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Cite This: ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX
Direct Measurement of the Intrinsic Sharpness of Magnetic Interfaces Formed by Chemical Disorder Using a He+ Beam Grace L. Causer,*,†,‡ David L. Cortie,† Hanliang Zhu,‡ Mihail Ionescu,‡ Gary J. Mankey,§ Xiaolin L. Wang,† and Frank Klose‡,∥ †
Institute for Superconducting and Electronic Materials, University of Wollongong, Wollongong, New South Wales 2500, Australia Australian Nuclear Science and Technology Organisation, Lucas Heights, New South Wales 2234, Australia § Department of Physics and Astronomy, University of Alabama, Tuscaloosa, Alabama 35487, United States ∥ Guangdong Technion-Israel Institute of Technology, Shantou 515063, P. R. China ‡
ABSTRACT: Using ion beams to locally modify material properties and subsequently drive magnetic phase transitions is rapidly gaining momentum as the technique of choice for the fabrication of magnetic nanoelements. This is because the method provides the capability to engineer in three dimensions on the nanometer length scale. This will be an important consideration for several emerging magnetic technologies (e.g., spintronic devices and racetrack and random-access memories) where device functionality will hinge on the spatial definition of the incorporated magnetic nanoelements. In this work, the fundamental sharpness of a magnetic interface formed by nanomachining FePt3 films using He+ irradiation is investigated. Through careful selection of the irradiating ion energy and fluence, room-temperature ferromagnetism is locally induced into a fractional volume of a paramagnetic (PM) FePt3 film by modifying the chemical order parameter. A combination of transmission electron microscopy, magnetometry, and polarized neutron reflectometry measurements demonstrates that the interface over which the PM-to-ferromagnetic modulation occurs in this model system is confined to a few atomic monolayers only, while the structural boundary transition is less well-defined. Using complementary density functional theory, the mechanism for the ion-beaminduced magnetic transition is elucidated and shown to be caused by an intermixing of Fe and Pt atoms in antisite defects above a threshold density. KEYWORDS: magnetic nanostructures, chemical disorder, ion irradiation, polarized neutron reflectometry, density functional theory
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irradiating A1 FePt films with He+ ions to obtain L10 FePt.10 Paramagnetic (PM) precursors, such as Fe60Al40 and FePt3, which display an attractive interplay between chemical order and magnetism have also been explored, because they constitute ideal systems for investigating the phenomenon of chemical disorder-driven ferromagnetism.3,11−13 In these model architectures, impinging ions act as mixing agents to knock magnetic and nonmagnetic atoms from their chemically ordered PM sites. The vacancies created during the irradiation process are filled randomly by the diffusion of atoms to form a nonequilibrium solid solution accommodating a modified electronic structure. Ferromagnetism is then locally generated in the disordered regions because of a change in the magnetic exchange interactions, which are tied to the Fe−Fe interatomic distances and bond angles. For the specific case of facecentered-cubic L12 FePt3 (with a lattice constant (a) of a = 3.87 Å), this process transforms the chemically ordered structure, which has a Néel temperature (TN) of 160 K and Fe atoms
INTRODUCTION Artificial magnetic nanostructures synthesized by ion beams exhibit novel properties, which are envisaged to become centrally important to the future functionality of ultrahigh density magnetic recording devices and to the scalability of magnetic logic gates in the spintronics paradigm.1−7 Aiding a move away from the traditional nanofabrication technique (i.e., lithography), focused ion irradiations have demonstrated an aptitude for crafting superior-quality three-dimensional nanoelements through a single-step, dopant-free fabrication process which preserves film topography. This advent has fostered a revolution in nanostructure design. It is now possible to tune electronic properties and artificially drive phase transitions though the introduction of chemical disorder in a wide range of materials with nanoscale selectivity. Cybart et al. used He+ ions to pattern thin insulating barriers into YBa2Cu3O7−δ films to form high-transition-temperature Josephson superconducting tunnel junctions.8 Ramirez et al. suppressed long-range electronic correlations to tune the metal-to-insulator transition in V2O3 via disorder-induced O+ irradiation.9 Bernas et al. reported that the converse is also true in certain systems chemical disorder can be driven to chemical orderafter © XXXX American Chemical Society
Received: February 23, 2018 Accepted: April 18, 2018
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DOI: 10.1021/acsami.8b03197 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX
Research Article
ACS Applied Materials & Interfaces occupying magnetic moments of 3.3 μB, into a ferromagnetic (FM) chemically disordered structure below a Curie temperature of 425 K.14−16 In this way, by actively tailoring the chemical properties of nonmagnetic frameworks, ion beams provide an avenue for direct stylus-type “writing” of nanocompact FM dot arrays, characteristic of ultrahigh-density magnetic recording devices. Ultimately, the maximum achievable FM dot density of these devices will be limited by the intrinsic nature of the ion beam, which spreads out (or straggles) in the host material. This character, which is fundamental to all low-energy ions, is thought to have the potential to blur the intended definition of the resulting magnetic bit, which could hinder applications. Moreover, as there is generally a dissociation between ion-induced magnetic and structural property changes,17 this, in turn, raises further questions about the overall magnetic uniformity of the irradiated volume. Recently, a semiempirical model was developed to account for the inhomogeneous depth penetration and lateral scattering of ions in matter.3 The model revealed an acute sensitivity of the ion-induced FM moment on the number of displacements undergone by an atom during the irradiation, that is, the displacements per atom (DPA). As an outcome, it was modeled that nearly-discrete magnetic elements with non-depth-varying saturated magnetization profiles could be obtained by ion irradiation. A typical FM/ PM interface sharpness on the order of 10 nm was predicted and found to be in line with experimental observations.3,12 However, as the collected data were affected by a marginal experimental depth resolution and also by the mechanical precision of the resist layer used, the ultimately achievable sharpness of interface transitions between ion-beam-designed FM and PM regions remains yet to be experimentally demonstrated. In this work, we present direct evidence of a superior, fundamental magnetic interface quality (i.e., 1 T) region is likely caused by antiferromagnetic (AFM) spins partially rotating into the field direction or due to exchange coupling between FM and AFM domains, where the latter acts as pinning centers causing the moment to increase reversibly with the field. This picture is supported by the high-temperature (300 K) magnetometry data recorded well above the AFM TN of FePt3. Here, the slope of the high-field region reduces to zero, and both samples can be fully saturated in external fields of 1 T. Typical of normal FM behavior, both the coercivity and magnetization of each sample are also shown to increase upon cooling. Polarized Neutron Reflectometry. PNR is used to investigate the temperature and depth dependence of the induced ferromagnetism within the ion-irradiated sample, in comparison to the as-grown sample. PNR probes small
Figure 2. (a) Cross-sectional TEM image of the as-grown sample, showing the 280 nm FePt3 film thickness. Dark-field TEM images measured with a (001) diffracted beam of the (b) as-grown sample and (c) within the irradiated volume of the ion-irradiated sample. The density of dark contrast areas (encircled in red), corresponding to regions of chemical disorder, is increased in the ion-irradiated sample. (d) High-resolution TEM image measured within the irradiated volume of the ion-irradiated sample in the [110] zone axis. Inset: corresponding SAED pattern displaying single-crystalline diffraction peaks.
the FePt3 film is found to be 280 nm (Figure 2a). Dark-field TEM images, formed with a (001) diffracted beam, measured for the as-grown sample and within the irradiated microstructure of the ion-irradiated sample are shown in Figure 2b,c, respectively. In these two images, bright and dark regions correspond to areas occupying significant degrees of chemical order and disorder, respectively. The ion-irradiated sample has a greater areal density of dark contrast domains (compared to the as-grown sample), indicating a higher percentage of chemically disordered regions. Further, the as-grown sample displays both relatively large chemically ordered and chemically disordered domains extending up to ∼500 Å along the crosssectional surface. This is in contrast to the ion-irradiated volume, where smaller chemically ordered domains (∼150 Å) are apparent. The high-resolution TEM image in Figure 2d illustrates the atom arrangement in the [110] zone axis of the ion-irradiated sample. The lattice parameter of the FePt3 crystal C
DOI: 10.1021/acsami.8b03197 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX
Research Article
ACS Applied Materials & Interfaces
Figure 4. PNR data (open symbols) and best fits (solid curves) to the R+ and R− channels of the (a) as-grown and (c) ion-irradiated samples, each measured at 300 and 5 K. For clarity, the 5 K PNR data sets and best fits are shifted on the ordinate. Corresponding nuclear and magnetic SLD profiles of the (b) as-grown and (d) ion-irradiated samples, obtained from the best fits.
to the critical edge positioned at QZ = 0.007 Å−1, the FePt3 layer was modeled with an SLD of 6.28 × 10−6 Å−2 (note: the SLD of bulk FePt3 is 6.59 × 10−6 Å−2). The SLD for the layer is reduced by ∼5% relative to that of bulk FePt3, which is a common feature in non-perfectly lattice matched epitaxial films and usually indicates a slightly reduced mass density or minor deviations in stoichiometry. Further, as there are no known structural phase transitions to occur in the materials over the temperature ranges considered here, the nuclear SLD profile was held as a constant during the subsequent fitting procedure performed at lower temperatures. At 5 K, a slight difference in the spin-dependent reflectivity is seen for the as-grown sample in the high-QZ region of Figure 4a. The asymmetry most likely originates from a small net moment provided by a minority of chemically disordered domains, in the otherwise nominally chemically ordered FePt3 film, at low temperatures. The distribution of magnetic splitting, which is minimal in the region surrounding the critical edge and maximum at QZ = 0.06 Å−1, is suitably described by a model which allocates a moment of 0.24 μB per Fe atom homogenously throughout the FePt3 layer, with the exception of a magnetically modified region close to the Pt buffer layer interface. In the latter region, the FePt3 possesses an increased magnetic moment of 0.76 μB per Fe atom. The resulting magnetic SLD profile of the as-grown sample at 5 K used to obtain the best fit to the experimental data is shown in Figure 4b. Apart from changes to the magnetic structure, no chemical modulation is assigned to the FePt3 film after ion irradiation. This is verified by the horizontal profile of the nuclear SLD (black profile) extending from the Pt buffer to the Pt capping layer in Figure 4d, where the densities and scattering lengths of the FePt3 layer remain homogenous throughout both the
momentum transfers at low angles of incidence, making the technique sensitive to magnetic periodicities on the nanometer length scale. As such, the technique is capable of depth profiling layer-averaged FM moments but incapable of investigating AFM order on the atomic scale. Figure 4 displays the 5 and 300 K PNR data collected for the as-grown (Figure 4a) and ionirradiated (Figure 4c) samples. Open symbols show the specular reflected neutron spin-up (R+) and spin-down (R−) data which were collected with the incident neutron polarization oriented parallel (+) and antiparallel (−) with respect to a 1 T external magnetic field. The R+ and R− reflectivites encode information regarding the nuclear and magnetic scattering length density (SLD) profiles of each film. The simulated SLD profiles obtained from the best fits to the PNR data are shown in Figure 4b,d, respectively, for the as-grown and ion-irradiated samples. In Figure 4a, no significant asymmetry is observed between the two spin-dependent reflectivity channels (R+ and R−) for the as-grown sample at 300 K. Accounting for the magnetization detection limit of the PNR instrument (see the Methods section), this indicates that the sample is very weakly FM, in agreement with magnetometry. The best-fit structural model determined from fitting the data, shown by the black nuclear SLD profile in Figure 4b, consists of a MgO substrate (gray) with a 25 Å Cr (yellow) seed layer and a 54 Å Pt (red) buffer layer, upon which a 2777 Å layer of FePt3 (blue) sits below a 95 Å Pt (red) capping layer. Each layer is assigned a nuclear rootmean-square roughness between 1 and 2 monolayers, aside from the FePt3 layer which has a surface roughness of 20 Å. The thickness of the FePt3 layer obtained from fitting the PNR data agrees well with the cross-sectional TEM result of Figure 2a. Bulk scattering lengths and atomic number densities were used to model the MgO, Cr and Pt layers. To achieve a fitting D
DOI: 10.1021/acsami.8b03197 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX
Research Article
ACS Applied Materials & Interfaces
Figure 5. Four potential long-range magnetically ordered configurations of the chemically ordered FePt3 supercell. The relative internal energies (Δ) are given for (a) (1/2 1/2 1/2) g-type AFM, (b) FM, (c) (1/2 0 0) a-type AFM, and (d) (1/2 1/2 0) c-type AFM structures in order of descending energy. Top: the simplified structure shows only the Fe spins corresponding to the different propagation vectors used in the calculations. Bottom: a detailed representation of the corresponding spin-density isosurfaces indicates patches of small quadrupolar and induced moments surrounding the Pt sites, together with a much larger component localized on the Fe sites. The isosurface is drawn at a value of 5% of the maximum difference density.
The suppression of the ion-induced moment at the sample surface is attributed to ion channeling along the crystallographic axes of the film. In addition, the thickness of the irradiated layer obtained from PNR analysis is slightly higher than the maximum ion-penetration depth determined from TRIM simulations. This suggests, first, that the FePt3 film may have a lower density than the bulk FePt3 density (ρ = 18.37 g cm−3) used to simulate the damage profile (see Figure 1gray DPA profile). This picture is supported by the ∼5% reduction from the bulk SLD (i.e., 6.59 × 10−6 Å−2) used to model the FePt3 layer during the PNR analysis. Performing TRIM calculations based on the reduced number density of the FePt3 layer extracted from the PNR fits, a new DPA profile is calculated (see Figure 1yellow DPA profile), which displays an increased ion-penetration depth compared to the DPA profile modeled using a bulk density FePt3 target. Second, it is noted that TRIM calculations neglect ion-beam channeling effects in crystalline materials, which have, in our case, significantly increased the projected ion range compared to the modeled amorphous structure. As a result of the significant influence of ion channeling in our single-crystalline FePt3 film, we are prevented from experimentally providing a direct numerical relationship between DPA and the ion-induced FM moment. Upon cooling the ion-irradiated sample to 5 K, a greater spin asymmetry and contrast in the Kiessig fringes develops, as shown in Figure 4c. During the fitting procedure, the magnetic SLD throughout the nonirradiated region of the FePt3 layer was fixed to the model determined for the as-grown sample and only the magnetic moment throughout the irradiated region was allowed to vary. This can be seen by comparing the magnetic SLD profiles in Figure 4b,d, in which the nonirradiated portion (>1900 Å from the sample surface) of the SLD profile of the FePt3 layer is identical to the as-grown sample at the corresponding temperature. The best fit to the data was obtained by the model shown in Figure 4d, in which a bulk FePt3 FM moment of 3.3 μB per Fe atom is carried by the majority of the layer. This moment decreases slightly to a value of 3 μB per Fe atom in proximity to the Pt capping layer, which is again suggestive of ion-beam channeling. Collectively, the results of our PNR investigations confirm that magnetically saturated FM elements which exhibit distinct interfaces can indeed be formed by low-energy ion irradiation and successfully probed with high-resolution experimental techniques.
irradiated and nonirradiated regions. As such and for consistency, the chemical model established for the as-grown sample was likewise adopted to fit the PNR data for the ionirradiated sample. As shown in Figure 4c and in contrast to the as-grown sample (Figure 4a), a clear spin splitting between the R+ and R− channels is found at the critical-edge location for the ion-irradiated film at 300 K, suggestive of a highly FM sample. In addition, the appearance of Kiessig fringeswhich result from the constructive and destructive interference of neutrons reflected from the AFM/FM interfaceprovides a clear indication that magnetic contrast now exists within the chemically homogenous FePt3 layer. The period of the Kiessig fringes is determined by the thickness of the FM layer relative to the total film thickness, whereas the QZ dependence of the amplitude of the Kiessig fringes is determined by the quality (i.e., roughness) of the AFM/FM interface. From analyzing the data, it is revealed that in order to observe Kiessig fringe oscillations in Figure 4c at QZ = 0.045 Å−1, the roughness of the magnetic interface between the ion-irradiated and nonirradiated regions of the FePt3 layer must be no greater than several atomic monolayers (i.e.,
