Competing Pathways for Nucleation of the Double Perovskite

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Article

Competing Pathways for Nucleation of the Double Perovskite Structure in the Epitaxial Synthesis of LaMnNiO 2

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Steven R. Spurgeon, Yingge Du, Timothy Droubay, Arun Devaraj, Xiahan Sang, Paolo Longo, Pengfei Yan, Paul G Kotula, Vaithiyalingam Shutthanandan, Mark E. Bowden, James M LeBeau, Chongmin Wang, Peter V Sushko, and Scott A. Chambers Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.6b00829 • Publication Date (Web): 10 May 2016 Downloaded from http://pubs.acs.org on May 12, 2016

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Chemistry of Materials

Competing Pathways for Nucleation of the Double Perovskite Structure in the Epitaxial Synthesis of La2MnNiO6 Steven R. Spurgeon,† Yingge Du,† Timothy Droubay,† Arun Devaraj,† Xiahan Sang,‡ Paolo Longo,¶ Pengfei Yan,§ Paul G. Kotula,k Vaithiyalingam Shutthanandan,§ Mark E. Bowden,§ James M. LeBeau,‡ Chongmin Wang,§ Peter V. Sushko,† and Scott A. Chambers∗,† †Physical and Computational Sciences Directorate, Pacific Northwest National Laboratory, Richland, WA, USA ‡Department of Materials Science and Engineering, North Carolina State University, Raleigh, NC, USA ¶Gatan Inc., Pleasanton, CA, USA §Environmental and Molecular Sciences Laboratory, Pacific Northwest National Laboratory, Richland, WA, USA kSandia National Laboratories, Albuquerque, NM, USA E-mail: [email protected]

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Abstract Double perovskites of the form R2 BB ′ O6 (where R is a rare earth cation and B and B ′ are chemically distinct transition metal cations with half-filled and empty eg orbitals, respectively) are of significant interest for their magnetoelectric properties. La2 MnNiO6 is particularly attractive because of its large expected ferromagnetic moment per formula unit (5 µB f.u.−1 ) and its semiconducting character. If the ideal structure nucleates, superexchange coupling can take place via the B — O — B ′ bonds that form, and the moment per formula unit can attain its maximum theoretical value. However, we show that even in the case of layer-by-layer deposition via molecular beam epitaxy, the system can follow multiple reaction pathways that lead to deviations from the double perovskite structure. In particular, we observe a spatially extended phase in which B-site cation disorder occurs, resulting in Mn — O — Mn and Ni — O — Ni antiferromagnetic domains, as well as the formation of quasi-epitaxial, antiferromagnetic NiO nanoscale inclusions, surrounded by a Mn-rich double perovskite. The coexistence of the double perovskite and secondary phases in oxygen deficient conditions is supported by first-principles modeling. However, extended annealing in air restores long-range B-site order and begins to dissolve the NiO inclusions, yielding an ideal structure and an enhanced ferromagnetic moment. This study reveals fundamental structure-property relationships that may not be apparent during the design phase of a multi-element crystalline solid and illustrates how to engineer a synthetic path to a desired product.

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Introduction Over the past decades a confluence of advanced synthesis techniques, 1,2 data-driven characterization, 3 and robust computational capabilities 4,5 has given rise to guided materials engineering, in which the prevailing hierarchical view of materials is combined with insight into complex interactions across length scales. 6,7 The “materials-by-design” approach, while radically transforming the development of multicomponent systems, tends to overlook the complex kinetic pathways that define solid state materials synthesis. Although we are able to envision almost limitless materials combinations, our ability to synthesize them in practice is constrained by existing characterization and modeling approaches that often fail to capture the inherent complexity of solid state systems. There is currently a disconnect between highly local structural characterization and macroscale properties measurements, resulting in oversimplified or incomplete structure-property models. The design of multicomponent systems in particular represents a grand challenge, 8 spanning materials classes as diverse as Heusler alloys, 9 topological insulators, 10 and complex oxides. 2 These systems often possess complex energy landscapes and varying kinetic pathways that can deviate significantly from equilibrium structures, demanding a novel approach to materials synthesis, characterization, and modeling. A detailed understanding of chemical and structural ordering is particularly important for the design of complex oxides, since these properties govern a wealth of useful electronic, magnetic, and optical functionalities. 11 The development of atomically-precise deposition techniques, such as molecular beam epitaxy (MBE), now enables the synthesis of high-quality oxide heterostructures. 2,12,13 However, the characterization of oxides still relies heavily on non-local X-ray scattering techniques, complemented by highly localized electron microscopy studies—a combination that is often insufficient to adequately correlate atomic-level and macroscale structure. Moreover, in the absence of complementary chemical mapping, bright and dark field imaging alone are often unable to fully characterize non-idealities in oxide thin films. Our inability to connect nano- and macroscale properties, both experimentally 3



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and theoretically, has yielded large uncertainties and hindered our understanding of these technologically significant materials. La2 MnNiO6 (LMNO) exemplifies the challenges and controversies associated with structureproperty engineering of multicomponent systems. LMNO is a ferromagnetic (FM) semiconductor with a relatively high FM Curie temperature of ∼280 K, making it attractive for thermoelectric and spintronic applications. 14 In a perfectly ordered LMNO crystal (P 21 /n space group), Ni and Mn occupy the cation sublattice in a rocksalt-like configuration, giving rise to FM behavior mediated by the superexchange interaction. 15,16 While it is thought that the cation species tend to attain Mn4+ and Ni2+ valences, 17 the dynamics of the ordering process, the role of antisite / antiphase defects, and the presence of secondary phases are still poorly understood phenomena. Antisite defects can promote antiferromagnetic (AF) correlations, giving rise to regions of magnetic disorder, 18 while antiphase boundaries at larger length scales are known to greatly affect the magnetism of other double perovskites. 19 Previous work has shown that low oxygen fugacity during growth can generate extended defects and degrade the magnetic properties of LMNO. 20 Past studies have considered the role and type of such defects, 18,21,22 but characterization has tended to rely on either local or sampleaveraged information and is thus unable to adequately describe the complexity emerging at the mesoscale. 16,19,20,23 The recent development of atomically-resolved energy dispersive X-ray spectroscopy (STEM-EDS) provides an excellent and currently unexploited tool to the study emergence of order in these systems. 24 More broadly, advanced analysis tools have now made it possible to extract quantifiable and statistically-significant information from larger volumes in the pursuit of true multiscale materials characterization. 25,26 A comprehensive approach, leveraging insights from synthesis, characterization, and modeling, is needed to disentangle the web of competing effects present in multicomponent systems. We show that LMNO can serve as a model system in which multiscale methods are combined with insights from first-principles calculations to consistently interpret macroscale magnetic behavior. Here we employ a suite of atomic-level and nanoscale electron microscopy,

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three-dimensional (3D) tomographic analysis, and macroscale properties characterization to explore structure-property relationships across multiple length scales. Aberration-corrected scanning transmission electron microscopy (STEM) offers a simultaneous high-resolution probe of atomic structure and chemistry, but projection issues make it difficult to extend this information to three dimensions. We therefore combine STEM with the emerging technique of laser-assisted atom probe tomography (APT), 27–29 a method to reconstruct 3D volumes of a material at the single-atom level that is being increasingly used in nanoscale analysis of oxide materials. 30 We show that APT provides invaluable complementary information about structure and composition, brilliantly illuminating a network of secondary phases. Finally, we discuss phase stability in light of ab initio modeling and propose a defect model underlying LMNO’s experimentally measured structural and magnetic properties. This array of techniques yields unprecedented insight into the structure and chemical ordering of a model multicomponent system, effectively bridging the gap between the nano- and macroscale.

Results and Discussion In the present work we employ molecular beam epitaxy to grow high-quality, epitaxial thin films. While other methods, such as pulsed laser deposition (PLD) and sputter deposition, offer the possibility of a high oxygen partial pressure during growth and can produce excellent films, they suffer from drawbacks such as macroparticle ejection as well as the presence of energetic species, which can lead to interlayer mixing and extended lattice defects that degrade film quality. 2 PLD in particular has been extensively used to synthesize LMNO films, leading to a range of reported properties and a small epitaxial growth window. 20,22,31–33 While MBE is not subject to the aforementioned effects, the technique is constrained by a rather low oxygen pressure during growth and limited LMNO films have been reported using this synthesis approach. 34 We have grown nominally 40 nm-thick La2 MnNiO6 films on SrTiO3 (STO) (001) sub-

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strates, as described in the methods section. Selected samples were sectioned to four 5 × 5 mm pieces to permit comparison studies that involve post-growth annealing, structural characterization, and property measurements. Pieces were sectioned and then post-annealed at 750 ◦ C for two hours in a flowing air tube furnace to promote chemical ordering. Nonlocal characterization, including X-ray photoelectron spectroscopy (XPS), X-ray diffraction (XRD), and Rutherford backscattering spectrometry (RBS), indicates no detectable change in composition, structure, or chemistry, as shown in supporting information Figure S1. Figure 1 shows bulk in-plane vibrating sample magnetometry (VSM) results from the asgrown and annealed films measured at -173 ◦ C. We observe a significant increase in saturation magnetization (MS ) upon annealing, from 1.3 µB f.u.−1 to 3.5 µB f.u.−1 , accompanied by a slight reduction in coercivity. A variety of mechanisms, including cation ordering and phase separation, have previously been invoked explain the observed increase in moment upon post-growth annealing. 31,32 In particular, prior investigations have focused on possible spin structure changes associated with cation/charge ordering; for instance, a charge disproportionation of the form Ni3+ (d7 : t62g e1g ) + Mn3+ (d4 : t32g e1g ) → Ni2+ (d8 : t62g e2g ) + Mn4+ (d3 : t32g e0g ) may give rise to a transition from an intermediate to high-spin state. 31,35 The theoretical high-spin moment of 5.0 µB f.u.−1 is typically used as an indirect benchmark for chemical/charge ordering; 31 within this framework we may estimate the degree of cation ordering as 3.5/5.0 ≈ 70%. However, such a simple analysis belies assumptions about average film structure, thickness, and composition, which are typically taken from non-local XRD techniques. A ferromagnetic ground state may be expected based purely on superexchange considerations, 36 but should such behavior still hold when chemical inhomogeneities and defects are present? While there has been some discussion of octahedral distortions associated with disorder, 37,38 direct mapping of local order and secondary phases has never been reported in this system. We have conducted extensive local structural characterization of this system, beginning with STEM imaging and chemical mapping. Figure 2.A shows a cross-sectional high-angle

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Figure 1: Macroscale magnetization. Bulk in-plane VSM measurements of the as-grown and annealed (750 ◦ C) films, showing a significant increase in saturation magnetization upon annealing. These measurements were conducted at -173 ◦ C. annular dark field (STEM-HAADF) image of the as-grown film, overlaid with a high magnification image of the film-substrate interface. Because of the similar atomic numbers of the two B-site cations, the image contrast in the LMNO layer is difficult to interpret without the aid of chemical mapping. Using energy dispersive X-ray spectroscopy we are able to measure composition with single atomic-column precision and correlate it to broader-scale chemical features and macroscale magnetic properties. Figure 2.B shows an atomically-resolved STEM-EDS map of the Ni K / Mn K peak ratio in a 4 × 4 nm2 region of the as-grown film. The data have been processed using an automated multivariate statistical analysis (MSA) routine, which both de-noises and extracts pure elemental components from which the relevant counts can be recovered. 39,40 These data show that there is no clear separation of Ni and Mn onto alternating columns, in contrast to what is expected for an ordered film in this projection (see supporting information Figure S2). Figure 2.C shows a very different situation after annealing, marked by the emergence of clearly ordered Mn and Ni planes along the diagonal pseudocubic [111] direction. Quantitatively, we find that the signal on

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a pure Mn or Ni column never exceeds 70% of the maximum, nor does it drop below 30% off the column; however, this is less an indication of disorder than of complex electron-beam delocalization and channeling effects. 24 Previous work on fully-ordered Y2 Ti2 O7 pyrochlore films has found a similar 80% / 20% on / off column behavior, supporting our claim of significant local Ni and Mn ordering. 39 While this result agrees well with the increase in moment shown in Figure 1, the final value is still well short of the theoretical 5.0 µB f.u.−1 , suggesting that the ordering is incomplete or that other factors must influence the magnetic properties.

Figure 2: Chemical mapping of local B-site cation ordering. (A) Representative STEM-HAADF cross-section of the LMNO / STO heterostructure, inset with a highmagnification image of the interface, showing excellent epitaxy and crystallinity. (B) Representative atomic-level STEM-EDS map of the as-grown film taken along the [110] pseudocubic zone-axis. The colormap corresponds to the Ni / Mn ratio, as shown on the inset scale. No long-range pattern in the ratio is evident. (C) Atomic-level STEM-EDS map of the annealed film, overlaid with the theoretical P 21 /n crystal structure in the lower right. In this case there is a clear ordering of the B-sites onto Mn and Ni cation planes along the diagonal pseudocubic [111] direction. While our STEM-EDS results effectively capture the cation disorder-order transition after post-growth annealing, such a highly localized picture can overlook the true nature of the sample, which may possess defects at larger length scales. In particular, by directly measuring pseudocubic interaxial angles over tens of nm, we can look for broad indications of structural disorder. In a projection of the ideal P 21 /n double perovskite along the pseudocubic [100] zone-axis, we expect a pattern of large (red > 92◦ ) and small (blue < 88◦ ) La – La – 8


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La interaxial angles, as shown in Figure 3.A (for additional details see supporting Figure S3). Distortion-corrected HAADF (RevSTEM 41 ) measurements of the annealed film enable us to search for this pattern, which is a signature of ordering in the annealed film. The interaxial angles were then directly measured using the atom columns located from fitting 2D Gaussian peaks to an atom column matrix representation. 42 At high-magnification (see Figure 3.B), we observe a binomial distribution of large and small angles centered around 92.6 ± 0.1◦ and 87.5 ± 0.1◦ , respectively, in excellent agreement with the cation ordering present in STEM-EDS measurements. In contrast, at lower-magnification (see Figure 3.C) we observe extensive regions of disrupted interaxial angles that are closer to 90◦ on average. Prior XRD work has revealed an intimate connection between these angles and A and Bsite occupancy; 38,43 the disordered regions may therefore result from incomplete chemical ordering, antiphase boundaries, 19 or even the presence of secondary phases. 44 These regions will effectively disrupt superexchange coupling and reduce the overall magnetic moment of the system.

Figure 3: Interaxial angle mapping of structural order at extended length scales. (A) Projection of the ideal P 21 /n crystal along the [100] pseudocubic zone-axis, overlaid with the expected large (red ∼ 92◦ ) and small (blue ∼ 88◦ ) La – La – La pseudocubic interaxial angles. (B) High-magnification RevSTEM interaxial angle map of the annealed film, showing a clear layering of long and short angles into pseudocubic (001) planes, commensurate with an ordered P 21 /n structure. (C) Intermediate-magnification RevSTEM interaxial angle map of the same film showing extensive disordered regions at larger (> 10 nm) length scales.

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Progressing from a highly local, atomic-level picture of chemical ordering to several tens of nm, we see that a hierarchy of complex structural factors begins to emerge. At the most basic level we observe cation ordering upon annealing, which promotes ferromagnetism, but our interaxial angle measurements show that these ordered regions are disrupted at larger length scales. This raises the possibility of the existence of disordered or secondary phases, which could limit the saturation magnetization. Parts of our STEM-HAADF images exhibit stripe-like contrast fluctuations, as shown in Figure 4.A. These stripes are qualitatively similar to unidentified defects observed during low PO2 PLD growth. 20 STEM-EDS maps of such a region (see Figure 4.B) indicate that these are Ni-rich regions, which are associated with a noticeable reduction in Mn and La concentrations (see supporting Figure S4 for more information); however, their small 2–5 nm size makes it difficult to quantify their composition and distribution. Moreover, the fine dispersion of these regions renders them invisible to conventional XRD, necessitating another approach to effectively distinguish them from the matrix. We now turn to laser-assisted atom probe tomography to reconstruct the three-dimensional spatial distribution of these regions at high-resolution. APT is widely used to study the properties of metallic alloys and site-specific features such as grain boundaries, 45 but it is considerably more difficult to apply to oxides, where needle specimen frequently fracture due to inhomogeneous evaporation rates. 27 However, APT makes an excellent complement to STEM, providing unique insight into the local environment of the Ni-rich regions and the surrounding matrix and adding a crucial third dimension to projected microscope images. 46 Figure 4.C shows the atom probe reconstruction of a representative region of the annealed film, with Ni-rich regions highlighted using green isocomposition surfaces at 15 at % Ni. This striking image reveals the 3D distribution of a higher density of Ni-rich regions than might be expected from STEM images alone; because of their small size and possible overlap in a 2D projection, many of these regions would otherwise be difficult to detect in a 30–40 nm-thick foil. These regions represent a sizable volume fraction of the film; for an isocomposition

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Figure 4: Complementary STEM / APT mapping of NiO secondary phase morphology and composition. (A, B) STEM-HAADF micrograph and corresponding STEMEDS map of the Ni K peak, revealing the presence of several extended, Ni-rich regions (marked with arrows). Images taken along the [110] pseudocubic zone-axis. (C) APT volume reconstruction of a portion of the annealed film; there is a clear distinction between the NiO regions highlighted by the 15 at % Ni isocomposition surface (green) and the annealed LMNO matrix (colored points). (D, E) Composition profiles across the secondary phase-matrix boundary for both the as-grown and annealed films, respectively, showing a statistically-significant interface broadening after annealing.

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surface defined at 15 at % Ni, as shown in Figure 4, this amounts to a nearly 20–30% volume disruption in the local film matrix. The choice of iso-surface will, of course, dictate the volume fraction of the secondary phase; we have chosen the isocomposition surface value of 15 at % based on the inflection point of the Ni concentration profile across the interface with the matrix, given in the proximity histograms of Figures 4.D-E. The Nirich regions tend to assume a columnar morphology, running from the film substrate to its surface, although isolated structures are occasionally present. Considering both elemental and molecular species, APT allows us to plot the film composition across these boundaries using a proximity histogram 47 without the projection issues present in STEM-EDS, as shown in Figures 4.D-E. For comparison, RBS measurements (shown in supporting Figure S1) establish a nominal volume-averaged cation ratio of La:Mn:Ni = 2:1:1. In contrast, APT measurements of the as-grown sample show that the average elemental composition of the LMNO matrix is 20.9 ± 0.9 La, 14.8 ± 0.8 Mn, 10.5 ± 0.7 Ni, and 53.8 ± 1.1 O at %, while the secondary phase composition is 6.0 ± 1.6 La, 4.5 ± 1.6 Mn, 42.3 ± 7.6 Ni, and 47.2 ± 7.7 O at %. For the annealed sample we find that the average elemental composition of the matrix is 22.9 ± 1.1 La, 14.1 ± 0.9 Mn, 10.2 ± 0.8 Ni, and 52.8 ± 1.2 O at %, while the secondary phase composition is 11.4 ± 1.6 La, 8.0 ± 1.6 Mn, 29.7 ± 4.0 Ni, and 50.9 ± 4.3 O at %. It is known that oxygen content is typically underreported during laser-assisted APT analysis of oxide materials, 30,48 so we have renormalized the La, Mn, and Ni compositions independently of O. For a nominal La2 MnNiO6 composition, we expect the fraction of total cations,

A , (La+Mn+Ni)

to be 0.5, 0.25, and 0.25 for La, Mn, and Ni, respectively. In the as-

grown sample matrix we measure La 0.45, Mn 0.32, and Ni 0.23, while in the as-grown secondary phases we measure La 0.11, Mn 0.08, and Ni 0.80. In the annealed sample matrix we measure La 0.48, Mn 0.30, and Ni 0.22, while in the annealed sample secondary phases we measure La 0.23, Mn 0.16, and Ni 0.60. These measurements show that the as-grown secondary phases exhibit nearly four times the Ni content of the matrix, while the La and Mn content is reduced by half, indicating the formation of a NiO phase. At the same time the

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surrounding matrix is enriched in Mn, compared to the nominal average film composition as measured by RBS. This finding is also supported by STEM measurements, shown in supporting Figure S5, where a local NiO epitaxial structure is directly imaged. Similar structures have been observed before in LaNiO3 / LaAlO3 superlattices, 44 but their presence has not been reported in LMNO double perovskites, which is not unexpected considering they are nearly impossible to detect via conventional non-local methods typically employed. Inspection of the NiO / matrix interface composition also shows an enhancement of La and Mn content on post-growth annealing, consistent with a homogenization process during which these atoms are reincorporated into the NiO regions. Our APT results reveal an extensive network of secondary phases, providing fundamental insight into defect evolution during film growth and annealing. The presence of the NiO phase raises questions about the relative stability of mixed oxide phases, the mechanisms of their formation, and their impact on magnetization. To this end, we have investigated the thermodynamic stability of ordered La2 MnNiO6 and phase-separated systems—NiO and defect-containing LMNO—using ab initio simulations (additional details are given in the supporting information). We have calculated the Gibbs free energies with respect to the binary oxides La2 O3 , MnO2 , and NiO, under the assumption that La:Mn:Ni = 2:1:1, which corresponds to the overall chemical composition of the macroscopic sample measured by RBS. These free energies are plotted as functions of the oxygen chemical potential ∆µO in Figure 5. We consider three cases that would lead to NiO phase separation: 1. The formation of LMNO with a missing NiO (i.e. a vacancy pair): La2 MnNi1−x O6−x and NiO (x = 0.125) 2. The formation of a Mn-rich phase with Ni vacancies, a Ni-rich phase and NiO: La2 Mn1+x Ni1−2x O6 , La2 Mn1−y Ni1+y O6−y and NiO (x = y = 0.125)

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3. The formation of a Mn-rich phase without any B-site vacancies and NiO: La2 Mn1+x Ni1−x O6 and NiO (x = 0.125 & 0.250) Case 1 can be thought of as a perfectly abrupt interface between the NiO secondary phase and the Mn-rich La2 MnNi1−x O6−x matrix. For this case we find that the cost of forming LMNO with a missing pair of neighboring Ni and O atoms is ∼3 eV. Case 2 corresponds to a more realistic, intermixed interface, as observed by STEM and APT; it is described by the formation of La2 Mn1+x Ni1−2x O6 , La2 Mn1−y Ni1+y O6−y , accompanied by NiO. For x = y = 0.125 this reaction is slightly more favorable than the formation of a Ni and O di-vacancy (see Figure 5); however, it is still quite energetically unfavorable in comparison to the formation of ordered LMNO. In general, the stabilities of these phases depend on the values of x and y and their corresponding volume fractions. While it is very challenging to extract accurate values for x and y from our datasets, we find that the existence of a transition region with changing x and y more accurately models the experimental situation. Finally, case 3 corresponds to another type of transition region between NiO phase and the Mn-rich La2 MnNi1−x O6−x matrix. In contrast to case 2, this model assumes the existence of Mn3+ ions and no B-site vacancies. Electron energy loss spectroscopy (EELS) measurements (shown in supporting Figures S6-S7) suggest the presence of Mn3+ and Mn4+ ions in the vicinity of the NiO, supporting this model. For the chemical potential most consistent with our experimental growth conditions (∆µO ≈ −1.6 eV), case 3 is in very close competition with the ordered La2 MnNiO6 phase for both x = 0.125 and x = 0.250. The energy difference between ordered LMNO and case 3 is so small that the latter is entirely possible, particularly once kinetics are considered. Thus, the results of ab initio modeling lend strong support to the formation of a Mn-rich matrix, separated from a NiO phase by an intermixed transition region, in good agreement with our experimental results. We now consider the consequences for magnetism associated with forming the NiO phase and the octahedral environment in the surrounding LMNO matrix. Previous synthesis studies using PLD have concluded that ordered double perovskites are only stable 14


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Figure 5: Ab initio calculations of defect scenarios and proposed defect model. Left: Calculated free energy diagram for several defect scenarios. The ordered LMNO phase is given by the gold colored line, while three primary defect cases are marked (green, navy, and black), as described in the main text. Red corresponds to the highly unfavorable case of Ni vacancy formation. The experimental growth conditions are shaded in blue. Right: Illustration of the proposed defect model, informed by both experiment and theory, where the nucleation of NiO regions during growth disrupts the LMNO double perovskite ordering.

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within a very narrow growth window. 31,49 Outside of this window, the system can tend to form a mix of ordered, disordered, and secondary phases, as has been observed via XRD in (Ba0.9 La0.9 )NiMnO6 . 50 NiO secondary phases have also previously been observed in the growth of LaNiO3 / LaAlO3 on STO, where their presence was attributed to a polar discontinuity. 44 As already mentioned, RBS shows that the films possess excellent average stoichiometry; however, our results clearly illustrate that non-local characterization methods can overlook even extensive compositional partitioning in a thin film system. At the low pressures used during MBE growth, it is possible that NiO seeds may stochastically form at the substrate interface; once established, the NiO phase propagates through the film in a columnar fashion and is kinetically locked into place. Our APT and STEM-EDS results reveal a Mn-rich matrix surrounding the NiO, which was expelled during the formation of the NiO. The high density of interface regions between the matrix and NiO can give rise to octahedral distortions, similar to those observed in our interaxial angle measurements. Through these distortions and misfit dislocations it is possible for the system to accommodate the strain of the additional NiO phase, as shown schematically in Figure 5. Away from these regions, the interplay of MnO6 and NiO6 octahedra can still result in ordering of the double perovskite structure in spite of a slight Mn enrichment, as confirmed by STEM-EDS. This heterogeneous environment will affect the magnetic properties of the system in several ways. First, since the NiO phase is antiferromagnetic, it will act as a magnetically “dead” volume of material. Neglecting exchange interactions between the NiO and surrounding matrix, a 20 − 30% volume fraction of NiO will reduce MS by 1.0 – 1.5 µB f.u.−1 , which accounts for a sizable portion of the reduced moment we observe in Figure 1. In practice these values will be slightly smaller, owing to the decreased moment of the Mn-rich matrix relative to stoichiometric LMNO. Second, the interface region around the NiO regions will likely exhibit a suppressed superexchange interaction owing to extensive anti-site disorder, again disrupting FM ordering. As mentioned earlier, EELS measurements also indicate that the interface region is associated with a reduced Mn valence, which can induce a transition

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from high- to intermediate-spin state, leading to a further decrease in magnetic moment. 31 However, away from the NiO a slightly Mn-rich matrix is still expected to possess FM order at low temperatures. 34 Finally, the NiO phase acts to pin magnetic domain wall motion, leading to a higher coercivity. The post-annealing process has a twofold effect: it both promotes cation ordering in the matrix, thereby increasing the overall magnetic moment, and it facilitates interdiffusion of Ni and Mn, homogenizing the composition across the secondary phase-matrix boundary and reducing the volume fraction of AF regions, as supported by our APT results. Moreover, annealing at high temperatures corresponds to a higher effective oxygen chemical potential, where our ab initio calculations show that an ordered LMNO is energetically preferred. Further annealing is likely to homogenize the structure even more, at the risk of introducing film-substrate intermixing. Our calculations suggest that the formation of NiO secondary phases may be unavoidable at low oxygen fugacity, imposing a fundamental limit on the synthesis of this class of materials. However, our results show that it may be possible to engineer nanocomposite systems in which secondary phases are used to impart useful functionalities.

Conclusions Our study shows that the nature of multicomponent materials synthesis can be much more complex than might be expected from non-local XRD or even high-resolution STEM characterization. Using a combination of aberration-corrected electron microscopy and atom probe tomography, informed by theory calculations, we are able to bridge the gap between atomicscale structure and macroscale properties. We find considerable evidence for the formation of NiO secondary phases, which are surrounded by a defective transition region leading to a Mn-rich matrix. Ab initio calculations support the conclusion that these AF regions are an intrinsic feature of films grown at low oxygen fugacity, which prevents the system from attaining its theoretical magnetic performance without post-annealing. This result illustrates

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the important role played by the energetics of the film growth process and suggests that similar features may be present in many other oxide systems, not just double perovskites. This understanding can directly inform synthesis and modeling efforts, which often incorporate simplistic assumptions about film growth and neglect thermodynamic constraints. More broadly, these results point the way toward designer materials, in which kineticallylimited secondary phases are used to engineer desirable functionalities in multicomponent systems. Taken together, our results open the door to a more comprehensive and predictive understanding of complex materials systems, guided by complementary multiscale synthesis, characterization, and modeling techniques.

Experimental Methods Thin Film Growth and Characterization As-received STO substrates (10 × 10 × 0.5 mm, MTI Corporation) were etched in buffered HF and annealed to ensure a TiO2 terminated surface. 51 They were subsequently cleaned by heating in the MBE chamber at 650 ◦ C for 20 min in an oxygen partial pressure of 6.0 × 10−6 Torr prior to film growth. The LMNO thin films were grown on SrTiO3 (001) substrates in a custom MBE system. La, Mn, and Ni were co-evaporated from high-temperature effusion cells and the evaporation rates were calibrated using a quartz crystal microbalance positioned at the substrate position prior to each growth. The substrate temperature during growth was varied from 600–750 ◦ C, and 650 ◦ C was determined to be the optimal condition for the growth of high quality films reported in this work. An activated oxygen plasma beam (with O2 partial pressure in the chamber set at ∼ 1 × 10−5 Torr) was incident on the sample during deposition. In situ RHEED was used to monitor the overall morphology and surface structure. After deposition, the substrate temperature was reduced at a rate of 30 ◦ C min−1 under the same oxygen environment. In situ high-resolution XPS using monochromatic Al Kα X-rays (hν = 1.486.6 eV) and a VG/Scienta SES 3000 electron energy analyzer in an 18


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appended chamber was carried out to determine cation valence, electronic properties, and film composition by calibrating with Rutherford backscattering spectrometry (RBS). The total energy resolution was 0.5 eV. 52 Lattice parameters and overall crystal quality were determined using high-resolution XRD and reciprocal space mapping with a Philips X’Pert diffractometer equipped with a Cu anode (λ = 1.54065 ˚ A) operating at 45 kV and 40 mA, a hybrid monochromator consisting of four-bounce double crystal Ge (220), and a Cu X-ray mirror. The cation compositions of the films were determined by RBS using 2 MeV He ions and SIMNRA for peak fitting. One sectioned piece of a 40 nm-thick LMNO film was post-annealed at 750 ◦ C (3 ◦ C hr−1 ramp rate up/down) for 2 hours in a tube furnace to promote cation ordering.

Magnetic Properties Bulk magnetic properties were obtained using a Quantum Design Physical Property Measurement System (PPMS) with Vibrating Sample Magnetometer (VSM). In-plane hysteresis loops were measured at -173 ◦ C and background contributions from the holder were subtracted.

Scanning Transmission Electron Microscopy Samples were prepared for STEM with a standard lift-out approach along the STO [100] and [110] zone-axes, using an FEI Helios DualBeam Focused Ion Beam operating at 2–30 keV and 5–7◦ incidence angle. STEM-HAADF images were acquired on JEOL ARM-200CF and FEI Titan probe-corrected microscopes, operating at 200 and 300 keV, respectively. Low-magnification STEM-EDS maps were collected on a JEOL ARM-200CF probecorrected microscope operating at 200 keV with a 34.4 mrad probe semi-convergence angle and 90 mrad probe semi-collection angle. The maps were processed using the ThermoFisher NORAN System 7 software; background subtraction, peak deconvolution, and COMPASS analysis were done within the program. Atomic-resolution maps were collected on a probe19



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corrected FEI Titan G2 80–200 microscope with an 18 mrad probe semi-convergence angle and equipped with four silicon-drift X-ray detectors (SuperX) with a combined solid angle of 0.7 sr. The probe used was ∼ 0.12 nm with 200 pA of current. The X-ray spectral images were acquired for a total dwell time of 1 msec pixel−1 at a magnification of 7.2 MX (300 × 300 pixels at 12 pm pixel spacing). The data were binned 3 × 3 spatially (36 pm) and 2× spectrally (20 eV) to improve the signal-to-noise ratio prior to MSA. RevSTEM was conducted with a probe-corrected FEI Titan G2 60–300 microscope operated at 200 kV with a beam current of ∼ 50 pA, a probe semi-convergence angle of ∼ 13.5 mrad, and a HAADF detector inner semi-angle of ∼ 77 mrad. In each RevSTEM dataset, twenty 1024 × 1024 pixel frames were acquired with 90◦ rotation between frames and a 2 µs pixel−1 dwell time. The time-dependent drift rate and direction were then measured from the angle modulation across the twenty frames. The final RevSTEM image was averaged from aligning and summing all the frames after drift distortion correction. STEM-EELS maps were collected on a JEOL ARM-200CF probe-corrected microscope operating at 200 keV with a 27.5 mrad probe semi-convergence angle, 43 mrad probe semicollection angle, and a 2.5 mm EELS entrance aperture. Spectra were acquired using a 0.25 eV ch−1 dispersion with a nominal 1.35 ˚ A probe size, a ∼120 pA probe current, 0.07 nm pixel size, and 30 ms pixel−1 dwell time.

Atom Probe Tomography APT needle samples were prepared using a FIB based lift-out process on FEI Helios and Quanta DualBeam Microscopes operating at 2–30 keV. 53 The APT analysis was conducted using a CAMECA LEAP 4000XHR system with a 355 nm wavelength pulsed UV laser, using a 100 pJ laser pulse energy and 100 KHz pulse repetition rate, while maintaining the sample temperature at 50–60 K and evaporation rate at 0.003 atoms per pulse. The APT results were analyzed using the IVAS 3.6.8 software.

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Density Functional Theory All calculations were performed using the periodic model. The ordered La2 MnNiO6 and its disordered phases were represented using an 80-atom supercell (La16 Mn8 Ni8 O48 ), corresponding to a 2 × 2 × 4 ABO3 crystallographic cell. The same supercell was used to model non-stoichiometric phases La2 Mn1+x Ni1−x O6 , La2 Mn1+x Ni1−2x O6 , La2 Mn1−x Ni1+x O6−x , and La2 MnNi1−x O6−x that participate in reactions resulting in the NiO phase formation. For every composition we considered all non-equivalent arrangements of the B-site cations and oxygen vacancies and used the lowest-energy one in further analysis. Averaging over all Mn / Ni configurations with the Boltzmann factor corresponding to the growth temperature does not affect our conclusions. The calculations were performed using the Vienna Ab Initio Simulation Package (VASP). 54,55 The projected augmented wave (PAW) method was used to approximate the core electron potential. 56 Exchange-correlation effects were treated within the Perdew-Burke-Ernzerhoff (PBE) functional form of the GGA, modified for solids (PBEsol) 57 within the GGA+U scheme 58 and UMn = 5.25 eV and UNi = 5.77 eV. 59 The plane-wave basis with a 500 eV cutoff was used. The calculations were performed using a 2×2×4 Monkhorst-Pack k-point mesh with its origin at the Γ-point. The charge and spin density distributions were analyzed using the Bader method. 60 The total energy was minimized with respect to the lattice parameters and internal coordinates. The energies of self-consistent calculations were converged to 10−6 eV cell−1 , and the lattice and atomic positions were relaxed until the forces on the ions were less than 0.02 eV ˚ A−1 .

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Acknowledgement S.R.S. thanks Drs. Yuanyuan Zhu and Quentin Ramasse for many constructive discussions. This work was supported by the U.S. Department of Energy, Office of Science, Division of Materials Sciences and Engineering under award #10122. MBE growth, VSM, RBS, STEM, STEM-EDS, STEM-EELS, APT, and DFT were performed in the Environmental Molecular Sciences Laboratory, a national science user facility sponsored by the Department of Energy’s Office of Biological and Environmental Research and located at Pacific Northwest National Laboratory. Pacific Northwest National Laboratory (PNNL) is a multiprogram national laboratory operated for DOE by Battelle. Additional STEM-EDS measurements were performed at Sandia National Laboratories, a multiprogram laboratory managed and operated by Sandia Corporation, a wholly owned subsidiary of Lockheed Martin Corporation, for the U.S. Department of Energy’s National Nuclear Security Administration under contract #DE-AC04-94AL85000 (P.G.K.). J.L. and X.S. gratefully acknowledge support from the National Science Foundation (#DMR-1350273). RevSTEM measurements were performed at the Analytical Instrumentation Facility (AIF) at North Carolina State University, which is supported by the State of North Carolina and the National Science Foundation. XAS measurements were performed at beamline 4-ID-C of the Advanced Photon Source, a U.S. Department of Energy (DOE) Office of Science User Facility operated for the DOE Office of Science by Argonne National Laboratory under Contract #DE-AC02-06CH11357.

Supporting Information Available Details about volume-averaged structural characterization, electron microscopy, and DFT calculations are provided in the supporting information section. This material is available free of charge via the Internet at http://pubs.acs.org/.

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References (1) Gautier, R.; Zhang, X.; Hu, L.; Yu, L.; Lin, Y.; Sunde, T. O. L.; Chon, D.; Poeppelmeier, K. R.; Zunger, A. Prediction and accelerated laboratory discovery of previously unknown 18-electron ABX compounds. Nat. Chem. 2015, 7, 308–316. (2) Martin, L.; Chu, Y.-H.; Ramesh, R. Advances in the growth and characterization of magnetic, ferroelectric, and multiferroic oxide thin films. Mater. Sci. Eng. R Reports 2010, 68, 89–133. (3) Kalinin, S. V.; Sumpter, B. G.; Archibald, R. K. Bigdeepsmart data in imaging for guiding materials design. Nat. Mater. 2015, 14, 973–980. (4) Curtarolo, S.; Hart, G. L. W.; Nardelli, M. B.; Mingo, N.; Sanvito, S.; Levy, O. The high-throughput highway to computational materials design. Nat. Mater. 2013, 12, 191–201. (5) Fischer, C. C.; Tibbetts, K. J.; Morgan, D.; Ceder, G. Predicting crystal structure by merging data mining with quantum mechanics. Nat. Mater. 2006, 5, 641–646. (6) Olson, G. B. Computational Design of Hierarchically Structured Materials. Science 1997, 277, 1237–1242. (7) Gleiter, H. Nanostructured materials: basic concepts and microstructure. Acta Mater. 2000, 48, 1–29. (8) Committee, B. E. S. A. Directing Matter and Energy: Five Challenges for Science and the Imagination; 2007. (9) Miura, Y.; Nagao, K.; Shirai, M. Atomic disorder effects on half-metallicity of the full-Heusler alloys Co2(Cr1xFex)Al: A first-principles study. Phys. Rev. B 2004, 69, 144413.

23



1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

(10) He, L.; Kou, X.; Wang, K. L. Review of 3D topological insulator thin-film growth by molecular beam epitaxy and potential applications. Phys. Status Solidi - Rapid Res. Lett. 2013, 7, 50–63. (11) Ramesh, R.; Keramidas, V. G. Metal-Oxide Heterostructures. Annu. Rev. Mater. Sci. 1995, 25, 647–678. (12) Rondinelli, J. M.; May, S. J. Oxide interfaces: Instrumental insights. Nat. Mater. 2012, 11, 833–834. (13) Chambers, S. A. Epitaxial growth and properties of doped transition metal and complex oxide films. Adv. Mater. 2010, 22, 219–48. (14) Rogado, N. S.; Li, J.; Sleight, A. W.; Subramanian, M. A. Magnetocapacitance and magnetoresistance near room temperature in a ferromagnetic semiconductor: La2 NiMnO6 . Adv. Mater. 2005, 17, 2225–2227. (15) King, G.; Woodward, P. M. Cation ordering in perovskites. J. Mater. Chem. 2010, 20, 5785–5796. (16) Guo, H.; Gupta, A.; Varela, M.; Pennycook, S.; Zhang, J. Local valence and magnetic characteristics of La2NiMnO6. Phys. Rev. B 2009, 79, 2–5. (17) Dass, R.; Goodenough, J. Multiple magnetic phases of La2CoMnO6-δ (0˜δ˜0.05). Phys. Rev. B 2003, 67, 1–9. (18) Menéndez, N.; Garc´ıa-Hernández, M.; Sánchez, D.; Tornero, J. D.; Mart´ınez, J. L.; Alonso, J. a. Charge transfer and disorder in double perovskites. Chem. Mater. 2004, 16, 3565–3572. (19) Egoavil, R.; H¨ uhn, S.; Jungbauer, M.; Gauquelin, N.; Béché, A.; Van Tendeloo, G.; Verbeeck, J.; Moshnyaga, V. Phase problem in the B-site ordering of La 2 CoMnO 6 : impact on structure and magnetism. Nanoscale 2015, 7, 9835–9843. 24


Page 24 of 30

Page 25 of 30

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


(20) Guo, H. Z.; Burgess, J.; Ada, E.; Street, S.; Gupta, A.; Iliev, M. N.; Kellock, A. J.; Magen, C.; Varela, M.; Pennycook, S. J. Influence of defects on structural and magnetic properties of multifunctional La2NiMnO6 thin films. Phys. Rev. B 2008, 77, 174423. (21) Zhao, S.; Shi, L.; Zhou, S.; Zhao, J.; Yang, H.; Guo, Y. Size-dependent magnetic properties and Raman spectra of La2NiMnO6 nanoparticles. J. Appl. Phys. 2009, 106, 123901. (22) Singh, M. P.; Grygiel, C.; Sheets, W. C.; Boullay, P.; Hervieu, M.; Prellier, W.; Mercey, B.; Simon, C.; Raveau, B. Absence of long-range NiMn ordering in ferromagnetic La2 NiMn O6 thin films. Appl. Phys. Lett. 2007, 91, 89–92. (23) Langenberg, E.; Rebled, J.; Estrad, S.; Daumont, C. J. M.; Ventura, J.; Coy, L. E.; Polo, M. C.; Garc´ıa-Cuenca, M. V.; Ferrater, C.; Noheda, B.; Peiró, F.; Varela, M.; Fontcuberta, J. Long-range order of Ni2+ and Mn4+ and ferromagnetism in multiferroic (Bi0.9La0.1)2NiMnO6 thin films. J. Appl. Phys. 2010, 108, 3–6. (24) Allen, L. J.; DAlfonso, A. J.; Freitag, B.; Klenov, D. O. Chemical mapping at atomic resolution using energy-dispersive x-ray spectroscopy. MRS Bull. 2012, 37, 47–52. (25) TMS, Modeling Across Scales: A Roadmapping Study for Connecting Materials Models and Simulations Across Length and Time Scales; 2015. (26) Rohrer, G. S.; Affatigato, M.; Backhaus, M.; Bordia, R. K.; Chan, H. M.; Curtarolo, S.; Demkov, A.; Eckstein, J. N.; Faber, K. T.; Garay, J. E.; Gogotsi, Y.; Huang, L.; Jones, L. E.; Kalinin, S. V.; Lad, R. J.; Levi, C. G.; Levy, J.; Maria, J.-P.; Mattos, L.; Navrotsky, A.; Orlovskaya, N.; Pantano, C.; Stebbins, J. F.; Sudarshan, T. S.; Tani, T.; Scott Weil, K. Challenges in Ceramic Science: A Report from the Workshop on Emerging Research Areas in Ceramic Science. Journal of the American Ceramic Society 2012, 95, 3699–3712.

25



1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

(27) Kelly, T. F.; Larson, D. J.; Thompson, K.; Alvis, R. L.; Bunton, J. H.; Olson, J. D.; Gorman, B. P. Atom Probe Tomography of Electronic Materials. Annu. Rev. Mater. Res. 2007, 37, 681–727. (28) Oberdorfer, C.; Stender, P.; Reinke, C.; Schmitz, G. Laser-Assisted Atom Probe Tomography of Oxide Materials. Microsc. Microanal. 2007, 13, 342–346. (29) Chen, Y.; Ohkubo, T.; Hono, K. Laser assisted field evaporation of oxides in atom probe analysis. Ultramicroscopy 2011, 111, 562–566. (30) Devaraj, A.; Gu, M.; Colby, R.; Yan, P.; Wang, C. M.; Zheng, J. M.; Xiao, J.; Genc, A.; Zhang, J. G.; Belharouak, I.; Wang, D.; Amine, K.; Thevuthasan, S. Visualizing nanoscale 3D compositional fluctuation of lithium in advanced lithium-ion battery cathodes. Nat. Commun. 2015, 6, 8014. (31) Singh, M. P.; Truong, K. D.; Jandl, S.; Fournier, P. Long-range Ni/Mn structural order in epitaxial double perovskite La2 NiMnO6 thin films. Phys. Rev. B 2009, 79, 1–6. (32) Kitamura, M.; Ohkubo, I.; Kubota, M.; Matsumoto, Y.; Koinuma, H.; Oshima, M. Ferromagnetic properties of epitaxial La2NiMnO6 thin films grown by pulsed laser deposition. Appl. Phys. Lett. 2009, 94, 2007–2010. (33) Guo, H.; Burgess, J.; Street, S.; Gupta, a.; Calvarese, T. G.; Subramanian, M. a. Growth of epitaxial thin films of the ordered double perovskite La 2NiMnO 6 on different substrates. Appl. Phys. Lett. 2006, 89, 19–22. (34) Tanaka, H.; Okawa, N.; Kawai, T. Magnetic exchange interactions in perovskite LaMnO3/LaMO3 (M=Ni, Co, Cr, Fe) superlattices. Solid State Commun. 1999, 110, 191–196. (35) Truong, K. D.; Singh, M. P.; Jandl, S.; Fournier, P. Influence of Ni/Mn cation order on

26


Page 26 of 30

Page 27 of 30

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


the spin-phonon coupling in multifunctional La2NiMnO6 epitaxial films by polarized Raman spectroscopy. Phys. Rev. B 2009, 80, 134424. (36) Saha-Dasgupta, T. Magnetism in Double Perovskites. J. Supercond. Nov. Magn. 2013, 26, 1991–1995. (37) Das, H.; Waghmare, U. V.; Saha-Dasgupta, T.; Sarma, D. D. Electronic structure, phonons, and dielectric anomaly in ferromagnetic insulating double pervoskite La2NiMnO6. Phys. Rev. Lett. 2008, 100, 1–4. (38) Sánchez, M. C.; Garc´ıa, J.; Blasco, J.; Sub´ıas, G.; Perez-Cacho, J. Local electronic and geometrical structure of LaNi1xMnxO3+δ perovskites determined by x-ray-absorption spectroscopy. Phys. Rev. B 2002, 65, 144409. (39) Kotula, P. G.; Klenov, D. O.; von Harrach, H. S. Challenges to Quantitative Multivariate Statistical Analysis of Atomic-Resolution X-Ray Spectral. Microsc. Microanal. 2012, 18, 691–698. (40) Kotula, P. G.; Keenan, M. R.; Michael, J. R. Automated analysis of SEM X-ray spectral images: a powerful new microanalysis tool. Microsc. Microanal. 2003, 9, 1–17. (41) Sang, X.; LeBeau, J. M. Revolving scanning transmission electron microscopy: Correcting sample drift distortion without prior knowledge. Ultramicroscopy 2014, 138, 28–35. (42) Sang, X.; Oni, A. A.; LeBeau, J. M. Atom Column Indexing: Atomic Resolution Image Analysis Through a Matrix Representation. Microsc. Microanal. 2014, 20, 1764–1771. (43) Alonso, J. A.; Mart´ınez-lope, M. J.; Casais, M. T.; Macmanus-driscoll, J. L.; de Silva, P. S. I. P. N.; Cohen, L. F.; Fernández-d´ıaz, M. T. Non-stoichiometry, structural defects and properties of LaMnO3+δ with high δ values (0.11δ0.29). J. Mater. Chem. 1997, 7, 2139–2144. 27



1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

(44) Detemple, E.; Ramasse, Q. M.; Sigle, W.; Cristiani, G.; Habermeier, H.-U.; Benckiser, E.; Boris, a. V.; Frano, A.; Wochner, P.; Wu, M.; Keimer, B.; van Aken, P. A. Polarity-driven nickel oxide precipitation in LaNiO3-LaAlO3 superlattices. Appl. Phys. Lett. 2011, 99, 211903. (45) Miller, M.; Forbes, R. Atom probe tomography. Mater. Charact. 2009, 60, 461–469. (46) Diercks, D. R.; Gorman, B. P.; Manerbino, A.; Coors, G. Atom Probe Tomography of Yttrium-Doped Barium-Cerium-Zirconium Oxide with NiO Addition. J. Am. Ceram. Soc. 2014, 97, 3301–3306. (47) Hellman, O.; Vandenbroucke, J.; Ruesing, J.; Isheim, D.; Seidman, D. Analysis of Three-dimensional Atom-probe Data by the Proximity Histogram. Microsc. Microanal. 2000, 6, 437–444. (48) Devaraj, A.; Colby, R.; Hess, W. P.; Perea, D. E.; Thevuthasan, S. Role of Photoexcitation and Field Ionization in the Measurement of Accurate Oxide Stoichiometry by Laser-Assisted Atom Probe Tomography. J. Phys. Chem. Lett. 2013, 4, 993–998. (49) Singh, M. P.; Truong, K. D.; Jandl, S.; Fournier, P. Multiferroic double perovskites: Opportunities, issues, and challenges. J. Appl. Phys. 2010, 107, 09D917. (50) Langenberg, E.; Varela, M.; Garc´ıa-Cuenca, M.; Ferrater, C.; Polo, M.; Fina, I.; Fàbrega, L.; Sánchez, F.; Fontcuberta, J. Epitaxial thin films of (Bi0.9La0.1)2NiMnO6 obtained by pulsed laser deposition. J. Magn. Magn. Mater. 2009, 321, 1748–1753. (51) Qiao, L.; Droubay, T. C.; Varga, T.; Bowden, M. E.; Shutthanandan, V.; Zhu, Z.; Kaspar, T. C.; Chambers, S. A. Epitaxial growth, structure, and intermixing at the LaAlO3 / SrTiO3 interface as the film stoichiometry is varied. Phys. Rev. B 2011, 83, 085408.

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(52) Zhang, K. H. L.; Sushko, P. V.; Colby, R.; Du, Y.; Bowden, M. E.; Chambers, S. A. Reversible nano-structuring of SrCrO3-δ through oxidation and reduction at low temperature. Nat. Commun. 2014, 5, 4669–1–7. (53) Thompson, K.; Lawrence, D.; Larson, D.; Olson, J.; Kelly, T.; Gorman, B. In situ site-specific specimen preparation for atom probe tomography. Ultramicroscopy 2007, 107, 131–139. (54) Kresse, G.; Joubert, D. From ultrasoft pseudopotentials to the projector augmentedwave method. Phys. Rev. B 1999, 59, 1758–1775. (55) Kresse, G.; Furthm¨ uller, J. Efficient iterative schemes for ab initio total-energy calculations using a plane-wave basis set. Phys. Rev. B 1996, 54, 11169–11186. (56) Blöchl, P. E. Projector augmented-wave method. Phys. Rev. B 1994, 50, 17953–17979. (57) Perdew, J. P.; Ruzsinszky, A.; Csonka, G. I.; Vydrov, O. A.; Scuseria, G. E.; Constantin, L. A.; Zhou, X.; Burke, K. Restoring the Density-Gradient Expansion for Exchange in Solids and Surfaces. Phys. Rev. Lett. 2008, 100, 136406. (58) Dudarev, S. L.; Botton, G. A.; Savrasov, S. Y.; Humphreys, C. J.; Sutton, A. P. Electron-energy-loss spectra and the structural stability of nickel oxide: An LSDA+U study. Phys. Rev. B 1998, 57, 1505–1509. (59) Floris, A.; de Gironcoli, S.; Gross, E. K. U.; Cococcioni, M. Vibrational properties of MnO and NiO from DFT+U-based density functional perturbation theory. Phys. Rev. B 2011, 84, 161102. (60) Bader, R. Atoms in Molecules A Quantum Theory; Oxford University Press: Oxford, 1990.

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Competing Pathways for Nucleation of the Double Perovskite

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