Polarized neutron reflectometry

Thin film samples of Si/Pd (50 nm)/AlO x (1 μm)/GdO x (2 nm)/Co (15 nm)/Pd (20 nm) were grown by sputtering and evaporation. The thin GdO x layer is crucial, and experiments performed without the 2 nm GdO x layer were unsuccessful. In contrast to previous studies, this thickness of Co is expected to give rise exclusively to an in-plane magnetic easy axis. The samples were E+T conditioned by heating the sample to 230 °C and applying +40 V to the top Pd film, while the buried Pd film was held at ground for 15 min; the reverse treatment applied −40 V to the top contact, also at 230 °C for 15 min (see the ‘Methods’ section). PNR measurements of the as-grown sample and each conditioned state are shown in Fig. 1a. The R++ and R−− reflectivities show sensitivity to the nuclear and magnetic depth profiles, evident by spin-dependent oscillations. The difference in the R++ and R−− is approximately proportional to the quotient of the magnetization and the nuclear SLD (see the ‘Methods’ section). Thus, the magnetic contribution to the data is highlighted by plotting the spin asymmetry (SA=(R++−R−−)/(R+++R−−)), as shown in Fig. 1b. The oscillation amplitude first decreases after conditioning in +40 V then increases after conditioning in −40 V, suggesting a decrease of the saturation magnetization, M S and/or a change in the structure, followed by a partial recovery towards the initial state. The nuclear and magnetic depth profiles from the converged model, shown in Fig. 1c, confirm these trends. Figure 1 is reconstructed into individual panels in Supplementary Fig. 1 to aid in visualization. The SLD profiles are a cross-sectional average of the nuclear and magnetic SLD at particular depths due to the neutron coherence distribution23 and, as such, measure the average contribution of all atoms at that depth.

Figure 1: PNR results of E+T-conditioned sample. (a) Fitted PNR data, with R++ (R−−) identified by solid circles (open triangles), scaled by q4 and (b) spin asymmetry for the sample as-grown and after E+T conditioning. (c) Depth-dependent real and imaginary nuclear SLD (ρ N and ρ imag ), and magnetic SLD (ρ M ) extracted from the PNR. In all panels black, red and blue lines identify the as-grown, +40 V conditioned, and +/−40 V conditioned samples, respectively. In c the solid and dashed lines identify ρ N and ρ M , respectively, and the green line identifies ρ imag . Background colours in c represent (red) Pd, (purple) AlO x , (green) GdO x and (yellow) Co, respectively. In a and b the experimental data are shown as symbols, and the lines are fits corresponding to the depth profile shown in c. Error bars in a and b correspond to ±1s.d.; error bars for c are shown in Supplementary Fig. 2. Arrows in a and b identify the bulk ρ N for Co (2.25 × 10−4 nm−2) and CoO (4.29 × 10−4 nm−2). Full size image

The extracted depth profile of the as-grown sample accurately reproduces the designed structure, both in terms of thickness and nuclear SLD, ρ N . Our fits show excellent agreement between the measured and expected ρ N values of Co (2.27 × 10−4 nm−2), Pd (4.02 × 10−4 nm−2), and GdO x (2.74 × 10−4 nm−2) (refs 24, 25). However, the measured SLD of the thick AlO x base layer is substantially lower than the expected bulk value (5.67 × 10−4 nm−2), suggesting the presence of significant voids or an oxygen-deficient stoichiometry. The GdO x —a neutron absorber—can be identified explicitly by the imaginary SLD in Fig. 1c. After conditioning the sample at +40 V, the nuclear SLD of the Co layer, ρ N Co, increases by 34%, approaching that of CoO (4.29 × 10−4 nm−2). Simultaneously, the measured GdO x /Co interface becomes much broader, increasing from a width of 3.3 to >10 nm and extending well into the AlO x . These results are consistent with the electric field removing oxygen from the GdO x and AlO x , decreasing their SLD, and depositing it into the Co, increasing its SLD and resulting in an apparent broadening of the interface. After switching the voltage polarity to −40 V, the GdO x /Co interface width is reduced to 1.9 nm and ρ N Co decreases, demonstrating an induced migration of oxygen from the CoO x into and through the GdO x . The recovery of ρ N Co occurs predominantly within the 10 nm nearest the GdO x interface, while the top 5 nm, near the Co/Pd interface, remains unchanged from the +40 V conditioned state. Thus, we observe oxygen ion migration throughout the thickness of the Co layer, but reversibly only within the 10 nm closest to the GdO x /Co interface. However, at the GdO x interface after this second conditioning, ρ N Co is still 32% larger than the as-grown sample, demonstrating that the oxygen migration is only semi-reversible. The uncertainty associated with the presented fits is shown in Supplementary Fig. 2, and discussed in Supplementary Note 1, and is much smaller than the changes between conditioning steps.

Trends in the magnetic depth profile (dashed lines in Fig. 1c) agree with those observed in the nuclear profile. Specifically, the as-grown sample has a sharp step-function like GdO x /Co (magnetic) interface and strong magnetic scattering, as indicated by its large magnetic SLD, ρ M , of 3.45 × 10−4 nm2 in the Co layer, corresponding to a bulk magnetization of 1,180 emu cm−3 (1 emu cm−3=1 kA m−1). Strikingly, the shape of the magnetic profile changes after conditioning in +40 V, with ρ M reduced significantly—by 80%—at the GdO x /Co interface and 38% in the bulk. Subsequent conditioning at −40 V recovers the original shape and ρ M increases to 92% of the as-grown value.

Control experiments were performed on a sample grown side-by-side with the E+T-conditioned sample following the same thermal treatment but without an electric field. PNR measurements of the as-grown sample, shown in Fig. 2, share a similar structure to the as-grown sample in Fig. 1, with small differences in ρ N likely due to sample aging (30 days between measurements in Figs 1 and 2). Similar to the E+T sample, the first thermal treatment also increases ρ N and reduces ρ M in the Co layer, but to a much lesser degree and with no significant changes in the GdO x /Co interface shape and extent. Quantitative comparison shows that ρ N Co increases by 17% and ρ M Co decreases by 12% after thermal treatment alone compared with the as-grown sample. This is much less than the E+T-treated sample, which showed an increase in ρ N Co of 34% and a decrease in ρ M Co of 38% after the +40 V treatment. After a second 15-min thermal conditioning the nuclear and magnetic profiles of the control sample do not change appreciably, suggesting a saturation effect or depletion of easily diffusible oxygen. These results confirm the role of the electric field in enhancing the oxidation of the Co layer during the +40 V conditioning and reducing it during the −40 V treatment.

Figure 2: PNR results of thermally conditioned control sample. (a) Fitted PNR data, with R++ (R− −) identified by solid circles (open triangles), scaled by q4 and (b) spin asymmetry for the sample as-grown and after thermal-only conditioning. (c) Depth-dependent nuclear and magnetic SLD, ρ N and ρ M , extracted from the PNR. In all panels black, red and blue lines identify the as-grown, once-treated (15 min) and twice-treated (2 × 15 min) samples, respectively. In c the solid and dashed lines identify ρ N and ρ M , respectively, and the green line identifies ρ imag . Background colours in c represent (red) Pd, (purple) AlO x , (green) GdO x and (yellow) Co. In a and b the experimental data are shown as symbols, and the lines are fits corresponding to the depth profile shown in c. Error bars in a and b correspond to ±1 s.d. Arrows in a and b identify the bulk ρ N for Co (2.25 × 10−4 nm−2) and CoO (4.29 × 10−4 nm−2). Full size image

Magnetometry

Magnetic hysteresis loops of the samples as-grown, after sequential +/−40 V treatment (E+T) and after two thermal-only treatments are shown in Fig. 3a. The Co M S in the as-grown sample was measured to be 1,230 emu cm−3, in good agreement with the PNR value of 1,180 emu cm−3. Further, M S decreased by 10% in the E+T-treated sample and 7% in the thermal-only sample compared with the as-grown sample, similar to the PNR data which show a reduction of 10%. The good agreement between the magnetometry and PNR results supports the validity of the model used to fit the data. Sample coercivity and remanent magnetization also decreased compared with the as-grown sample by 68% and 55% in the E+T sample and 54 and 11% in the thermal sample, respectively, indicating significant changes in the magnetic characteristics.

Figure 3: Magnetometry and FORC investigations. (a) Combined major hysteresis loops for the sample (black) as-grown, (red) after +/−40 V conditioning and (blue) after thermal-only conditioning. Family of FORCs and FORC distributions for the sample (b,e) as-grown, (c,f) after +/−40 V conditioning and (d,g) after thermal-only conditioning. Color bar in e identifies the max (red) and min (blue) value of the FORC distribution. Full size image

Details of the magnetization reversal have been investigated by the FORC method26,27,28,29. The family of FORCs and FORC distribution for the as-grown sample are shown in Fig. 3b and e, respectively. The family of FORCs show the minor loops fill the major loop, and the calculated FORC distribution shows only a single feature, centred at (local coercivity μ 0 H C =4.9 mT, bias field μ 0 H B =0 mT, where μ 0 is the vacuum permeability 4π × 10−7 N A−2), indicating the sample is comprised of a single magnetic phase29,30. After the combined +/−40 V electric field treatment (E+T) the family of FORCs, Fig. 3c, still fill the major loop, but the FORC distribution, Fig. 3f, now shows two features. The main feature is centred at (μ 0 H C =2.6 mT, μ 0 H B =0.33 mT) and is circularly symmetric, again indicative of an irreversible (that is, hysteretic) phase. The shift in H C relative to the as-grown sample indicates the coercivity is significantly reduced. The non-zero bias suggests a finite interaction experienced by this phase and may be the result of exchange bias with residual antiferromagnetic CoO with an enhanced Néel temperature31. A second phase is identified by the elongated FORC ridge along the μ 0 H B axis centred at μ 0 H C =0 mT. This ridge represents a reversible phase with an internal demagnetizing interaction of 6 mT at saturation, identified by the spread of the feature along the μ 0 H B axis27. A negative feature centred at (μ 0 H C =0.8 mT, μ 0 H B =−1.7 mT) identifies reversal events which are present on FORC branches that start at H R 1, M(H,H R 1), but absent on FORCs that begin at more negative H R 2, M(H,H R 2) with H R 2<H R 1 (ref. 29). In this case, the negative feature in Fig. 3f aligns in H R with the peak of the reversible feature, and in H with the irreversible feature. This behaviour indicates that once the irreversible switching event occurs, the reversible phase changes its upswitching field due to magnetic coupling between the reversible and irreversible phases.

The family of FORCs for the thermally treated sample, Fig. 3d, is significantly different, with the minor loop protruding outside of the major loop. This indicates that the domain structure evolves more easily under fields applied along the major loop that increase from the saturated state, than under fields that increase from the mixed multi-domain state32. This result further underscores the role of the electric field in determining the oxygen distribution. Similar to the E+T sample, the FORC distribution for the thermal-only sample, Fig. 3g, also exhibits reversible and irreversible phases. The main FORC feature is centred at (μ o H C =2.6 mT, μ o H B =0.17 mT) and is no-longer circularly symmetric, but rather has a 90° bend with symmetries along the +H and -H R axes, typical of a domain nucleation/growth reversal mechanism28,33. The FORC distribution for the thermal-only sample shows the same negative feature that again identifies magnetic coupling, and a new feature associated with the observed major loop protrusion. Integrating the FORC features gives a magnetic phase fraction30: the reversible phase contributes to 0%, 31% and 24% of the magnetization in the as-grown, E+T and thermal-only samples, respectively.

Reversible phases exhibit no hysteresis, and therefore are manifested in the FORC distribution along the μ o H C =0 mT axis, for example, when the phase has essentially zero coercivity or when the magnetic field is applied along the magnetic hard axis. Major hysteresis loops measured in the out-of-plane direction (see Supplementary Fig. 3 and Supplementary Note 2) for these samples display little hysteresis, implying that the out-of-plane direction remains the hard axis. These results suggest that oxygen migrates in the film after E+T or thermal-only conditioning and segregates to grain boundaries in the Co layer, thus disrupting long-range magnetic correlations and effectively breaking down the affected Co films into isolated grains. Once the coupling between the grains becomes weaker, their respective magnetocrystalline anisotropies in confined grains play a large role in determining the magnetic orientation resulting in much reduced coercivity and remanent magnetization.

X-ray absorption and circular dichroism

Oxidation of the cobalt after both +/−40 V E+T and thermal-only treatments is confirmed in the XA and XMCD measurements (see the ‘Methods’ section) shown in Fig. 4a. Oxidation of the Co layer34 is identified directly by the emergence of peaks at E=779.2 eV and 776.8 eV, which are not present in the as-grown profile. The peak at 779.2 eV is largest in the thermal-only sample, indicating significantly increased oxidation relative to the E+T sample. This trend is supported by the XMCD spectra, which shows that the as-grown sample has the largest dichroism, indicating the largest magnetization. The E+T sample has the second largest dichroism, and the thermal-only sample has the smallest. The different ordering in the dichroism, compared with the bulk magnetometry, may identify variation in the depth-dependent oxygen-binding behaviour. This is consistent with the XA results, which showed a larger oxidation peak in the thermal-only treatment than the E+T sample. XMCD signal from the Gd, shown in Fig. 4b, shows no dichroism for all three samples, indicating a negligible contribution to the magnetization. Interestingly, the XA signal for the Gd shows no significant change for any of the samples, suggesting a relatively constant Gd oxidation state, regardless of oxygen migration into or out of the Co. While the Gd XAS is expected to be less sensitive to the oxidation state than the Co XAS, the amount of oxygen necessary to facilitate the observed changes in the Co would constitute a significant change if it had come from the 2-nm thick Gd layer, and would be observable in the Gd XAS.

Figure 4: X-ray absorption and XMCD spectra. XA (solid) and XMCD (dashed) spectra for (a) cobalt and (b) gadolinium in samples as-grown (grey), after +/−40 V conditioning (E+T, red) and thermal-only (blue) conditioning. Arrows indicate 779.2 eV and 776.8 eV. Full size image

The PNR, magnetometry and X-ray results clearly demonstrate that E+T conditioning can drive oxygen semi-reversibly into a thick (15 nm) Co film, profoundly changing its magnetic properties. Depth profiling with PNR indicates that while these effects are most prominent at the GdO/Co interface, they also extend throughout the entire 15-nm thick Co film. Reversing the polarity of the applied voltage drives oxygen out of the Co, partly restoring ρ N Co and ρ M Co to their original values at the GdO x /Co interface, but leaving ρ N Co and ρ M Co unchanged near the Co/Pd interface. Thermal conditioning of the control sample also promotes oxidation of the Co layer, but the supply of highly mobile oxygen that can be moved by entropy-driven diffusion is clearly limited. In the following discussion we determine the oxygen stoichiometry from the nuclear scattering profile and consider the underlying mechanics of the oxygen migration.

The role of the electric field and entropy-driven oxygen migration is seen qualitatively by comparing the profiles for the +40 V E+T-treated sample with the thermally treated sample (Figs 1c and 2c, respectively). The thermal treatment is shown to scale the magnetic depth profile relative to the as-grown sample, but not change its shape. In comparison, the magnetic depth profile for the +40 V sample strongly deviates from that in the as-grown sample, suggesting that the electric field drives oxygen into the film, while the thermally activated, entropy-driven oxygen migration is relatively uniformly distributed.

Using the neutron coherent scattering length, b, for cobalt (2.49 fm) and oxygen (5.81 fm) (ref. 35), the CoO x stoichiometry can be directly calculated. Specifically, the nuclear SLD is calculated as: where N Co(O) is the total number of cobalt (oxygen) atoms within the volume, V, of the Co film. Assuming the as-grown film is pristine Co, which is supported by the good agreement with the referenced bulk ρ N Co value, and that the cobalt number density remains constant during treatment, a lower-limit to the oxygen profile can be calculated: . The stoichiometry is then defined as CoO N(O)/N(Co) . The nuclear depth profiles suggest that the +40 V conditioning forms CoO 0.18±0.01 , compared with CoO 0.07±0.01 in the thermally treated sample. Treatment with a reversed polarity removes about one-third of the absorbed oxygen, leaving CoO 0.12±0.01 . The fact that the magnetization measurements show only a 10% reduction in M S illustrates an indirect correspondence (for example, not one-to-one) of the magnetization and nuclear composition and suggests some of the O2− may be forming magnetic compositions other than CoO, such as Co 2 O, or may remain as interstitial oxygen. Integrating the oxygen profiles for the conditioned samples shows conservation of oxygen within this system to within 3%. Since we do not expect any external sources of oxygen, this agreement is another good support for the validity of the presented models. With additional assumptions we can further determine a depth-resolved oxygen profile (see Supplementary Fig. 4 and Supplementary Note 3) to highlight the interface and bulk migration explicitly.

Interestingly, the depth profiles (especially the magnetic profile) indicate that the oxygen is semi-reversibly driven out of the GdO x /Co (0–10 nm) interface after treatment in −40 V, while the remaining oxygen is left trapped deeper within the Co. We suggest that as the GdO x /Co interface becomes depleted of oxygen, becoming more metallic. The oxidized region deeper within the film gets surrounded by conductive layers above and below, screening the electric field; without an electric field the oxygen does not migrate, resulting in the observed trapping effect. Thus, we suggest that the observed 10 nm thickness of the Co presents a practical limit for the electrically driven oxygen migration within these otherwise metallic films.

Mechanics of oxygen migration

Brief considerations of the underlying mechanisms quickly reveal that this effect cannot be justified with bulk chemistry alone. First, considering the thermally treated sample, the initial treatment both suppresses the magnetism and increases the nuclear SLD, consistent with inclusion of oxygen. Further treatment only weakly changed these parameters, indicating that all of the easily diffusible oxygen had migrated during the first treatment. Since the available thermal energy (k B T=43 meV) is not enough to reduce Gd 2 O 3 (enthalpy of formation, ΔQ=18.9 eV), Al 2 O 3 (ΔQ=17.4 eV) or CoO (ΔQ=2.5 eV), we suggest the source of the mobile oxygen is likely interstitial, for example, from trapped sputtering gas located at grain boundaries or voids in the GdO x and AlO x , which diffuses and reacts with the Co (Co+O→CoO, ΔQ=−2.5 eV or 2Co+O 2 →2CoO, ΔQ=−0.2 eV) (refs 36, 37).

Next, we consider the role of the electric field on oxygen mobility. Defect sites in oxygen-rich transition metal oxide films have been previously shown to act as p-type dopants38, giving them an effective negative charge. In the presence of an electric field with the anode on the Pd cap surface and cathode on the buried Pd seed film, excess O2− defects in the GdO x and AlO x will be pulled towards the cobalt, while vacancies are pulled towards the bottom Pd electrode. The chemical potential within each of the AlO x , GdO x and Co films is expected to be uniform, and thus O2− defects are expected to be highly mobile within each17. The application of a static electric field then causes field-induced ion migration39. However, at the boundary between two layers there will exist a difference in the chemical potential, which may cause an accumulation of oxygen at, for example, the GdO x /Co interface. Considering this issue, the enthalpy of formation is calculated in each layer and compared with the electric potential energy available to overcome this interfacial barrier. Assuming a lattice-site hopping model40,41, the electric potential energy at the interface can be calculated to be 24 meV (qEΔx, where q is the oxygen charge of 2e−, Δx is the atomic site spacing of 3 Å and E is the electric field of 400 kV cm−1). The sum of the electric potential energy and thermal energy defines the scale of the energy landscape available to drive the reversible oxygen migration back and forth across the interface. First considerations of an ideal system were that the migrating oxygen is moved from chemically stable Gd 2 O 3 to the Co (Gd 2 O 3 +3Co→3CoO+2Gd, ΔQ=+11.4 eV) (refs 36, 37). Clearly the available energies (67 meV) cannot drive this reaction; stability of the Gd 2 O 3 oxidation state is supported by the XA in Fig. 4. An alternative scenario suggested above is that the trapped oxygen react with the Co to form CoO (2Co+O 2 →2CoO, ΔQ=−0.2 eV). However, without a charge, the O 2 experiences no net force from the electric field, suggesting this reaction is not responsible for the observed effects induced by the electric field. Similar reactions with elemental oxygen (Co+O→CoO, ΔQ=−2.5 eV) can occur along the forward reaction, but will not be reversible as it again costs too much energy. Other considered reactions are presented in Supplementary Note 4, but no bulk energy calculation was able to support the observed results.

We present a possible alternative energy consideration, treating the grains as nanoscale clusters, which provides a reasonable energy landscape for the observed migration. In the picture illustrated in Fig. 5, oxygen ions are bound to the surface of the nanocrystalline grains. These O2− ions are off-stoichiometry defects, and have a different binding energy than the core of the grain. Since the grains were all fabricated at the same time by sputtering in an oxygen-rich environment, the surface structure and hence energy landscape, is expected to be relatively uniform. Thus applying an electric field moves the surface O2− towards the anode and into the Co layer, forming CoO x . Once on the grain surfaces these O2− ions experience chemically driven diffusion, which occurs at a rate of ∼10–20 nm h−1 at these temperatures42, preferentially oxidizing the surface, but likely also the core of the small cluster-like Co grains. Reversing the applied field drives O2− to leave the CoO x film. Using a cluster-like formation approach, the binding energy per atom for GdO x (ref. 43) and CoO x (ref. 44) can be determined by modelling to be 4.50 and 4.48 eV per atom near nominal stoichiometry, respectively. Since the energy landscape is isotropic to within 20 meV, one only has to overcome the small barrier and the activation energy to move oxygen ions, destroy old bonds and create new ones. Thus oxygen freely leaves the CoO x , but due to the slow bulk diffusion in the CoO x grains, a gradient oxygen profile is expected to form within the grains, with the surface being oxygen deficient. Once the grain surface loses enough oxygen to become conductive, the internal electric field is again screened, trapping the oxygen within the core. Thus, both at the film-length mesoscale and grain-length nanoscale, the oxygen migration induces screening effects which limits the oxygen mobility. We propose that, as a result, the conditioned films have a highly defective structure, with a complex mix of CoO x core-shell-surface grains. This model presents a mechanism which is semi-reversible and nearly recovers the saturation magnetization, thus consistent with the PNR and magnetometry. Perhaps more remarkably, this model also has a disrupted long-range ordering, consistent with the two-phase construction identified by the FORC measurement. Finally, this model also agrees well with the XA spectra; the oxygen migration in the Co grains occurs both on the surface and throughout the bulk. In contrast, only the oxygen on the surface of the GdO x grains is likely mobile with the bulk of the grain remaining stoichiometry balanced.

Figure 5: Illustration of oxygen migration mechanism. Cross-section view of (a) the as-grown film, (b) a single grain during +40 V treatment, (c) the film after +40 V treatment, (d) a single grain during subsequent −40 V treatment, and (e, f) cross section and grain after −40 V treatment. Colours identify AlO x (red), GdO x (green), metallic FM Co (light blue), insulating non-FM CoO x (blue) and interstitial oxygen (orange); large gold arrow indicates the electric field. Illustration emphasizes fast surface migration, identified by the red arrows, and slow bulk diffusion, indicated by the grey arrows. Full size image