is seen in thesmall kink at 105 K in Fig. 2A. This kink is always present in theM(T) data (SI Appendix, Fig. S8). The SMR temperature pre-cedes the STO transition temperature, so the bulk STO transitionis not the origin of the SMR, but the LSMO film amplifies thesubtle bulk transition.At room temperature, the observed SMR can be viewed as a

magnetic switch in the monolithic film. As shown in Fig. 2C, in

the absence of a magnetic field, the magnetic moments align in adirection that we label negative, while applying a magnetic fieldaligns all of the magnetic moments in the positive direction.Removing the magnetic field reverses the magnetic moments,resembling an exchange spring. The presence/absence of amagnetic field, which can be as small as 10 Oe, controls theswitching, showcasing the potential applications of this materialfor ultrathin devices.Fig. 3A shows the magnetic hysteresis loops at different tem-

peratures measured using the same procedure as in Fig. 2A. Thehysteresis loop at 25 K represents the normal hysteresis for a softFM with coercivity of 11 Oe (15, 50, 51). At 75 K, the left andright coercive fields start to move toward zero, and IH begins toappear. Further increase of temperature leads to a partial IH. Mreverses to a negative (positive) value in the presence of a pos-itive (negative) magnetic field, giving rise to a negative magneticcoercivity. From 100 K to 275 K, the negative coercivity increaseswith increasing temperature and saturates at higher tempera-tures. Negative coercivity persists up to 325 K, well above roomtemperature. It is noteworthy that no detectable vertical shift inthe hysteresis was observed, pointing to the fact that the biasinglayer is very small. It is well known that the area of the hysteresisloop corresponds to the work done by the magnetic field, and itcannot be negative. Fig. 3B illustrates that the internal interac-tions can make the hysteresis appear to be negative in the regimewhere exchange coupling between interface and the film, JAFM,is greater than both magnetic field (Zeeman energy) and mag-netic anisotropy, Ku. When measuring the hysteresis, if only thesoft layer (the bulk of the film) changes its direction, the hys-teresis will be positive (the case in normal FM or exchange-biasedsystems). However, JAFM can reverse the soft layer, while H is

Fig. 2. (A) Magnetization vs. temperature. Small kink is seen at 105 K. Inset shows a schematic process for SMR. (B) AC susceptibility measured in the absenceof magnetic field. (C) Two-state switch in 30-unit cell LSMO at room temperature under different magnetic fields. The momentary (permanent) ON (OFF) stateis in the presence (absence) of magnetic field.

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strong enough to reverse the hard layer as well. In this case, thetotal energy is conserved, and total hysteresis will be positive (42).The complete IH has been observed in bulk Ni45Fe55 where fine-grain boundaries are formed between a very small embeddedNi3Fe phase and the primary phase. Similar to our findings, thetwo phases couple antiferromagnetically at interface that leads tocomplete inversion of hysteresis loop (52). For our film, themagnetic signal of the bulk of the thin film is dominated, making itdifficult for us to estimate the thickness of the biasing layer. Thisrare phenomenon occurs only when JAFM > H and Ku (43, 53).

DFT Calculations and Explanation of the Structural andElectronic DataWe performed DFT calculations for both stoichiometric andseveral Sr-rich interfaces on a structural model depicted in Fig.4A (SI Appendix, Computational Details). Neither the stoichio-metric heterostructure nor the heterostructure with an Sr-richinterfacial LSMO region exhibits any kind of divergent internalelectric field. Instead, the calculated electrostatic potential, asshown in Fig. 4B, is flat in both cases (SI Appendix, Fig. S10). Atthe same time, the DFT calculations allow us to examine thecharge density and check for charge transfer at the interface.Plots of the charge density around each Mn atom (both in-planeand OOP) at the interface and the bulk of thin film are identical(SI Appendix, Fig. S11). To obtain a quantitative comparison, weshow, in Fig. 4C, the radially integrated charge density aroundMn atoms at the LSMO/STO interfacial region (Sr-rich andstoichiometric interface), and in layer 3 of the LSMO thin film.The three curves are identical, demonstrating that there is nophysical charge transfer or charge accumulation near the in-terface. A similar plot for oxygen shows that the radially in-tegrated charge is the same on oxygen atoms at the interface andin the bulk of the LSMO film, again confirming the absence ofany charge transfer (SI Appendix, Fig. S12). The flat electrostaticpotential and the absence of charge transfer at the interfaceindicate that the polar discontinuity model for polar−nonpolarinterfaces (polar catastrophe) is not applicable to metal−insulatorpolar−nonpolar interfaces.In the absence of physical charge transfer, we need to address

the reason for the observed change in the EELS-derived Mnoxidation state at the interface. EEL spectra are known to de-pend on the available unoccupied states, which depend on the

local environment and bond lengths (40, 54). SI Appendix, Fig.S13 shows the density of states projected on individual atomsat the interface and the bulk of the film, demonstrating thatthe EEL spectra can be different at the interface withoutnecessarily any charge transfer (see SI Appendix for furtherdetails).The DFT results for the OOTs are shown in Fig. 4D along with

the experimental data. The agreement between experimentaland theoretical values is excellent on the STO side. The theo-retical values of OOTs in bulk LSMO are significantly largerthan the measured values, and the same persists in the LSMO/STO heterostructure. This behavior is probably caused by ourchoice of DFT functional, which we have not optimized. Nev-ertheless, the theoretical OOT values reproduce the observedtrend of a dip in the OOTs near the interface.The DFT results for the OOP lattice constants are compared

with the experimental values in Fig. 4E. The only significantdisagreement is that the data show an elongation of the OOPlattice constant at the interface, whereas the theoretical resultsshow a shortening. The same elongation was found in ref. 55,where a case was made that it is caused by the presence of Ovacancies. Indeed, we find that the formation energy of O va-cancies is lower in the interface by 0.64 eV compared with thebulk of the film and that the presence of just 2% of O vacanciescan reverse the OOP lattice constant shortening and produceabout 2% elongation, as observed.

DFT Magnetic Calculations and Atomic-Scale Explanation ofthe DataTo elucidate the atomic-scale mechanism that underlies theobserved magnetic properties, we used DFT calculations to de-termine the magnetic moments on individual Mn atoms near theinterface and in the bulk of the LSMO film (layer 3). We first rancalculations for different Sr configurations in bulk LSMO andfound no significant sensitivity of the magnetic results. Thesecalculations informed us that, in the lowest energy configura-tions, the Sr atoms are farthest from each other. We employedthis criterion in constructing the heterostructure supercell. Wefound that all Mn spins align ferromagnetically (SI Appendix, Fig.S14). Such FM ordering in the ground state does not constitute aplatform that can explain the complex magnetic behavior thathas been observed. The STEM-EELS results shown in Fig. 1

Fig. 3. (A) FM hysteresis loop at different temperatures. The green solid (red open) symbols show leftward with green arrow (rightward with red arrow) fieldsweep. (B) Schematic showing magnetic moment arrangement throughout the FM hysteresis, where Ku and JAFM are magnetic and exchange anisotropy,respectively. To better see the effect, the scales in the figure are exaggerated.

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reveal that the interface is not stoichiometric. We, therefore,turn our attention to calculations on a heterostructure with anSr-rich interfacial region (SI Appendix, Computational Details).The Sr-rich interface indeed features a complex magneticordering that serves as a foundation for explaining the ob-servations and sheds light on the microscopic picturethat underlies the phenomenological scenario described inMagnetic Characterization.Fig. 5 shows the calculated layer-by-layer magnetic moment

alignment in an Sr-rich heterostructure. In the supercell used inthe calculations (Fig. 4A), each layer consists of eight Mn atoms.In the first layer, Mn atoms exhibit intralayer FM coupling. Thesecond layer serves as a transition layer, with about 50% of themagnetic moments in each direction. The magnetic moments in

the remaining layers are aligned antiparallel to those in the firstlayer. Thus, we find a complex magnetic ordering that has a uniquekind of AFM coupling of the FM film to the first two interfacialatomic layers that have different stoichiometries. This type ofcoupling at the interface is not affected by the LSMO surfacetermination (SI Appendix, Fig. S14).We find that, in addition to the Sr-rich interfacial stoichiom-

etry, the suppression of MnO6 octahedral tilts at the interface bythe cubic STO substrate is critical for antiparallel alignment ofmagnetic moments in the LSMO first layer relative to the rest ofthe thin film (Fig. 5). We considered LSMO with no octahedraltilts, i.e., cubic LSMO on cubic STO. The calculations show thatan FM configuration is energetically favorable by 91 meV/f.u.(formula unit) over the complex magnetic ordering described

Fig. 4. (A) Structural model used in DFT calculations. (B) In-plane and microscopic averages of the electrostatic potential across LSMO/STO (001) interface. BulkLSMO electrostatic potential is shown for comparison. (C) No change in the physical valence charge Q(R) on interfacial Mn atoms in stoichiometric interface and100% Sr-rich interface in comparison with Mn atoms in bulk LSMO. (D) Oxygen octahedra rotation across the interface. (E) OOP lattice constant across the in-terface. The good agreement between experimental and computational data is illustrated by overlaid green experimental data points in D and E.

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above. Hence, the suppression of octahedral tilts alone does notfavor the complex magnetic ordering. Instead, the gradient in theoctahedral tilts in the LSMO thin film caused by their suppres-sion at the interface along with a nonstoichiometric interface arecritical for the complex magnetic ordering at the interface. Layer-projected densities of states show spin channel conversion, illus-trating AFM coupling (SI Appendix, Fig. S13). Contributions tothe energy bands in the immediate vicinity of the Fermi level (i.e.,∼1 eV above and below the Fermi level) are mainly from 3d and3p orbitals of interfacial Mn atoms and 2p orbitals of O atoms. It isclear that the LSMO layer 3 shows bulk-like FM ordering, whileLSMO layer 2 has almost equal up and down spins acting as atransition layer. The first LSMO layer has spin in the oppositedirection from the ones in the bulk of the film.The above results of DFT calculations for magnetic moment

alignments of course correspond to T = 0 K. Finite-temperature,time-dependent DFT calculations or even noncollinear calcula-tions to check for magnetic anisotropy energy are not practical.Nevertheless, we can use the T = 0 K DFT results to deduce apossible atomic-scale mechanism for the observed SMR as fol-lows. Initially, at T = 2 K, the external magnetic field aligns allmagnetic moments in the same direction. When the magneticfield is turned off, and the temperature is gradually raised, thetotal magnetization starts to decrease, suggesting that thermalspin-flipping begins to occur. Normally, the flipping of spins isgoverned by the coercivity, but that refers to the resistance of amaterial to elimination of its preferred magnetic ordering by anexternal magnetic field. Here we have the opposite: When theexternal magnetic field is turned off, the system finds itself in ahigh-energy ordering, and spin flips would lower the energy.Although we have no way to estimate the energy barrier for spinflips, the DFT results (Fig. 5) suggest the following. Since FMordering is the energetically preferred option, there is no reasonfor the spins in layers 3, 4, 5, and beyond (Fig. 5) to flip at thesevery low temperatures, whereas those of layers 1 and 2 wouldlower the energy of the system if they flip. The thermal fluctuations

that occur even at 2 K and as the temperature is increased wouldmediate this process. We can then infer that the two interfaciallayers are the first to flip to their T = 0 K alignments, since thatwould lower the total energy.The immediate question then is what is going to happen as the

temperature is raised further. The energetically easiest scenariois to flip more magnetic moments in the second layer, which isalready almost 50% flipped, driving it either to align with thefirst layer or with the remainder of the film. We performed thecorresponding two calculations and found that the total enthalpyis lower by 68 meV/f.u. if the second-layer magnetic momentsalign with the first layer as opposed to the second layer aligningwith the remainder of the film. This result suggests that the in-terfacial transition layer triggers a domino scenario, flipping eachsubsequent layer in the film, and gradually reversing the mag-netization of the entire film as observed. As this process goespast the third atomic layer, however, we are faced with two ex-change energy costs due to spin misalignments: the traveling onethat gradually flips the rest of the LSMO film and a mis-alignment at the interface. As the temperature is raised further,spin flips occur at the interface to eliminate the extra energy cost.Thus, the first interfacial LSMO layer recovers the orientationthat it had at 2 K, which explains why a single atomic layer canbehave as a hard magnet in the phenomenological scenario.

SummaryWe have investigated the origin of intriguing magnetic propertiesof LSMO/STO heterojunction employing atomically resolvedSTEM-EELS and DFT calculations. Our findings show that theinterface is controlling the magnetic properties of the entire thinfilm, exhibiting rare magnetic behaviors such as SMR and IH.The atomically resolved STEM-EELS data are key to our abilityto carry out pertinent DFT calculations that account for theobservations. We find that the concept of polar mismatch doesnot underlie the behavior of nominally polar−nonpolar metal−insulator interfaces, and that a change in the formally definedoxidation state is not equivalent to charge transfer. Using themeasured local structure and composition, DFT shows that thegradient in the octahedral local geometry as well as local com-position are crucial for understanding the observed SMR. TheSMR and IH are features that can be used for potential deviceapplications, taking advantage of room-temperature functional-ity, reduced thin-film dimensionality, and ease of fabrication withone-step growth process.

ACKNOWLEDGMENTS. This work is primarily supported by the US Depart-ment of Energy (DOE) under Grant DOE DE-SC0002136. The electronic mi-croscopic work done at Brookhaven National Laboratory is sponsored by theUS DOE Basic Energy Sciences, Materials Sciences and Engineering Divisionunder Contract DE-AC02-98CH10886. The transmission electron microscopysample was prepared at the Center for Functional Nanomaterials, which is aUS DOE Office of Science Facility, at Brookhaven National Laboratory underContract DE-SC0012704. XRD measurements were performed at Shared In-strumentation Facility at Louisiana State University. S.K. and S.T.P. weresupported by DOE Grant DE-FG02-09ER46554. Computation time was fur-nished by the Extreme Science and Engineering Discovery Environment (56),which is supported by National Science Foundation Grant ACI-1548562 underAllocation TG-DMR180037.

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