We report on the perpendicular magnetic anisotropy (PMA) behavior of heavy metal (HM)/Fe alloy/MgO thin film heterostructures when an ultrathin HfO2 passivation layer is inserted between the Fe alloy and MgO. This is accomplished by depositing one to two atomic layers of Hf onto the Fe alloy before the subsequent rf sputter deposition of the MgO layer. This Hf layer is fully oxidized during the subsequent deposition of the MgO layer, as confirmed by X-ray photoelectron spectroscopy measurements. The HfO2 insertion generates a strong interfacial perpendicular anisotropy energy density without any post-fabrication annealing treatment, for example, 1.7 erg / cm 2 for the Ta/Fe60Co20B20/HfO2/MgO heterostructure. We also demonstrate PMA even in Ni80Fe20/HfO2/MgO structures for low-damping, low-magnetostriction Ni80Fe20 thin films. Depending on the choice of the HM, further enhancements of the PMA can be realized by thermal annealing to at least 400 o C . We show that ultra-thin HfO2 layers offer a range of options for enhancing the PMA properties of magnetic heterostructures for spintronics applications.

The realization of robust perpendicular magnetic anisotropy (PMA) in heavy metal (HM)/Fe alloy/MgO thin-film heterostructures,1,2 where typically the Fe alloy is Fe80-xCoxB20 (FeCoB), has enabled a pathway for the implementation of high density memory elements based on the spin transfer torque switching of perpendicularly magnetized tunnel junctions (MTJs).2–4 Strong PMA is also required to create the perpendicularly magnetized nanowire structures needed to enable the manipulation of domain walls and novel magnetic chiral structures such as skyrmions by the spin Hall effect.5–8 At present, the only viable FM/oxide combination that yields the strong PMA and low damping required for practical devices is Fe80-xCoxB20 (FeCoB)/MgO where the PMA originates from the strong spin-orbit interaction in the hybridized 3d Fe-2p O bonding at the FeCoB/MgO interface.9,10 However, obtaining significant PMA in the previous work requires an annealing step1–4 that can compromise the layers in the magnetic heterostructure due, for example, to materials intermixing at the interface or to longer distance diffusion.

In a previous paper, we inserted an ultrathin oxidized Hf layer at the FeCoB/MgO interface of HM/FeCoB/MgO heterostructures to study how changing the oxide interface modifies the field-like spin-orbit torque.11 Here, we report a systematic study that this addition to the surface of FeCoB of as little as 0.2 nm of Hf “dusting,” which is oxidized to HfO2 during the subsequent MgO deposition process, can yield strong PMA without any post-fabrication annealing treatment. Depending on the base HM, the system can also, if that is desired, be annealed to at least 400 o C to further enhance the PMA. This simple Hf dusting technique not only improves the performance of FeCoB/MgO structures but also allows for the PMA devices to be made from the low-damping, low-magnetostriction alloy permalloy (Ni80Fe20) and likely other Fe alloys. This technique therefore substantially expands the options for engineering magnetic thin film heterostructures for spintronics.

All the samples in this paper were prepared via standard direct current (DC) sputtering (with RF magnetron sputtering for the MgO layer), with a base pressure of < 4 × 10 8 Torr . The DC sputtering conditions were 2 mTorr Ar pressure and 30 watts power. To form the interfacial HfO2, an ultrathin Hf dusting layer was first sputtered on FeCoB with a low deposition rate of 0.01 nm / s , and the MgO layer was then sputtered on the Hf layer with a growth rate of 0.005 nm / s (at 100 watts power and 2 mTorr Ar pressure), a process that oxidized the Hf layer. In each case, the top Ta film serves as a capping layer to protect the underlayers from degradation due to atmospheric exposure. The heterostructures were fabricated into Hall cross bar structures with a lateral dimension of 5 × 60 μ m 2 for measurements. All the samples were baked at 115 ° C for 1 min twice during the standard photolithography fabrication for photoresist treatment.

We will first discuss the results obtained from a set of samples (series A) of Si/SiO2/Ta(6)/FeCoB( t FeCoB )/HfO2( t Hf )/MgO(2)/Ta(1) with a range of HfO2 thicknesses and from two control series, one with a Ta dusting layer (series B), Si/SiO2/Ta(6)/FeCoB( t FeCoB )/TaOx( t Ta )/MgO(2)/Ta(1), and another without any dusting layer (series C), Si/SiO2/Ta(6)/FeCoB( t FeCoB )/MgO(2)/Ta(1) (The numbers in parentheses are the thicknesses in nm). Since the complete oxidation of the insulator at the Fe alloy/oxide interface is held to be critical for the formation of PMA in HM/Fe alloy/oxide heterostructures,10 we performed x-ray photoelectron spectroscopy (XPS) measurements on an as-grown Ta(6)/FeCoB(1.2)/HfO2(0.2)/MgO/Ta series A sample. Ion etching was used to remove most of the Ta capping layer before performing XPS. As shown in Fig. 1(a), the HfO2 4f7/2 and 4f5/2 peaks are clearly displayed at 17.1 eV and 18.8 eV, respectively, with only a very small sub-oxide peak at ∼16.0 eV and no evidence for the Hf metallic 4f7/2 peak at 14.3 eV. To achieve strong PMA, it is also required that the Fe alloy should not be oxidized beyond the interfacial Fe-O bonds. Fig. 1(b) shows that for the same sample, the XPS 2p3/2 peak of Fe is at 706.0 eV, which can be fitted quite well with the narrow asymmetric spin-split peak function characteristic of metallic Fe.12 For a series C sample without the Hf dusting layer, the Fe 2p3/2 peak is much broader with a high energy tail indicative of substantial oxidation of the Fe surface during the direct deposition of MgO by rf sputtering13 (Fig. 1(b)). We also examined the Fe XPS signal for a series B sample with Ta as the dusting layer (0.3 nm), which yielded a metallic signal indistinguishable from the Hf case, again indicating the protection of the ferromagnetic surface from significant oxidation. However, the magnetic characteristics of these heterostructures are quite different.

FIG. 1.

XPS spectra from (a) HfO2 4f and (b) Fe 2p spectral regions for the as-grown samples Ta(6)/FeCoB(1.2)/HfO2(0.2)/MgO/Ta(1) and Ta(6)/FeCoB(1.3)/MgO/Ta(1).

FIG. 1.

XPS spectra from (a) HfO2 4f and (b) Fe 2p spectral regions for the as-grown samples Ta(6)/FeCoB(1.2)/HfO2(0.2)/MgO/Ta(1) and Ta(6)/FeCoB(1.3)/MgO/Ta(1).

Close modal

In Fig. 2(a), we present the magnetic moment per area as a function of t FeCoB for a set of series A samples (Hf dusting) and also for series B samples (Ta dusting), as measured by vibrating sample magnetometry (VSM). (The average thickness provided for the dusting layer throughout this paper is that of the deposited metal, before oxidation.) The fit to the Hf dusting series (A) gives a saturation magnetization of M s = 1260 emu / cm 3 and a very small apparent “dead layer” thickness t d 0.1 nm ; both are consistent with previous results from as-deposited Ta/FeCoB/MgO structures.1,11,14 In contrast, the series B samples indicate t d 0.8 nm and a much larger M s = 1800 emu / cm 3 ; the results are comparable to some previous studies of annealed (∼ 300 ° C ) Ta/FeCoB/MgO samples where the dead layer15 has been attributed to the undesirable diffusion of Ta into FeCoB, perhaps to the ferromagnet/oxide interface.16 Thus, we tentatively attribute the thick dead layer in series B samples to the intermixing of Ta and FeCoB during the deposition of the Ta dusting layer.

FIG. 2.

(a) VSM measurements of magnetization, (b) effective anisotropy energy density K eff determined from anomalous Hall measurements as a function of the in-plane magnetic field, and (c) anomalous Hall measurements as a function of the out-of-plane magnetic field, for the as-grown samples Ta(6)/FeCoB( t FeCoB )/HfO2(0.2)/MgO/Ta and Ta(6)/FeCoB( t FeCoB )/TaOx(0.2)/MgO/Ta. The solid and dashed straight lines are linear fits to the data. (d) The perpendicular anisotropy fields of Ta(6)/FeCoB(0.8)/HfO2( t Hf )/MgO/Ta samples as deposited and after different post-fabrication annealing treatments.

FIG. 2.

(a) VSM measurements of magnetization, (b) effective anisotropy energy density K eff determined from anomalous Hall measurements as a function of the in-plane magnetic field, and (c) anomalous Hall measurements as a function of the out-of-plane magnetic field, for the as-grown samples Ta(6)/FeCoB( t FeCoB )/HfO2(0.2)/MgO/Ta and Ta(6)/FeCoB( t FeCoB )/TaOx(0.2)/MgO/Ta. The solid and dashed straight lines are linear fits to the data. (d) The perpendicular anisotropy fields of Ta(6)/FeCoB(0.8)/HfO2( t Hf )/MgO/Ta samples as deposited and after different post-fabrication annealing treatments.

Close modal

While Ta/FeCoB/MgO structures with a thin FM layer typically only exhibit, at most, a weak perpendicular magnetic anisotropy (PMA) in the as-deposited state,3,14,15 we obtained robust PMA behavior in as-deposited structures with the HfO2 dusting layer. For example, in Figure 2(b), we plot the PMA energy density K eff μ 0 H a M s / 2 as a function of the effective thickness t FeCoB eff = t FeCoB t d of FeCoB for the series A Ta/FeCoB( t FeCoB )/HfO2(0.2)/MgO samples. Here, H a is the perpendicular magnetic anisotropy field as determined from the measurement of the first harmonic response of the anomalous Hall voltage to an oscillating current while a non-oscillatory external field is applied,17,18 and we use the values of M s determined from the VSM measurements discussed above. Since we can expect, at least approximately, that when t FeCoB eff is sufficiently large K eff · t FeCoB eff = ( K v 2 π M s 2 ) · t FeCoB eff + K s , where K v ( K s ) is the bulk (interfacial) anisotropy energy density, we can use a linear fit to this plot to determine the surface anisotropy K s = 1.74 ± 0.09 erg / cm 2 . For Ta and Hf base layer systems without the Hf dusting, comparable surface anisotropies, and effective anisotropies, can only be obtained via high temperature (≥ 200 C) annealing.19–21 In the inset of Fig. 2(b), we also show the effective PMA energy density K eff for the series B samples with a 0.2 nm Ta dusting layer. Notice that here K eff for Ta dusting is an order smaller than that for the Hf dusting.

Consistent with the strong Ha of the HfO2 passivated samples, the coercive field Hc of those PMA structures is relatively high, typically ≥300 Oe, in comparison to quite low values ( < 20 Oe ) for the Ta dusting samples. Examples of the field switching that is obtained with an external field applied normal to the film surface are provided in Fig. 2(c) for Ta(6)/FeCoB(1.1)/HfO2(0.2)/MgO/Ta(1) and Ta(6)/FeCoB(1.1)/TaOx(0.2)/MgO/Ta(1) samples. Since Hc of such PMA samples depends on both the anisotropy field and its uniformity, which together act to set the depinning field for magnetic reversal, further enhancement in Hc might be possible with improvements in the smoothness and uniformity of such heterostructures.

We measured the perpendicular anisotropy fields H a as a function of HfO2 thicknesses in a different set of samples Ta(6)/FeCoB(0.8)/HfO2( t Hf )/MgO/Ta with t Hf 0.2 0.4 nm , as indicated in Fig. 2(d). For the as-grown samples, H a increases with the HfO2 thickness and grows above 1 T when t Hf 0.3 nm . This is likely due to the fact that a more completely continuous HfO2 layer being formed at the FeCoB/MgO interface as t Hf is increased over this range and hence a higher Fe-O-Hf hybridized bond density that enhances the interfacial PMA.

Previously, high temperature post-fabrication annealing treatment has been considered to be necessary for the achievement of robust PMA in HM/FeCoB/MgO heterostructures. There are generally two important functions of this annealing process: (i) avoidance of the over-oxidation of the FeCoB surface that occurs during MgO deposition13,22 and (ii) promotion of the out-diffusion of boron from the initially amorphous FeCoB 13,19 to obtain a more ordered, crystalline FeCo/MgO interface. Our results here indicate that the first function is quite important. We also speculate that another possibility, which merits further research, is that the Fe-O-Hf hybridized bonds result in a stronger spin-splitting of the orbitals than the Fe-O-Mg bonds do.

Obtaining strong PMA in HM/Fe alloy/oxide systems without the necessity of thermal annealing may facilitate important applications as this could avoid complications such as material diffusion/intermixing during high temperature excursions. On the other hand, since many applications of PMA heterostructures do require high temperature processing, both for integration with Si circuits and to attain a high tunneling magnetoresistance (TMR) with MTJs, we also studied how different heat treatments affect the PMA of our HfO2 structures. We show in Fig. 2(d) that after annealing at 210 ° C for 1 h, H a increases for every HfO2 thickness studied, while the general dependence of H a on t Hf remains unchanged. However, after annealing at 300 ° C for 1 h, the PMA deteriorates, with a much weaker PMA retained only for t Hf 0.3 nm . We attribute this deterioration to the diffusion of Ta through the ferromagnetic layer to the MgO interface since such diffusion has been known to damage the interfacial PMA in the Ta based PMA systems.16 

We have also examined whether this Ta in-diffusion problem can be avoided by the use of other heavy metal base layers, especially those with strong spin Hall effects, e.g., W and Pt. In Fig. 3(a), we show the values of Ha obtained from a set of W(4)/FeCoB(0.8)/HfO2( t Hf )/MgO(1.6)/Ta samples as a function of t Hf for the as-deposited case, after 1 h at 300 ° C , and after 1 h at 410 ° C . Here, the W is in the high resistivity beta-W phase. The anisotropy increases with annealing temperature, and with 410 ° C vacuum annealing, we obtain Ha > 1.6 T for a sufficiently thick HfO2 passivation layer, indicative of an interfacial anisotropy energy density ≥ 1.5 ergs/cm2. When we use a 1 nm Ta seeding layer before the deposition of the W layer, it results in the W being smoother and being in the lower resistivity alpha-phase. As shown in Fig. 3(b), relatively high anisotropy fields are obtained after 300 ° C annealing of such Ta(1)/W(4)/FeCoB(0.8)/HfO2( t Hf )/MgO(1.6)/Ta samples for t Hf ≥ 0.1 nm but annealing at 410 ° C degrades Ha, particularly for the heterostructures with thinner HfO2, likely due to in-diffusion of Ta from the bottom seeding layer.

FIG. 3.

The perpendicular anisotropy fields of (a) beta-W(4)/FeCoB(0.8)/HfO2( t Hf )/MgO/Ta and (b) Ta/alpha-W(4)/FeCoB(0.8)/HfO2( t Hf )/MgO/Ta after different post-fabrication annealing treatments. (c) Anomalous Hall measurements of the as-grown samples Ta(6)/NiFe(1.4)/HfO2(0.2)/MgO/Ta and Ta(6)/Hf(0.5)/NiFe(1.5)/HfO2(0.2)/MgO/Ta as a function of the in-plane magnetic field. (d) The perpendicular anisotropy fields of MgO(1.6)/FeCoB( t FeCoB )/HfO2(0.3)/MgO(0.8)/Ta samples after different post-fabrication annealing treatments.

FIG. 3.

The perpendicular anisotropy fields of (a) beta-W(4)/FeCoB(0.8)/HfO2( t Hf )/MgO/Ta and (b) Ta/alpha-W(4)/FeCoB(0.8)/HfO2( t Hf )/MgO/Ta after different post-fabrication annealing treatments. (c) Anomalous Hall measurements of the as-grown samples Ta(6)/NiFe(1.4)/HfO2(0.2)/MgO/Ta and Ta(6)/Hf(0.5)/NiFe(1.5)/HfO2(0.2)/MgO/Ta as a function of the in-plane magnetic field. (d) The perpendicular anisotropy fields of MgO(1.6)/FeCoB( t FeCoB )/HfO2(0.3)/MgO(0.8)/Ta samples after different post-fabrication annealing treatments.

Close modal

While most PMA heterostructure research currently utilizes either Ta/FeCoB/MgO or Pt/Co/Oxide multilayers, where in the latter case, the PMA is concluded to originate largely from spin orbit effects at the Pt/Co interface, other magnetic layers with attractive properties, such as Ni80Fe20, could be of possible interest and value if samples of sufficiently strong anisotropy can be produced. A previous study with Ni80Fe20 has reported that when it is grown on an MgO underlayer, PMA can be achieved with proper capping materials (i.e., with a MgO/Ni80Fe20/capping layer structure).23 Here, we find that significant interfacial anisotropy can also be obtained with a suitable combination of HfO2 and Ni80Fe20, e.g., with Ta/Ni80Fe20/HfO2/MgO and with Ta/Hf(0.5)/Ni80Fe20/HfO2/MgO multilayers. Our best results have been obtained with an amorphous Hf(0.5) spacer between the Ta base layer and NiFe, which presumably helps to accommodate the crystalline mismatch between Ta and NiFe. In Fig. 3(c), we show anomalous Hall measurements as a function of an in-plane magnetic field for as-deposited Ta based NiFe(∼1.5)/HfO2(0.2)/MgO samples with and without the Hf spacer at the Ta/NiFe interface. Ha for the structure without the Hf spacer was 1.1 kOe, while for the sample with the 0.5 nm Hf spacer, Ha is doubled to 2.1 kOe, larger than the previous results.23 In particular, the coercive field of Ta/Hf(0.5)/Ni80Fe20(1.5)/HfO2/MgO is around 100 Oe (not shown here). We have also obtained PMA in Pt/Hf(0.5)/FeCoB/HfO2(0.2)/MgO multilayers both as deposited and after 300 ° C annealing. We conclude that the combination of a HfO2 passivation layer at the Fe alloy/oxide interface and a thin Hf spacer layer between the HM and the Fe alloy (when needed due to crystalline mismatch between the HM and the Fe alloy) can be a robust strategy for engineering the PMA of a range of thin-film magnetic heterostructures.

Finally, a recent development in MTJ technology for spin transfer torque applications is to include a second, thinner MgO layer on the other side of the FeCoB free layer, opposite to the MgO tunnel barrier interface.24,25 This enhances K e f f of the free layer, permitting the use of a thicker layer with more thermal stability, and also suppresses the magnetic damping enhancement that would otherwise occur via spin pumping to the adjacent normal metal contact. We have examined a modification of this approach by depositing multilayer stacks of MgO(1.6)/FeCoB( t FeCoB )/HfO2(0.2)/MgO(0.8)/Ta onto oxidized Si substrates. Before annealing, these samples do not exhibit PMA, which we attribute to the deleterious effect of chemisorbed oxygen on the surface of the sputter deposited MgO after deposition13 that, we presume, oxidizes the lower surface of the subsequently deposited FeCoB. After annealing, there is strong PMA. In Fig. 3(d), we show a plot of Ha of such samples as a function of tFeCoB. Quite strong anisotropy fields are obtained at high values of t FeCoB , particularly for the samples annealed at 370 ° C . Field modulated ferromagnetic resonant studies of such a heterostructure with t FeCoB  = 1.6 nm yielded a magnetic damping parameter α = 0.009 , while a Ta/FeCoB(1.6)/MgO(1.6)/Ta sample showed α = 0.02 , consistent with earlier work.1,26

An important question in terms of applications is whether MTJs with a HfO2 passivation layer at the tunnel barrier-free layer interface can provide sufficiently high TMR that is useful for STT and other spintronics applications. As reported previously,27 a TMR of 80% has been achieved with an in-plane magnetized Pt/Hf/FeCoB (1.6)/MgO(1.6)/FeCoB/Ru/Ta MTJ structure annealed at 300 C, where analytical STEM reveals substantial Hf within the tunnel barrier and the greatly reduced demagnetization field, ≈ 4 kOe, indicates a substantial K s . As will be reported elsewhere,28 we are now obtaining similar results with W base layer MTJs, and so, hybrid HfO2MgO tunnel barriers may provide sufficient TMR to be of technology interest.

In summary, we have demonstrated that perpendicular magnetic anisotropy in HM/Fe alloy/MgO heterostructures can be dramatically strengthened by incorporating a very thin HfO2 dusting layer at the Fe alloy/MgO interface. In HM/FeCoB/MgO devices, the dusting layer enables strong PMA even in the absence of the post-deposition annealing step that has previously been necessary. When annealing is desired, the dusting layer allows the PMA to remain strong for annealing temperatures even above 400 ° C , provided that a proper base layer is utilized, a much higher limit than for some current spin-transfer torque magnetic random-access memory prototype technologies. This can allow easier integration with Si circuitry. The HfO2 dusting can also create robust PMA using magnetic materials beyond just FeCoB, thereby expanding the portfolio of magnetic materials available for PMA technologies. In particular, we have demonstrated PMA with thin-film Ni80Fe20, a material that is attractive for its low damping and low magnetostriction. Overall, the strengthening of PMA with the use of HfO2 dusting layers has great promise both for enhancing the performance of spin-transfer-torque magnetic memory based on PMA magnetic tunnel junctions and also for improving the control of chiral domain walls and skyrmion structures within PMA HM/Fe alloy/MgO structures.5,6,29–31

We thank Darren Dale for the assistance with the XPS measurements and analysis. This research was supported by ONR and by NSF/MRSEC (DMR-1120296) through the Cornell Center for Materials Research (CCMR) and by NSF through the use of the Cornell Nanofabrication Facility (CNF)/NINN (ECCS-1542081) and the CCMR facilities.

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