masters thesis physics/optics
strength
the hysteresis loop is narrower when the field is applied perpendicular to the nanowire axis. In contrast to these reports, our results of Fe/ MgO/Fe nanowires grown dense demonstrate hysteresis loop corre- sponding to a coherent type reversal magnetization for all orientation of the applied field and abrupt steps are absent.
As shown in Fig. 4 the in-plane value of Hc of nanowires of Fe/ MgO/Fe is 54% higher than thin film of Fe/MgO/Fe and the in-plane Ms of nanowires is 173% higher than the thin film, although both sets of samples were synthesized at 100 °C. The difference between in-plane and out-of-plane values of Ms for nanowires is 25.5% whereas for films it is 76.5%. As the orientation of magnetic field with respect to the sample surface changes the variation in films is three times greater than in nanowires. In both nanowires and films most of the magnetiza- tion variation occurs by domain wall motion which is confirmed by small change in the values of Hc as the angle between the sample surface and applied field change.
TMM measurements for nanowires and planar film of Fe/MgO/Fe synthesized at 100 °C are shown in Fig. 5(a) and (b) respectively. In both cases a pronounced two-fold symmetry in the rotational hysteresis loop appeared for an applied field of 500 Oe and higher. Torque on the film is 20 times higher than on nanowires for the same applied field. In Fig. 5(c), we plot the ratio of Mr/Ms for the planar samples synthesized at different temperatures, as measured by longitudinal magneto- optic Kerr effect (MOKE). In order to investigate the asymmetry of the data, further a set of low- field Kerr magneto-optic loops were recorded. A combined longitudinal and transversal Kerr and domain observations of 20 nm thick Fe on MgO(001) show different switching depending on the relative orientation of the magnetic field with respect to the Fe crystallographic axes [21]. This is demonstrated by coercive field measurements using VSM at several angles with respect to Fe crystal- lographic axes as shown in Fig. 5(d). Strong dependence on angle was observed much more in films than in nanowires indicating domination of uniaxial anisotropy over magneto crystalline anisotropy.
6. Conclusions
In summary this comparative study of the structural and magnetic properties of Fe/MgO/Fe nanowires grown inside substrate supported carbon nanotubes for the first time and planar thin films MTJs showed several similarities, such as two-fold magnetic symmetry and ratio of orbital moment to spin moment. Deposition conditions in particular substrate temperature influences both structural and magnetic proper- ties of the Fe/MgO/Fe planar films. Nanowires of Fe/MgO/Fe prepared at 100 °C substrate temperature showed improved magnetic property compared to planar thin films, and we attribute this to the shape anisotropy of nanowires. The saturation magnetization is higher by 173% for nanowires compared to planar thin films prepared at the same substrate temperature. The coercive field of nanowires of Fe/ MgO/Fe is 54% higher than thin film of Fe/MgO/Fe both synthesized at 100 °C of substrate temperature. Nanowires of Fe/MgO/Fe showed higher saturation magnetization by a factor of 2.7 compared to planar thin films of Fe/MgO/Fe at 1.5 kOe. These enhanced magnetic proper- ties may result from shape anisotropy of the nanowires as well as hybridization that occurs between the π-electronic states of carbon and 3d-bands of the Fe-surface. The small change in the coercive field as the angle between the sample surface and the applied field is changed in both nanowires and films confirmed that most of the magnetization variation occur by domain wall motion. Strong dependence on angle was observed on the values of coercive field much more in films than in nanowires indicating domination of uniaxial anisotropy over magneto
crystalline anisotropy. The well-defined XMCD spectra of both nano- wires and films synthesized at substrate temperature of 100 °C indicate the Fe film is magnetically active. Previous predictions of magnetically dead layers were based on the study of Fe films that were synthesized at room temperature.
Acknowledgments
The corresponding author D.S acknowledges the support of NSF- MRI-1337339, NSF-MRI-R2-0958950, ARL-W911NF-12-2-0041, and BNL-NSLS-PASS-27168 for XAS /XMCD beam time.
References
[1] S. Wolf, D. Awschalom, R. Buhrman, J. Daughton, S. Von Molnar, M. Roukes, A.Y. Chtchelkanova, D. Treger, Spintronics: a spin-based electronics vision for the future, Science 294 (2001) 1488–1495.
[2] W.H. Butler, X.G. Zhang, T.C. Schulthess, J.M. MacLaren, Spin-dependent tunneling conductance of Fe–MgO–Fe sandwiches, Phys. Rev. B 63 (2001).
[3] D. Waldron, V. Timoshevskii, Y. Hu, K. Xia, H. Guo, First principles modeling of tunnel magnetoresistance of Fe/MgO/Fe trilayers, Phys. Rev. Lett. 97 (2006) 226802.
[4] J. Mathon, A. Umerski, Theory of tunneling magnetoresistance in a disorderedFe? MgO?Fe(001)junction, Phys. Rev. B 74 (2006).
[5] Y. Fan, K.J. Smith, G. Lupke, A.T. Hanbicki, R. Goswami, C.H. Li, H.B. Zhao, B.T. Jonker, Exchange bias of the interface spin system at the Fe/MgO interface, Nat. Nanotechnol. 8 (2013) 438–444.
[6] S. Yuasa, T. Nagahama, A. Fukushima, Y. Suzuki, K. Ando, Giant room-temperature magnetoresistance in single-crystal Fe/MgO/Fe magnetic tunnel junctions, Nat. Mater. 3 (2004) 868–871.
[7] S.S. Parkin, C. Kaiser, A. Panchula, P.M. Rice, B. Hughes, M. Samant, S.-H. Yang, Giant tunnelling magnetoresistance at room temperature with MgO (100) tunnel barriers, Nat. Mater. 3 (2004) 862–867.
[8] A. Newman, S. Khatiwada, S. Neupane, D. Seifu, Nanowires of Fe/multi-walled carbon nanotubes and nanometric thin films of Fe/MgO, J. Appl. Phys. 117 (2015) 144302.
[9] F. Lpez-Uras, E. Munoz-Sandoval, M. Reyes-Reyes, A. Romero, M. Terrones, J. Morán-López, Creation of helical vortices during magnetization of aligned carbon nanotubes filled with Fe: theory and experiment, Phys. Rev. Lett. 94 (2005) 216102.
[10] B. Wei, R. Vajtai, Y. Jung, J. Ward, R. Zhang, G. Ramanath, P. Ajayan, Assembly of highly organized carbon nanotube architectures by chemical vapor deposition, Chem. Mater. 15 (2003) 1598–1606.
[11] C. Li, A. Freeman, Giant monolayer magnetization of Fe on MgO: a nearly ideal two-dimensional magnetic system, Phys. Rev. B 43 (1991) 780.
[12] P. Luches, S. Benedetti, M. Liberati, F. Boscherini, I.I. Pronin, S. Valeri, Absence of oxide formation at the Fe/MgO(001) interface, Surf. Sci. 583 (2005) 191–198.
[13] S. Yang, H.K. Park, J.S. Kim, J.Y. Kim, B.G. Park, Magnetism of ultrathin Fe films on MgO(001), J. Appl. Phys. 110 (2011) 093920.
[14] C.M. Boubeta, C. Clavero, J.M. Garca-Martn, G. Armelles, A. Cebollada, L. Balcells, J.L. Menéndez, F. Peir, A. Cornet, M.F. Toney, Coverage effects on the magnetism of Fe? Mg O (001) ultrathin films, Phys. Rev. B 71 (2005) 014407.
[15] C. Chen, Y. Idzerda, H.-J. Lin, N. Smith, G. Meigs, E. Chaban, G. Ho, E. Pellegrin, F. Sette, Experimental confirmation of the X-ray magnetic circular dichroism sum rules for iron and cobalt, Phys. Rev. Lett. 75 (1995) 152.
[16] R. Wu, A. Freeman, Limitation of the magnetic-circular-dichroism spin sum rule for transition metals and importance of the magnetic dipole term, Phys. Rev. Lett. 73 (1994).
[17] M.K. Niranjan, C.-G. Duan, S.S. Jaswal, E.Y. Tsymbal, Electric field effect on magnetization at the Fe/MgO(001) interface, Appl. Phys. Lett. 96 (2010) 222504.
[18] P.K.J. Wong, T.L.A. Tran, P. Brinks, W.G. van der Wiel, M. Huijben, M.P. de Jong, Highly ordered C60 films on epitaxial Fe/MgO(001) surfaces for organic spin- tronics, Org. Electron. 14 (2013) 451–456.
[19] Y. Peng, H.-L. Zhang, S.-L. Pan, H.-L. Li, Magnetic properties and magnetization reversal of ?-Fe nanowires deposited in alumina film, J. Appl. Phys. 87 (2000) 7405–7408.
[20] K. Nielsch, R. Hertel, R. Wehrspohn, J. Barthel, J. Kirschner, U. Gösele, S. Fischer, H. Kronmüller, Switching behavior of single nanowires inside dense nickel nanowire arrays, IEEE Trans. Magn. 38 (2002) 2571–2573.
[21] J.L. Costa-Krämer, J.L. Menéndez, A. Cebollada, F. Briones, D. Garc?a, A. Hernando, Magnetization reversal asymmetry in Fe/MgO (001) thin films, J. Magn. Magn. Mater. 210 (2000) 341–348.
D. Aryee, D. Seifu Journal of Magnetism and Magnetic Materials 429 (2017) 161–165
165
Magnetically Engineered Spintronic Sensors and Memory STUART PARKIN, SENIOR MEMBER, IEEE, XIN JIANG, CHRISTIAN KAISER, ALEX PANCHULA, KEVIN ROCHE, AND MAHESH SAMANT
Invited Paper
The discovery of enhanced magnetoresistance and oscillatory interlayer exchange coupling in transition metal multilayers just over a decade ago has enabled the development of new classes of magnetically engineered magnetic thin-film materials suitable for advanced magnetic sensors and magnetic random access memo- ries. Magnetic sensors based on spin-valve giant magnetoresistive (GMR) sandwiches with artificial antiferromagnetic reference layers have resulted in enormous increases in the storage capacity of magnetic hard disk drives. The unique properties of magnetic tunnel junction (MTJ) devices has led to the development of an advanced high performance nonvolatile magnet random access memory with density approaching that of dynamic random-access memory (RAM) and read-write speeds comparable to static RAM. Both GMR and MTJ devices are examples of spintronic materials in which the flow of spin-polarized electrons is manipulated by controlling, via magnetic fields, the orientation of magnetic moments in inhomogeneous magnetic thin film systems. More complex devices, including three-terminal hot electron magnetic tunnel transistors, suggest that there are many other applications of spintronic materials.
Keywords—Field sensor, giant magnetoresistance (GMR), magnetic engineering, magnetic random-access memory (MRAM), magnetic recording, magnetic tunneling junction (MTJ), magnetic tunneling, magnetoelectronics, magnetoresistance, oscillatory interlayer coupling, read head, spin-dependent transport, spin valve, spintronics.
I. INTRODUCTION
Conventional magnetic materials have long been used for various near-ubiquitous applications including electric mo- tors and magnetic compassesand sensors [1]. In recentyears a new class of magnetic materials has emerged based on the microscopic generation and manipulation of spin-polarized electrons in magnetic multilayered thin-film structures [2].
Manuscript received February 15, 2003; revised March 10, 2003. This work was supported in part by the Defense Advanced Research Projects Agency. The authors are with the IBM Almaden Research Center, San Jose, CA
95120-6099 USA (e-mail: [email protected]). Digital Object Identifier 10.1109/JPROC.2003.811807
In particular, these materials can act as extremely sensitive magnetic fieldsensors,because theirelectrical resistancecan change in the presence of magnetic fields at room temper- ature by factors much larger than are possible with conven- tionalmagneticmaterials. In thisarticlewewill focuson two classes of novel materials, metallic magnetic multilayered structures and magnetic tunnel junctions, and their applica- tions to magnetic information storage in the form of mag- netic hard disk drives and magnetic random access memory [3].Wewill alsobrieflydiscussmorecomplexdevicesbased on these materials, in particular the magnetic tunnel tran- sistor (MTT). In order to make technologically useful de- vices, these materials have to be magnetically engineered so as to control and tune their response to magnetic fields. The discoveryofoscillatory interlayercoupling in1989 in transi- tion metal based magnetic multilayers [4] plus the phenom- enon of exchange biasing discovered much earlier in 1959 [5] together form the basis of the magnetic engineering of many of today’s most useful magnetic nanostructures.
II. SPIN-VALVE MAGNETIC RECORDING READ HEADS
In a magnetic recording hard disk drive information is stored by magnetizing regions within a magnetic thin film (see Fig. 1). The transitions between these regions repre- sent “bits” which are detected, via their fringing magnetic fields, by the readsensor.The readsensor ispart of amerged read–write recording head which has a separate writing el- ement. The recording head is attached to a small ceramic “slider”which is flownonanairbearingabove the recording medium at a height of just a few nanometers. The number of magnetic bits per unit area, the areal density, has increased at compound growth rates (CGRs) exceeding 100% for the past several years (see Fig. 2). These astounding increases in storage capacity have been driven in large part by a new generation of magnetic recording read heads based on the phenomenonofgiantmagnetoresistance(GMR),whichwere
0018-9219/03$17.00 © 2003 IEEE
PROCEEDINGS OF THE IEEE, VOL. 91, NO. 5, MAY 2003 661
Fig. 1. Fundamentals of magnetic recording. Schematic diagram of a hard disk drive. (a) Information is stored by magnetizing regions of a thin magnetic film on the surface of a disk. Bits are detected by sensing the magnetic fringing fields of the transitions between adjacent regions as the disk is rotated beneath a magnetic sensor. As the area of the magnetized region has decreased, the read element has had to scale down in size accordingly. (b) The read element is incorporated into a merged read/write head, which is mounted on the rear edge of a ceramic slider flown above the surface of the rapidly spinning disk via a cantilevered suspension. A hard drive unit usually consists of a stack of several such head–disk assemblies plus all the motors and control electronics required for operation (for more details see, for example, [7]).
firstdevelopedandintroducedbyIBMinlate1997[2].GMR headsarenowfoundinvirtuallyallharddiskdrivesproduced today. From the earliest days of magnetic recording more than
50 years ago [6] there have been a succession of predictions of the ultimate achievable recording density but so far these have always been surpassed. As the areal density has increased and the bit size correspondingly decreased many
technological challenges have been met. These challenges [7] range from mechanical issues such as reduced flying height and contamination sensitivity, servoing and head stiction; read issues such as magnetic shielding and sensor sensitivity; write issues concerning field strength and speed; media issues including noise and data stability as bit size approaches the super-paramagnetic limit. One of the most important technologicalhurdleshasbeenandcontinues tobe
662 PROCEEDINGS OF THE IEEE, VOL. 91, NO. 5, MAY 2003
Fig. 2. Increases in areal density and shipped capacity of magnetic storage over time. Since the invention of magnetic disk recording in the 1950s, the areal density (bits stored per square inch) of disk drives has increased rapidly. Fueled by the development of the anisotropic magnetoresistance (AMR) and GMR read sensors, in recent years its compound growth rate (CGR) has exceeded 100%. Developments in the magnetic media have kept pace with those of the read sensor and the write head. For example, antiferromagnetically coupled (AFC) recording media allow the writing of narrower tracks on media which would otherwise be susceptible to the super-paramagnetic effect. The inset plot shows the total capacity of hard drives shipped per year; in 2002 that shipped capacity was 10 EB worth of data. [Data provided by Ed Grochowski (private communication)].
the need to extract ever more read signal from diminishing volumes of magnetic material. The read signal from a GMR recording head is one to two
orders of magnitude bigger than that from prior generation state-of-the-art read heads [8] which were based on the phe- nomenon of anisotropic magnetoresistance (AMR) [9]. The resistanceofconventional ferromagneticmetalsvariesasap- proximately where is the angle between the direc- tion of the magnetic moment of the ferromagnet and the di- rection of current flow. The fringing fields from transitions in the magnetic media are then detected by changes in resis- tance of the read sensor element as the media is rotated at great speed under the sensor. These fields are quite small in the rangeof 10–100Oe.Theactivecomponentof anAMR sensor is essentially an ultrathin magnetoresistive ferromag- netic layer (typically permalloy, ) whose plane is arranged to be orthogonal to that of the magnetic media. The sensinglayer issandwichedbetweenandelectrically isolated fromtwomuch thickerultrasoftmagnetic layerswhichactas shields toensure that thesense layermeasures fluxfromonly one transition. The separation of these shields plus the thick- ness of the sense layer determines the spatial resolution of the read head. AMR, usually quite small—just a few percent at room
temperature, is a consequence of bulk scattering, and moreover, in thin films typically decreases with decreasing film thickness as scattering from the surfaces of the film
becomes more important [10]. This modest magnitude of AMR led to predictions that areal densities would be limited to Gb/in . In contrast to AMR, GMR is a much larger effect of up to more than 100% at room temperature [11] and is dominated by interface scattering [12], [13]. The GMR effect was originally discovered in MBE (molecular beam epitaxy) grown epitaxial (100) oriented Fe/Cr/Fe sandwiches [14] and Fe/Cr multilayers [15] but the effects were quite modest at room temperature. Shortly afterwards it was discovered that similar effects could be found in polycrystalline sputtered Fe/Cr multilayers [4] and subse- quently very large room temperature magnetoresistance was found in Co/Cu and related multilayers [11], [16]. These latter materials form the basis of GMR sensors and storage devices today. AMR and GMR are compared with other classes of structures and materials which display significant magnetoresistance in Fig. 3. GMR is a result of spin-dependent scattering in inho-
mogeneous magnetic metallic systems (for recent reviews see, for example, [17] and [18]). As originally discussed by Mott in the 1930s [19], [20] current in 3d transition fer- romagnetic metals is carried independently by spin-up and spin-down electrons. According to this two-current model [21], [22] thescattering ratescanbequitedifferent for electronswithin these twochannels. In themostnaïvemodel the current is considered to be carried predominantly by the low-mass electrons but the heavier electrons, which
PARKIN et al.: MAGNETICALLY ENGINEERED SPINTRONIC SENSORS AND MEMORY 663
Fig. 3. Types of magnetoresistance. (a) AMR results from bulk spin-polarized scattering within a ferromagnetic metal; it is manifested as a dependence of the resistance on the angle between an applied external field and the direction of current flowing through the material. (b) Colossal magnetoresistance (CMR) results from interactions predominantly between adjacent atoms in certain crystalline perovskites. (c) GMR results from interfacial spin-polarized scattering between ferromagnets separated by conducting spacers in a heterogeneous magnetic material, such as a magnetic multilayer or granular alloy. (d) Tunneling magnetoresistance (TMR) in magnetic tunnel junctions results from spin filtering as spin-polarized electrons tunnel across an insulating barrier from one ferromagnet to another. (e) Anomalous MR from domain wall effects has been observed in single-crystal ferromagnetic Fe whiskers and patterned magnetic wires. Spin-polarized scattering of electrons as they cross from one domain to another leads to increased resistance in a GMR-like model, although both increased and decreased resistance has been observed in lithographically patterned magnetic elements in the presence of domain walls. (f) Ballistic MR (BMR) is another type of domain wall effect in the limit of very narrow constrictions where the conductance may be quantized. Very large BMR effects have been observed but their origin, whether from spin-polarized transport, or from magnetostrictive effects is controversial.
give rise to the ferromagnetism, provide a spin-dependent reservoirofemptystates intowhich thesespelectronscanbe scattered. Since the density of states (DOS) of and elec- trons at the Fermi energy are quite distinct this rationalizes
the spin-dependent scattering rates. In an inhomogeneous magnetic system where the magnetic moment direction varies spatially the scattering rates for and electrons will also vary in space. In particular, as shown schematically in
664 PROCEEDINGS OF THE IEEE, VOL. 91, NO. 5, MAY 2003
Fig. 3(c), magnetic multilayers will typically have a lower resistance when the magnetic moments of the individual layers are parallel (P) than when antiparallel (AP) [23]. Defining the magnetoresistance as the largest GMR effects at room temperature of 110% have been found in sputter deposited Co/Cu multilayers [11]. For GMR the MR varies as the cosine of the angle between the moments of adjacent ferromagnetic layers. The largest GMR effects in magnetic multilayers are ob-
served for structures containing the thinnest possible mag- netic and nonmagnetic layers. This is because the GMR ef- fect is dominated by spin-dependent scattering at the mag- netic/nonmagnetic interfaces [12], [13] and, especially for current flow in the plane of the multilayer, thickening these layers largely results in shunting of current away from these interfacial regions [24]. In the early days of GMR the pre- ponderance of groups argued that GMR was largely a con- sequence of spin-dependent bulk scattering within the inte- rior of the ferromagnetic layers [25] and the dominant role of the interfacial contributionwasnot appreciated.Ofcourse in structures with very thick ( 100) ferromagnetic layers, when the GMR effects are quite small, contributions from bulk scattering may compare with those from interface scat- tering. However, for technological applications the interfa- cial origin of GMR makes the effect much more useful be- cause only very thin ferromagnetic layers are needed to sup- port large GMR effects. In particular, demagnetizing fields associated with the magnetic moments of the ferromagnetic layers within these devices will increase with the magnitude of these magnetic moments so it is usually very important to minimize the volume of magnetic material. For devices of microscopicdimensions theselfdemagnetizingfields,which depend in detail on the shape of the device, will increase the magnetic field required to change the magnetic state of such devices. For sensors this will eventually limit the smallest fields which such devices can detect. For memory applica- tions demagnetizing fields may lead to interactions between closely packed neighboring devices which will eventually limit the density of such memories. The understanding and control of demagnetizing fields in magnetic nanostructures is key to the fabrication of useful devices.
III. MAGNETIC ENGINEERING
Magnetic multilayers such as those shown schematically in Fig. 3(c) are not by themselves useful for most sensor or memory applications because the fields required to generate large changes in the resistance of such multilayers are typi- cally large and, moreover, are very sensitive to the thickness of the nonmagnetic spacer layer. In addition, especially for sensing applications, the maximum sensitivity to small mag- netic fields isobtained forparticularorientationsof the mag- neticmoments within the sensor. ForAMRreadsensors [see Fig. 4(a)] the moment of the magnetic layer must be set at an angle of from the direction of the current flowing through the layer.When themoment isparallelorperpendic- ular to thecurrentdirection, for smallchanges in themoment
direction, therewill be little change in resistance.Various in- genious schemesweredeveloped tomaximize the sensitivity of AMR sensors by controlling the current flow through the device and the quiescent moment direction [26]. Usually de- magnetizing fields from a second soft ferromagnetic layer separated from the sense layer by a nonmagnetic conducting layer (e.g., Ta) would be used to help bias the magnetic mo- ment of the sense layer at the correct angle. The moment of the bias layer was itself set by the self-field of the sense current passing through the device. The moment direction of the sense layer was then determined by a combination of the bias field, the magnetic anisotropy and shape demagnetizing fieldsof the sense layerand theself-fieldof thesensecurrent [see Fig. 4(b)]. The magnitude of the magnetic fields required to saturate
the change in resistance of magnetic multilayers with thin layers are very high and are determined by the intrinsic ex- change coupling mediated through the nonmagnetic spacer layers [27]. The coupling strength was discovered [4] to oscillate between ferromagnetic (F) and antiferromagnetic (AF) coupling with increasing spacer layer thickness for almost all nonmagnetic transition metal spacer layers [28]. Companion oscillations in the magnitude of the GMR effect were also found [4]. These latter oscillations simply reflect the fact that for ferromagnetically coupled magnetic layers in symmetric multilayer structures no relative change in magnetic orientation of adjacent magnetic layers will result with the application of magnetic field. By taking advantage of the oscillatory interlayer coupling it was shown that multilayer structures could be “spin engineered” [29] or constructed in such a way that ferromagnetically coupled neighboring magnetic layers could be arranged to become antiparallel in intermediatemagnetic fields.This allowsboth the measurement of the ferromagnetic coupling strength between these layers and the observation of a GMR effect from these layers [29]. For antiferromagnetically coupled magnetic multilayers
the moments of adjacent magnetic layers naturally lie an- tiparallel inzeromagnetic field. Inastrongenoughmagnetic field these moments will eventually align with the field be- coming parallel to one another and so resulting in a change in the resistance of the multilayer. The saturation field is di- rectly related to the strength of the antiferromagnetic inter- layer exchange coupling energy. The magnitude of the oscil- latory interlayer coupling energy was found to vary system- atically for magnetic multilayers with spacer layers formed from the 3d, 4d, and 5d transition metals in the periodic table, increasing in strength with d band filling and from the 5d to 4d to 3d elements [28]. For technological applications one of the most useful elements is Ru which not only dis- plays relatively high interlayer oscillatory exchange energy but also, for certain ferromagnetic materials, displays anti- ferromagnetic coupling in the limit of ultrathin Ru layers as thin as 2 to 3 [4], [29]. For GMR sensors, where additional metallic layerswill shunt andsodilute [24] themagnitudeof the effect, using a very thin AF coupling layer is important. Moreover,Ruhasotheruseful properties includingexcellent thermal stability and growth habits.
PARKIN et al.: MAGNETICALLY ENGINEERED SPINTRONIC SENSORS AND MEMORY 665
Nanowires of Fe/multi-walled carbon nanotubes and nanometric thin films of Fe/MgO Alexander Newman, Suman Khatiwada, Suman Neupane, and Dereje Seifu Citation: Journal of Applied Physics 117, 144302 (2015); doi: 10.1063/1.4917051 View online: http://dx.doi.org/10.1063/1.4917051 View Table of Contents: http://scitation.aip.org/content/aip/journal/jap/117/14?ver=pdfcov Published by the AIP Publishing Articles you may be interested in Effect of hydrogen plasma irradiation of catalyst films on growth of carbon nanotubes filled with iron nanowires J. Vac. Sci. Technol. A 32, 02B102 (2014); 10.1116/1.4827822 Optical and magnetic properties of porous anodic alumina/Ni nanocomposite films J. Appl. Phys. 113, 244305 (2013); 10.1063/1.4812466 A study of the effect of iron island morphology and interface oxidation on the magnetic hysteresis of Fe-MgO (001) thin film composites J. Appl. Phys. 112, 013905 (2012); 10.1063/1.4730630 Electrical and magnetic properties of nanosized Mg0.2Mn0.5Ni0.3AlyFe2-yO4 ferrites AIP Conf. Proc. 1447, 1099 (2012); 10.1063/1.4710390 Spatial control of magnetic anisotropy for current induced domain wall injection in perpendicularly magnetized CoFeB|MgO nanostructures Appl. Phys. Lett. 100, 192411 (2012); 10.1063/1.4711016
[This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 158.103.0.1 On: Tue, 14 Apr 2015 17:32:33
Nanowires of Fe/multi-walled carbon nanotubes and nanometric thin films of Fe/MgO
Alexander Newman, Suman Khatiwada, Suman Neupane, and Dereje Seifua)
Department of Physics, Morgan State University, Baltimore, Maryland 21251, USA
(Received 30 December 2014; accepted 27 March 2015; published online 9 April 2015; corrected 10 April 2015)
We observed that nanowires of Fe grown in the lumens of multi-walled carbon nanotubes required four times higher magnetic field strength to reach saturation compared to planar nanometric thin films of Fe on MgO(100). Nanowires of Fe and nanometric thin films of Fe both exhibited two fold magnetic symmetries. Structural and magnetic properties of 1-dimensional nanowires and 2-dimensional nanometric films were studied by several magnetometery techniques. The h-2h x-ray diffraction measurements showed that a (200) peak of Fe appeared on thin film samples deposited at higher substrate temperatures. In these samples prepared at higher temperatures, lower coercive field and highly pronounced two-fold magnetic symmetry were observed. Our results show that maximum magnetocrystalline anisotropy occurred for sample deposited at 100 !C and it decreased at higher deposition temperatures. VC 2015 AIP Publishing LLC. [http://dx.doi.org/10.1063/1.4917051]
I. INTRODUCTION
Nanowires of Fe were grown using magnetron DC sput- tering in the lumens of carbon nanotubes vertically grown on SiO2 substrate. Nanometric films of Fe on MgO(100) were synthesized at several deposition temperatures using magne- tron DC sputtering. All samples were capped with Cu at room temperature without breaking vacuum to deter oxida- tion. Samples were studied using x-ray diffraction (XRD), scanning electron microscope (SEM), scanning transmission electron microscope (STEM), vibrating sample magnetome- ter (VSM), torque magnetometer (TMM), and magneto-optic Kerr effect (MOKE). The system, Fe/MgO(100), despite its simplicity is family of systems that reveal unique properties with attractive effects for technological applications such as giant and tunneling magneto-resistance (TMR)1–5 as well as antiferromagnetic,6 oscillatory,7 and biquadratic exchange8
couplings. Nanometric ultrathin films exhibit an out-of-plane uniaxial surface anisotropy sufficient to overcome demagnet- izing field.1 This feature is important for higher density mag- netic media. This system in addition to its enormous potential for technological applications it is an attractive research object in nanomagnetism.9 The substrate MgO is an ideal substrate to grow quasi free standing metal structures, on which effects of reduced dimensionality of metal can be studied because of the very weak interaction between Fe and MgO predicted theoretically by full-potential linearized augmented-plane-wave-total-energy method10 and demon- strated experimentally by x-ray photoelectron spectroscopy (XPS) and x-ray absorption spectroscopy (XAS) showing the interface between Fe(001) and MgO(100) to be stable for temperatures up to 670 K.11
The importance of Fe/MgO system is also demonstrated by the several experimental and theoretical studies on
Fe/MgO (100) nanometric films.10–25 Ab-initio calculations yielded optimistic TMR ratio in excess of 100%12 for Fe/ MgO/Fe. XAS and photoemission experiments observed changes in 3d band due to evolution of Fe local atoms coordi- nation from a bulk-line situation to a configuration where low dimensionality effects are significant.13 Salvador et al.14
showed that for uncapped nanometric thin films of Fe, the magnetocrystalline anisotropy of the films increased with dep- osition temperature indicating improved crystalline structures. The saturation value of the magnetocrystalline anisotropy (550 Oe corresponding to bulk Fe) was reached at a deposition temperature of 300 !C was further shown. Our study on nano- metric thin films of Fe/MgO(100) also showed saturation magnetization was a maximum at deposition temperature of 200 !C. Maximum coercive field occurred at 100 !C, Figure 3(a), indicative of higher magnetocrystalline anisotropy. The coercive field showed a large decrease at deposition tempera- tures of 200 !C and 300 !C. This corresponds with the appear- ance of Fe(200) peak at h¼65! in the XRD for samples deposited at these two high temperatures, Figure 1.
Density functional calculations have shown that the application of electric field has significant effect on the inter- facemagnetization and magnetocrystalline anisotropy. This is due to change in relative occupancy of 3d-orbitals of Fe atoms on Fe/MgO interface. This suggests possible applica- tion of Fe/MgO systems for electrically controlled magnetic data storage, multi-ferroic device.15
In this study, structural and magnetic properties of nano- metric Fe thin films on MgO(100) and nanowires of Fe pre- pared in the lumens of multi-walled carbon nanotubes (MWCNTs) using magnetron DC-sputtering were studied using XRD, SEM/STEM, VSM, TMM, and MOKE measure- ments. Magnetic measurements using VSM showed that samples prepared at 100 !C exhibit the highest coercive field (Hc¼176 Oe), while samples prepared at 50 !C show high remnant magnetization (Mr¼119 emu/g) and samples
a)Author to whom correspondence should be addressed. Electronic mail: [email protected].
0021-8979/2015/117(14)/144302/5/$30.00 VC 2015 AIP Publishing LLC117, 144302-1
JOURNAL OF APPLIED PHYSICS 117, 144302 (2015)
[This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 158.103.0.1 On: Tue, 14 Apr 2015 17:32:33
prepared at 200 !C show the highest saturation magnetization (Ms¼147 emu/g), which is 68% the saturation field of bulk Fe measured by a force method at room temperature.20 As shown in Figure 3(a), the three parameters extracted from a B-H loop Hc, Mr, and Ms are at their maximum values not at one particular rather at different substrate growth tempera- tures. In our study, nanowires of Fe capped with Cu in the lumens of MWCNTs vertically grown on SiO2 exhibited an enhanced magnetic property of 137 Oe coercivity. This result is comparable with our recent result of single walled carbon nanotubes (SWCNTs) coated with Fe2O3 nanoparticles
26 and exceeds that of graphene coated with Fe2O3 nanoparticles.
27
In the past, magnetic nanoparticles were filled inside and outside MWCNTs and SWCNTs28–33 using chemical meth- ods. This work illustrates an alternate method of filling sub- strate supported by MWCNTs to synthesize magnetic nanowires using magnetron sputtering.
II. EXPERIMENT
Vertically aligned MWCNTs were filled with Fe and capped with Cu in this experiment were grown using a ther- mal CVD method34 on SiO2 substrate. This method involves exposing silica structures to a mixture of ferrocene and xy- lene at 770 !C for 10 min. The furnace is pumped down to #200 mTorr in argon bleed and then heated to a temperature of 770 !C. The solution of ferrocene dissolved in xylene (#0.01 g/ml) is pre-heated in a bubbler to 175 !C and then passed through the tube furnace. The furnace is then cooled down to room temperature. The open ended MWCNTs tips were filled with Fe and capped with Cu using DC magnetron sputtering method at a substrate temperature of 100 !C. This substrate temperature yielded in planar samples the highest value of coercive field amongst several other substrate tem- peratures, as shown in Figure 3(a).
Nanometric thin films epitaxially grown on MgO (100) substrate (5 mm $ 5 mm $ 50 lm) (from MTI company) using magnetron DC sputtering (AXXIS tool from K. J. Lesker Company) at several temperatures. All substrates were degassed at 350 !C in vacuum of 10%7 Torr for 1800 s and samples were pre and post annealed at a pre-selected deposition temperature for 1800 s in vacuum. All samples
were capped with 5 nm of Cu to deter oxidation. The source substrate distance was kept fixed at 30 cm and the substrate was kept at 45!, while being rotated at a constant rate of 20 rpm for uniform deposition. With these condition, epitax- ial Fe grows on MgO (100) due to a good lattice match of MgO (a¼4.213 A!) and Fe (a¼2.866 A!), and weak inter- face interaction10,11 free standing Fe is formed. The deposi- tion rate for Fe was 0.17 nm/s as calibrated by the deposition time versus the thickness measurements for Fe films several hundred thick.
Thin film samples of Fe/MgO(100) were characterized by thin film X-ray diffraction (Rigaku D/max Ultima II 40 kV/40 mA) using CuKa radiation in h-2h geometry.
Surface morphologies of pristine MWCNTs and nano- wires of Fe grown in the lumens of carbon nanotubes were characterized by a Hitachi S-5500 field emission SEM/ STEM. For STEM imaging, some nanotubes were scrapped off SiO2 substrate and dispersed in dimethylformamide, the resulting solution was dripped on holey carbon coated car- bon TEM grid.
VSM measurements were carried out using Vector Magnetometer Model 10 VSM system from MicroSense equipped with 3T electromagnet. Magnetic torque was meas- ured using EV7 TMM system equipped with a 2T electro- magnet. For MOKE measurements, an in-house built instrument was used. The operating principles of MOKE consist of measuring changes in polarization of light reflected from magnetic sample. The MOKE setup consists of a monochromatic light source (HeNe laser, 632 nm wave- length, and 5 mW output power), polarizer, analyzer, photo- diode, electromagnet, lock-in amplifier, pre-amplifier and an optical chopper, or photoelastic modulator. The reflected beam from a magnetic sample is passed through a second po- larizer, which is the analyzer in order to select the compo- nent of the E-field perpendicular to the plane of incidence. This will make the normalized intensity detected propor- tional to the component of the magnetization of the sample parallel to the applied magnetic field.
III. RESULTS AND DISCUSSIONS
It is well known that the structural property is important in magnetism.1 For this reason, one of the goals of this study is to correlate the structure of films to their magnetic prop- erty. The structures were characterized by x-ray diffraction using CuKa radiation.
As shown in Figure 1, all samples prepared at several different temperatures peaked at the same angle except for samples prepared at higher substrate temperatures beyond
FIG. 1. XRD patterns of Fe/MgO(100) at various growth temperatures.
TABLE I. Magnetization values of Fe grown on MgO (100) by magnetron
DC sputtering at several deposition temperatures. Measurements were taken using VSM at room temperature.
T ( !C) Hc (Oe) Ms (emu/g) Mr (emu/g) S¼Mr/Ms
50 103 132 119 0.906
100 176 91.8 76.4 0.832
200 41.6 147 104 0.707
300 8.30 122 94.6 0.773
144302-2 Newman et al. J. Appl. Phys. 117, 144302 (2015)
[This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 158.103.0.1 On: Tue, 14 Apr 2015 17:32:33
100 !C. These samples exhibited a (200) sharp peak at 65!. This change in crystalline structure and a sharp decrease in coercive field were observed in samples prepared at higher substrate temperatures. This is due to a decrease in magneto- crystalline anisotropy in samples prepared at substrate tem- perature higher than 100 !C. This result is in direct contrast to the finding of Salvador et al.14 that magneto-crystalline anisotropy increases monotonously with temperature up to 350 !C.
As shown in Figure 1, XRD indicates that Fe(200) peak is present at higher deposition temperature, this can be
explained qualitatively as follows: an Fe atom reaching the substrate at an arbitrary site will not stay at that site for depo- sition temperature high enough. It will have enough thermal energy to move to an appropriate site to form a crystalline structure. XRD of the film deposited at 100 !C does not show a peak at Fe(200), it possesses a maximum coercive field, as shown in Figure 3(a). A change in crystalline structure and a sharp decrease in coercive field were observed in samples prepared at higher substrate temperatures, Table I.
In Figure 2, SEM/STEM images of pristine and nano- wires of Fe/MWCNTs on SiO2 substrates. As compared to
FIG. 2. (a) SEM image of pristine MWCNTs. (b) SEM images of nano- wires of Fe/MWCNTs (inset: high magnification). (c) Bright field STEM images of Fe/MWCNTs. (d) Dark field STEM images of Fe/MWCNTs.
FIG. 3. (a) Coercive field Vs growth temperature of Fe/MgO(100) measured with VSM. Coercive field of Fe/ MWCNTs nanowires synthesized at 100 !C indicated by “*.” (b) Magnetic hysteresis loops for Fe/MgO(100) films for various growth temperatures as measured by VSM at room tempera- ture. All loops are measured in-plane with the applied field perpendicular to the normal vector to the sample’s sur- face h¼0!. (c) Magnetic hysteresis loops for Fe/MgO(100) film prepared at 100 !C at various angles between the normal vector to the surface and the applied magnetic field. The purple loop is for in-plane (IP) measurement h¼0!. (d) Magnetic hysteresis loops for Fe/MWCNT wire prepared at 100 !C at various angles between the normal vector to the sample’s surface and the applied magnetic field.
144302-3 Newman et al. J. Appl. Phys. 117, 144302 (2015)
[This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 158.103.0.1 On: Tue, 14 Apr 2015 17:32:33
pristine MWCNTs, Fe/MWCNTs show a three-fold increase in diameter from (10.6 6 1.66) nm to (37 6 8) nm. The bright contrast at tips of MWCNTs on Figure 2(b) is due to Fe grown on MWCNTs. The high-magnification image in the inset shows that Fe forms a uniform coating creating Fe/ MWCNTs coaxial structures. The continuous bright spots on bright-field STEM image (Figure 2(c)) are obtained due to the uniform coating of Fe along the entire length of MWCNTs. The dark-field STEM image in Figure 2(d) also shows the formation of nanowires of Fe/MWCNT with ho- mogeneous deposition of Fe.
Figure 3(a) depicts coercive field (Hc), saturation and remnant magnetization (Ms and Mr) Vs temperature for pla- nar Fe/MgO films, as well as nanowire of Fe/MWCNTs. Figure 3(b) VSM measurements of M Vs H for samples syn- thesized at several substrate temperatures. In Figures 3(c) and 3(d), M Vs H loops of planar nanometric thin film and
nanowire synthesized at a substrate temperature of 100 !C at several angles between the applied field and normal to the surface. The remnant and saturation magnetizations are min- ima at growth temperature of 100 !C, as shown in Figure 3(a). The saturation magnetization, Ms, has a maximum value for films deposited at 200 !C, for which the coercive field is a minimum. The squareness S¼Mr/Ms has a mini- mum value for films grown at this deposition temperature. This variation is depicted in Figure 3(b), where hysteresis loops for the various Fe/MgO(100) film prepared at several deposition temperatures measured using VSM at room tem- perature. For the sample with the highest value of Hc, Figure 3(c) depicts, hysteresis loops for various orientation of the magnetic field with the surface normal for Fe/MgO(100) film prepared at 100 !C measured using VSM at room tempera- ture. The values of Hc, Ms, and Mr for Fe/MgO(100) planar film and wires of Fe/MWCNTs deposited at 100 !C meas- ured at various angles between the normal vector to the sur- face and applied magnetic field are listed in Table II and shown in Figures 3(c) and 3(d). Hc was measured to be at its maximum value at 45!. Mr and Ms were maximum and Hc at its minimum at 0! and at in-plane orientation. This is another indication that the easy magnetization axis lie in plane and the hard axis lie out of plane. Szalkowski35 showed that the hysteresis loop of Fe-filled MWCNTs exhibits an anomalous narrowing of the loop at the zero magnetization axes. L!opez- Ur!ıas et al.25 reported the formation of helical spin configu- rations during magnetization of ferromagnetic nanowires encapsulated inside carbon nanotubes.
TABLE II. Magnetization values of Fe/MgO(100) film and Fe/MWCNTs/ SiO2 nano-wires deposited at 100
!C by magnetron DC sputtering at several angles between the applied magnetic field and the surface normal.
Measurements were taken using VSM at room temperature.
Angle
Hc (Oe) Ms (emu/g) Mr (emu/g)
Film Nano-Wire Film Nano-Wire Film Nano-Wire
0 176 137 91.8 38 76.4 27
45 243 115 57.9 35 46.2 24
90 179 59.8 13.2 28 88.6 19
FIG. 4. (a) Magnetic torque curves of Fe/MgO(100) films for various growth temperatures at room temperature measured at an applied field of 20 kOe. (b) Magneto-optical longitudinal Kerr loops of Fe/MgO(100) films for various growth temperatures at room temperature. (c) Magnetic torque curves of Fe/ MgO(100) films at growth temperature of 100 !C measured at several applied fields at room temperature measured. (d) Magnetic torque curves of Fe/ MWCNTs wires at growth temperature of 100 !C measured at several applied fields at room temperature measured.
144302-4 Newman et al. J. Appl. Phys. 117, 144302 (2015)
[This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 158.103.0.1 On: Tue, 14 Apr 2015 17:32:33
Maximum coercive field of 176 Oe was observed at synthe- sis temperature of 100 !C for films and 137 Oe for nanowires, a maximum remnant magnetization of Mr¼120.0 emu/g was observed at 50!C, a maximum saturation magnetization of Ms¼146.4 emu/g was observed at 200!C, and the squareness defined S¼Mr/Ms was 80% for 2D films and 70% for 1D wire.
Figure 4(a) depicts torque magnetometer measurements at an applied magnetic field of 20 kOe of samples synthe- sized at several substrate temperatures. In Figure 4(b), MOKE measurements are presented in polar plots. Hysteresis loops were measured using MOKE in longitudinal configuration, from the recoded hysteresis loop the coercive fields of samples were measured as a function of angular position of the sample relative to the applied magnetic field direction, Figure 4(b).
MOKE data indicates the coercive field is a maximum for samples with growth temperature of 100 !C in agreement with VSM. The magnetic torque curve, Figure 4, indicates pronounced two-fold magnetic symmetry at higher deposi- tion temperature. The torque curve in Figure 4(a) for sample prepared at 100 !C shows less pronounced two-fold symme- try compared to torque curve of samples prepared at higher substrate temperatures. In Figure 4(c), the 100 !C sample shows comparatively higher pronounced at 20 kOe than at lower field strengths, there is a transition between 1 kOe and 5 kOe. Figures 4(c) and 4(d) show torque magnetometer measurements of planar film of Fe/MgO(100) and nanowire of Fe/MWCNTs synthesized at 100 !C at several applied fields. Both show similar trend except at low fields, the loop for nanowire is off center. Figure 4(b) depicts coercive field measured using MOKE in polar plot for thin films synthe- sized at several growth temperatures. The same trend is observed in torque magnetometer measurements for lower fields, as depicted in Figure 4(c).
IV. CONCLUSION
In conclusion, we have synthesized Cu capped Fe/ MgO(100) nanometric thin films and Cu capped nanowires of Fe using MWCNTs as templates. Nanowires were grown at an optimized condition set by growing planar films at sev- eral deposition temperatures that showed the best magneto- crystalline property. Magnetic measurements showed that nanowires exhibited higher anisotropy requiring higher satu- ration field compared to planar thin films due to magnetic shape anisotropy, however, similar magnetic symmetry, two- fold, was observed in nanometric films and nanowires. The squareness of nanowires is by 10% less than planar nanomet- ric thin films.
ACKNOWLEDGMENTS
One of the authors, Dereje Seifu, would like to acknowledge the supports of NSF-MRI-R2 Grant (Award No. 0958950), NSF-MRI (Award No. 1337339), and the U.S. Army Research Laboratory through a Cooperative Research Agreement USARMY-W911NF-12-2-0041. We also thank Professor Ajayan’s group at Rice University for providing nanotubes used in this experiment.
1Ultrathin Magnetic Structures, edited by B. Heinrich and J. A. C. Bland (Springer, Berlin, 1994).
2S. Yuasa, T. Nagahama, A. Fukushima, Y. Suzuki, and K. Ando, Nature 3, 868 (2004).
3D. D. Djayaprawira, K. Tsunekawa, M. Nagai, H. Maehara, S. Yamagata, N. Watanabe, S. Yuasa, Y. Suzuki, and K. Ando, Appl. Phys. Lett. 86, 092502 (2005).
4S. Yuasa, A. Fukushima, T. Nagahama, K. Ando, and Y. Suzuki, Jpn. J. Appl. Phys., Part 2 43(4B), L 588–L 590 (2004).
5M. Bowen, V. Cros, F. Petroff, A. Fert, C. Mart!ınez Boubeta, J. L. Costa- Kr€amer, J. V. Anguita, A. Cebollada, F. Briones, J. M. de Teresa, L. Morell!on, M. R. Ibarra, F. G€uell, F. Peir!o, and A. Cornet, Appl. Phys. Lett. 79, 1655 (2001).
6P. Grunberg, R. Schreiber, Y. Pang, M. B. Brodsky, and H. Sowers, Phys. Rev. Lett. 57(19), 2442 (1986).
7S. S. P. Parkin, N. More, and K. P. Roche, Phys. Rev. Lett. 64(19), 2304 (1990).
8M. Ruhrig, R. Schafer, A. Hubekt, R. Mosler, J. A. Wolf, S. Demokrlrov, and P. Grunrer, Phys. Status Solidi A 125, 635 (1991).
9S. D. Bader, Rev. Mod. Phys. 78, 1 (2006). 10C. Li and A. J. Freeman, Phys. Rev. B 43, 780 (1991). 11P. Luches, S. Benedetti, M. Liberati, F. Boscherini, I. I. Pornin, and S.
Valeri, Surf. Sci. 583, 191 (2005). 12J. Mathon and A. Umerski, Phys. Rev. B 63, 220403 (2001). 13P. Luches, P. Torelli, S. Benedetti, E. Ferramola, R. Gotter, and S. Valeri,
Surf. Sci. 601, 3902 (2007). 14C. Salvador, T. Freire, C. G. Bezerra, C. Chesman, E. A. Soares, R.
Paniago, E. Silva-Pinto, and B. R. A. Neves, J. Phys. D: Appl. Phys. 41, 205005 (2008).
15M. K. Niranjan, C. G. Duan, S. S. Jaswal, and E. Y. Tsymbal, Appl. Phys. Lett. 96, 222504 (2010).
16C. M. Boubeta, C. Clavero, J. M. Garc!ıa-Mart!ın, G. Armelles, A. Cebollada, L. Balcells, J. L. Men!endez, F. Peir!o, A. Cornet, and M. F. Toney, Phys. Rev. B 71, 014407 (2005).
17J. Balogh, I. D!ezsi, C. Fetzer, J. Korecki, A. Kozioł-Rachwał, E. Mły!nczak, and A. Nakanishi, Phys. Rev. B 87, 174415 (2013).
18S. Yang, H.-K. Park, J.-S. Kim, J.-Y. Kim, and B.-G. Park, J. Appl. Phys. 110(9), 093920 (2011).
19J. L. Costa-Kramer, J. L. Men!endez, A. Cebollada, F. Briones, D. Garcia, and A. Hernando, J. Magn. Magn. Mater. 210, 341–348 (2000).
20J. Crangle and G. M. Goodman, Proc. R. Soc. London, Ser. A 321, 477 (1971).
21H. Fuke, A. Sawabe, and T. Mizoguchi, Jpn. J. Appl. Phys. 32, 1137–1140 (1993).
22K. Postava, J. F. Bobo, M. D. Ortega, B. Raquet, H. Jares, E. Snoeck, M. Goiran, A. R. Fert, J. P. Redoules, J. Pistora, and J. C. Ousset, J. Magn. Magn. Mater. 163, 8 (1996).
23G. Sch€utz, W. Wagner, W. Wilhelm, P. Kienle, R. Zeller, R. Frahm, and G. Materlik, Phys. Rev. Lett. 58, 737 (1987).
24B. Sinković, P. D. Johnson, N. B. Brookes, A. Clarke, and N. V. Smith, Phys. Rev. Lett. 65, 1647 (1990).
25F. L!opez-Ur!ıas, E. Mu~noz-Sandoval, M. Reyes-Reyes, A. H. Romero, M. Terrones, and J. L. Mor!an-L!opez, Phys. Rev. Lett. 94, 216102 (2005).
26S. Neupane, S. Khatiwada, C. Jaye, D. Fisher, H. Younes, H. Hong, S. Karna, S. Hirsch, and D. Seifu, ECS J. Solid State Sci. Technol. 3(8), M39–M44 (2014).
27S. Neupane, H. Hong, L. Giri, S. P. Karna, and D. Seifu, “Enhanced mag- netic properties of graphene coated with Fe2O3 nanoparticles,” J. Nanosci. Nanotechnol. (to be published).
28G. Korneva, H. Ye, Y. Gogotsi, D. Halverson, G. Friedman, J. C. Bradley, and K. G. Kornev, Nano Lett. 5, 879 (2005).
29J. P. Tessonnier, O. Ersen, G. Weinberg, C. Pham-Huu, D. S. Su, and R. Schlogl, ACS Nano 3, 2081 (2009).
30X. P. Gao, Y. Zhang, X. Chen, G. L. Pan, J. Yan, F. Wu, H. T. Yuan, and D. Y. Song, Carbon 42(1), 47–52 (2004).
31F. Tan, X. Fan, G. Zhang, and F. Zhang, Mater. Lett. 61(8–9), 1805–1808 (2007).
32D. Jain and R. Wilhelm, Carbon 45, 602–606 (2007). 33D. Seifu, Y. Hijji, G. Hirsch, and S. P. Karna, J. Magn. Magn. Mater. 320,
312–315 (2008). 34B. Q. Wei, R. Vajtai, Y. Jung, J. Ward, R. Zhang, G. Ramanath, and P. M.
Ajayan, Chem. Mater. 15, 1598–1606 (2003). 35Carbon Nanotubes: Select Army Research Laboratory Studies, edited by
P. Szalkowski (NOVA, 2013), Chap. 4.
144302-5 Newman et al. J. Appl. Phys. 117, 144302 (2015)
[This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP: 158.103.0.1 On: Tue, 14 Apr 2015 17:32:33
- Exchange bias of the interface spin system at the Fe/MgO interface
- Current understanding of exchange bias
- The ferromagnet/oxide interface
- Selectively probing the interface and bulk spin systems
- Atomic-scale model of the interface spin system
- Controlling the interface exchange bias
- Conclusions
- Methods
- Figure 1 Model of TMR structure and of atomic moments near the Fe/MgO interface.
- Figure 2 Measurement geometry and MOKE/MSHG data.
- Figure 3 Classic model of exchange bias.
- Figure 4 TEM images of sample interfaces.
- Figure 5 Magnitude of exchange bias field HE versus applied field direction.
- Figure 6 Temperature dependence of the exchange bias.
- References
- Acknowledgements
- Author contributions
- Additional information
- Competing financial interests