Electric-Field Control of Ferromagnetism in a Nanocomposite via a

Letter pubs.acs.org/NanoLett

Electric-Field Control of Ferromagnetism in a Nanocomposite via a ZnO Phase Thomas Fix,*,† Eun-Mi Choi,† Jason W. A. Robinson,† Shin Buhm Lee,† Aiping Chen,‡ Bhagwati Prasad,† Haiyan Wang,‡ Mark G. Blamire,† and Judith L. MacManus-Driscoll† †

Department of Materials Science, University of Cambridge, 27 Charles Babbage Road, Cambridge CB3 0FS, United Kingdom Department of Electrical and Computer Engineering, Texas A&M University, College Station, Texas 77843-3128, United States

‡

S Supporting Information *

ABSTRACT: La2CoMnO6 (LcmO)−ZnO nanocomposite thin films grown on SrTiO3 and Nb−SrTiO3 (001) are investigated. The films grow in the form of self-assembled epitaxial vertically aligned structures. We show that, at 120 K, an electric field applied across the nanocomposite reversibly alters magnetic properties of LcmO. The effect is consistent with charge-mediated coupling between magnetism and an electric field that can be induced by changes in ion valences.

KEYWORDS: Nanocomposite, magnetoelectric, ferromagnetism, thin film, oxide

C

stress. We use vertical composite heterostructures which are promising22,23 because the interfacial area is large and the influence of clamping to the substrate is reduced so that larger ME coupling is expected. For instance, the BiFeO3−CoFe2O4 system has allowed control of magnetization reversal with an electric field.24 However, so far there has been a lack of macroscopic magnetic measurements which are necessary for the applicability of these materials. The composite chosen in this study is a vertical heteroepitaxial nanostructure based on La2CoMnO6 (LcmO) (magnetostrictive and ferromagnetic below 230 K25) and ZnO (piezoelectric and semiconducting). The motivation to use LcmO is that it is a rare example of a ferromagnetic insulator26 so that large electric fields can be applied vertically to the nanostructure (see Figure 1). At the same time, there is potential for surface charging of the ZnO and hence to charge doping of the LcmO. In ref 9, Molegraaf et al. found electronic charge-modulation of magnetism in a bilayer of PbZr0.2Ti0.8O3 (PZT)/4 nm La0.8Sr0.2MnO3 (LSMO). LSMO had to be not thicker than a few nanometers in order for the PZT to charge dope it effectively. Therefore, by increasing the amount of interface, effects should be obtained on a larger scale with thicker magnetic films, which is achieved here through vertical nanostructures. Microstructural properties of the LcmO−ZnO nanostructures was examined by X-ray diffraction (XRD). Figure 2 shows a θ−2θ XRD scan of a LcmO−ZnO film grown on STO (001).

urrent designs for magnetic random access memory (MRAM) require that the magnetic state is controlled either by the application of local magnetic fields or sufficiently large currents to induce spin-transfer torque switching,1 both of which result in significant on-chip dissipation. Replacing magnetic fields with electric fields would reduce power consumption and help miniaturize devices, and new concepts such as four-state memories are also possible.2 Some materials that are magnetically and electrically polarizable can exhibit magnetoelectric (ME) coupling, that is, a change of magnetization induced by an applied electric field or change of polarization induced by an applied magnetic field. Large ME coupling values should be found in ferromagnetic ferroelectrics because ferroelectrics can provide large dielectric coefficients and ferromagnets large magnetic permeabilities.3 However, most ferromagnets are conductive (with partially filled d-bands), while ferroelectrics are insulating (with filled bands); hence ferromagnetic ferroelectrics are scarce, especially at room temperature. There have been four main different approaches to obtain ME coupling in thin films: (a) the use of (anti)ferromagnetic ferroelectrics;4,5 (b) the use of a magnetic layer on a nonmagnetic ferroelectric layer or substrate;6−10 (c) the use of a magnetic layer on an antiferromagnetic ferroelectric layer (e.g., coupling through magnetic exchange bias);11−14 or even (d) the use of a magnetic layer on an oxide or electrolyte (e.g., electron filling control with an electric field).15−18 Interesting reviews on ME coupling can be found in refs 19 and 20. The approach in our study, as first proposed by Suchtelen et al.21 in 1972, is the use of composites of magnetostrictive and piezoelectric phases where the ME coupling originates from © 2013 American Chemical Society

Received: July 26, 2013 Revised: October 25, 2013 Published: November 27, 2013 5886

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can already be seen in Figure 3b where some plates seem to start disappearing. Conductive tip AFM corresponding to Figure 3c is shown in Figure 3d and indicates that the LcmO is insulating as expected while leakage comes from the ZnO plates. To confirm the XRD results and to understand the interplay between the LcmO and ZnO phases, a transmission electron microscopy (TEM) study was conducted, and representative cross-section TEM images are shown in Figure 4. They show that there are vertical nanostructures of ZnO with lateral size of ∼20 nm embedded in the LcmO matrix, consistent with the AFM images. Both ZnO and LcmO are epitaxial as shown in the high-resolution TEM image of Figure 4b). From the diffraction patterns (not shown) the epitaxial relationships are determined to be out-of-plane LcmO (001)∥ZnO (112̅0)∥STO (001) and in-plane LcmO [1̅10]∥STO [100] and ZnO [0001]∥STO [110] or ZnO [0001]∥STO [1̅10]. An energy dispersive X-ray analysis (EDX) in a line scan (not shown) did not detect any Zn substitution in LcmO. The magnetic properties of the nanocomposites were measured as a function of temperature without applied electric field. Figure 5 shows magnetization loops at 10, 100, and 200 K. The evolution is similar to that of the pure LcmO films in which spin-glass behavior has been identified.25 At 10 K the coercive field is ∼10.5 kOe, and at 200 K it becomes around 50 times smaller. The inset of Figure 5 shows the evolution of the magnetization vs temperature, showing a Curie temperature (TC) at around 214 K, which is again close to previous values of pure LcmO films.25 It is well-known that ordering affects the TC. A TC of 150 K has been found to correspond to ∼48% ordering and 225 K to ∼81%.29 Since here the TC is 214 K, it is reasonable to assume that there is a high level of ordering. The evaluation of the valences (see Supporting Information) is compatible with Co2+ and Mn4+, and this further supports ferromagnetic ordering.25 We now discuss the effect of an electric field applied across the nanocomposite structure using the geometry illustrated in Figure 1. Figure 6 shows the magnetization hysteresis loops for different positive electric field values at 120 K. From 0 to 289 kV/cm, a reduction of the coercive field by 35% is observed. The effect goes further for larger electric fields. At 305 kV/cm the coercive field is around 430 Oe, eight times lower than the value of 3417 Oe at 0 kV/cm. The effect is almost symmetrical between positive and negative electric fields. Also, when the electric field is removed, the coercive field goes back to high values, as shown in Figure 6. Another effect of the electric field is visible in Figure 6. The magnetization value is decreasing at higher electric fields: from 0 to 289 kV/cm there is a reduction of 14%, and from 0 kV/cm to 305 kV/cm the reduction is 60%. The latter value gives a ME coefficient of ΔM/ΔE ∼ 3 Oe·cm· kV−1 at 120 K. Another important observation is that neither a large change of coercive field or magnetization can be observed in pure LcmO films contacted and measured in the same way as before, far away from the Curie transition temperature. This shows that the magnetic changes vs electric field observed are caused by ZnO. We finally investigated how the magnetization can be tuned by sweeping the electric field with or without a magnetic field as shown in Figure 7. The measurement procedure is as follows: first we saturate at 10 kOe, and then either we remove the magnetic field in Figure 7a or leave a field of 500 Oe in Figure 7b. The electric field is varied (1) from 0 to 305 kV/cm, then (2) to −278 kV/cm. We see in a that at step (1) the remanent

Figure 1. Schematic showing how the electric field is applied across the LcmO−ZnO nanocomposite. ZnO is symbolized as plates.

Figure 2. XRD pattern showing a 90 nm thick LcmO−ZnO nanostructure on STO (001) using the Bruker diffractometer and in the inset using high-resolution XRD with the Philips diffractometer.

ZnO (112̅0) can be identified, while the LcmO (004) reflection is only seen in the inset from high resolution XRD due to its proximity to STO (002). The full-width at half-maximum of the rocking curve on LcmO (004) is ∼0.1°, similar to the value for the substrate. This shows that LcmO has excellent epitaxy on the STO substrate despite the presence of a ZnO phase. From the θ−2θ scan and the reciprocal space map around (103) STO (not shown), we estimate the lattice parameters aLcmO = 0.5517 ± 0.0005 nm and cLcmO = 0.7758 ± 0.0005 nm, with the following epitaxial relationship (001) [1̅10] LcmO∥(001) [100] STO. The introduction of ZnO did not lead to extra parasitic phases indicating that LcmO and ZnO do not react together. Also the fact that the LcmO lattice parameters are close to bulk values27 indicates little incorporation of Zn in LcmO. While Zn does substitute for Mn in La0.7Sr0.3MnO3 at high temperatures, Zn exolves in the form of ZnO at lower temperatures as a result of spinodal decomposition.28 A similar process may occur here. To understand the nanostructure better, atomic force microscopy (AFM) was performed on a LcmO film without ZnO (Figure 3a) compared to 90 and 180 nm thick films of LcmO−ZnO (Figure 3b,c). The extra features observed in Figure 3b,c are attributed to ZnO, with ZnO forming longer plates for the thicker sample. In the initial growth there is more nucleation of ZnO (i.e., more plates, Figure 3b). Then some plates stop their growth, and there is further growth on the remaining plates that expand mainly in length (Figure 3c). This 5887

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Figure 3. AFM topography of (a) 90 nm thick LcmO film, (b) 90 nm thick LcmO−ZnO nanostructure, and (c) 180 nm thick LcmO−ZnO nanostructure. (d) A current map corresponding to c.

Figure 4. Cross-sectional TEM images of LcmO−ZnO nanostructures in (a) low magnification (target phase ratio 1:4) and (b) high resolution (target phase ratio 1:1), of a 80 nm thick film on STO (001) in the [100] STO zone axis. The arrows indicate the presence of ZnO in the composite. High-resolution TEM (b) clearly shows the highly epitaxial growth of ZnO and LcmO nanostructures.

Figure 6. Magnetization curves at 120 K for different electric fields applied to a 180 nm thick LcmO−ZnO nanostructure on Nb-STO contacted as in Figure 1. The electric field was varied in the following way: 0 → 222 → 289 → 305 → 0 kV/cm as shown in the legend. The magnetic field is applied in plane to the sample along STO [100]. The magnetization in emu has been divided by the total volume of the film.

Figure 5. Magnetization curves of a 90 nm thick LcmO−ZnO nanostructure on STO (001). Inset: magnetization vs temperature curve at 1 kOe of LcmO−ZnO nanostructure on Nb-STO (001). The magnetic field is applied in plane to the sample along STO [100]. The magnetization in emu has been divided by the total volume of the film.

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account for the phenomena observed, although local heating cannot be completely excluded. The second reason is the fact that ZnO is piezoelectric and that the electric field applied induces strain in the ZnO which then strains the LcmO phase. In this case, assuming the inverse magnetostrictive coefficient of LcmO is constant one would expect a progressive evolution of M−H loops as the electric field increases which is not the case as shown in Figure 6. However first-principles calculations predict that strain in LcmO can switch the orbital occupancy of Co2+ and modify the electronic structure of the material inducing a metal−insulator transition,30 possibly with a nonlinear response to strain. Experiments by piezoresponse force microscopy and XRD with varying electric field could not show evidence of piezoelectricity in the nanocomposite; therefore the second reason is also excluded (see Supporting Information). It should also be noted that experiments using Nb-STO (110) substrates where ZnO is (0001) oriented here did not lead to stronger electric field dependence, while stronger piezoelectricity would be expected in this orientation (see Supporting Information). The third reason is related to the magnetic sensitivity of strongly correlated oxides to charge.19 Some examples with manganites show charge-mediated coupling between magnetism and electric field by inducing changes in valences.9,31,32 In particular, work on bulk LcmO33 has demonstrated that valence changes driven by changes to the oxygen vacancy concentration lead to shifts toward a different TC magnetic phase. We believe that this explanation is the most plausible to account for the phenomena observed. When two structurally incompatible oxides are contacted, the chemical bonding is modified at the interface due to the strain and structural mismatch.34 In the case of our composite, this will lead to a high density of oxygen vacancies at the vertical interfaces of LcmO and ZnO. These vertical interfaces are more conducting than the LcmO matrix, as shown by conductive tip AFM (Figure 3d). The electrons, which conduct through the vertical interfaces, will get trapped or detrapped on oxygen vacancies35 leading to changes to the valence states of Mn and Co cations in the LcmO.33 We have ignored any effect of oxygen vacancy migration under electric field because of the very low measurement temperature of 120 K,36 but it is reasonable to suppose that any vacancy migration will add to this affect. Compared with LcmO−ZnO nanostructures, there are few conducting pathways in the insulating LcmO thin film, and so charge can much less easily be transported; so, as observed, there is no electric-field control of the magnetism in the pure material. To conclude, we have shown the potential of novel piezoelectric−ferromagnetic nanocomposites obtained by selfassembly by pulsed laser deposition (PLD) growth. The influence of the electric field on the magnetic properties of LcmO−ZnO is large and over a significant sample volume. Although the effects observed here are at 120 K, the results provide a strong stimulus for further research pointing the way to future simple self-assembled thin film composite device designs and to eventual room temperature operation. Experimental Section. Films were grown on single-crystal STO and Nb-STO (001) substrates by PLD using a KrF laser with a 248 nm wavelength. The growth conditions were 10 Hz repetition rate and a fluence of around 1 J·cm−2 on the target at a substrate target distance of 120 mm; the heater temperature was 750 °C, and the oxygen pressure while depositing and cooling down was around 10−2 mbar and 0.5 mbar of O2, respectively. LcmO−ZnO targets with a 1:1 molar composition

Figure 7. Magnetization at 120 K at remanence (a) and 500 Oe (b) vs electric field of a 180 nm thick LcmO−ZnO nanostructure on NbSTO contacted as in Figure 1. The magnetic field is applied parallel to the sample plane along STO [100]. The magnetization in emu has been divided by the total volume of the film. The arrows and numbers next to them indicate how the electric field was swept.

magnetization is first constant and then decreases as the electric field increases, consistently with Figure 6. Then for (2) the magnetization remains at low values. A different scenario happens for Figure 7b where a field of 500 Oe is maintained while sweeping the electric field. Step (1) is similar to before, but then step (2) shows that the magnetization increases again while decreasing the electric field from 305 to 0 kV/cm. Symmetrically with step (1), step (2) shows a decrease of the magnetization from 0 to −278 kV/cm. We have shown from XRD and TEM results that the LcmO−ZnO is a fully epitaxial self-assembled nanocomposite. The M−H and M−T measurements indicate that the LcmO magnetic properties are not degraded by the ZnO. More interestingly, Figure 7 shows that the remanent magnetization can be tuned in a nonvolatile manner. Indeed, in Figure 7a, once that the remanent magnetization has been decreased to a low value by applying an electric field, it does not recover to the initial values by decreasing the electric field, unless a magnetic field is applied simultaneously as shown in Figure 7b. A strong dependence of the magnetic properties (Figure 6) with an electric field applied across the sample is observed. This could happen for several reasons. The first one is a ferromagnetic−paramagnetic transition induced by heating generated by the electric field in a leaking sample. To eliminate this possibility, many growth runs were undertaken to reduce this effect. While this reason could have been true in early leaking samples measured close to the Curie temperature, leakage was reduced considerably (0.25 mA for 1 V) by increasing the nanostructure thickness (180 nm), and samples were measured at 120 K. Furthermore, a rise of temperature of the sample induced by the electric field from 120 K to around 190 K would be necessary to account for solely heating effects. However, changes of temperature on the sample with application of electric field, measured with a Pt resistor, were found to be limited to several kelvin for the highest electric field applied, which is significantly less than the 70 K rise which would be necessary to give the measured change in Mr if this occurred solely because of heating effects. Thus, heating effects are not the most plausible mechanism to 5889

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(unless specified otherwise) were made by milling, presintering at 900 °C for 6 h, pressing, and sintering at 1100 °C for 6 h a stoichiometric mixture of high purity (99.99%) La2O3, MnO2, Co3O4, and ZnO powders. The orientation, crystallinity, and lattice parameters of the films were determined by XRD with a Philips X’Pert PRO and a Bruker D8 Advance diffractometer (Cu Kα) and TEM using a FEI Tecnai G2 F20 analytical microscope operated at 200 kV. AFM images were obtained at room temperature either with a Veeco Dimension 3100 or an Agilent SPM 5500. Conductive tip AFM was performed using a Pt-coated tip biased at 2 V. Magnetization measurements (M−T and M−H) were performed using a Princeton Micromag 3900 vibrating sample magnetometer (VSM) and a superconducting quantum interference device (SQUID) (Quantum Design, MPMS). For VSM measurements with an electric field, a Pt electrode was deposited on 5 × 5 mm2 films, and then contacts were made using silver paste.

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ASSOCIATED CONTENT

S Supporting Information *

Composition analysis, additional curves of magnetization vs electric field, X-ray absorption spectroscopy, and piezoresponse force microscopy. This material is available free of charge via the Internet at http://pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Author

*Address: ICube laboratory (Université de Strasbourg and CNRS) 23 rue du Loess BP 20 CR, 67037 Strasbourg Cedex 2, France. E-mail: thomas.fi[email protected]. Present Address

A.C.: Center for Integrated Nanotechnologies (CINT), Los Alamos National Laboratory, Los Alamos, New Mexico 87545, United States. Author Contributions

T.F. is responsible for the experiments and preparation of the paper; E.M.C. performed the SQUID measurements; B.P. contributed to the film deposition; A.P.C. and H.W. performed the TEM analysis; J.W.A.R., S.B.L., M.G.B., and J.L.M. provided guidance for the experiments and for editing and proofreading the paper. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS T.F. thanks S. Kar-Narayan, A. Kursumovic, and F. Schoofs for technical assistance and discussion. J.W.A.R. is funded by the Royal Society. This research was partially funded by the ERC Advanced Investigator Grant (no. 247276, NOVOX).

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REFERENCES

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