J Buchwald
email: joerg.buchwald@ufz.de

M Hennes

Adsorption and diffusion of Au, Pt, and Co adatoms on SrTiO 3 (001) surfaces: A density functional theory study

Keywords: SrTiO 3, Surface diffusion, DFT

We use Density Functional Theory (DFT) and Hubbard-model-based DFT+U calculations to determine preferential adsorption sites and energies of single, isolated Au, Pt, and Co atoms on planar SrTiO 3 (001) surfaces. Based on these results, we employ a nudged elastic band (NEB) approach to calculate relevant diffusion energy barriers for the three transition metal species on SrO and TiO 2 terminated surfaces. This provides valuable quantitative input for future experimental and simulation studies and a sound basis to guide research aiming at controlling the microstructure of transition-metal-doped SrTiO 3 thin films.

Introduction

Transition metal atom doping is a versatile approach that can be used to tailor the optical, catalytic, magnetic, and electrical transport properties of SrTiO 3 . While stoichiometric SrTiO 3 crystallizes in a cubic perovskite structure, which, at room temperature, behaves as a diamagnetic, paraelectric insulator with a large band gap, addition of small amounts of impurity atoms can have a dramatic impact on the physical and chemical properties of this functional oxide. Prominent and technologically relevant examples are the creation of n-type conductivity via introduction of Nb [START_REF] Tomio | Control of electrical conductivity in laser deposited SrTiO 3 thin films with Nb doping[END_REF][START_REF] Zhao | Highly conductive Nb doped SrTiO 3 epitaxial thin films grown by laser molecular beam epitaxy[END_REF], formation of ferromagnetic thin films by addition of Co [START_REF] Lee | Magnetic properties of Co-and Mn-implanted BaTiO 3 , SrTiO 3 and KTaO 3[END_REF][START_REF] Fix | Ferromagnetism in Co-doped (La,Sr)TiO 3[END_REF][START_REF] Kim | Magnetoelastic effects in SrTi 1-x MxO 3 (M = Fe, Co, or Cr) epitaxial thin films[END_REF], Ni [START_REF] Schuler | Huge metastable axial strain in ultrathin heteroepitaxial vertically aligned nanowires[END_REF] or Mn [START_REF] Middey | Evidence of oxygen-vacancyinduced ferromagnetic order in single crystal Mn-doped SrTiO 3[END_REF], doping for memresistive applications [START_REF] Janousch | Andreasson, Role of Oxygen Vacancies in Cr-Doped SrTiO 3 for Resistance-Change Memory[END_REF][START_REF] Muenstermann | Coexistence of Filamentary and Homogeneous Resistive Switching in Fe-Doped SrTiO 3 Thin-Film Memristive Devices[END_REF] or the use of small amounts of Cr [START_REF] Maruska | Transition-metal dopants for extending the response of titanate photoelectrolysis anodes[END_REF][START_REF] Wang | Photophysical and Photocatalytic Properties of SrTiO 3 Doped with Cr Cations on Different Sites[END_REF], Rh and Ir [START_REF] Konta | Photocatalytic activities of noble metal ion doped SrTiO 3 under visible light irradiation[END_REF][START_REF] Kawasaki | Elucidation of Rh-Induced In-Gap States of Rh:SrTiO 3 Visible-Light-Driven Photocatalyst by Soft X-ray Spectroscopy and First-Principles Calculations[END_REF][START_REF] Kawasaki | Electronic structure and photoelectrochemical properties of an Ir-doped SrTiO3 photocatalyst[END_REF] for photoelectrochemical water splitting in the visible spectrum.

Molecular beam epitaxy (MBE) and pulsed laser deposition (PLD) based approaches have been routinely employed to obtain doped SrTiO 3 thin films with tunable composition and good structural quality. In contrast to ion bombardment, adding the dopant species during growth does not damage the thin films nor require subsequent annealing steps [START_REF] Moretti | Niobium implantation effects in BaTiO 3 and SrTiO 3[END_REF][START_REF] Thevuthasan | Accumulation and thermal recovery of disorder in Au 2+irradiated SrTiO 3[END_REF]. However, the introduction of metal atoms during synthesis raises questions about their spatial distribution in the surrounding oxide. Depending on the growth conditions, the system might exhibit phase separation tendencies, and the added transition metals can form metallic precipitates, which will impact the final functionality of the thin film. For some applications, this clustering of metal atoms might be seen as a detrimental side effect, that needs to be avoided. However, it can also be sought after, when one is interested in the synthesis of SrTiO 3 thin films with nanocomposite structure [START_REF] Schuler | Huge metastable axial strain in ultrathin heteroepitaxial vertically aligned nanowires[END_REF][START_REF] Kawasaki | Photoelectrochemical water splitting enhanced by selfassembled metal nanopillars embedded in an oxide semiconductor photoelectrode[END_REF][START_REF] Hennes | Growth of vertically aligned nanowires in metaloxide nanocomposites: kinetic Monte-Carlo modeling versus experiments[END_REF][START_REF] Weng | Ultrathin Ni nanowires embedded in SrTiO 3 : Vertical epitaxy, strain relaxation mechanisms, and solid-state amorphization[END_REF][START_REF] Lippmaa | Noble metal clustering and nanopillar formation in an oxide matrix[END_REF].

In order to guide experimental research and obtain thin films with the desired microstructure, knowledge of the diffusion properties of dopant adatoms on SrTiO 3 surfaces is required. If the in-plane mobility during synthesis is low and the thin film growth velocity is high enough, the system will very likely incorporate the dopants as isolated atoms. On the contrary, fast diffusion will increase the propensity for phase separation to occur. To address this problem quantitatively, one requires information about the adsorption sites and diffusion pathways of metal atoms on SrTiO 3 surfaces. Surprisingly, this has not been studied in detail yet. While computational data are available describing the adsorption and diffusion of a variety of elements on other prototypical oxides, like MgO for example [START_REF] Yudanov | Systematic Density Functional Study of the Adsorption of Transition Metal Atoms on the MgO(001) Surface[END_REF][START_REF] Hu | Initial stage of Ag deposition on regular MgO(001) surface: A DFT study[END_REF][START_REF] Shim | Ab initio study of early stage Nb growth on MgO(001)[END_REF], only few studies on SrTiO 3 have been performed yet. Adsorption energies were calculated for Pt and Au on SrTiO 3 , but the problem has not been addressed for Co adatoms [START_REF] Wei | Au adsorption and Aumediated charge transfer on the SrO-termination of SrTiO3 (001) surface[END_REF][START_REF] Asthagiri | DFT study of Pt adsorption on low index SrTiO 3 surfaces: SrTiO 3 (100), SrTiO 3 (111) and SrTiO 3 (110[END_REF]. With regard to diffusion, the emphasis was predominantly put on pure SrTiO 3 growth [START_REF] Guhl | Oxygen adatoms at SrTiO 3 (001): A density-functional theory study[END_REF][START_REF] Hong | Surface diffusion on SrTiO 3 (100): A temperature accelerated dynamics and first principles study[END_REF]. Recently, an attempt has been made to perform calculations for a Ni atom on SrTiO 3 [START_REF] Hennes | Growth of vertically aligned nanowires in metaloxide nanocomposites: kinetic Monte-Carlo modeling versus experiments[END_REF].

To fill this gap, we scrutinize Co, Pt and Au adatom diffusion on planar strontium titanate (STO) (001) surfaces using a computational approach based on density functional theory (DFT), with particular emphasis on the relevance of the (DFT+U) formalism. We will start the following sections by giving an overview of our computational methods and then proceed with a description of our results. These will consist of a detailed presentation of adsorption sites and energy barriers for both terminations. We will then close this work with a brief discussion of the results before turning to the summary and conclusion.

Computational details

All DFT calculations were carried out using the Quantum Espresso software package [START_REF] Giannozzi | QUANTUM ESPRESSO: a modular and open-source software project for quantum simulations of materials[END_REF] with ultrasoft pseudopotentials and PBE-GGA exchange correlation [START_REF] Perdew | Generalized Gradient Approximation Made Simple[END_REF][START_REF] Zhang | Comment on "Generalized Gradient Approximation Made Simple[END_REF]. The kinetic energy cutoff was set to 35 Ry, while the charge density cutoff was chosen to be 500 Ry. The integration over the Brillouin zone was performed with a 3×3×1 Monkhorst-Pack mesh [START_REF] Monkhorst | Special points for Brillouin-zone integrations[END_REF]. By optimizing the unit cell and atom positions, we computed a value of a STO 0 = 3.923 Å as bulk lattice constant which is in good agreement with earlier reported experimental and computed values for the PBE functional [START_REF] Piskunov | Bulk properties and electronic structure of SrTiO3, BaTiO3, PbTiO3 perovskites: an ab initio HF/DFT study[END_REF]. The simulation cell employed in our study consisted of a single STO-cube with volume 3a STO 0 ×3a STO 0 ×3a STO 0 and periodic boundary conditions in-plane, combined with 3 a STO 0 vacuum inserted in z direction to create a (001) surface. The SrTiO 3 surface was structurally relaxed using the BFGS algorithm [START_REF] Broyden | The Convergence of a Class of Double-rank Minimization Algorithms 1. General Considerations[END_REF][START_REF] Fletcher | A new approach to variable metric algorithms[END_REF][START_REF] Goldfarb | A family of variable-metric methods derived by variational means[END_REF][START_REF] Shanno | Conditioning of quasi-Newton methods for function minimization[END_REF] until the force components on every atom were less than 10 -3 Ry/a 0 2 , and changes in the total energy were found smaller than 10 -4 Ry. Single Au, Co, and Pt atoms were placed at various canonical positions (e.g., on top of oxygen atoms or at bridge positions between two oxygen atoms, as shown in Fig. 1) corresponding to a surface coverage θ = 1/9. Keeping the lower atom layer fixed, we performed further geometry optimization on the system. For all cells containing magnetic Co atoms, we 2 Here, a 0 denotes the bohr radius. used a spin-polarized calculation scheme. Adsorption energies were calculated using

E a = E total -E STO -E m (1)
where E total is the energy of the adsorbed system, E STO , the energy of the clean STO surface and E m the energy of a single metal atom (i.e., Au, Pt, Co). To compute the height of the energy barriers, we employed the nudged elastic band (NEB) method with a Climbing Image (CI) scheme as implemented in Quantum Espresso to assess the height of the energy barrier. For the open-shell 3d and 5d metals Co and Pt, the results were additionally computed by means of DFT+U calculations using a Hubbard U correction from linear response [START_REF] Cococcioni | Linear response approach to the calculation of the effective interaction parameters in the LDA + U method[END_REF] to overcome the problem of over-delocalization in these systems. Due to different system configurations of the adatom sites, different values of U were obtained. For Co, we found U to be within 4.6-4.9 eV and for Pt between 2.8 eV and 3.3 eV. In this paper, we present the energies referring to maximum values of U only, as all energy differences were found to be less than 0.05 eV.

Results and Discussion

For both common surface terminations, we explored typical binding sites for Au, Pt, and Co atoms in order to get appropriate minimal energy configurations as starting points for the NEB calculations (Fig. 1). In Tables 123we list the computed adsorption energies as well as the corresponding next neighbor distances. A detailed description of these results and comparison with literature data will be provided in the remainder of this section.

SrO-terminated surface

For the SrO-terminated surface, we identified the top oxygen positions (TO) as local minimal energy points for all three transition metal species. For Au, the direct TO position is nearly equally favorable to a distorted TO position, forming an O-O-Au angle of 66.4 degrees, a configuration which has not been reported so far [START_REF] Wei | Density Functional Theory Study of Ag Adsorption on SrTiO 3 (001) Surface[END_REF]. The Pt position appears also slightly distorted in the TiO plane, forming an O-O-Pt angle of 61 degrees with the nearby oxygen, which also contrasts with earlier results [START_REF] Asthagiri | First principles study of Pt adhesion and growth on SrO-and TiO 2 -terminated SrTiO 3 (100)[END_REF], where no such distortion was reported. This might be related to the symmetric monolayer arrangements used in [START_REF] Asthagiri | First principles study of Pt adhesion and growth on SrO-and TiO 2 -terminated SrTiO 3 (100)[END_REF] and the tiny activation barrier we found, between the minimal position and the high symmetry TO configuration, mainly when no +U correction is applied. In contrast, we find that the Co adatom sits directly on top of the oxygen atom. From a quantitative perspective, our results are in reasonable agreement with available literature data. In the TO position, the adhesion energy of Au and the Au-O bond length (Table 1) are close to the results given by Wei et al.: E ads = 1.26 eV and d Au-O = 2.17 Å [START_REF] Wei | Au adsorption and Aumediated charge transfer on the SrO-termination of SrTiO3 (001) surface[END_REF]. Note that the difference in bond length is around 1%, while the binding energies differ by more than 8%. The discrepancy is more pronounced for Pt (Table 2), where Asthagiri and Sholl report E ads = 2.77 eV and d Pt-O = 1.97 Å for a TO site position in their GGA calculations [START_REF] Asthagiri | DFT study of Pt adsorption on low index SrTiO 3 surfaces: SrTiO 3 (100), SrTiO 3 (111) and SrTiO 3 (110[END_REF]. This corresponds to a difference of roughly 2% in bond length and up to 30% difference in the binding energies. We emphasize, however, that those earlier calculations were performed for significantly larger surface coverage of the Pt atoms (θ = 0.25). In fact, for Au as well as for Pt on SrTiO 3 , it has been shown that a decrease of θ correlates with higher adsorption energy [START_REF] Wei | Au adsorption and Aumediated charge transfer on the SrO-termination of SrTiO3 (001) surface[END_REF][START_REF] Asthagiri | DFT study of Pt adsorption on low index SrTiO 3 surfaces: SrTiO 3 (100), SrTiO 3 (111) and SrTiO 3 (110[END_REF]. In addition to the TO positions, we identified BrO sites for the Au and Pt adatoms as minimum energy positions (Table 1-2). Surprisingly, for Au, the energy of this bridge site is slightly lower than the TO position. This clearly challenges earlier findings, where the TO placement of the Au adatom was identified as the global minimum of the system [START_REF] Wei | Au adsorption and Aumediated charge transfer on the SrO-termination of SrTiO3 (001) surface[END_REF]. It is interesting to notice that in the same study, the authors analyzed a similar adsorption configuration, substituting a single subsurface Ti atom by Nb. This gave rise to BrO minimum energy configurations, identical to the one observed in the present case. Considering that both configurations are very close in energy (∆E = 40 meV), we might attribute the observed discrepancy to a different relaxation of the surface. This, again, might be traced back to the amount of Au surface coverage. For the Pt atom, we also identified BrO sites as local minima. However, in contrast to Au, their energy is larger by approximately 0.9 eV. Thus, as already described in earlier studies, we conclude that the TO position represents the preferential binding site [START_REF] Asthagiri | First principles study of Pt adhesion and growth on SrO-and TiO 2 -terminated SrTiO 3 (100)[END_REF]. Finally, we find a qualitatively different behavior for the Co-adatom, which tends to bind to both near-by oxygen atoms. This gives rise to an additional minimum located between the TO and the BrO positions, which is very similar to what has been observed for Ni adsorption on SrO terminated SrTiO 3 surface [START_REF] Hennes | Growth of vertically aligned nanowires in metaloxide nanocomposites: kinetic Monte-Carlo modeling versus experiments[END_REF].

Having identified preferential adsorption sites, we now turn to the description of diffusion barriers on SrO -terminated surfaces. These were computed using NEB-CI calculations with one to five intermediate images. For paths containing more than two relevant minimal points, the path was split into segments and reassembled after- wards. The resulting lowest energy paths are given in Fig. 234. For the Au adatom, the minimum energy diffusion path connects two BrO positions via a TO site. The activation barrier between the BrO and the TO position is found to be small E act = 0.1 eV and lies on the same energy level as the distorted 'TO' positions (I1) ( Fig. 2). While these were also found to be minimal positions, they are not identified as such in the NEB diagram. We thus conclude that the activation barrier I1→BrO is negligible compared to the global barrier of BrO→TO. For Pt, as already emphasized, the minimal state consists in a slightly shifted TO position (Fig. 3). A change of orientation with respect to the oxygen atom is related to the crossing of a small barrier of 0.03 eV (GGA+U: 0.13 eV) from TO1 to TO2. To jump from one oxygen atom to another, a barrier of 1.0 eV (GGA+U: 1.04 eV) has to be overcome. The BrO position serves as an intermediate minimum. Similar to the case of Au diffusion, the Pt adatom can escape this minimum by crossing an energy barrier of 0.15 eV (GGA+U: 0.05 eV). Finally, the behavior of the Co atom is qualitatively different. While the diffusion path connects two TO positions, intermediate steps involve twofold coordinated sites. The reaction barrier height equals E act = 0.75 eV (GGA+U: 1.24 eV) from the one-oxygen coordinated TO to the two oxygen coordinated I2 minimum position (Fig. 4). The transition between the two equivalent two-coordinated states I2 and I3 is reached by crossing a barrier of E = 0.7 eV (GGA+U: 0.4 eV), passing over the BrO position, which, unlike for Au and Pt does not represent a local minimum. To eventually return to the original minimum state, i.e., dissolving from one oxygen atom, a barrier of E act = 1.05 eV (GGA+U: 0.76 eV) has to be overcome.

TiO 2 -terminated surface

The adsorption behavior of Au, Pt, and Co on TiO 2 differs markedly from what we found on the SrO surface. In all three cases, a steep energy minimum was observed at the fourfold hollow site H. Note that such a fourfold coordinated minimum has also been reported for Ag and Ni adsorption on SrTiO 3 [START_REF] Hennes | Growth of vertically aligned nanowires in metaloxide nanocomposites: kinetic Monte-Carlo modeling versus experiments[END_REF][START_REF] Wei | Density Functional Theory Study of Ag Adsorption on SrTiO 3 (001) Surface[END_REF]. With respect to the TO position that acts as a transition point between cells we find an energy difference of 0.29 eV for Au, 1.28 eV (GGA+U: 1.38 eV) for Pt and 2.75 eV (GGA+U: 1.75 eV) for Co between both states. For Au and Co, this position indeed corresponds to a global minimum (Table 1-3). In contrast for Pt, our GGA calculations suggest another, even lower-lying site. Pt atoms coordinated to two Oatoms yield an overall energy lower by 90 meV (Table 2). As will be shown later, this picture changes when using the more advanced DFT+U formalism. It is noteworthy that our results conflict with those presented in [START_REF] Asthagiri | DFT study of Pt adsorption on low index SrTiO 3 surfaces: SrTiO 3 (100), SrTiO 3 (111) and SrTiO 3 (110[END_REF], where the TO site is claimed to be energetically preferred with an adsorption energy of 3.47 eV, while our calculations indicate E a = 2.87 eV for the same site using the GGA approach.

As shown in Fig. 2 and Fig. 4, the diffusion paths of Au and Co adatoms exhibit a simple structure and are characterized by a single saddle point. The hopping from one STO unit cell to the adjacent one happens via the TO position for Au and Co (Fig. 2 and4), a transition that is reminiscent of the diffusion path of Ni on TiO 2 -terminated STO [START_REF] Hennes | Growth of vertically aligned nanowires in metaloxide nanocomposites: kinetic Monte-Carlo modeling versus experiments[END_REF]. For Au, the barrier height is small and amounts to 0.29 eV while we obtain a substantial barrier of 2.75 eV (GGA+U: 1.75 eV) for Co. A similarly high value was observed in the case of Ni diffusion on the same planar surface [START_REF] Hennes | Growth of vertically aligned nanowires in metaloxide nanocomposites: kinetic Monte-Carlo modeling versus experiments[END_REF]. Qualitative differences in the energy landscape were found for Pt, where the diffusion path appears to be significantly more complex. In addition to the hollow site H, another minimum is observed at an intermediate state I2 in which Pt is linearly coordinated to two oxygen atoms, pulling one of them out of the surface. The jump between these two states requires to overcome a barrier of 1.19 eV (1.27 eV). Once trapped in the twofold coordinated state, the Pt atom can climb into a TO position by overcoming a barrier of 1.72 eV (GGA+U: 1.42 eV). In contrast to Au and Co, this configuration represents a local minimum with short dwell time, being separated from the adjacent minimum by a small barrier of 0.32 eV (0.33 eV). Unfortunately, data on the diffusion of Pt on SrTiO 3 is scarce. In references [START_REF] Jin | Effects of single metal atom (Pt, Pd, Rh and Ru) adsorption on the pho-tocatalytic properties of anatase TiO 2[END_REF] and [START_REF] Zhou | Growth of Pt particles on the anatase TiO 2 (101) surface[END_REF], a detailed study of Pt n cluster diffusion on anatase (101) TiO 2 is presented. For the case of n = 1, the authors identified minimal energy configurations corresponding to equally stable O-O bridge and O-Ti-O sites [START_REF] Zhou | Growth of Pt particles on the anatase TiO 2 (101) surface[END_REF]. Diffusion barrier heights reported in this study ranged from 0.86 eV to 1.41 eV, which is surprisingly close to our findings despite the difference in the surface structure of the two systems.

DFT calculations vs. DFT+U

While DFT+U results have been given alongside DFT calculations throughout this paper, we close this section by providing additional remarks concerning the two approaches.

The use of the DFT+U approach appears relevant to us in the present context, as it efficiently compensates the effect of over-delocalization. A stronger impact is thus to be expected for highly localized 3d-when compared to 5d-metals, due to their more extended radial wave functions [START_REF] Lan | A linear response approach to determine Hubbard U and its application to evaluate properties of Y 2 B 2 O 7 , B = transition metals 3d, 4d and 5d[END_REF]. This is fully compatible with our simulation results. Indeed, adsorption energy diffusion calculations for Co and Pt reveal that the Hubbard U correction amounts to 20 percent in the adsorption energies of Pt while it can become as high as 40 percent in the case of Co. A clear impact on the height of the diffusion barriers is also evidenced. In the case of Pt, the correction is marginal but suggests that the H site is indeed the global minimum of the system. The DFT+U approach seems to have a substantial effect on the adsorption energies of Pt amounting to a difference of half an electronvolt and to the energy of the two-coordinated intermediate state while the structure stays nearly the same (GGA:

d(O 1 Pt) = 2.05 Å, d(O 2 Pt) = 2.0 Å, O 1 PtO 2 = 174.45 • , GGA+U: d(O 1 Pt) = 2.06 Å, d(O 2 Pt) = 2.0 Å, O 1 PtO 2 = 174.48 • ).
In contrast, for Co adatom diffusion, we find pronounced differences between both approaches: While the DFT+Ucalculations reveal only a small effect on the barrier heights of paramagnetic Pt, a significant impact on the diffusion energies of the magnetic Co adatom is unraveled, leading to a shift of global minimum positions on the SrO surface as well as a reduction of about 1 eV in the activation energy on the TiO 2 terminated surface (Fig. 4).

Summary and conclusion

The diffusion behavior of different transition metals (Au, Pt, Co) on STO(001) was investigated by means of first principle DFT calculations. Minimal positions for the adatoms were determined on both possible planar surface terminations, and representative diffusion paths were found using the NEB-CI method. We observe that (a) the choice of the SrTiO 3 -termination has a large impact on the diffusion behavior of the adatoms, with enhanced mobility observed on the SrO surface, (b) Au diffuses much faster on SrTiO 3 , than Co and Pt. With an activation barrier smaller than 0.3 eV, it can be surmised that Au atoms will more likely aggregate into clusters at the surface of the perovskite, even when deposited at moderate temperatures. And finally, (c) our study unravels the need to compensate over-delocalization effects to obtain a more reliable description of the diffusion properties of 3d metals such as Co.

Although doping of SrTiO 3 has evolved into a mature research field, computational approaches that are able to grasp the initial stages of thin film growth remain scarce. Coarse-grained kinetic Monte Carlo (KMC) computation schemes can serve as a viable tool to gain a better understanding of growth phenomena at oxide surfaces and unravel whether or not, aggregation of dopant atoms will take place. However, such studies require detailed information on the hopping probabilities of the surface atoms, which is difficult to assess experimentally. The data presented herein can thus provide a solid starting point for further computational investigations, ultimately aiming at obtaining a better understanding of processing parameters relevant to SrTiO 3 thin film doping.

Acknowledgements

Computing time provided by the Center for Information Services and High Performance Computing (ZIH), TU Dresden is gratefully acknowledged. M. H. acknowledges financial support from the French Embassy in Berlin (Service pour la Science et la Technologie) and Campus France.

Figure 1 :

 1 Figure 1: Binding sites on a (a) SrO-and (b) TiO 2 -terminated STO surface: Sr atoms are shown in yellow (top-Sr sites labeled TSr), O atoms in red with top-oxygen sites (labeled TO), Ti atoms are shown in pink (top-Ti positions labeled TTi). Bridge oxygen positions (BrO) on the SrO termination and fourfold hollow sites (H) on the TiO 2 surface are shown as well (color online).

Figure 2 :

 2 Figure 2: a) TO minimal configuration of an Au adatom on SrOterminated STO. b) H minimal configuration of Au on TiO 2terminated STO. c) Lowest energy paths of Au adatoms on a SrO (blue) and a TiO 2 (red) -terminated STO surface.

Figure 3 :

 3 Figure 3: a) TO minimal configuration of Pt adatom on SrO terminated STO. b) H minimal configuration of Pt on TiO 2 termnated STO. c) Lowest energy paths of Pt adatoms on a SrO (blue) and TiO 2 (red) -terminated STO surface.

Figure 4 :

 4 Figure 4: a) TO minimal configuration of Co adatom on SrO terminated STO. b) H minimal configuration of Co on TiO terminated STO. c) Lowest energy paths of Co adatoms on a SrO (blue) and TiO 2 (red) -terminated STO surface.

Table 1 :

 1 Adsorption energies and next-neighbor distances computed for different Au adatom arrangements on SrO-and TiO 2 -terminated surfaces.

			SrO termination
	site E a (eV) d Au-O ( Å) d Au-Sr ( Å) d Au-Ti ( Å)
	TO	1.37	2.19	3.32	-
	BrO	1.41	3.4	3.2	-
			TiO 2 termination
	site E a (eV) d Au-O ( Å) d Au-Sr ( Å) d Au-Ti ( Å)
	H	0.6	2.21	3.49	3.27

Table 2 :

 2 Adsorption energies and next-neighbor distances computed for different Pt adatom arrangements on SrO-and TiO 2 -terminated surfaces.

			SrO termination
	site	E a (eV) E +U a (eV) d Pt-O ( Å) d Pt-Sr ( Å) d Pt-Ti ( Å)
	TO	3.66	3.01	1.93	3.16	-
	BrO	2.74	2.02	2.12	2.7	-
			TiO 2 termination
	site	E a (eV) E +U a (eV) d Pt-O ( Å) d Pt-Sr ( Å)	d Pt-Ti
	H	4.16	3.52	2.05	2.91	2.94
	TO	2.87	2.14	1.97	-	2.84
	O-PT-O	4.27	3.23	2.03	3.21	2.43

Table 3 :

 3 Adsorption energies and next-neighbor distances computed for different Co adatom arrangements on SrO-and TiO 2 -terminated surfaces.

			SrO termination
	site	E a (eV) E +U a (eV) d Co-O ( Å) d Co-Sr ( Å) d Co-Ti ( Å)
	TO	1.64	1.03	1.80	3.56	-
	O-Co-O	1.93	0.65	1.73	3.19	3.53
			TiO 2 termination
	site	E a (eV) E +U a (eV) d Co-O ( Å) d Co-Sr ( Å) d Co-Ti ( Å)
	H	3.81	3.03	1.97	3.03	3.04


Dr Guillaume Radtke
email: guillaume.radtke@upmc.fr

Dr Marcel Hennes

Dr Matthieu Bugnet

Dr Quentin M Ramasse

Dr Xiaorong Weng

Dr Dominique Demaille

Benoît Gobaut

Dr Philippe Ohresser

Dr Edwige Otero

Dr Fadi Choueikani

Benoit Gobaut

Amélie Juhin

Dr Philippe Sainctavit

Dr Yunlin Zheng

Franck Vidal
email: franck.vidal@insp.jussieu.fr

Dr A Juhin

Dr B Gobaut

Atomic-Scale Study of Metal-Oxide Interfaces and Magnetoelastic Coupling in Self-Assembled Epitaxial Vertically Aligned Magnetic Nanocomposites

Keywords: interfaces, magnetic nanocomposites, electron microscopy, electron energy loss spectroscopy, x-ray magnetic dichroism

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Introduction

Vertically aligned nanocomposites (VANs) form a class of hybrid thin films, [1,2] that has attracted significant interest since the discovery of self-assembled vertical nanopillar microstructures in the La0.67Ca0.33MnO3:MgO system. [3] Possible applications for these heterostructured materials have been proposed in the fields of nanomagnetism, [4] spintronics, [5,6] multiferroics, [7,14] catalysis and ionic conductors. [15] The versatility of VANs results from their peculiar nanoarchitecture: they are composed of an array of nanopillars vertically embedded and epitaxially coupled to a surrounding single crystal matrix, which is depicted schematically in Figure 1 for the case of the particular combination of materials studied in this work. While initial work on the growth of VANs reported the successful synthesis of vertical metallic nanostructures in an oxide host (Fe:LaSrFeO4, [16] ) most of the attention turned to oxide/oxide VANs in the following years, driven by the wide range of functionalities available in oxide materials and the good structural compatibility between spinel and perovskite structures. In particular, ferroelectric/ferrimagnetic VANs were intensively studied with the aim of coupling the ferroic properties in an artificial multiferroic material. [7,14] Recent years however have seen the emergence of a variety of studies on metal/oxide VANs coupling transition or noble metals with oxide host matrices. The successful growth of nanocomposites consisting of elemental Fe, Ni, Co, Au, Pt, Ir nanorods, as well as of CoxNi1-x alloys, epitaxied in SrTiO3 (STO), BaTiO3, BaZrO3, or CeO2 thin films was demonstrated. [17,26] Like their oxide/oxide counterparts, metal/oxide VANs are of interest in many areas: Au:BaTiO3 VAN systems are appealing in the field of optical metamaterials, [23] Pt:STO and Ir:STO VANs were demonstrated to exhibit enhanced activity in photoelectrochemical water splitting, [22] VANs incorporating Fe, Ni or Co transition metals display interesting properties such as large perpendicular magnetic anisotropy potentially appealing for the development of materials for data storage. [19,20,24] The large interface/volume ratio is a salient feature of the VAN nanoarchitecture and the vertical interface between the pillars and the surrounding matrix plays a prominent role in determining the final functionality of the nanocomposite. For example, the vertical interface can provide additional conduction channels in ionic conductors; the photocarrier transport can be more efficient due to the Schottky junctions around metallic pillars embedded in a semiconducting matrix; efficient magneto-elastic coupling through epitaxial strain is intimately related to the structure of the vertical interface. [27,28] The metal/oxide interface results from the immiscibility of the two phases, but grows in out-ofequilibrium conditions. Its structure/composition may thus reflect the fact that the system is stuck in a metastable state. In addition, the creation of interfacial compounds might also be envisioned. As a consequence, the interfacial region may not be as sharp as in oxide/oxide systems. Despite the crucial role played by the vertical interface, studies addressing its structure/composition at the atomic scale are scarce, possibly due to the difficulty to access the interface compared to more conventional planar heterostructures.

In this work, we therefore present an atomic-scale analysis of metal/oxide interfaces using Ni:SrTiO3 VANs (Ni:STO) as a model system. Ni:STO VANs have been extensively studied recently, from both a theoretical and an experimental point of view. The initial growth stages of this composite were simulated using a kinetic Monte-Carlo approach based on density functional theory. [25] The interfaces observed in this study were sharp, but these results must be considered with care, as the underlying coarse-grained model might not be able to grasp the atomic details of the interface structure. In a more recent study, a detailed analysis of the structure of the nanocomposite was given, unveiling a tilting relaxation mechanism as well as a possible amorphization of the embedded metallic structure upon diameter reduction. [26] No experimental data on the chemical composition of the interface is available yet. In the present paper, we combine scanning transmission electron microscopy (STEM), atomic resolution electron energy-loss spectroscopy (EELS) with x-ray magnetic circular dichroism (XMCD) and conventional magnetometry measurements to show that the interfaces do not present any sign of oxidation and show only very weak interdiffusion, two essential prerequisites to eventually obtain pronounced magneto-elastic coupling effects.

Results and discussion

A global view of the sample, observed along the [001] zone axis of the STO matrix, is displayed in Figure 2(a). The High Angle Annular Dark Field (HAADF) contrast reveals a set of bright, well separated Ni nanowires of diameter close to 2 nm in the midst of the STO matrix, distant from each other by about 10 to 20 nm. A number of smaller diameter, shorter wires do not reach the surface of the cross-sectional sample, such that the STO and Ni lattices appear superimposed.

The faceted nanowire visible at the center of the image is presented at higher magnification in Figure 2(b), and was used as a prototypical object for the atomic-scale spectroscopic investigations presented in the remainder of this work. The metal/oxide interface is structurally abrupt with no visually detectable phase reaction between the two materials. Epitaxial strain indicative of a partially coherent interface between the nanowire and the matrix is visible through the deformation of the Ni lattice planes. Atomic-resolution elemental maps acquired on this nanowire are presented in Figure 2(d), (e) and (f) for Ti, O and Ni, respectively. A composite map summarizing these results is shown in (c). Chemical profiles extracted from these maps and obtained by integrating the signal in a vertical box shown in inset are presented in Figure 3. These profiles reveal a chemical width of the interface of about 0.6 nm when integrated over the thickness of the sample. The chemical interface is therefore larger that the apparently more abrupt structural interface contrast observed on the HAADF images. A sizeable intermixing is therefore clearly detected at the interface even if one accounts for the inelastic delocalization of the EELS signal caused by the long-ranged Coulomb interaction, estimated between 0.1 and 0.2 nm for edges of a few hundreds of eV, [START_REF] Egerton | Electron Energy Loss Spectroscopy in the Electron Microscope[END_REF] and which leads inevitably to a broadening of the profiles. Very interestingly, a slight difference of 0.4 nm (i.e. 0.2 nm per interface) is observed between the global widths of the Ti and the O profiles. This is indicative of an increase of the Ti/O ratio at the interface with respect to the STO matrix and thus reveals a Ti-rich, or O-deficient, interface. Note that this effect is observed all along the interface and therefore, cannot be attributed to a geometrical tilt of the nanowire along a particular direction. As shown in Figure S1, this effect is indeed also visible on the To go beyond this elemental analysis, spectrum images recorded with a higher energy resolution (∼160 meV) have been acquired at the Ni-L2,3, Ti-L2,3 and O-K edges to monitor the spatial variation of their fine structures at the nanowire/matrix interface. The results are summarized in Figure 4, where the different regions of the spectrum images shown in (a) have been used to extract the spectra displayed in (b), (c) and (d). Beginning at the bottom of Figure 4(c), (d), we first present reference edges (black solid lines) of bulk STO recorded in the matrix. The Ti-L2,3 edges are characteristic of Ti 4+ (3d 0 ) in octahedral environment, [START_REF] De Groot | [END_REF] where the 2p hole spin-orbit coupling splits the edge into a lower-lying L3 and a higher-lying L2 components. These two edges are further split into two t2g and eg components as a consequence of the octahedral crystalfield acting on the Ti-3d unoccupied states. This crystal-field splitting is echoed in peaks a and b at the onset of the O-K edge, [33][34][35] which reflects the O-p component of the t2g and eg hybridized states dominating the bottom of the conduction band of STO. Diffuse O-p orbitals further hybridize with the high-lying Sr-4d states leading to the formation of the double structure c at about 7-8 eV above the edge onset.

The first series of spectra extracted from region A, on a TiO2 atomic plane located at one matrix unit cell (∼ 3.9 Å) from the interface, still displays all the characteristic features of the bulk electronic structure of STO without any detectable Ni. On the contrary, spectra from region B, located on the last atomic plane which can still be considered as belonging to the matrix from the HAADF contrast, are largely modified. The symmetry breaking induced by the presence of the interface as well as the pronounced covalent character of the bonding with neighboring metal atoms contrasting with the predominantly ionic bonding in STO, lead to a strong broadening of theTi-L2,3 edges fine structure where the t2g/eg splitting is barely resolved. The fine structure of the O-K edge still exhibits the signature of an hybridization with Ti-3d states through the presence of a sharp peak a and a broader structure b. Interestingly, the missing SrO plane on the left side of the interface leads to a considerable loss of spectral weight in the region of the double structure c, primarily associated with the Sr-4d empty states. Note that a weak Ni signal is already detected in this region, consistently with the elemental profiles shown in Figure 3. As already discussed, although this signal might be partly related to the delocalization inherent to EELS, the presence of Ni in the last atomic planes of the matrix cannot be ruled out.

Note that while the modification of the Ti-L2,3 edges fine structure is particularly abrupt along this [100]STO||[100]Ni interface, it occurs in a more gradual manner for the [110]STO|| [110]Ni facets where the Ti environment in the last atomic plane of the matrix is qualitatively closer to the bulk. Figure S2 presents a comparison of the spectra obtained for these two types of facets in the vicinity of the interface with the nanowire, and shows that in the latter case, the crystalfield splitting of the Ti-3d states is still clearly visible whereas it almost disappeared in the former.

Region C corresponds to the first atomic plane, which can be considered structurally as a part of the nanowire from the HAADF contrast. Consequently, a large increase of the Ni signal is observed with respect to region B. A comparison with a reference spectrum from bulk Ni [36,37] shown in Figure S3(a) reveals that, despite the relatively low signal-to-noise ratio of these spatially resolved data, a substantial broadening of the L3 edge is clearly observed. This effect might be attributed to fluctuations of Ni oxidation state and local environment within the thickness of the specimen both in terms of geometry due to the presence of the metal/oxide interface and to the local disorder associated with the diffusion of Ti and O in the first atomic planes of the nanowire. The Ti-L2,3 edges in this area is also further broadened to reach a shape which closely resembles that of the Ti-L2,3edges in bulk Ti or NiTi alloy. [36] A comparison between this Ti-L2,3 and a reference spectrum [37] recorded in bulk Ti is presented in Figure S3(b). The O-K edge also displays broader structures when compared to region B, however retains its global shape without displaying any particular new features.

In region D, the Ni-L2,3 edges are essentially that of bulk Ni (see Figure S3(a)), while very weak Ti and O signals are still detectable. Finally, regions E and F are characteristic of pure bulk Ni. The Ni-L3 to L2 ratios extracted from these data are presented in Figure 4(e), the procedure employed to obtain these values is detailed in Figure S4. Although the reduced signal-to-noise ratio inherent to these spatially resolved data limits the accuracy of the results, no substantial deviation from the bulk Ni value is observed in this series of spectra.

In summary, no reaction phase could be detected at the metal/oxide interface. While structurally abrupt, spectroscopic investigations at the atomic-scale indicate however the presence of a weak interdiffusion occurring over a few atomic planes only. This leads essentially to a metallization of the last TiOx (x< 2) atomic plane of the matrix and to the formation of a few atomic planes of a non-stoichiometric mixed Ni-Ti oxide on the external part of the nanowire. Despite the presence of detectable O in these planes, the metal L2,3 edges closely resemble those observed in bulk Ti or Ti intermetallics. This is consistent with the O-K edge fine structure at the interface, which does not display any characteristic features typically found in NiO. [38] Unlike in the case of a previous study carried out on a Ni thin film epitaxially grown on a STO(001) surface, [39] no evidence of the presence of a stoichiometric NiO layer could be found at the interface here. The results rather evidence the formation of metallic Ni-Ti bonds at the interface between the nanowire and the matrix. In previous work, a phenomenon was reported whereby some structural disorder could be detected within wires with diameters reduced below 2 nm. [26] In light of the present results, we suggest that the interdiffusion of Ti at the Ni:STO interface could contribute to this disordering process in ultrathin nanowires.

While the EELS results presented so far provide clear evidence for the preservation of the metallic character of the embedded Ni nanostructures, the analysis is limited to a very small number of nanowires, due to the intrinsic working principle of STEM-EELS as a highly local probe. To exclude any bias, Ni-STO VANs were thus further analyzed with x-ray absorption spectroscopy at the Ni-L2,3 edges, providing insight into the nature of chemical bonding of Ni atoms on a global, sample-wide scale. Figure 5(a) shows the x-ray absorption data of a Ni-STO VAN containing nanowires with mean diameter d=5 nm. The spectrum is very similar to those of earlier studies of bulk Ni, [42,43] with an absence of spectral features related to nickel oxide formation. Note for example the presence of a broad satellite peak located 6 eV above the L3 maximum, which has been reported in other studies on metallic Ni. [42] The calculated Ni-L3 to L2 ratio (ratio of the integrated intensities) is 3.1 and matches, within error bars, the one determined previously by EELS measurements of the VAN with 2.4 nm diameter wires. It is also in good agreement with values reported in the literature. [43] This spectrum thus confirms the EELS data and, considering the surface sensitivity of the total electron yield (TEY) mode, also shows that an ultrathin AlOx capping layer can be efficiently used to prevent oxidation at the apex of the wires.

The results show that NiO is efficiently reduced at the TiO2-terminated interface with STO, leading to the formation of metallic Ni during the growth under vacuum. This can be attributed to the higher reactivity of the TiO2-terminated surface of STO with respect to the more stable SrO-terminated surface. [44] This is a consequence of the fact that Sr levels are located deeper below the Fermi level than Ti and O levels that form the valence and conduction bands of STO. [45] This leads to a less reactive SrO surface with lower surface energy. In the case of Pt growth on SrTiO3(001), the SrO-terminated surface has been shown to give rise to a weaker bonding with Pt than the TiO2-terminated surface. [46] As NiO reduction at TiO2-terminated interface with STO is a key step in the formation of the VAN, growth has to be performed under vacuum in order to favor this reaction. Under such conditions, we anticipate that using a mixed We finally close this article by highlighting the importance of Ni metallicity conservation and interface sharpness with regard to strain engineered magnetic properties in VANs. As put forward in the introduction, nanocolumnar composites are an ideal test bed to scrutinize magneto-elastic effects, due to their huge heterointerface. This especially holds true for transition-metal/oxide VANs, where the pronounced lattice mismatch between the oxide and metal species can induce large axial nanowire strains εzz, as recently demonstrated for Ni:STO. [26] In this study, a decreasing linear relationship between εzz and the nanowire radius r was evidenced. Thus, in principle, the reduction of r provides a simple strategy to achieve large magneto-elastic effects. However, it also results in an increase of the relative amount of interfacial Ni atoms, which can have a negative impact on the magnetic properties of the metallic phase, via oxidation or compound formation. [47] Similarly, partial interfacial mixing between lattice mismatched phases might induce lattice parameter accommodation, a detrimental effect, considering the quest for maximum strain. Efficient strain engineering in metal/oxide composites thus requires growth strategies that preserve the metallicity of the embedded wires and provide sharp metal/oxide interfaces. Both criteria are met in selfassembled Ni:STO columnar composites, and large values of the axial strain can be reached, as a consequence of the vertical epitaxy. The mean axial (εzz) and radial strain (εrr) within the 5nm-diameter Ni nanowires were determined by x-ray diffraction measurements: εzz=2.4% and εrr=-0.3 %. This very large tensile strain along the wires axis should give rise to strong magnetoelastic effects, as will be examined in the remainder of the paper. define θ as the angle between the applied field and the axis of the wires. Figure 5(c) shows magnetic hysteresis loops acquired with the field applied along the nanowires axis (θ=0°) and perpendicular to this axis (θ=90°). Comparison of the cycles reveals that the remanence is larger for θ=90° than for θ=0°. Furthermore, the saturation field is weaker when the field is applied perpendicular to the wires axis. This indicates that the axis of the wires is not the magnetic easy axis. This is indeed confirmed by the results displayed in Figure 5(d), where the variation of the remanent magnetization as a function of θ is given. The remanent magnetization is a monotonous function of θ with a minimum value for θ=0. This indicates that the axis of the nanowires is a hard magnetic axis.

To analyze the magnetic properties of the present composite, two different anisotropy contributions have to be taken into consideration. First, there is a shape-induced term resulting from the high aspect ratio of the wires. This term is usually dominating in conventional magnetic nanowire arrays, typically obtained via electrodeposition into cylindrical nanopores, giving rise to an easy axis magnetic direction along the wire axis. [48] In the present system, the magnitude of this uniaxial shape anisotropy contribution is given by:

𝐾𝐾 𝑚𝑚𝑚𝑚 = 𝜇𝜇 0 4 𝑀𝑀 𝑚𝑚 2 (1 -3𝑃𝑃)
where Ms is the saturation magnetization and P is the porosity (volume fraction occupied by the wires) of the NWs assembly. [49] Using a typical value of the porosity corresponding to samples studied in this work, P=10%, gives Kms=5.3 10 4 J.m -3 . In addition, a strain induced anisotropy contribution has to be considered. Bulk Ni has a negative magnetostriction coefficient λ001. [50] In the case of uniaxial strain εzz along [001], the uniaxial magnetoelastic contribution is given by: [51] 𝐾𝐾 𝑚𝑚𝑚𝑚 = 3 2 𝜆𝜆 001 (𝑐𝑐 11 -𝑐𝑐 12 )(𝜀𝜀 𝑧𝑧𝑧𝑧 -𝜀𝜀 𝑟𝑟𝑟𝑟 )

where cij are elastic constants. [52] Neglecting εrr, a reasonable approximation in the present case where the axial strain is much larger than the radial strain, the magnetoelastic contribution becomes:

𝐾𝐾 𝑚𝑚𝑚𝑚 ≅ 3 2 𝜆𝜆 001 (𝑐𝑐 11 -𝑐𝑐 12 )𝜀𝜀 𝑧𝑧𝑧𝑧
Because λ001 is negative in Ni, the total uniaxial anisotropy, K=Kms+Kme, is the sum of two terms with opposite signs. Kms is positive and favors a magnetic easy axis along the wires axis.

Kme is negative and favors a magnetic hard axis along the wires axis. The two terms compete and the wires axis becomes a hard axis if K<0, i.e. if Kms<|Kme|. This corresponds to the situation where εzz exceeds the critical value defined by:

𝜀𝜀 𝑐𝑐𝑟𝑟𝑐𝑐𝑐𝑐 = � 𝜇𝜇 0 𝑀𝑀 𝑚𝑚 2 (1 -3𝑃𝑃) 6𝜆𝜆 001 (𝑐𝑐 11 -𝑐𝑐 12 ) � ≅ 0.6%
Since εzz exceeds εcrit, the system exhibits a hard axis oriented along the wire axis, as clearly demonstrated by the results in Figure 5(c-d). Thus, these results confirm that the total uniaxial anisotropy has changed sign compared to the unstrained case. This provides clear experimental evidence for strong magneto-structural coupling in the Ni:STO VAN.

Conclusion

The structure and chemical composition of the vertical interface in a Ni:SrTiO3 VAN VANs the axis of the Ni wires is the hard magnetic axis due to strong magneto-elastic effects.

This establishes Ni:SrTiO3 as a model system for the study of strain-engineering of the magnetic anisotropy in VANs.

Experimental Section

Sample growth: Samples were grown by pulsed laser deposition (PLD) under high vacuum (base pressure 2.10 -6 mbar) using NiO and SrTiO3 (STO) targets and a quadrupled Nd:YAG laser (wavelength 266 nm) operating at 10 Hz with fluence in the 1-3 mJ.cm -2 range. Growth was performed on SrTiO3(001) single crystalline substrates held at 650°C. A sequential deposition scheme was adopted, as described in previous works. [20,26] NiO and STO targets are ablated sequentially, resulting in a succession of submonolayer deposition step. This leads to the formation of the vertically aligned Ni nanowires embedded in the epitaxial SrTiO3/SrTiO3(001) matrix. The mean diameter of the nanowires can be tuned by adjusting the ratio of laser shots on NiO and STO targets during a sequence. [26] After completion of the sequential deposition steps, the samples were capped in-situ with an ultrathin AlOx layer to Sample analysis: Samples were thinned using a two-step method. The back side of the sample was first mechanically thinned until the thickness becomes less than a few tens of microns.

Then, Ar-ion milling was used in order to reach the electron transparency. HAADF STEM images and STEM-EELS spectrum-images were obtained using a Nion Ultra-STEM MC HERMES 100 microscope operating at 60 kV, equipped with a C5 Nion QO probe corrector, a monochromator [40] and a Gatan Enfinium spectrometer. EELS elemental maps are shown after principal component analysis [41] of a spectrum-image with dimensions 68×68×2048, acquired with an energy resolution of 350 meV, a beam current of 120pA, a dwell time of 0.2s/pixel, and a dispersion of 0. X-ray absorption spectroscopy and x-ray magnetic circular dichroism experiments were carried out on the DEIMOS beamline of SOLEIL synchrotron. [53][54][55] Data were acquired in total electron yield mode. The samples were kept under ultrahigh vacuum at 4.2 K during measurements. The applied magnetic field was collinear to the x-ray propagation vector. The incidence on the sample was set through a rotation of the sample by an angle θ. θ=0° refers to normal incidence. XAS and XMCD spectra were acquired at the Ni-L2,3 edges by scanning the photon energy in the 840-880 eV range. Element-specific hysteresis loops were acquired by scanning the magnetic field at constant x-ray energy, set to the maximum of the XMCD signal, for left and right circular polarization. For the acquisition of XMCD spectra at remanence, the following scheme was used: the sample is submitted to a magnetic field exceeding the saturation field in the positive direction, the field is then set to zero and absorption spectra are recorded with left and right circular polarizations. The same procedure is then applied with the saturating field in the opposite direction. The XMCD spectra at remanence are then obtained by properly averaging data recorded for both direction of application of the field. Magnetic hysteresis loops were also acquired using a vibrating sample magnetometer (VSM).

Supporting Information

Additional data extracted from EELS spectrum-images are provided. Figure S1 shows the chemical profiles extracted through the nanowire along a direction oriented at 45° with respect to the one used in Figure 3. Figure S3 compares the Ti-L2,3 and Ni-L2,3 edges fine structure obtained in the first atomic plane of the nanowire with bulk references obtained respectively in metallic Ti and Ni. Figure S4 illustrates the method employed to extract the Ni-L3 to L2 ratio from atomically resolved EELS data.

Supporting Information is available from the Wiley Online Library or from the author. 3) and reference spectra obtained on bulk Ni and bulk Ti, respectively. The points correspond to the experimental data while the lines are just guides for eye.

[

  110]STO||[110]Ni facets oriented at 45°.

  NiO-STO target or a Ni metallic target instead of NiO would lead to Ni:STO VAN formation as well. Other possible pairs of transition metal/oxide matrix, such as Co and CoNi alloys in SrTiO3 and in BaTiO3, may lead to VAN formation through the same mechanisms for which reduction at TiO2 interface is crucial.

Figure 5 (

 5 Figure 5(b) shows magnetic hysteresis loops acquired by measuring x-ray magnetic circular

  could be determined by STEM and EELS measurements with atomic resolution. The obtained results demonstrate the absence of NiO and the metallic character of the last external atomic planes of the nanowires embedded in SrTiO3. The full preservation of Ni metallicity is further confirmed by x-ray absorption spectroscopy, and magnetic measurements indicate that in Ni:SrTiO3

  prevent surface oxidation upon exposure to air. The structural properties of the samples were characterized by x-ray diffraction and transmission electron microscopy, in order to determine the epitaxial relationships and dimensions of the nanowires. Ni is found to grow epitaxially within the STO matrix with (001)Ni||(001)STO and [100]Ni||[100]STO. In the samples studied in the present work, the Ni nanowires have mean diameters of 2.4 nm and 5 nm.

  37 eV/ch. Qualitative elemental profiles were obtained by integrating EELS intensity maps across the nanowire. Energy-loss near-edge structure spectrum-images were acquired with an energy resolution of 160 meV, a beam current of 55 pA, a dispersion of 0.05 eV/ch, and dwell times of 0.20 s/pixel for the Ti-L2,3 edges, 0.35 s/pixel for the O-K edge and 0.4 s/pixel for the Ni-L2,3 edges. High signal to noise in images of Figure 2(b), Figure 4(a), S1(inset), and S2(a) were obtained after alignment of a stack of 42 images, using the built-in stack-alignment plugin in Gatan Digital Micrograph software. The convergence angle was ≈ 32 mrad, the HAADF collection angular range ≈ 85-195 mrad, and the EELS collection semi angles ≈ 44 mrad.

  Figure S2 shows a comparison of the Ti-L2,3 and O-K edges fine structure obtained at the [100]STO||[100]Ni and the [110]STO||[110]Ni nanowire-matrix interfaces.

Figure 2 .

 2 Figure 2. (a) Global HAADF top view, along [001] zone axis, of the sample showing a few Ni nanowires in the STO matrix. The STO and Ni lattices appear superimposed for short nanowires, as indicated by arrows, in which case the top or bottom surface of the thinned sample does not intersect the nanowires. Detailed analysis of a typical nanowire of diameter 2 nm: (b) highresolution HAADF image obtained after alignment of a stack of 42 images, showing the atomic structure of the nanowire; (c) composite chemical map of the nanowire built from the elemental Ti map shown in (d), O map in (e) and Ni map in (f) extracted from the Ti-L2,3 edges (∼457eV), the O-K edge (∼531eV) and the Ni-L2,3 edges (∼855eV) respectively. All maps are displayed after noise removal through principal component analysis (PCA). Variations of the intensity with the STO lattice periodicity is visible on the Ti and O maps but not in the Ni nanowire due to the smaller lattice spacing in the metal.

Figure 3 .

 3 Figure 3. Ni, Ti and O chemical profiles across the nanowire extracted from elemental maps, obtained by integrating EELS signal in the box shown in inset. Intensities are normalized to the same maximum height. According to the Ni profile, the projected diameter of the nanowire is close to 2nm. The slightly wider O profile compared to Ti indicate a Ti-rich interface. The points correspond to the experimental data while the lines are just guides for eye.

Figure 4 .

 4 Figure 4. Atomic and electronic structure of the nanowire-matrix [100]STO||[100]Ni interface. (a) High-resolution HAADF image of the interface between the Ni nanowire and the SrTiO3 matrix, boxes indicate the spectrum-image integration areas used to obtain the spectra shown in (b) for the Ni-L2,3 edges, (c) for the Ti-L2,3 edges and (d) for the O-K edge. (e) Ni-L3 to L2 ratio obtained from the spectra displayed in (b). The points correspond to the experimental data while the lines are just guides for eye.

Figure 5 .

 5 Figure 5. X-ray absorption spectroscopy and magnetic measurements of Ni nanowires assembly in STO, measured at the Ni-L2,3 edges. (a) X-ray absorption spectrum. (b) Magnetic hysteresis loops acquired with the magnetic field applied along the Ni nanowires axis as measured by XMCD at the Ni-L3 edge (symbols) and using a vibrating sample magnetometer (line). (c) Magnetic hysteresis loops acquired with the magnetic field applied along the Ni nanowires axis (thin line) and perpendicular to the axis (thick line) using a vibrating sample magnetometer. (d) Triangle: XMCD at the Ni-L3 edge, measured at remanence, as a function of the angle θ between the nanowires axis and the applied magnetic field (θ=0 corresponds to a field applied along the axis of the nanowires). Disk: value of the normalized remanence for θ=90° deduced from vibrating sample magnetometer measurements.

Figure S2 .

 S2 Figure S2. Ti-L2,3 and O-K ELNES differences between [100]STO||[100]Ni and the [110]STO||[110]Ni nanowire-matrix interfaces. While the crystal-field splitting is almost completely lost in B due to the symmetry breaking at the interface and the covalent interaction with neighboring Ni atoms, it is still partially visible in G where the Ti local environment is slightly less affected by the interface. The O-K edge still shows the fingerprint of the hybridization with the Ti-3d states but the spectral weight of the higher lying Sr-4d states is strongly reduced in B and G compared to bulk STO. The points correspond to the experimental data while the lines are just guides for eye.

Figure S3 .

 S3 Figure S3. (a) Comparison between (a) the Ni-L2,3 and (b) the Ti-L2,3 edges obtained at the Ni-STO matrix interface (area C in Figure3) and reference spectra obtained on bulk Ni and bulk Ti, respectively. The points correspond to the experimental data while the lines are just guides for eye.

Copyright WILEY-VCH Verlag GmbH & Co. KGaA, 69469 Weinheim, Germany, 2018.

Acknowledgements

Scanning Transmission Electron Microscopy (STEM) experiments were carried out at Super-STEM, the EPSRC National Research Facility for Advanced Electron Microscopy. M. Bugnet and G. Radtke are grateful to SuperSTEM for hospitality and access to the microscopes. We acknowledge SOLEIL for provision of synchrotron radiation facilities and thank the staff of the MPBT platform of Sorbonne Université for their support. The authors also gratefully thank F.

Breton for his assistance and design of the PLD control system. This work was supported by French state funds managed by the ANR within the Investissements d'Avenir programme under reference ANR-11-IDEX-0004-02, and more specifically within the framework of the Cluster of Excellence MATISSE led by Sorbonne Universités. M. Bugnet acknowledges funding from INSA Lyon through the BQR fund THERMOS. M. Hennes acknowledges financial support from the French Embassy in Berlin (Service pour la Science et la Technologie) and Campus France. X. Weng gratefully acknowledges financial support of the Chinese Government Scholarship from China Scholarship Council (CSC No. 201608070062). This work has benefited from a funding from LabEx PALM (ANR-10-LABX-0039-PALM).

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