In: Laser Ablation: Fundamentals, Methods and Applications ISBN: 978-1-63482-589-4 Editors: C. Gerhard, S. Wieneke and W. Viöl © 2015 Nova Science Publishers, Inc.
Chapter 3
LASER ABLATION FOR MULTIFERROIC HETEROSTRUCTURES Devajyoti Mukherjee*, Sarath Witanachchi and Pritish Mukherjee Center for Integrated Functional Materials & Department of Physics, University of South Florida, Tampa, Fl, US
ABSTRACT Multiferroic heterostructures of ferroelectric (FE) and ferromagnetic (FM) materials have attracted considerable attention for their multifunctional properties and potential applications in magnetoelectric memory devices. Pulsed laser deposition (PLD), although quite successful in the growth of complex-oxide thin films, is complicated in the synthesis of multiferroic heterostructures. This is primarily associated with the dissimilar ablation conditions of the insulating FE, the conducting FM phases in such heterostructures and intrinsic differences in their material properties. In this chapter, we address some challenges in the laser ablation of epitaxial multilayered thin films: the FE perovskite PbZr0.52Ti0.48O3 (PZT) with the FM oxide La0.7Sr0.3MnO3 (LSMO). We show that a careful optimization of the laser ablation parameters can lead to the growth of PZT/LSMO heterostructures with superior FE/FM properties. The major issue of the preferential evaporation of Pb during laser ablation of PZT that leads to Pb-deficient nonFE thin films has been investigated using laser-target interaction studies and optical ICCD imaging of the laser-ablated plumes. Using the optimized growth parameters in PLD, epitaxial PZT/LSMO thin films have been deposited on single-crystal SrTiO3 (100) and MgO (100) substrates. Detailed structural analyses using X-ray diffraction, atomic force microscopy and transmission electron microscopy reveal the single crystalline nature, the smooth particulate-free surface morphologies of the layers and the sharp and flat interfaces in the heterostructures. Magnetic and polarization measurements show simultaneous room temperature FE and FM behaviors in PZT/LSMO heterostructures. Further, the incorporation of a hard-magnetic CoFe2O4 (CFO) sandwich layer in PZT/LSMO heterostructures was achieved using the dual-laser ablation technique that combines a KrF excimer and a CO 2 laser outputs. The optimum coupling of the laser *
Corresponding Author: 4202 E. Fowler Avenue, ISA 2019, Tampa, Fl 33620, USA. Email:
[email protected]
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Devajyoti Mukherjee, Sarath Witanachchi and Pritish Mukherjee energies in dual laser ablation caused higher ionization of the ablated species and enhanced gas phase reaction resulting in highly crystalline film deposition. Structural characterization of PZT/CFO/LSMO heterostructures deposited using dual laser ablation showed their single-crystalline nature, defect-free interfaces and smooth surface morphologies. With the introduction of the CFO sandwich-layer, the low magnetic coercivity of PZT/LSMO as well as the remanent polarization increased; making PZT/CFO/LSMO desirable for multiferroic device applications. The results presented here provide a detailed recipe for the pulsed laser deposition of multiferroic heterostructures, which is crucial for the coherent design of future memory devices based on these composite systems.
Keywords: Magnetoelectric, epitaxial thin films, laser-target interactions, dual-laser ablation
INTRODUCTION In the past decade, multiferroic materials with coexisting and intimately coupled ferroelectric (FE) and ferromagnetic (FM) orders have attracted considerable attention for their vast application potentials and the fundamental interest in the magnetoelectric (ME) effect in these systems [1-4]. Mutual control of the ferroelectricity and magnetism promises the design of multifunctional devices ranging from ultrasensitive magnetic sensors/transducers, solid state transformers, energy harvesting and magneto-electro-optic devices to non-volatile memories [5,6]. The scarcity of naturally occurring single-phase multiferroic materials and their weak ME coupling coefficients [7,8] has directed the research towards combining the electronic and magnetic properties of diverse FE and FM materials to create “artificial” multiferroic composite systems, exhibiting ME coefficients that are orders of magnitude larger than those typical of single-phase multiferroics [9,10]. The preparation of multiferroic FE/FM heterostructures is challenging due to the intrinsic differences in the properties of FE and FM materials, often making them chemically incompatible [11]. While FE materials are mostly insulators with unfilled d-orbitals, existence of transition metal delectrons is essential for magnetic ordering in FM materials, which complicates their combination into ME composites. In this respect, horizontal heterostructures consisting of alternate layers of FE and FM phases have been amongst the most widely investigated composites due to their ease of fabrication and effective separation of the FE and FM phases [9-11]. Epitaxial multilayered thin films provide an additional degree of freedom compared to bulk samples through the ‘strain engineering’ between the substrate and the film, and achieve strain-mediated ME coupling in these systems [11-13]. Over the years, several multilayered heterostructures comprising of FE perovskites (viz. Pb(Zr,Ti)O3, BaTiO3) with magnetic spinel (viz. Fe3O4, CoFe2O4, Ni0.8Zn0.2Fe2O4) or FM manganite phases (viz. LaxSr1−xMnO3) have been investigated and evaluated for their potential ME applications [14-18]. In particular, ME heterostructures of the state-of-the-art FE ceramic PbZr0.52Ti0.48O3 (PZT) when combined with the FM half-metallic oxide La0.7Sr0.3MnO3 (LSMO) [19-22] have been given special attention due to the demonstration of charge-mediated ME effect in PZT/LSMO thin films for novel spin-based logic devices in the context of spintronics [23, 24]. PZT/LSMO thin films have already found applications in non-volatile data-storage [25], field-effect transistors [26], magnetic tunnel junctions (TJs) [27], and FE TJs [28]. The use of the conducting LSMO layer as the bottom electrode during the polarization of PZT/LSMO
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capacitors has shown to improve their FE and fatigue endurance properties as compared to metallic or other oxide electrodes [29,30]. Among the various thin film deposition processes such as sputtering [31], molecular beam epitaxy (MBE) [20], chemical solution deposition [32], and sol-gel processing [33, 34] that have been used to deposit PZT/LSMO thin films, pulsed laser deposition (PLD) has undoubtedly been the most widely used technique [19, 26, 35, 36]. PLD offers unique advantages in terms of in-situ growth of multilayered thin films by sequential ablation of multiple target-materials and preservation of complex stoichiometries in the deposited films [37]. Still, there are several challenges in stoichiometric growth of PZT/LSMO thin films using PLD [38]. Laser ablation of PZT thin films is complicated by the high volatility of Pb in PZT that leads to its preferential evaporation during the ablation process and consequently non-stoichiometric film deposition [39, 40]. Pb deficiency in PZT films is responsible for the coexistence of a meta-stable pyrochlore (non-ferroelectric) phase along with the perovskite PZT structure, thereby degrading the FE properties of the films. The compositions of the asdeposited PZT films depend strongly on the laser fluence (energy density) and ambient O2 pressure (pO2). Generally, a high pO2 (~200-400 mTorr) and a high laser fluence (> 3.5 J/cm2) are used during deposition, which exacerbate the particulate ejection from the target and lead to nonuniform, particulate-laden films [41]. On the other hand, laser ablation of LSMO thin films exhibiting smooth surface morphologies, large magnetization and high conductivity is quite complicated and still in progress [42-44]. The physical properties of LSMO thin films are highly sensitive to the laser ablation parameters and ambient O2 pressure (typically, 10 < pO2 < 50 mTorr) [45]. Other detrimental effects during the growth of laser-ablated LSMO thin films include formation and precipitation of a secondary Mn3O4 phase as a result of off-stoichiometric deposition and Sr segregation towards the surface of LSMO thin films [46]. Due to the chemical compatibility of Mn3O4 with LSMO, the presence of micro-precipitates of the impurity Mn3O4 phase in LSMO thin films does not affect its stoichiometry, making it difficult to detect using standard techniques [47]. However, the occurance of Mn3O4 impurites degrade the FM properties of LSMO thin films since Mn3O4 is a room temperature (RT) paramagnetic material [47]. In this chapter, we address some of these challenges in the stoichiometric growth of highquality epitaxial PZT/LSMO multiferroic heterostructures using PLD. We present a comprehensive study using laser-target interactions and optical imaging of laser-ablated plumes that led to the optimization of ablation parameters in the epitaxial growth of these heterostructures. The as-deposited PZT/LSMO thin films exhibited high crystalline quality, smooth surface morphologies and RT multiferroic behavior. As a technological advancement, the incorporation of a hard-magnetic CoFe2O4 (CFO) sandwich-layer in PZT/LSMO heterostructures was achieved using the dual-laser ablation technique, which combines a KrF excimer and a pulsed CO2 laser outputs [48]. Significant improvement of the figures of merit, such as the magnetic coercivity and remanent polarization, was observed in the PZT/CFO/LSMO thin films making them desirable for potential device fabrication. The results are essential both for the fundamental understanding of the laser ablation process for the growth of multiferroic heterostructures and for the coherent design of multifunctional devices based on these materials.
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Laser-Target Interactions during Ablation of PZT and LSMO The laser fluence as well as the optical, topological and thermal properties of the target material influence the mechanisms that lead to the ablation of materials [49]. Investigations of the surface irregularities formed on the target surface after irradiation with the KrF excimer laser pulses (typically used in PLD) reveal the thermodynamic processes that occur during ablation.
Figure 1. Time-integrated ICCD images of the total visible emission from the laser-ablated plasma plumes and the corresponding SEM images of the laser-target interaction sites after ablation of (a) LSMO and (b) PZT targets at varying laser fluences from 1 J/cm2 to 4 J/cm2.
Here, the total visible emissions from the laser-ablated plasma plumes were captured using an intensified charge-coupled detector (ICCD) camera (PI-MAX:512 UNIGEN) system aligned normal to the plume propagation direction (see Refs. 38 and 39 for details). After ablation, the laser target interaction sites on the target surfaces were examined using a
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scanning electron microscope (Jeol, JSM-6389LV). Figures 1 (a) and 1 (b) show the timeintegrated ICCD plume images under varying laser fluences of 1-4 J/cm2 and the corresponding SEM images of the laser-target interaction sites after ablation of the (a) LSMO and (b) the PZT targets under pO2 of 10 mTorr and 500 mTorr, respectively. From the plume images in Figure 1, it is observed that the lateral dimensions of the plasma near the target resemble the laser spot size (left side of the images), while the axial dimension (perpendicular to the target) increases drastically with increase in the laser fluence [38,39]. Under low pO2, the most intense zone (higher counts) of the plasma is centered near the target (Figure 1 (a)). However, under high pO2, due to the surrounding pressure from the ambient O2, the most intense plasma region shifts away from the target, as clearly seen in the high fluence (3-4 J/cm2) plume images in Figure 1 (b). Since higher counts in the ICCD plume images represent higher ionization of the ablated species, spatial variation in ICCD plume images (as observed in Figure 1 under varying laser fluences or ambient O2 pressure) can affect the crystallinity and morphology of the deposited films. From the SEM images in Figure 1(a) , it is clear that the structural features on the irradiated target surfaces change drastically with the laser fluences. At low fluence of 1 J/cm2, the LSMO target exhibits a partially-melted granular surface (Figure 1 (a)) while at 2 J/cm2 the surface is indicative of complete melt-texturing. Ablation at higher fluences of 3-4 J/cm2 creates distinct cavities on the target surface, indicative of non-congruent target erosion. LSMO films deposited using laser fluence of 1 J/cm2 were non-stoichiometric, while those deposited using laser fluence of 3-4 J/cm2 showed rough surfaces with high particulate densities, making them undesirable for device applications. Sub-micron particulates in the LSMO films drastically decrease their electrical conductivity since they act as scattering sites for electrons during their propagation. From the foregoing laser-target interaction studies, a fluence of 2 J/cm2 was considered optimum for the ablation of LSMO thin films. At 2 J/cm2 laser fluence, a target to substrate distance of 4 cm was found to be optimum for deposition (see ICCD plume image in Figure 1(a)). SEM images in Figure 1 (b) show that the surface features on the PZT target change from distinct conical structures at low fluences of 1-2 J/cm2 to irregular spherical aspirations at high fluences of 3-4 J/cm2, indicative of varied ablation conditions.
Preferential Evaporation of Pb during Ablation of PZT Targets When the PZT target is ablated under high pO2 of 500 mTorr using low and high KrF laser fluences of 1 J/cm2 and 5 J/cm2, as shown in the ICCD plume images in Figures 2 (a) and 2 (b), respectively, diverse surface morphologies are observed on the irradiated PZT target surface. The conical structures formed at low fluence as shown in Figure 2 (c) are reminiscent of severe melt texturing along with smooth, globular cone-tips and distinct conebodies underneath. Energy dispersive spectroscopy (EDS) measurements reveal that the cone tips are extremely Pb deficient as compared to their bodies. A formation of such conical structures is indicative of the non-stoichiometric evaporation of PZT, due to the preferential ablation of Pb [38,39]. PZT films deposited under these ablation conditions are non-stoichiometric and highly Pb deficient [39,50]. The structural aspirates formed at high fluence as shown in Figure 2 (d) are associated to the surface
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instabilities of the melted target surface after absorption of the high energy laser pulse [51,52]. However, such features are not attributed to a non-congruent target erosion. EDS analysis showed that the preferential Pb loss from the target was minimized at high laser fluences (5 J/cm2) where the high energy facilitated the congruent evaporation of materials from the target surface. PZT films deposited under these conditions exhibited standard FE properties [39]. SEM images of PZT films deposited at 1 J/cm2 and 5 J/cm2 as shown in Figs. 2 (e) and 2 (f), respectively, reveal a drastic increase in particulate density as the KrF laser fluence is increased. Laser ablation of PZT using high fluence of 5 J/cm2 excerbates the particulate ejection from the PZT target surface during the ablation, as shown in the ICCD image in the inset to Figure 2 (f). Particulate-laden films, as shown in Figure 2 (f), are undesirable in FE/FM layered structures, since such micron-sized particulates in the FE PZT layer can cause a shorting of the electrodes during their polarization measurements. The micron-sized spherical particulate shown in the inset to Figure 2 (e) suggests its ejection from the tip of the conical structures formed on the irradiated PZT target. From the above analysis, during PLD of PZT films under high pO2 (500 mTorr) a laser fluence of 3 J/cm2 was considered optimum, where both the non-congruent ablation of PZT and particulate ejection were minimized.
Figure 2. Time-integrated ICCD images of the total visible emission from the laser-ablated plasma plumes after ablation of PZT target under high pO2 of 500 mTorr, using low and high fluences of (a) 1 J/cm2 and (b) 5 J/cm2, respectively [38]. SEM images of (c and d) conical structures formed at the lasertarget interaction sites on the PZT target surface and (e and f) surfaces of deposited PZT films after ablation using low (1 J/cm2) and high (5 J/cm2) laser fluences, respectively [39]. The insets in Figs. 2 (d) and 2 (e) show the details of one of the cone tips and a particulate on the deposited PZT film, respectively. The inset to Figure 2 (f) show an ICCD image depicting severe particulate ejection from the PZT target surface during ablation. Reprinted with permissions from Mukherjee et al. [38] and Mukherjee et al. [39]. Copyright [2012], AIP Publishing LLC.
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Laser Ablation of PZT/LSMO Multiferroic Heterostructures Epitaxial PZT/LSMO heterostructures were grown on single-crystal SrTiO3 (STO) and MgO (100) substrates using optimized ablation parameters in PLD [38]. In a typical synthesis process, an initial layer of LSMO (thickness ~ 100 nm) was deposited at 800 ˚C under pO2 of 10 mTorr and using a KrF excimer laser (Lambda Physik, 256 nm, 10 Hz) at a fluence of 2 J/cm2. Subsequently, a PZT layer (thickness ~ 300 nm) was deposited at 550 ˚C under a pO2 of 500 mT and using a laser fluence of 3 J/cm2. A shadow mask was used during PZT deposition to preserve an open access to the LSMO bottom electrode. After the PZT layer deposition, top LSMO electrodes (≈ 100 μm pads) were deposited using a shadow mask at 550 ˚C under a pO2 of 10 mTorr and using a laser fluence of 2 J/cm2. The prototype PZT/LSMO heterostructure shown schematically in Figure 3 (a) was structurally characterized using X-ray diffraction (XRD, Bruker D8 Focus Diffractometer), atomic force microscopy (AFM, Digital Instruments DI III), and transmission electron microscopy (TEM, FEI Tecnai F 20 S-Twin). Figure 3 (b) shows the cross-sectional TEM image of the PZT/LSMO heterostructure. The image reveals sharp and flat interfaces with uniform thicknesses of the individual PZT and LSMO layers grown on the STO substrate.
Figure 3. (a) Schematic diagram of the PZT/LSMO heterostructures grown on STO and MgO (100) substrates using laser ablation. (b) Cross-sectional TEM image of PZT/LSMO/STO heterostructure. (c and d) AFM images of the surfaces of the LSMO and the PZT layers, respectively [38]. (e) XRD θ-2θ patterns for epitaxial PZT/LSMO thin films grown on STO (100) and MgO (100) substrates [38]. Inset to Figure 3 (e) shows the details of the PZT (002), STO (200) and LSMO (200) peaks. Reprinted with permission from Mukherjee et al. [38]. Copyright [2012], AIP Publishing LLC.
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The surface morphologies of the bottom LSMO layer and the top PZT layer as shown in the AFM images in Figs. 3 (c) and 3 (d) reveal smooth surfaces with no indication of microcracks or defects (pitting or pores), suggesting the uniform deposition of the layers. Clearly, the LSMO bottom layers (Figure 3 c) show smaller grain sizes (D ≈ 55 nm) and roughness values (Rrms = 2.5 nm) as compared to the PZT top layers (Figs. 3 d) (D ≈ 170 nm, Rrms = 18 nm) [38]. A possible cause for the higher roughness values in the PZT layers could be associated with the higher pO2 used during the deposition of the PZT layer. High pO2 reduces the kinetic energy of the ablated species due to increased ionic collisions, which result in lower ad-atom mobility and higher surface roughness of the deposited films [49].The singlecrystalline nature of the PZT/LSMO heterostructures grown on STO and MgO substrates is evidenced from their XRD θ-2θ patterns as shown in Figure 3 (e). In both cases, only (00l) (l = 1, 2, and 3) diffraction peaks of the tetragonal PZT phase (JCPDS 01-070-4060) along with the (l00) (l = 1, 2, and 3) peaks of the pseudo-cubic perovskite LSMO phase (JCPDS 01-0894461) are observed, indicating the epitaxial growth of the individual layers on the singlecrystal substrates [38]. The small lattice mismatch between LSMO (pseudo-cubic perovskite, a = 0.387 nm) and the STO (cubic perovskite, a = 0.3905 nm) substrate resulted in the close proximity of their XRD peaks (as shown in the inset to Figure 3 (e)) [53]. No secondary phase formations are observable within the resolution limits of XRD in Figure 3 (e). XRD azimuthal (ϕ) scans and rocking curves (not shown here) confirmed the in-plane epitaxial relationships in the heterostructures, similar to our earlier reports [39]. The magentic properties of PZT/LSMO thin films were measured at 300 K using a commercial vibrating-sample-magnetometer (VSM, Quantum Design) with magnetic fields upto 10 kOe [38]. Figure 4 (a) shows the magnetization hysteresis (M-H) loops for PZT/LSMO heterostructures grown on STO and MgO substrates. The magnetic field has been applied parallel to the film plane as shown schematically in the inset to Figure 4 (a). The observed saturation magnetization (Ms ≈ 263 emu/cm3) and coercivity (Hc ≈ 30 Oe) for the PZT/LSMO heterostructure grown on STO substrate match well with those previously reported for PZT/LSMO/STO (100) heterostructures [19, 38]. The measured magnetizations for both the PZT/LSMO/STO and PZT/LSMO/MgO heterostructures are close to their corresponding LSMO single-layered films [38,43,44] which suggests that atomic inter-diffusion did not deteriorate the magnetic properties of the LSMO bottom layers with the high temperature growth of the PZT top layers under the optimized growth conditions using PLD [38]. The low values of Hc (≈ 30 Oe) in PZT/LSMO films indicate the low density of surface defects, such as high surface roughness, grain boundaries and other point defects, that could act as domain pinning sites to increase the Hc [38]. The lower Ms value for PZT/LSMO/MgO (Ms ≈ 224 emu/cm3) as compared to PZT/LSMO/STO (Ms ≈ 263 emu/cm3) is probably associated with the larger tensile strain of the LSMO layer in PZT/LSMO/MgO as seen in the XRD strain analysis, similar to our earlier report on epitaxial LSMO thin films grown on MgO and STO substrates [44]. For the FE property measurement of PZT/LSMO heterostructures, a thin film capacitor structure was fabricated using the conducting LSMO layer as the bottom electrode and depositing LSMO top electrodes using PLD, as shown schematically in the inset to Figure 4 (b).
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Figure 4. (a) Magnetic (M-H) [38] and (b) polarization (P-E) hysteresis loops measured at 300 K for PZT/LSMO heterostructures grown on STO and MgO (100) substrates using laser ablation. Insets to Figure 4 (a) show schematically the magnetic field applied parallel to the film plane and the enlarged portion of the M-H loops at low field ranges. Inset to Figure 4 (b) shows schematic diagram of the PZT/LSMO FE capacitor structure using the conducting LSMO layer as the bottom-electrode and LSMO top-electrodes. Reprinted with permission from Mukherjee et al. [38]. Copyright [2012], AIP Publishing LLC.
The polarization was measured using a commercial Ferrotester (Radiant Tech. Inc) using an applied voltage of 9 V. Figure 4 (b) shows the polarization hysteresis (P-E) loops at 300 K for PZT/LSMO heterostructures grown on STO and MgO substrates. The well-saturated and square P-E loops in Figure 4 (b) can be attributed to the high crystalline quality and stoichiometric PZT film deposition under the optimized growth conditions reported here. The remanent polarization (Pr) and coercive field (Ec) values for PZT/LSMO/STO (Pr ≈ 52 μC/cm2 at Ec ≈ 36 kV/cm) and PZT/LSMO/MgO (Pr ≈ 43 μC/cm2 at Ec ≈ 50 kV/cm) match well with those reported earlier for epitaxial PZT films (of the same composition of PZT) and epitaxial PZT/LSMO films [19,20,54,55]. The P-E loop for PZT/LSMO/MgO shown in Figure 4 (b) is slightly asymmetric along the abscissa (+Ec > - Ec), which could be associated with a built-in field in the structure, consistent with earlier reports [19]. To summarize, epitaxial PZT/LSMO thin films were grown on STO (100) and MgO (100) substrates using PLD. The ablation conditions were optimized using laser-target interactions and optical plume diagnostics. XRD measurements and TEM/AFM images revealed a single-crystalline nature and defect-free interfaces with smooth surface morphologies for the thin films. PZT/LSMO heterostructures exhibited RT multiferroic behavior while the individual layers retained their FE and FM properties.
Dual-Laser Ablation of PZT/CFO/LSMO Multiferroic Heterostructures In terms of memory applications, the low magnetic coercivity of PZT/LSMO heterostructures as seen earlier in Figure 4 (a) is a deterrent, since external stray fields could potentially flip their magnetic moments and destroy the stored information. It is possible that
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the incorporation of a hard magnetic material such as CoFe2O4 (CFO) [56,57] in PZT/LSMO heterostructures could improve their intrinsically soft FM properties, and consequently the ME device performance [58]. However, an incorporation of CFO layers (≈ 50-100 nm thick) in PZT/LSMO layered thin films introduces huge dielectric losses which consequently degrade the FE polarization of the PZT layer in the these heterostructures [59-61]. Recently, we reported on the enhanced FE and FM properties observed in PZT/CFO/LSMO heterostructures grown on MgO (100) substrates using the dual-laser ablation technique which combines a KrF excimer and a pulsed CO2 laser outputs [62]. In the dual-laser ablation process, the KrF excimer and the pulsed CO2 laser pulses are spatially and temporally overlapped onto the target surface as shown schematically in Figure 5 (a). The interpulse delay between the two lasers is temporally synchronized (Figure 5 b) to achieve optimum coupling of the laser energies. Initially, the target is heated by the CO 2 laser (wavelength 10.6 μm) pulse to produce a shallow transient molten layer, from which a slightly time-delayed KrF excimer laser pulse initiates the ablation [48, 63].
Figure 5. (a) Schematic diagram depicting the spatial overlap of the KrF excimer and the CO2 laser beams in dual-laser ablation [63]. (b) Waveforms showing the interpulse peak to peak (p-p) delay between KrF and CO2 laser pulses in dual-laser ablation process [63]. (c) Time-integrated ICCD plume images for the total intensities of the laser-ablated plumes in single (left) and dual-laser (right) ablation processes captured under the same conditions [63]. Reprinted with permission from Mukherjee et al. [63]. Copyright [2012], AIP Publishing LLC.
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Over the years, it has been confirmed that ablation from a momentarily liquid target results in drastic reduction of particulates in the deposited films [64-69]. Moreover, the effective heating of the KrF excimer ablated plume by the tail end of the CO2 pulse (Figure 5 b) produces a higher excited plume with a broader angular distribution, compared to the plume generated by an excimer pulse alone, as shown in the ICCD images of the ablated plumes in single KrF ablation and dual-laser ablation using the same KrF laser fluence in Figure 5c. The higher excitation of the ablated species leads to an enhanced gas phase reaction and consequently enhanced surface morphology and crystallinity in the deposited films. Compared to single-laser ablation, the broader transverse plume expansion results in the deposition of uniform films over larger areas making them desirable for device fabrication. Earlier, we have reported that dual-laser ablation minimizes the preferential evaporation of Pb during ablation of PZT targets, resulting in stoichiometric PZT films with smooth surface morphologies and enhanced FE properties as compared to those using traditional PLD [39]. Using the dual-laser process, PZT/CFO/LSMO heterostructures were grown on singlecrystal MgO (100) substrates. The details of the experimental set-up have been reported in [62]. In the optimized synthesis process, an initial layer of LSMO (thickness ~ 200 nm) was deposited on MgO substrate at 800˚C under a background oxygen pressure (pO2) of 10 mTorr, followed by the CFO layer (thickness ~ 50 nm) at 450 ˚C with pO 2 of 10 mTorr, and then the PZT layer (thickness ~ 400 nm) at 550˚C under a pO2 of 500 mTorr. An open access to the LSMO bottom electrode was kept throughout the process using shadow masks. Finally, top LSMO electrodes of 100 μm diameter were deposited using a shadow mask at 550 ˚C under a pO2 of 10 mTorr, as shown schematically in Figure 6 (a). For comparison, PZT/LSMO heterostructures were also grown on MgO substrates under the same conditions. The cross-sectional TEM image in Figure 6 (b) of the PZT/CFO/LSMO heterostructure reveals the sharp and flat interfaces with uniform thicknesses of the individual PZT, CFO, and LSMO layers grown on the MgO substrate, along with the LSMO top-electrode layer (over PZT layer) and a protective Pt-layer (deposited during TEM sample preparation). Figures 6 (c, d, and e) show AFM images of the surface morphologies of the LSMO, CFO and PZT layers, respectively, in the PZT/CFO/LSMO heterostructure. They exhibit particulate-free, flat and smooth surfaces with uniform grain growth and no indication of micro-cracks or defects (pitts or pores). The grain sizes (D) and surface roughness (Rrms) increase from the LSMO layer (D ≈ 50 nm, Rrms ≈ 2 nm) to the CFO layer (D ≈ 80 nm, Rrms ≈ 10 nm) and finally to the PZT layer (D ≈ 50 nm, Rrms ≈ 2 nm) [62]. The single-crystalline natures of the PZT/CFO/LSMO and PZT/LSMO heterostructures on MgO substrates are evidenced from their XRD θ-2θ patterns as shown in Figure 6 (f). In all cases, only (00l) (l = 1, 2, and 3) diffraction peaks of the tetragonal PZT phase (JCPDS 01070-4060) along with the (400) peak of the face-centered cubic (fcc) CFO phase (JCPDS 00022-1086) and the (l00) (l = 1, 2, and 3) peaks of the pseudo-cubic perovskite LSMO phase (JCPDS 01-089-4461) are observed, indicating the epitaxial growth of the individual layers on the single-crystal MgO (100) substrates, consistent with previous reports [44, 62,70,71]. Due to the small lattice mismatch between MgO (face-centered-cubic fcc, 2 x lattice parameter = 8.42 Å) and CFO (fcc, lattice parameter = 8.391 Å), the MgO (200) and CFO (400) peaks occur in close proximity to each other in the θ-2θ spectra, as shown in details in the inset to Figure 6 (f) [70,71].
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Figure 6. (a) Schematic diagram of the PZT/CFO/LSMO heterostructure grown on MgO (100) substrate using dual-laser ablation technique. (b) Cross-sectional TEM image of PZT/CFO/LSMO/MgO heterostructure. (c, d, and e) AFM images of the surface morphologies of the LSMO, CFO and PZT layers in the PZT/CFO/LSMO heterostructure, respectively [62]. (f) XRD θ-2θ patterns for the PZT/CFO/LSMO and PZT/LSMO heterostructures on MgO substrates [62]. (g) XRD rocking curves performed about CFO (400), PZT (002) and LSMO (200) planes in the PZT/CFO/LSMO heterostructure [62]. The inset to Figure 6 (e) shows the details of the PZT (002), STO (200) and LSMO (200) peaks. Reprinted with permission from Mukherjee et al. [62]. Copyright [2014], AIP Publishing LLC.
No secondary phase formations are observable within the resolution limits of XRD in Figure 6 (f). The small full-width-at-half-maximum (FWHM) values in the XRD rocking curves shown in Figure 6 (g) performed about CFO (400), PZT (002) and LSMO (200) planes confirmed the excellent in-plane orientation of the layers in PZT/CFO/LSMO heterostructure.
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Figure 7. (a) Magentic (M-H) and (b) polarization (P-E) hysteresis loops measured at 300 K for the PZT/CFO/LSMO and the PZT/LSMO heterostructures [62]. Insets to Figure 7 (a) schematically show the magnetic field applied parallel to the film plane and the enlarged portion of the M-H loops at low field ranges. The inset to Figure 7 (b) shows a schematic diagram of the PZT/CFO/LSMO FE capacitor structure using the conducting LSMO layer as the bottom-electrode and LSMO top-electrodes. Reprinted with permission from Mukherjee et al. [62]. Copyright [2014], AIP Publishing LLC.
Similar to the PZT/LSMO thin films, PZT/CFO/LSMO heterostructure exhibited simultaneous FM and FE behavior as observed in the magnetization (M-H) and polarization (P-E) hysteresis loops shown in Figs. 7 (a) and (b), respectively. Enhanced Msat from 244 emu/cm3 in PZT/LSMO to 288 emu/cm3 in PZT/CFO/LSMO is observed in Figure 7 (a). With the introduction of the hard magnetic CFO sandwich-layer in PZT/LSMO, the magnetic coercivity (Hc) is enhanced from 0.1 kOe to 14 kOe [62]. Higher magentic figures of merit (i.e., squareness and coercivity) make the PZT/CFO/LSMO prototype device more desirable for device applications as compared to PZT/LSMO thin films. On the other hand, PZT/CFO/LSMO heterostructures show well-saturated FE hysteresis loops (Figure 7 b) with enhanced remanent polarization (Pr ≈ 69 μC/cm2 at Ec of 88 kV/cm) as compared to PZT/LSMO thin films. An enhanced polarization in PZT/CFO/LSMO is associated with its high crystalline quality and the defect-free interfaces of the individual layers under the optimized growth conditions in the dual-laser process.The enhanced FM and FE properties in the PZT/CFO/LSMO with a thin CFO sandwich layer demonstrate the effectiveness of the dual-laser ablation process in growing high-quality multiferroic heterostructures for device fabrication.
CONCLUSION Multiferroics heterostructures exhibit stronger ME coupling than intrinsic multiferroics and promise novel functionalities that can be optimized for next-generation electronic devices. Still, much remains to be investigated in terms of materials optimization,
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Devajyoti Mukherjee, Sarath Witanachchi and Pritish Mukherjee
development and characterization of advanced growth techniques, and understanding of the interfacial processes mediating the ME coupling in these systems. PLD has been the technique most widely used for the growth of multilayered thin films of the FE ceramic PZT and the FM metallic-oxide LSMO. However, the intrinsic differences in the material properties of the insulating PZT and the conducting LSMO make the task challenging. Here, we have shown that laser-target interaction studies and optical plume imaging techniques can be used to evaluate the (non-)congruent ablation of PZT and LSMO targets, which allows a proper optimization of the laser ablation parameters. Epitaxial PZT/LSMO thin films deposited under optimized conditions showed singlecrystalline nature, smooth surface morphologies and defect-free interfaces. Coexisting FE and FM properties were observed at room temperature in PZT/LSMO thin films. In order to enhance the soft FM properties of PZT/LSMO to achieve better device functionality, the incorporation of a hard-magnetic CFO sandwich layer was envisioned. Using the technologically advanced dual-laser ablation technique, PZT/CFO/LSMO heterostructures were grown on MgO substrates. While structural characterization revealed the high-quality epitaxial growth of PZT/CFO/LSMO, magnetic and FE measurements showed enhanced coercive fields and remanent polarization in PZT/CFO/LSMO as compared to PZT/LSMO. The enhanced figures of merit achieved in PZT/CFO/LSMO multiferroic heterostructures make them desirable for future data-storage applications.
ACKNOWLEDGMENTS This work was partially supported by the United States Army (Grant No. W81XWH1020101/3349) and the Florida Cluster for Advanced Smart Sensor Technologies (FCASST).
REFERENCES [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14]
N. A. Spaldin, and M. Fiebig, Science 309, 391 (2005). M. Fiebig, J. Phys. D: Appl. Phys. 38, R123 (2005). R. Ramesh and N. A. Spaldin, Nat. Mater., 6, 21 (2007). S. W. Cheong, and M. Mostovoy, Nat. Mater. 6, 13 (2007). Y. Zhang, Z. Li, C. Y. Deng, J. Ma, Y. H. Lin, and C. W. Nan, Appl. Phys. Lett. 92, 152510 (2008). M. Bibes and A. Barthelemy, Nat. Mater. 7, 425 (2008). N. A. Hill, J. Phys. Chem. B 104, 6694 (2000). D. I. Khomskii, Bull. Am. Phys. Soc. C 21, 002 (2001). Y. Wang, J. Hu, Y. Lin, and Ce-Wen Nan, NPG Asia Materials 2, 61 (2010). J. Ma, J. M. Hu, Z. Li, and C. W. Nan, Adv. Mater. 23, 1062 (2011). L.W. Martin, Y.-H. Chu, and R. Ramesh, Mater. Sci. Eng. R 68, 89 (2010). J. P. Zhou, Z. C. Qiu, and P. Liu, Mater. Res. Bull. 43, 3514 (2008). M. Vopsaroiu, J. Blackburn, and M. G. Cain, J. Phys. D: Appl. Phys. 40, 5027 (2007). S. Ryu, J. H. Park, and H. M. Jang, Appl. Phys. Lett. 91, 142910 (2007).
Laser Ablation for Multiferroic Heterostructures
59
[15] M. Ziese, A. Bollero, I. Panagiotopoulos, and N. Moutis, Appl. Phys. Lett. 88, 212502 (2006). [16] Y. G. Ma, W. N. Cheng, M. Ning, and C. K. Ong, Appl. Phys. Lett. 90, 152911 (2007). [17] Y. Zhang, C. Deng, J. Ma, Y. Lin, and C.-W. Nan, Appl. Phys. Lett. 92, 062911 (2008). [18] T. Wu, M. A. Zurbuchen, S. Saha, R. V. Wang, S. K. Streiffer, and J. F. Mitchell, Phys. Rev. B 73, 134416 (2006). [19] S. Dussan, A. Kumar, J. F. Scott, and R. S. Katiyar, Appl. Phys. Lett. 96, 072904 (2010). [20] C. A. F. Vaz, Y. Segal, J. Hoffman, F. J. Walker, and C. H. Ahn, J. Vac. Sci. Technol. B 28(4), C5A6 (2010). [21] N. Zhang, G. Srinivasan and A. M. Balbashov, J. Mater. Sci. 44, 5120 (2009). [22] M. A. Zurbuchen, T. Wu, S. Saha, and J. Mitchell, and S. K. Streiffer, Appl. Phys. Lett. 87, 232908 (2005). [23] C. A. F. Vaz, J. Hoffman, Y. Segal, J. W. Reiner, R. D. Grober, Z. Zhang, C. H. Ahn, and F. J. Walker, Phys. Rev. Lett. 104, 127202 (2010). [24] J. M. Rondinelli, M. Stengel, and N. A. Spaldin, Nat. Nanotechnol. 3[1], 46 (2008). [25] J. Hoffman, X. Hong, and C. H. Ahn, Nanotechnology 22, 254014 (2011). [26] C. Thiele, K. D¨orr, W. -M. Lin, K. -H. M¨uller, and L. Schultz, Sensor. Actuat. A 129, 180 (2006). [27] K. Pradhan, R. Mundle, R. B. Konda, O. Yasar, F. Williams, and M. Bahoura, J. Mater. Res. 24, 3065 (2009). [28] L. Jiang, W. S. Choi, H. Jeen, S. Dong, Y. Kim, M. –G. Han, Y. Zhu, S. V. Kalinin, E. Dagotto, T. Egami, and H. N. Lee, Nano Lett. 13, 5837 (2013). [29] F. Chen, Q. Z. Liu, H. F. Wang, F. H. Zhang, and W. Wu, Appl. Phys. Lett. 90, 192907 (2007). [30] W. Wu, K. H. Wong, C. L. Choy, and Y. H. Zhang, Appl. Phys. Lett. 77, 3441 (2000). [31] S. C. Ray, M. Algueró, J. Ricote, M. L. Calzada, C. Prieto, A. de Andrés, and M. Garcia-Hernández, Mater. Lett. 60, 1714 (2006). [32] F. McNallya, J. H. Kima, and F. F. Langea, J. Mater. Res. 15, 1546 (2000). [33] J. -C. Yu, J. -X. Wu, T. H. Yeh, and C. -C. Chou, Ferroelectrics 383, 127 (2009). [34] H. -C. Pan, C. -C Chou, and H. -L. Tsai, Appl. Phys. Lett. 83, 3156 (2003). [35] W. Wu, K. H. Wong, and C. L. Choy, Appl. Phys. Lett. 85, 5013 (2004). [36] F. Mitsugi, Y. Yamagata, T. Ikegami, K. Ebihara, J. Narayan, and A. M. Grishin, Jpn. J. Appl. Phys. 39, 5418 (2000). [37] R. Eason, Pulsed Laser Deposition of Thin Films: Applications-Led Growth of Functional Materials (Wiley, Hoboken, New Jersey, 2007). [38] D. Mukherjee, M. Hordagoda, N. Bingham, H. Srikanth, P. Mukherjee and S. Witanachchi, J. Appl. Phys. 112, 064101 (2012). [39] D. Mukherjee, R. Hyde, P. Mukherjee, H. Srikanth, and S. Witanachchi, J. Appl. Phys. 111, 064102 (2012). [40] N. Izyumskaya, Y. I. Alivov, S. J. Cho, H. Morkoç, H. Lee, and Y. S. Kang, Crit. Rev. Solid State Mater. Sci. 32, 111 (2007). [41] T. J. Zhu, L. Lu, and M. O. Lai, Appl. Phys. A 81, 701 (2005). [42] H. Boschker, M. Huijben, A. Vailionis, J. Verbeeck, S. van Aert, M. Luysberg, S. Bals, G. van Tendeloo, E. P. Houwman, G. Koster, D. H. A. Blank, and G. Rijnders, J. Phys. D: Appl. Phys. 44, 205001 (2011).
60
Devajyoti Mukherjee, Sarath Witanachchi and Pritish Mukherjee
[43] D. Mukherjee, N. Bingham, M. H. Phan, H. Srikanth, P. Mukherjee, and S. Witanachchi, J. Appl. Physics 111, 07D730 (2012). [44] D. Mukherjee, N. Bingham, M. Hordagoda, M. H. Phan, H. Srikanth, S. Witanachchi, and P. Mukherjee, J. Appl. Phys. 112, 083910 (2012). [45] J. Dho, N. H. Hur, I. S. Kim, and Y. K. Park, J. Appl. Phys. 94, 7670 (2003). [46] T. Higuchi, T. Yajima, L. Fitting Kourkoutis, Y. Hikita, N. Nakagawa, D. A. Muller, and H. Y. Hwang, Appl. Phys. Lett. 95, 043112 (2009). [47] B. Vertruyen, R. Cloots, M. Ausloos, J. F. Fagnard, and P. Vanderbemden, Phys. Rev. B 75, 165112 (2007). [48] S. Witanachchi, K. Ahmed, P. Sakthivel, and P. Mukherjee, Appl. Phys. Lett. 66, 1469 (1995). [49] Pulsed Laser Deposition of Thin Films, edited by D. B. Chrisey, G. K. Hubler, (Wiley, New York, 1994). [50] S. R. Foltyn, R. C. Dye, K. C. Ott, E. Peterson, K. M. Hubbard, et. al., Appl. Phys. Lett. 59, 594 (1991). [51] R. K. Kelly and A. Miotello, Pulsed Laser Deposition of Thin Films, edited by D. B. Chrisey, G. K. Hubler (Wiley, New York, 1994) p. 55. [52] B. Brailovsky, S. V. gaponov, and V. I. Luchin, Appl. Phys. A 61, 81 (1995). [53] F. Yang, N. Kemik, M. D. Biegalski, H. M. Christen, E. Arenholz, and Y. Takamura, Appl. Phys. Lett. 97, 092503 (2010). [54] M. D. Nguyen, M. Dekkers, E. Houwman, R. Steenwelle, X. Wan, A. Roelofs, T. Schmitz-Kempen, and G. Rijnders, Appl. Phys. Lett. 99, 252904 (2011). [55] J. Schwarzkopf, and R. Fornari, Prog. Cryst. Growth Character. Mater. 52, 159 (2006). [56] D. Mukherjee, M. Hordagoda, R. Hyde, N. Bingham, H. Srikanth, S. Witanachchi, and P. Mukherjee, ACS Appl. Mater. Interfaces 5(15), 7450 (2013). [57] D. Mukherjee, T. Dhakal, M. Phan, H. Srikanth, P. Mukherjee, and S. Witanachchi, Physica B 406, 2663 (2011). [58] H. C. He, J. Ma, Y. H. Lin, and C. W. Nan, J. Phys. D: Appl. Phys. 42, 095008 (2009). [59] X. D. Zhang, J. Dho, S. Park, H. Kwon, J. Hwang, G. Park, D. Kwon, B. Kim, Y. Jin, B. G. Kim, D. Karpinsky, and A. L. Kholkin, J. Appl. Phys. 110, 064115 (2011). [60] W. Chen, Z. H. Wang, W. Zhu and O. K. Tan, J. Phys. D: Appl. Phys. 42, 075421 (2009). [61] N. Ortega, A. Kumar, R. S. Katiyar, and C. Rinaldi, J. Mater. Sci. 44, 5127 (2009). [62] D. Mukherjee, M. Hordagoda, P. Lampen, M. H. Phan, H. Srikanth, S. Witanachchi, and P. Mukherjee, J. Appl. Phys. 115, 17D707 (2014). [63] D. Mukherjee, S. Witanachchi, R. Hyde, and P. Mukherjee, AIP Conf. Proc. 1464, 325 (2012). [64] S. Witanachchi, and P. Mukherjee, J. Vac. Sci. Technol. A 13, 1171 (1995). [65] P. Mukherjee, J. Cuff, and S. Witanachchi, Appl. Surf. Sci. 127, 620 (1998). [66] P. Mukherjee, S. Chen, and S. Witanachchi, Appl. Phys. Lett. 74, 1546 (1999). [67] P. Mukherjee, S. Chen, J. B. Cuff, P. Sakthivel, and S. Witanachchi, J. Appl. Phys. 91, 1828 (2002). [68] P. Mukherjee, S. Chen, J. B. Cuff, and S. Witanachchi, J. Appl. Phys. 91, 1837 (2002). [69] P. Mukherjee, J. B. Cuff, and S. Witanachchi, Rev. Sci. Instrum. 72, 2380 (2001).
Laser Ablation for Multiferroic Heterostructures
61
[70] D. Mukherjee, T. Dhakal, R. Hyde, P. Mukherjee, H. Srikanth, and S. Witanachchi, J. of Phys. D: Appl. Phys. 43, 485001 (2010). [71] T. Dhakal, D. Mukherjee, R. Hyde, P. Mukherjee, M. H. Phan, H. Srikanth, and S. Witanachchi, J. Appl. Phys. 107, 053914 (2010).
