A perpendicular-anisotropy CoFeB2013MgO magnetic tunnel junction

A perpendicular-anisotropy CoFeB2013MgO magnetic tunnel junction

(Parte 1 de 2)

A perpendicular-anisotropy CoFeB–MgO magnetic tunnel junction

Magnetic tunnel junctions (MTJs) with ferromagnetic electrodes possessing a perpendicular magnetic easy axis are of great interest as they have a potential for realizing next-generation high-density non-volatile memory and logic chips with high thermal stability and low critical current for current-induced magnetization switching1–3. To attain perpendicular anisotropy, a number of material systems have been explored as electrodes, which include rare-earth/transition-

high thermal stability at reduced dimension, low-current current-induced magnetization switching and high tunnel magnetoresistance ratio all at the same time. Here, we use interfacial perpendicular anisotropy between the ferromagnetic electrodes and the tunnel barrier of the MTJ by employing the material combination of CoFeB–MgO, a system widely adopted to produce a giant tunnel magnetoresistance ratio in MTJs with in-plane anisotropy11–13. This approach requires no material other than those used in conventional in-plane-anisotropy MTJs. The perpendicular MTJs consisting of Ta/CoFeB/MgO/CoFeB/Ta show a high tunnel magnetoresistance ratio, over 120%, high thermal stability at dimension aslowas40nmdiameterandalowswitchingcurrentof49µA.

The three conditions that high-performance perpendicular

MTJs need to satisfy impose a stringent set of requirements on the materials to be used in the MTJ structure. First of all, the thermal stability factor E/kBT of the free (recording) layer needs to be more than 40 (ref. 14) for non-volatility, where E = MSHKV/2 is the energy barrier that separates the two magnetization directions; here,

MS is the saturation magnetization, HK the anisotropy field, kB the Boltzmann constant and T the temperature. Because the volume V of the free layer reduces as the junction dimension is reduced, the anisotropy energy density K = MSHK/2 needs to be high enough to ensure high thermal stability. A number of perpendicular- anisotropy materials such as FePt satisfy this first condition15.

However, the intrinsic threshold current IC0 for current-induced magnetization switching (CIMS) is proportional to E,

IC0 =α γe µBg MSHKV =2α γe µBg E (1) where α is the magnetic damping constant, γ the gyromagnetic ratio, e the elementary charge, µB the Bohr magneton and g a function of the spin polarization of the tunnel current and the angle between the magnetizations of the free and the reference layers16,17. Note that for in-plane-anisotropy MTJs E in equation (1) is

1Center for Spintronics Integrated Systems, Tohoku University, 2-1-1 Katahira, Aoba-ku, Sendai 980-8577, Japan, 2Laboratory for Nanoelectronics and Spintronics, Research Institute of Electrical Communication, Tohoku University, 2-1-1 Katahira, Aoba-ku, Sendai 980-8577, Japan, 3Hitachi Ltd, Advanced Research Laboratory, 1-280 Higashi-koigakubo, Kokubunji-shi, Tokyo 185-8601, Japan. *e-mail:sikeda@riec.tohoku.ac.jp; ohno@riec.tohoku.ac.jp.

replaced by the demagnetization energy Edemag, resulting in large E, whichisthereasonwhyperpendicularanisotropyisrequiredforthe reduction of switching current. This equation shows that low α is neededforlowswitchingcurrentforagivenE.However,commonly known perpendicular-anisotropy materials and structures use noble metals with high spin–orbit interaction18, which increases α (refs 3,19–21). For example, the typical α is larger than 0.1 for Co/Pt (ref. 19). In addition, there is no established material system that provideshightunnelmagnetoresistance(TMR)ratioapartfromthe well-known body-centred cubic (bcc) (001) CoFe(B)–MgO system. The crystal structures of perpendicular-anisotropy materials are usuallydifferentfrombcc,andonannealingtheinitiallyamorphous CoFeB tends to crystallize in structures other than the wanted bcc because they are deposited in direct contact with the perpendicularanisotropy materials10. In the following, we show that all three conditions for high-performance perpendicular MTJs can be met with the CoFeB–MgO standard material system that is widely used for in-plane-anisotropy MTJs.

All the stack structures in this study are prepared on thermally oxidized Si(001) substrate by RF sputtering at room temperature13. The MTJ structures consist of, from the substrate processed into circular devices with a 40 or 150nm diameter by electron-beam lithography and Ar-ion milling (Fig. 1b). For magnetization M and ferromagnetic resonance (FMR) measurements, two kinds of stack structure with a single CoFeB layer are prepared: CoFeB (1.0–20)/MgO (1.0), which corresponds to the bottom CoFeB layer in the MTJ, and the reversed structure, MgO (1.0)/CoFeB (0.5–3.0)/Ta (5), which corresponds to the top CoFeB layer in the MTJ. The former is deposited on a Ta/Ru/Ta buffer layer and the latter on a Ta buffer layer. The completed

MTJs/stacked structures are annealed at a temperature Ta ranging from 250 to 400◦C in a vacuum under a perpendicular magnetic field of 400mT for an hour.

Figure 2 shows the in-plane and out-of-plane magnetization versus external magnetic field (M–H) curves for annealed and tCoFeB = 1.3nm (Fig. 2b). The sample with tCoFeB = 2.0nm has an in-plane easy axis with out-of-plane saturation field much smaller than the saturation magnetization MS, consistent with earlier studies22–24 indicating the presence of a perpendicular- anisotropy component. The sample with tCoFeB = 1.3nm shows a clear perpendicular easy axis with in-plane saturation field

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LETTERS NATUREMATERIALSDOI:10.1038/NMAT2804 V

V 40nm

5nmRu

10nm Ru

5nmTa Al O

5nmTa

Si/SiO sub. 5 nm Ta t CoFeB t MgO

Cr/Au t CoFeB

Figure 1 | MTJ structure. a, Schematic of an MTJ device for TMR and CIMS measurements. b, Top view of an MTJ pillar taken by scanning electron microscope.

(µ0: permeability in free space). The saturation magnetization is 1.58T. The perpendicular-anisotropy energy density K at this CoFeB thickness, which determines the thermal stability, is 2.1 × 105 Jm−3, a value comparable to that of the Co–Pd perpendicular multilayers25 and high enough to secure good thermal stability at reduced dimensions (40nm diameter). To separate the bulk and interfacial contribution of the anisotropy, the tCoFeB dependence of K = Kb −MS2/2µ0 +Ki/tCoFeB is measured, as shown in the inset of Fig. 2b. Here, Kb is the bulk crystalline anisotropy and Ki the interfacial anisotropy. From the intercept, Ki is determined to be 1.3mJm−2. The bulk contribution is consistent with the demagnetization (−MS2/2µ0), indicating that

Kb is negligible, that is, that the perpendicular anisotropy in this systemisentirelyduetotheCoFeB–MgOinterfacialanisotropy26.

From FMR measurements, the information of HK and α can be obtained.WehavemeasuredFMRspectraatamicrowavefrequency of 9.0GHz for annealed CoFeB/MgO samples at Ta = 300◦C as a function of the angle θ between H and the normal axis to the sample surface as shown in Fig. 3a. The θ dependencies of resonant field HR and linewidth (full-width at half-maximum, FWHM)aresummarizedinFig. 3a,b,fromwhichwecandetermine

HK and α (ref. 19). Figure 3d,e shows the tCoFeB dependence of the obtained HK and α. The HK increases as thickness reduces and changes its sign reflecting the change of magnetic-easy-axis direction around tCoFeB =1.5nm. The tCoFeB dependence of K ·tCoFeB is plotted together with that obtained from M–H curves in the inset of Fig. 2, showing good correspondence between the two measurements. Although the magnitude of α steeply increases as thickness decreases below 2 nm, it is still smaller than those for materials including noble metals18. Full understanding of the origin of the increase is important to further reduce IC0. The interfacial perpendicular anisotropy between oxide and ferromagnetic metal (Fe/MgO) has been predicted by firstprinciples calculation and attributed to hybridization of Fe 3d and O 2p orbitals27. Although earlier experimental studies also indicated the presence of perpendicular anisotropy at the interface in Pt/Co/MOx (M = Al, Mg, Ta and Ru) trilayer structures28,29 and in MgO/CoFeB/Pt (ref. 30), these structures always contained Pt in direct contact with ferromagnetic transition metals to stabilize the perpendicular anisotropy, which made the origin of the anisotropy ambiguous. As demonstrated in the following, the interfacial anisotropybetweenMgOandCoFeBislargeenoughtorealizehighperformance perpendicular CoFeB–MgO MTJs and no addition of noble metal is necessary.

In plane Out of plane

0 (mJ m

K. t CoFeB tCoFeB = 2.0 nm tCoFeB = 1.3 nm

Figure 2 | In-plane and out-of-plane magnetization curves for

CoFeB/MgO. a, tCoFeB =2.0nm. b, tCoFeB =1.3nm. Inset: tCoFeB dependence of the product of K and tCoFeB, where the intercept to the vertical axis and the slope of the linear extrapolation of the data correspond to Ki and

Kb−MS2/2µ0. Circles and squares are obtained from magnetization and FMR measurements, respectively.

Now we turn to the TMR properties of perpendicular MTJs. Figure 4a,b shows junction resistance R as a function of H (R–H curves) of a 150-nm-diameter MTJ annealed at Ta = 300◦C, with two different magnetic-field directions. The top and bottom CoFeB electrodes of the MTJ have nominally identical tCoFeB of 1.3nm, and tMgO is 0.9nm. Reflecting the perpendicular anisotropy, the R–H curve shows a clear hysteresis with distinct high- (antiparallel

M configuration: AP) and low-R (parallel M: P) states (TMR ratio of 100%) when the magnetic field is applied out of plane, whereas the in-plane R–H curve shows virtually constant R.

The coercivity HC is much larger than those shown in Fig. 1b taken on a millimetre-size sample, most probably owing to the suppression of domain-structure formation. The obtained HC, however, is smaller than 2K/MS, suggesting that there remains a contribution of domain nucleation to H-induced magnetization reversal in these MTJ structures. The HC difference between the nominally identical electrodes may be due to different degrees of intermixing at the two Ta–CoFeB interfaces during sputtering31 and/or different areas of the two electrodes because of a taper of

MTJ pillar introduced during ion milling; the tCoFeB dependence of MS indicates an approximately 0.5-nm-thick magnetically dead layer in the CoFeB/Ta interface (corresponding to the top layer in the MTJ) and no signature of a dead layer for Ta/CoFeB/MgO (corresponding to the bottom layer) (not shown). Perpendicular magnetic anisotropy with a clear R–H hysteresis is obtained at Ta greater than 250◦C, and the TMR ratio increases monotonically with increasing Ta and reaches 121% after annealing at Ta =350◦C, as shown in Fig. 4c. It should be noted that 350◦C annealing is required for integration with complementary metal–oxide– semiconductor transistors. Further increase of Ta leads a decrease of the TMR ratio.

Next, to show the potential of this material system at reduced dimensions, circular 40-nm-diameter MTJs are fabricated. Figure 5a shows an R–H curve of such an MTJ. The MTJ has tCoFeB = 1.0 and 1.7nm for bottom and top CoFeB layers, respectively, and tMgO = 0.85nm. The MTJ is annealed at 300◦C. The TMR ratio is 124% with resistance–area product RA = 18 µm2. The minor loop of the top free layer (the free layer is identified from the CIMS measurement; see below) is shifted by 37mT with respect

722 NATURE MATERIALS | VOL 9 | SEPTEMBER 2010 | w.nature.com/naturematerials

NATUREMATERIALSDOI:10.1038/NMAT2804 LETTERS =

= 0.027= 0.28 T FWHM (T)

0 Derivative absorption

(arb. units)

α t = 1.3 nm t (nm) t(nm)H (T)μH μ b d

Figure 3|FMR spectra for CoFeB/MgO and obtained material parameters. a, Angle dependence of FMR spectra for CoFeB/MgO with magnetic fields, from which HK is determined. The solid line shows the fitting result. c, Angle dependence of linewidths (FWHM), from which α is determined. The solid line shows the fitting result. d, tCoFeB dependence of HK. e, tCoFeB dependence of α.

to H = 0, indicating the existence of dipolar interlayer coupling between the two CoFeB layers.

Finally, we show the results of CIMS in a 40-nm-diameter MTJ. Figure 5b shows the resistance versus applied current density

(R–J curves) measured at two current-pulse durations τP, 300µs and 1.0s, in the absence of an external magnetic field. We can see clear switching between high- and low-resistance states, whose magnitudes correspond to high- (AP) and low-resistance (P) states

R (k

R (k

(Parte 1 de 2)

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