Magnetization reversal and enhanced tunnel magnetoresistance ratio inperpendicular magnetic tunnel junctions based on exchange springelectrodesYi Wang, Xiaolu Yin, D. Le Roy, Jun Jiang, H. X. Wei et al. Citation: J. Appl. Phys. 113, 133906 (2013); doi: 10.1063/1.4798507 View online: http://dx.doi.org/10.1063/1.4798507 View Table of Contents: http://jap.aip.org/resource/1/JAPIAU/v113/i13 Published by the American Institute of Physics. Additional information on J. Appl. Phys.Journal Homepage: http://jap.aip.org/ Journal Information: http://jap.aip.org/about/about_the_journal Top downloads: http://jap.aip.org/features/most_downloaded Information for Authors: http://jap.aip.org/authors

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Magnetization reversal and enhanced tunnel magnetoresistanceratio in perpendicular magnetic tunnel junctions basedon exchange spring electrodes

Yi Wang,1,2 Xiaolu Yin,1 D. Le Roy,1 Jun Jiang,2 H. X. Wei,2 S. H. Liou,1 and X. F. Han2,a)

1Department of Physics and Astronomy, Nebraska Center for Materials and Nanoscience, University ofNebraska-Lincoln, Lincoln, Nebraska 68588, USA2Beijing National Laboratory for Condensed Matter Physics, Institute of Physics, Chinese Academy ofSciences, Beijing 100190, China

(Received 23 January 2013; accepted 13 March 2013; published online 2 April 2013)

The ½Co=Pt�n multilayer based perpendicular magnetic tunnel junction stacks with wedged

Co60Fe20B20 insertions up to 2 nm, and corresponding perpendicular magnetic tunnel junctions

were magnetically and electrically investigated. The focus is on the influence of CoFeB insertions

in the free and reference electrodes on the overall junction magnetization reversal and

magnetoresistance response. The exchange spring behavior was revealed as the Co60Fe20B20 spins

canting towards the in-plane direction in the ½Co=Pt�n=Co60Fe20B20 hard/soft perpendicular

magnetic electrodes. The broad range thickness of wedged Co60Fe20B20 insertion enables to reveal

the critical transition, in particular, from rigid coupling to exchange spring coupling. With the help

of 375 �C annealing under 10 kOe magnetic field, the recovery from distinct multi-domain

structure to nearly single domain structure was distinctly observed in the unpatterned perpendicular

magnetic tunnel junction (p-MTJ) films with CoFeB thickness tCFB � 1:5 nm. Meanwhile, for the

corresponding patterned perpendicular magnetic tunnel junctions with AlOx barrier, the tunnel

magnetoresistance (TMR) ratio exhibited an intense enhancement over 100%. The TMR results

and spin configurations were illustrated using an exchange spring model in both magnetic

electrodes. The presented study shows the benefit of using exchange spring magnetic electrodes in

perpendicular magnetic tunnel junction on their performance. VC 2013 American Institute ofPhysics. [http://dx.doi.org/10.1063/1.4798507]

I. INTRODUCTION

Over the last few years, magnetic tunnel junctions

(MTJ) with perpendicular magnetic anisotropy (PMA) have

been intensively investigated from both the fundamental and

technological viewpoints.1–10 Basically, perpendicular mag-

netic tunnel junctions (p-MTJs) possess some fascinating

advantages over conventional in-plane MTJs, as shown in

theoretical and experimental investigations,3,11–14 which illu-

minate the realization of the emerging spin transfer torque

magnetic random access memory (STT-MRAM) with high

performance. However, in the early studies, the tunnel mag-

netoresistance (TMR) ratio of p-MTJs was not as high as

that of in-plane MTJs15–19 for applications. Subsequently,

inspired by the wonderful performance of MgO tunnel bar-

rier for in-plane MTJs,19 some pioneer studies have been

performed to enhance the TMR ratio for p-MTJs via directly

replacing AlOx tunnel barrier with MgO barrier. However,

low TMR ratios at room temperature were obtained in

p-MTJs with ordered alloy or multilayer perpendicular mag-

netic electrodes.20,21 This was generally attributed to either

the large lattice mismatch between ordered magnetic electro-

des and MgO barrier (about 9.5% between L10 CoPt alloy

and MgO) or the fcc (111) polycrystalline structure of

as-grown magnetic multilayer electrodes (such as Co/Pt or

CoFe/Pd multilayers), which would not promote the forma-

tion of good bcc (001) MgO barrier after annealing. As a

consequence, spin coherent tunneling and spin filter effect

cannot be well realized and thus no high TMR ratio was

observed.22–24 These critical issues were tentatively

addressed in recent studies. A better crystallization of MgO

barrier and thus an over 100% TMR ratio could be achieved

by inserting Fe or CoFe thin layer between L10 CoPt (or

FePt) and MgO barrier.3 Meanwhile, for another category of

perpendicular magnetic electrodes such as Co/Pt (or Co/Pd,

CoFe/Pd) multilayer, the critical issue is how to diminish the

unwished effect of fcc (111) magnetic electrodes on the crys-

tallization of bcc (001) MgO barrier during the annealing

process. As we all know, the MgO barrier grown on amor-

phous CoFeB shows superior crystalline quality after anneal-

ing.19 Therefore, CoFeB insertions were introduced to the

conventional CoFe/Pd based p-MTJs by Ohno et al., leading

to a substantial increase of the TMR ratio, from 1.7% to

120% at room temperature.7,21 In addition, a high tempera-

ture annealing stability up to 350 �C was improved by intro-

ducing CoFeB high spin polarizing layer to MgO-based

p-MTJs with Co/Pd multilayers.25 Recently, Ikeda et al.have reported the realization of p-MTJs with single CoFeB

electrodes and an over 120% TMR ratio at room tempera-

ture.5 However, the phase diagram for CoFeB single film by

Wang et al. revealed that the optimal thickness of perpendic-

ular anisotropic CoFeB layer is around 1.0 nm actually after

taking into account of magnetic dead layer.26

a)Author to whom correspondence should be addressed. Email:

xfhan@aphy.iphy.ac.cn

0021-8979/2013/113(13)/133906/8/$30.00 VC 2013 American Institute of Physics113, 133906-1

JOURNAL OF APPLIED PHYSICS 113, 133906 (2013)

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To date, the efforts have been mostly focused on achiev-

ing good perpendicular magnetic electrodes for p-MTJs ei-

ther by rigidly coupled CoFeB/PMA layer or by single

CoFeB thin films, which would somewhat limit the CoFeB

thickness in p-MTJs. Hence, one might not realize the full

polarization of electrons or perfect crystallization of MgO

barrier to obtain TMR ratio as high as that of in-plane

MTJs.19 In other words, the TMR ratio of p-MTJs has poten-

tial to be enhanced if the thickness of CoFeB insertions could

be further increased, indeed. Meanwhile, the magnetic elec-

trodes with thicker CoFeB soft phase would exhibit

exchange spring (ES) behavior in the hard/soft perpendicular

ES (p-ES) system, in which the soft phase are exchange

coupled but not rigidly coupled with a perpendicular hard

phase at the interface.27–31 As reported by Sousa et al.,29 in

the FePt/Fe exchange spring bilayer at zero magnetic field,

the spins in soft Fe layer would start to exhibit continuous

rotation towards in-plane direction, with the tilt angle gradu-

ally increasing with increasing distance from the hard per-

pendicular FePt layer, even finally the toppest Fe spins

nearly lay in-plane orientation. Here, the tilt angle of spins

could also be flexibly tuned via selecting the hard phase

materials or the interfacial exchange coupling strength Jex

besides the soft-layer thickness.27,32,33 More attractively, the

addition of the soft phase does not weaken but even gain the

thermal stability of the total materials.34,35 As a conse-

quence, it should be of great interest to introduce the novel

perpendicular exchange spring electrodes into perpendicular

magnetic tunnel junctions for spintronics application.

In this paper, the p-MTJs were fabricated with core struc-

ture of ½Co=Pt�n=Co=Co60Fe20B20ðtCFBÞ=AlOx=Co60Fe20B20

ðtCFBÞ=Co=½Co=Pt�n consisting of wedged Co60Fe20B20 inser-

tions with large thickness variation. Via continuously varying

the Co60Fe20B20 thickness and the annealing temperature, we

could systematically investigate this kind of p-MTJs in terms of

magnetic domain structures, magnetization reversal, and corre-

sponding magnetotransport properties. It is greatly worth noting

that the reasons for adopting AlOx barrier rather than the cur-

rently popular MgO barrier are that (1) the present research can

be well performed on the basis of our previous research on

p-MTJs with AlOx barrier without CoFeB insertions,36 (2)

more importantly, very few research work reported have intro-

duced p-ES magnetic electrodes into p-MTJs,25 our focus here

however is on fundamentally clarifying the magnetization re-

versal mechanisms and spin configurations in both magnetic

electrodes after different temperature annealing. It could be bet-

ter deduced from the corresponding magnetoresistance in

p-MTJs with amorphous AlOx barrier than that with MgO bar-

rier, due to the different quality of bcc (001) crystallization of

MgO barrier after different temperature annealing would signif-

icantly influence the magnetoresistance and thus introduce dis-

turbance for the analysis of magnetization reversal process and

spin configurations. The results here can serve as a good basis

for further development of high TMR ratio in MgO based p-

MTJs.

II. EXPERIMENTAL DETAILS

The wedged p-MTJ stacks adopted hard/soft perpendicu-

lar magnetic multilayer electrodes with the structure of Seed

layer /[Co(0.4)/Pt(2)]10/Co(0.4)/Co60Fe20B20ðtCFBÞ/AlOxð1Þ/Co60Fe20B20ðtCFBÞ/Co(0.6)/[Pt(2)/Co(0.4)]5/cap layer (unit in

nm) were deposited on a 50 mm long Si=SiO2 wafer substrate

by using ULVAC ultrahigh vacuum sputtering system with a

base pressure better than 7� 10�7 Pa at room temperature, as

shown schematically in Fig. 1(a), where tCFB varies from 0 nm

to 2 nm. The wedged Co60Fe20B20 layer was achieved by a

moving shutter in the chamber. Learning from the magnetic

hysteresis loop (M-H loop), Co60Fe20B20ðtCFBÞ/Co(0.6)/[Pt(2)/

Co(0.4)]5 are top (free) electrode with lower coercivity. The

deposition rates of Co, Pt, and Co60Fe20B20 are 0.53, 1.05, and

0.44 A/s, respectively. As shown in Fig. 1(b), finally, the

whole wedged p-MTJ stack was cut into two strips, each of

them covering tCFB from 0 nm to 2 nm with CoFeB variation

of 0.0444 nm/mm. The first strip was cut into 11 small pieces,

in which the first 10 small pieces had equal length of 4.5 mm,

resulting in a Co60Fe20B20 (hereafter denoted as CoFeB)

thickness interval of 0.2 nm from one piece to a successive

one. And the last remained 5 mm length piece was the one

with tCFB ¼ 0nm. The second strips were patterned into p-

FIG. 1. (a) The schematic diagram of

wedged shape p-MTJ stacks (not to

scale), (b) two strips cut from the whole

wedged p-MTJ stack each covering tCFB

from 0 nm to 2 nm, one strip (left) is cut

into 11 pieces of films, the other strip

(right) is patterned into p-MTJs,

correspondingly.

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MTJs with rectangular junction of aspect ratio of 3 (i.e.,

12 lm by 4 lm) by contact UV lithography. The CoFeB thick-

ness could be accurately determined via the distance from

each junction to the edge of film, where tCFB ¼ 2nm. The p-

MTJ films were annealed at TA ¼ 200 �C–375 �C for 40 min

in a high vacuum furnace with a base pressure below 10�5 Pa

under an applied perpendicular magnetic field of 10 kOe. In

addition, another two reference samples are also prepared with

structures of Si=SiO2-sub/Tað10Þ=Ptð10Þ=½Coð0:4Þ=Ptð2Þ�10/

Co(0.4)/Ta(4) (referred as S1) and Si=SiO2 � sub=Tað10Þ=Ptð30Þ=AlOxð1Þ=Coð0:6Þ=½Ptð2Þ=Coð0:4Þ�5=Ptð10Þ (referred

as S2), respectively (unit in nm). The magnetization hysteresis

loops (M-H loops) and magnetoresistance curves (R-H curves)

were measured by Alternating Gradient Force Magnetometer

(AGFM), Superconducting Quantum Interference device mag-

netometer (SQUID), and standard four probe method at room

temperature in the perpendicular magnetic field, respectively.

The magnetic domain structures were probed by Magnetic

Force Microscope (MFM) measurement performed in tapping

mode with the high Hc and low Ms tip by a Dimension 3000

mode system. All the MFM images were obtained at rema-

nence state after out-of-plane saturation.

III. RESULTS AND DISCUSSION

A. Unpatterned p-MTJ films

Figure 2 shows the MFM images of p-MTJ films with

selected tCFBðsÞ at as-grown, 300 �C, and 375 �C annealing

temperature, respectively. For cases of both the as-grown

and 300 �C annealing films, the conventional transition from

single domain structure to multi-domain structure was dis-

tinctly observed at the critical thickness tCFB ¼ 0:9 nm (as-

grown) and tCFB ¼ 1:1 nm (300 �CÞ. The multi-domain

exhibits a typical maze-like domain pattern, where the dark

and light contrasts represent domains with magnetization

pointing out of and into the film plane, respectively. This

transition should be attributed to the competition of perpen-

dicular magnetic anisotropy energy (K?) from the hard Co/

Pt multilayer and the in-plane magnetic anisotropy energy

(Kk) arising from the soft CoFeB insertion layers in the hard/

soft typed magnetic electrodes.33 Basically, in the regions of

0 nm � tCFB < 0:9 nm (as-grown) and 0 nm � tCFB < 1:1 nm

ð300 �CÞ, the magnetic moments in the thin CoFeB layer are

pinned by the hard Co/Pt perpendicular magnetic multilayer

due to the exchange coupling at the interface. It results in a

single domain structure at remanence state. For tCFB �0:9 nm (as-grown) and tCFB � 1:1nm ð300 �CÞ, a multi-

domain structure was revealed and the amount of reverse

domains rose monotonically with increasing tCFB. The

increase of the total magnetization coming from CoFeB

insertions leads to a domination of the demagnetizing effects

over the interfacial perpendicular pinning and promotes the

formation of multi-domain structure. Moreover, the shift of

the critical tCFB from 0.9 nm to 1.1 nm due to the increased

annealing temperature under 10 kOe perpendicular magnetic

field revealed the enhancement of perpendicular magnetic

anisotropy of the p-MTJ films. More interestingly, for the p-

MTJ films after 375 �C annealing, significant changes

revealed in the MFM images. At tCFB ¼ 0:5 nm, the film

abruptly broke into multi-domain structure. Subsequently,

the amount of reverse domains did not increase after it

reached the maximum around tCFB ¼ 1:0 nm, in contrast, it

began to decrease and finally a completely different domain

patterns (the initial distinct perpendicularly oriented domain

patterns disappeared) were observed at tCFB � 1:5 nm,

although a very blurry magnetic contrast seems to be present

FIG. 2. The remanence state MFM images of p-MTJ films with different thickness of CoFeB insertions tCFB (denoted on top) and annealing temperature TA

(denoted on left)

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in these films. This will be interpreted later. It is also noted

that from the M-H loop at tCFB ¼ 0 nm (not shown here), the

coercivity increased in both top and bottom magnetic elec-

trodes. It suggests the good perpendicular magnetic anisot-

ropy resulting from either the improvement of the magnetic

electrode texture or the formation of a little amount of CoPtx

alloy at the Co and Pt interface.37

To clarify the observations and obtain good insight into

the magnetization evolution, the squareness of Mr=Ms and

the nucleation field Hn were extracted from M-H loops and

plotted with respect to both tCFB and TA, as shown in Figs.

3(a) and 3(b), accompanying with some representative M–Hcurves. Here, the Mr and Ms represent the remanent magnet-

ization and saturation magnetization for the total p-MTJ

films, respectively. As reported in some previous research

work,38,39 the magnetization reversal process of our wedged

Co/Pt based electrodes occurs in two steps, with a first nucle-

ation of reversed domains followed by an abrupt irreversible

magnetization switching. The magnetic field where rapid

decrease of magnetization starts is defined as Hn, which is

denoted by the arrow in the measured M-H loops. As deter-

mined by the reference samples S1 and S2, the K? of the bot-

tom magnetic electrode without CoFeB is about six times

larger than that of the top magnetic electrode. Moreover,

according to the reports on the analogous hard/soft struc-

ture,34,40 with the same thickness CoFeB soft layer in both

magnetic electrodes, the K? of bottom electrode remains the

larger one. Hence, Hn is reasonable to represent the nuclea-

tion field of total top magnetic electrode. It is worth noting

that a very thin enough CoFeB magnetic layer deposited on

insulating barrier layer typically can lead to the formation of

superparamagnetic nano-islands.41 However, it is a different

case in the present work, the CoFeB is deposited with a Co-

terminated Co/Pt multilayer. The superparamagnetism was

not observed because CoFeB is ferromagnetic coupled with

high quality perpendicular Co/Pt multilayers tightly and it

remains beyond the superparamagnetic threshold. Figures

3(a) and 3(b) show four distinguishable regions, namely,

regions A, B, C, and D.

For all the cases of TA � 350 �C, in the region A

(tCFB � 0:3 nmÞ, good squareness value of Mr=Ms larger

than 0.8 were observed, which means the CoFeB spins are

almost well pinned along perpendicular direction. However,

Hn drastically reduced with tCFB increasing. In the region B

(0:3 nm < tCFB < 0:9 nm), Mr=Ms decreased slightly with

tCFB increasing. As revealed in MFM images, the single do-

main structures still remained. Consequently, we could

deduce that a little amount of CoFeB spins began to tilt away

from perpendicular direction in p-MTJ films. This is consist-

ent with the reported spin configurations in hard/soft p-ES

system at remanence state.29 Thus, we could determine that

the transition of exchange coupling from rigid coupling to

exchange spring coupling was triggered in region B. It is

noted that the still large value of Mr=Ms implies that the tilt

angle is not very large. The Hn exhibited a reproducible

behavior as first increasing then decreasing. Such singular

behavior is likely attributed to either the magnetostatic inter-

action between the top and bottom magnetic electrodes or

the discontinuity of the CoFeB/Co-terminating bilayer. In

the regions C and D (0:9 nm � tCFB � 1:9 nmÞ, Mr=Ms

FIG. 3. The plots of (a) squareness value of Mr=Ms and (b) the nucleation field Hn as function of averaged thickness of CoFeB (tCFB from 0 nm to 1.9 nm) and

annealing temperature TA at as-grown, 200 �C, 300 �C, 350 �C, and 375 �C, respectively. And accompany with some representative M–H curves, the inset

shows the location of Hn. (Note: the magnetization of p-MTJ stack with tCFB¼ 1.9 nm after 375 �C annealing is not saturated under perpendicular magnetic

field of 3 kOe.)

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quickly decreased to its minimum value of about 0.15 with

further increasing tCFB. There are two aspects accounting for

it: one is that the films broken into multi-domain structures

as shown in corresponding MFM images, and the other is

within one domain the tilt angle of CoFeB spins towards in-

plane direction further increased. Meanwhile, the irreversible

magnetization switching in the top electrode would be fur-

ther promoted by the tilted CoFeB spins, as a consequence,

Hn monotonously decreased even change the sign to the min-

imum of �250 Oe.

For the case of TA ¼ 375 �C, the magnetization reversal

mechanism seems different from that of other annealing tem-

peratures, which is valuable for us to have an insight into it.

In the region A, Mr=Ms increased to near 1, which indicates

the soft CoFeB insertion is rigidly coupled to the Co/Pt mul-

tilayer very well as a single phase magnet (magnetization

reverses together) via exchange coupling at the interface.29

According to some previous work,42,43 the Hn for the hard/

soft system in the rigid coupling regime is given by

Hn ¼2ðthKh þ tsKsÞthMh þ tsMs

; (1)

where the th (ts), Kh (Ks), and Mh (Ms) are the thickness,

magnetic anisotropy, and magnetization of magnetically

hard (soft) layers, respectively. Normally, the thicker the soft

layer, the smaller the value of Hn. This perfectly accounts

for the first abrupt reduction of Hn with tCFB increasing. The

high value of Mr=Ms and positive value of Hn both recon-

firmed the single domain structure in the film. In the region

B, both Mr=Ms and Hn abruptly declined to the minimum

value at tCFB ¼ 0:9 nm, where the sign of Hn even changed

to be negative. It is indicated the reverse domains had

appeared at least in the top free electrode before reaching

zero magnetic field, which is in good agreement with the cor-

responding multi-domain MFM images in Fig. 2. These

results suggest that the perpendicular magnetic anisotropy

(K?) of the total film became weak after 375 �C annealing.

However, we have known the K? of p-MTJ film without

CoFeB became better. Hence, the drastic drop of Mr=Ms and

Hn should be attributed to the influence of CoFeB insertions.

More specifically, it is probably related to a small amount of

B diffusion from CoFeB insertions into the adjacent Co/Pt

multilayer after 375 �C annealing, which would lead to a

graded interface between CoFeB and Co/Pt multilayer with

graded K?. Indeed, after a high temperature annealing, the

partial crystallization of CoFeB into CoFe would be trig-

gered by the B diffusing into other layers beside

CoFeB.7,44,45 Surprisingly, in the regions C and D, Mr=Ms

and Hn stopped decreasing but exhibited significant increase.

Especially, in the region D, Mr=Ms recovered to be a high

value of about 0.7 and Hn suddenly increased to be a positive

value close to that in region A. It is indicated that the irre-

versible magnetization switching had not taken place in top

magnetic electrode at zero magnetic field. Actually, the more

diffusion of B atoms into adjacent Co/Pt multilayers

occurred, meanwhile, the partial crystallization of CoFeB

into CoFe also enhanced the demagnetization energy and

thus Kk. Under this condition, in order to minimize the total

Gibbs free energy, the CoFeB (with CoFe) spins first tilted

coupled together in-plane direction but not prefer to switch

down to the opposite direction and form the multi-domain

structures. Hence, the irreversible magnetization switching

was hindered via the exchange spring coupling at the inter-

face, resulting in the increase of Hn. This is in good agree-

ment with MFM images as shown in Fig. 2, which hardly

have distinct perpendicularly reverse domains. Note that the

very blurry magnetic contrast referred above just reveals the

cone state of spin structures with large in-plane component

reported by Fr€omter,46 which arises from the canted magnet-

ization both in the CoFeB (with CoFe) layer and exchange

coupled adjacent hard phase multilayer in the p-MTJ films.

B. Patterned p-MTJs

In order to get better understanding of the mechanisms

that govern the magnetization reversal and in particular the

interfacial spin configurations (i.e., the spin configurations

near the barrier interface), we performed magnetotransport

measurement on patterned p-MTJs. Indeed, it is well known

that this technique is highly sensitive to the interfacial spin

structures. Shown in Figs. 4(a)–4(d) are some representative

R-H curves for four p-MTJs with tCFB ¼ 0 nm, 0.4 nm,

0.95 nm, and 1.7 nm, and for each p-MTJ by the sequence of

as-grown, 300 �C annealing and 375 �C annealing. As pre-

sented in Fig. 4(a), R-H curves for the p-MTJ with

tCFB ¼ 0 nm showed a good perpendicular anisotropy and dis-

tinct parallel (antiparallel) state of magnetization illustrated

by the arrows. The TMR ratio increased from 8% to 12% by

50% after 300 �C annealing due to the improved amorphous

tunnel barrier and interface.47 Subsequently, the TMR ratio

abruptly dropped to around 5% after 375 �C annealing, which

would be attributed to the degradation of the tunnel barrier as

well as the interface.48 Shown in Fig. 4(b) are R-H curves for

the p-MTJ with tCFB ¼ 0:4 nm. For the as-grown p-MTJ, the

platform width shrinked due to the difference of Hn between

bottom and top magnetic electrodes decreased with CoFeB

insertions. Interestingly, upward curvature of the magnetore-

sistance in the ascending- (descending-) field branch in the

negative (positive) field range was distinctly observed, which

is attributed to the unpinned interfacial CoFeB spins in the

top electrode. As shown in Fig. 4(c), for the R-H curves of as-

grown p-MTJ with tCFB ¼ 0:95 nm, the upward curvature sig-

nificantly increased, which reveals a larger tilt angle of the

unpinned interfacial CoFeB spins away from perpendicular

direction due to the thicker CoFeB insertions. However, the

upward curvature decreased after 300 �C annealing, which is

due to the improvement of perpendicular anisotropy. Here, by

using the high sensitive magnetoresistance measurement tech-

nique, the exchange coupling transition from rigid to

exchange spring coupling in the top electrode was clearly

detected. As shown in Fig. 4(d), for the as-grown or 300 �Cannealing p-MTJ with tCFB ¼ 1:7 nm, there are no significant

changes in the R-H curves compared with that of as-grown or

300 �C annealing p-MTJ with tCFB ¼ 0:95 nm, only that the

contribution of CoFeB spins to the magnetoresistance became

more dominating. Nevertheless, after 375 �C annealing, a sig-

nificant change in the R-H curve was exhibited: the TMR ratio

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drastically increased from initial 15% to 32% by an enhance-

ment of 113%. It is worth noting that we also acquired a simi-

lar evolution of R-H curves for p-MTJs with tCFB ¼ 1:52 nm

and tCFB ¼ 1:85 nm. Surprisingly, the TMR ratio of AlOx

based MTJ should have decreased after 375 �C high tempera-

ture annealing,48,49 just as illustrated in Fig. 4(a) for the

p-MTJ without CoFeB insertions. On the contrary, it here

doubled. Therefore, the contribution to the distinct enhance-

ment of TMR ratio is manifested not from the AlOx barrier

but from the interfacial spin configurations of CoFeB inser-

tions in both magnetic electrodes.

According to the conductance theory of MTJ by

Slonczewski (Ref. 50), the Julliere model and the similar na-

ture of the top and bottom magnetic electrodes, the TMR ra-

tio of AlOx based p-MTJ can be written as

TMR ¼ P2 � P2 cos h1þ P2 cos h

; (2)

where P and angle h represent the spin polarization of the

magnetic electrodes and the angle of magnetization configu-

rations between top and bottom electrodes, respectively. To

date, the highest P of CoFeB electrodes experimental dem-

onstrated in the AlOx barrier based micro-sized MTJ at room

temperature is about 51%.17 By adopting this value, for our

p-MTJ with TMR ratio of 32%, an angle of 100�

between

the magnetic moments of both sides of the tunnel barrier

could be estimated. Actually, after 375 �C high temperature

annealing, the spin polarization P of magnetic electrodes

would be lower than 51% due to the partial crystallization of

CoFeB into CoFe and the combination with Co/Pt multilayer

with low spin polarization as well as the degradation of the

tunnel barrier and barrier interface. Consequently, after tak-

ing these factors into account, one can expect that the angle

h should be much larger than 100�

in our p-MTJ with TMR

ratio of 32%. Considering all the p-MTJs with tCFB � 1:5 nm

showed similar R-H curves, similar domain pattern and sig-

nificantly enhanced TMR ratio, one can expect the magnetic

moments of both sides of the tunnel barrier tend to be anti-

parallel at the point showing the higher TMR ratio.

In the following, we will present the magnetization re-

versal mechanisms for the higher TMR ratio we obtained.

Taking the p-MTJ with tCFB ¼ 1:5 nm, for example, six rep-

resentative points in the ascending-field branch of R-H curve

were denoted by A, B, C, D, E, and F, respectively, as shown

in Fig. 5. The spin configurations were illustrated in the cor-

responding schematic diagrams (bottom panel), where the

different colors represent different spin directions as shown

FIG. 5. (Top panel) Magnetoresistance

curve (R-H curve) of representative p-

MTJ with tCFB ¼ 1:52 nm, A, B, C, D,

E, and F in the ascending-field branch

are six representative points; the inset

shows the amplified rectangle shaped

junction with aspect ratio of 3. (Bottom

panel) the illustrations of corresponding

spin configurations, where the different

colors represent different spin directions

defined by the pseudo-colour table and

the small arrows represent the spins of

different magnetic layers.

FIG. 4. Magnetoresistance (R-H) curves

of four representative p-MTJs with (a)

tCFB ¼ 0 nm, (b) tCFB ¼ 0:4 nm, (c)

tCFB ¼ 0:95 nm, and (d) tCFB ¼ 1:7 nm,

respectively, for each p-MTJ the R-Hcurves at as-grown state, TA ¼ 300

�C

and TA ¼ 375�C are selected, respec-

tively, and the small arrows represent

the spins of the top and bottom magnetic

electrodes.

133906-6 Wang et al. J. Appl. Phys. 113, 133906 (2013)

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in the pseudo-colour table (right-top) and the small arrows

represent the spins of different magnetic layers. At point A

(the negative saturation field), all the spins in both of mag-

netic electrodes were driven along the direction of magnetic

field resulting in the parallel state and no TMR ratio. With

the magnetic field reaching to zero, magnetoresistance

increased to point B. Similar to the previous analysis, during

the decrease of magnetic field, the CoFeB (with CoFe) spins

could not be stabilized out of plane. More specifically, due to

the rectangle shaped junction with aspect ratio of 3 (i.e.,

12 lm by 4 lm) as shown in the inset, the large in-plane

shape anisotropy tends to retain the tilted spins along the

long axis (6x axis) of the junction. For the given remanent

magnetization and the configuration of the first tilted spins

adjacent to the barrier in the top electrode (along -x axis) as

illustrated in diagram B, the direction of the subsequently

tilted CoFeB (with CoFe) spins in the bottom electrode

might have two possible cases: either along x axis or �x

axis. In order to know the exact interfacial spin configura-

tions, the total Gibbs free energy for these two cases was

compared. According to the micromagnetism model, the

total Gibbs free energy could be written as

Etotal ¼ Eext þ Ea þ Eex þ Em; (3)

where Eext, Ea, Eex, and Em represent the Zeeman energy, the

magnetic anisotropy energy, the exchange energy, and the

magnetostatic energy, respectively. Considering the rela-

tively small magnetocrystalline anisotropy of CoFeB (with

CoFe) insertions and very weak exchange coupling with

non-magnetic amorphous AlOx barrier, for simplification,

we only take the two dominating terms of Eext and Em into

consideration. The term of Eext is given by

Eext ¼ �~Hext � ~M top � ~Hext � ~Mbottom; (4)

where ~Hext is the effective perpendicular magnetic field con-

sisting of the external magnetic field and the equivalent mag-

netic field from the perpendicular magnetized Co/Pt

multilayer. ~M top and ~Mbottom represent the magnetization of

the top and bottom electrodes, respectively. Each has three

components of Mtop-x (Mbottom-x), Mtop-y (Mbottom-y), and

Mtop-z (Mbottom-z) along x, y, and z axes, respectively. In the

perpendicular magnetic field, the in-plane magnetization

components have no contribution to the Zeeman energy and

the out-of-plane magnetization components have the same

contribution. Hence, the magnetostatic energy is dominating

and given by

Em ¼ �1

4p

ðððd3~r

ðððd3~r0 mð~rÞ ~Nð~r; ~r0 Þmð~r0 Þ: (5)

Here, due to the large shape anisotropy of our junction

and the very thin magnetic electrodes, we assume Mtop-y

¼ 0, Mbottom-y ¼ 0, and the insertion layers are approxima-

tively uniformly magnetized. It has been proved that

N13 ¼ N31 ¼ 0, N11 > 0, and N33 < 0 in the demagnetization

matrix for our junction. So the simplified demagnetization

energy is estimated to be

Em ¼ ðMtop-x Mtop-z ÞN11 N13

N31 N33

� �Mbottom-x

Mbottom-z

� �

¼ Mtop-xN11Mbottom-x þMtop-zN33Mbottom-z: (6)

When the Mtop-x and Mbottom-x are antiparallel and Mtop-z

and Mbottom-z are parallel, the system has low energy. So it

suggests that the dipole interaction arising from the first

tilted spins in the top electrode (along �x) would drive the

spins in bottom electrode to tilt along þx axis as illustrated

in diagram B. Note that according to the exchange spring

model and the competing energy in play, the tilt angle of

spins is expected to gradually increase as we come closer to

the barrier interface, so the analysis and spin configuration

illustrations above are just for simplification, but enough to

illustrate the physical nature. The magnetoresistance (TMR

ratio) kept increasing to point C indicating the further

increase of tilt angle after the magnetic field reversal.

Interestingly, just passing 450 Oe, the magnetoresistance

exhibited an abrupt increase to point D. Under the reverse

perpendicular magnetic field, most of the spins of

Coð0:6Þ=½Coð0:4Þ=Ptð2Þ�5 (unit in nm) layer in the top elec-

trode switched down followed by spins of adjacent CoFeB

(with CoFe) quickly rotating down through the exchange

spring coupling at the coupling interface. Thus, the interfa-

cial spin configurations between the top and bottom electro-

des would tend to be antiparallel as estimated for Fig. 4(d)

(the angle h is much larger than 100�), which would result in

the highest TMR ratio here. Then the magnetoresistance

(TMR) showed a decrease followed by a small abrupt drop

to point E, likewise, which suggests most of the spins of

½Coð0:4Þ=Ptð2Þ�10=Coð0:4Þ (unit in nm) layer in the bottom

electrode switched down followed by the spins of adjacent

CoFeB (with CoFe) quickly rotating down. Finally, the mag-

netic field reached to positive saturation field, all the spins in

both of the magnetic electrodes were driven along the direc-

tion of positive magnetic field resulting in no TMR ratio

again with the spin configurations illustrated in diagram F.

By now, the magnetization reversal mechanisms for p-MTJs

revealing high TMR ratio have been illustrated. If the per-

pendicular magnetic anisotropy of hard phase materials as

well as the exchange coupling strength between CoFeB and

hard phase materials are both improved,27,32 CoFeB thick-

ness could be increased accordingly. The presented mecha-

nisms indicate that a much higher TMR ratio is expected for

MgO based p-MTJs. Moreover, the thermal stability of p-

MTJs is expected to increase as well.

IV. CONCLUSIONS

In this paper, the evolutions of magnetization reversal

and magnetoresistance as functions of tCFB (from 0 nm to

2 nm) and TA (from as-grown to 375 �C) were investigated in

p-MTJ films and patterned p-MTJs with ½Co=Pt�n=CoFeB

composite magnetic electrodes. The transition from rigid cou-

pling to exchange spring coupling were distinguished. After

375 �C annealing and when CoFeB insertions were thick

enough, in particular at tCFB ¼ 1:7 nm, the initial distinct per-

pendicularly oriented domain patterns disappeared in the

p-MTJ film, meanwhile, the corresponding p-MTJ exhibited

133906-7 Wang et al. J. Appl. Phys. 113, 133906 (2013)

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an enhancement of TMR ratio from initial 15% to 32%. The

increased TMR ratio resulted from the antiparallel alignment

of tilted CoFeB spins in the perpendicular exchange spring

typed magnetic electrodes. Meanwhile, based on the R-Hcurves and exchange spring model, the corresponding magnet-

ization reversal mechanisms and spin configurations were

illustrated. This study has proposed that the thicker CoFeB

layers and a higher TMR ratio are promising to be achieved

by using the perpendicular exchange spring magnetic electro-

des in perpendicular magnetic tunnel junctions.

ACKNOWLEDGMENTS

The project was supported by the Department of Defense

(Army Research Office) under Award Nos. W911NF-09-2-

0039 and W911NF-10-2-0099, and by the National Science

Foundation MRSEC under Grant No. DMR-0820521, the State

Key Project of Fundamental Research of MOST of China [No.

2010CB934400] and National Natural Science Foundation of

China [NSFC, Grant Nos. 10934099, 51021061, 10904167,

and 11174341], and International collaborative research pro-

grams between NSFC and EPSRC of United Kingdom with

Foundation No. 10911130234 and between NSFC and ANR of

France with Foundation No. F040803.

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