Effect of target nickel stoichiometry on the synthesis of infinite-layer nickelates

Abstract

The discovery of superconductivity in infinite-layer nickelate films marks a groundbreaking addition to the family of unconventional superconductors, providing new insights into mechanism of unconventional high temperature superconductivity. However, synthesizing these superconducting nickelates presents significant challenges: they cannot be grown directly and instead require a two-step synthesis protocol involving initial deposition of a perovskite precursor phase (e.g., Nd_0.8Sr_0.2NiO₃) followed by topotactic reduction to the infinite-layer structure (Nd_0.8Sr_0.2NiO₂). This process is further complicated by the extreme sensitivity of both steps to synthesis conditions, necessitating stringent control over the crystallinity and stoichiometry of the parent phase. In this study, we uncover nickel deficiency during pulsed laser deposition (PLD) of the parent-phase Nd_0.8Sr_0.2NiO₃. By incorporating 15% excess nickel into the PLD target, we mitigate this loss, suppress secondary phase formation in the Nd_0.8Sr_0.2NiO₃ parent film, and ultimately obtain a phase-pure Nd_0.8Sr_0.2NiO₂ film exhibiting superconductivity after following reduction. Notably, we observe a doping-dependent insulator-to-superconductor transition in films synthesized from targets with varying nickel content after reduction. X-ray photoelectron spectroscopy (XPS) confirms that the Nd/Ni ratio in films derived from nickel-over-doped targets (15% excess) aligns closely with the ideal stoichiometry. These findings underscore the indispensable role of stoichiometric precision in stabilizing infinite-layer nickelates and establish a practical synthesis strategy for optimizing their superconducting performance.

Graphical abstract

Keywords

target / pulsed laser deposition / nickel deficiency / infinite-layer nickelates / superconductivity

1 Introduction

Since the discovery of copper-based high-temperature superconductors [1], researchers have persistently sought to find alternative complex oxide superconducting systems. Nickelates with Ni⁺ oxidation state have been regarded as particularly promising candidates owing to their electronic configuration closely resembling that of copper. Although extensive efforts were made decades ago to explore superconductivity in nickelate systems [2−5], no conclusive evidence was found. Until in 2019, the breakthrough discovery of superconductivity in an infinite-layer Nd_0.8Sr_0.2NiO₂ thin films reignited extensive exploration of nickel-based superconductors [6]. Subsequent (La/Pr)_1−x(Eu/Sr/Ca)_xNiO₂ [7−11] infinite-layer superconducting films have been discovered, and achieved higher T_c values successively [12]. Furthermore, superconductivity has also been demonstrated in Ruddlesden−Popper (RP) nickelates, including reduced Nd₆Ni₅O₁₂ films [13] and high-pressure La₃Ni₂O₇ [14], La₄Ni₃O₁₀ [15], La₂PrNi₂O₇ [16] crystals. And recently, superconductivity surpassing the McMillan limit (40 K) has been observed at ambient pressure in La₃Ni₂O₇ [17] and La₂PrNi₂O₇ [18] thin films constrained by a substrate. These findings establish nickelates as a new generation of unconventional high-temperature superconductors, significantly advancing the exploration of unconventional superconductivity mechanisms.

Among nickelate superconductors, infinite-layer nickelates exhibit many similar properties to cuprates and have attracted much attention [19−26]. However, due to the inability to directly obtain superconducting infinite-layer nickelate thin films through growth and the necessity of topotactic hydrogenation reduction of high-quality thin film precursors, the synthesis of superconducting infinite-layer nickelates poses a significant challenge. For the Sr-doped perovskite precursor Nd_1−xSr_xNiO₃, the Sr doping makes it more prone to generate some secondary phases during the growth process, greatly affecting the quality of the obtained precursor [27]. Therefore, compared to the undoped NdNiO₃, the growth conditions of Nd_0.8Sr_0.2NiO₃ are more demanding. Additionally, stoichiometry is a crucial parameter during thin film growth. Just like in NdNiO₃ thin films, the lack of nickel and oxygen will directly impact the metal-insulator transition temperature [28, 29]. After Sr doping, Nd_1−xSr_xNiO₃ thin films become more sensitive to this parameter. Even after a thorough reduction process, off-stoichiometric precursor thin films are difficult to generate superconducting properties, and may not even lead to an infinite-layer structure [30−32]. During the thin film deposition process, the stoichiometry of the target directly impacts the final stoichiometry of the thin films. If there is an element loss during the growth process, the elements provided by the target synthesized according to the molar ratio will no longer match the optimal stoichiometry, leading to the formation of nickel-deficient RP phase defects, accompanied by the enrichment of NiO_x particles on the surface [33, 34].

Here, we compensated for the nickel stoichiometric loss during the growth process by doping the target material with varying additional nickel content. We demonstrate that when the excess nickel compensation amount reached 15%, the secondary phase content in the thin film precursor significantly decreased, and after reduction, a well-formed infinite-layer phase was achieved, exhibiting superconductivity. XPS quantitative analysis indicated that with the increase in nickel compensation content, the relative content of the grown precursor thin film approached closer to the ideal value. Subsequent XAS tests showed that the excess nickel compensation did not have a significant impact on the electronic structure of nickel in the thin film. This strategy can more effectively avoid the off-stoichiometry of the parent thin films during the growth process and increase the success rate of achieving infinite-layer superconducting thin films.

2 Experiment section

2.1 Preparation of targets and thin films

Nd₂O₃, NiO, and SrCO₃ powders were mixed according to a stoichiometric ratio of Nd:Sr:Ni = 0.8:0.2:x (x = 1, 1.05, 1.1, 1.15) and ball-milled for 24 hours to ensure homogeneity. The mixture was first subjected to high-temperature sintering at 1200 °C for 12 hours to remove carbon, followed by another sintering step at 1350 °C for 12 hours to allow complete reaction. The resulting well-mixed uniform powder was pressed into cylindrical target form and sintered again at 1350 °C to obtain high-density Nd_0.8Sr_0.2Ni_xO₃ targets for laser ablation.

Nd_0.8Sr_0.2Ni_xO₃ thin films of approximately 12 nm (30 unit cells) were grown on SrTiO₃ substrates by pulsed laser deposition (PLD, KrF, λ = 248 nm), and no extra layer was capped. The growth process was monitored in situ by reflection high-energy electron diffraction (RHEED). The SrTiO₃ substrates were cleaned, etched by deionized water and HF, and then annealed at 950 °C for 1.5 hours to form a sharp TiO₂-terminated step surface. During film growth, the substrate temperature was maintained at 600 °C, with a laser energy density of 2.5 J/cm² (spot size is about 1.5 mm²), a repetition rate of 4 Hz, and an oxygen pressure of 0.1 mbar.

The as-grown films were cut into cut into four pieces with a size of 2.5 mm × 2.5 mm × 0.5 mm and then wrapped aluminum foil and placed in a sapphire crucible with 1g of CaH₂, avoiding direct contact. Then the crucible is placed in a vacuum chamber with a chamber pressure of approximately 10⁻³ mbar, and subsequently heated at 300−320 ℃ for 4−6 hours, the heating and cooling rates were 10 ℃/min. The adequacy of the reduction process is determined by the peak position of the film in the XRD spectrum, and a continuous reduction process is used to ensure that the film is fully reduced.

2.2 Composition analysis and structural characterization

The surface morphologies were acquired using an atomic force microscope (AFM) in contact mode. The crystal structures of the targets and thin films were measured by X-ray diffraction (XRD, Cu-Kα). The relative proportions of different elements in the target materials were semi-quantitative using scanning electron microscopy-energy dispersive spectroscopy (SEM-EDS).

The relative proportions of cationic elements in the thin films were quantitatively analyzed using X-ray photoelectron spectroscopy (XPS), with the XPS spectra obtained using Al-Kα radiation. Peak fitting was performed on the XPS spectra for Nd 3d, Ni 2p, and Sr 3d, using the Shirley method to subtract the background. The peak area ratios were controlled to be 3d_5/2:3d_3/2 = 3:2 and 2p_3/2:2p_1/2 = 2:1, the half-height width of each group of split peaks was restricted to the same value, the ratio of Lorentz and Gaussian functions in the fitting function was 80:20.

The main peak positions for Nd 3d were maintained at 1004 eV (Nd 3d_3/2) and 981.7 eV (Nd 3d_5/2), which can be attributed to Nd³⁺ final state [50]. The Ni 2p main peaks were divided into Ni²⁺ and Ni³⁺ components, with binding energy of 873.69 eV (Ni 2p_1/2, Ni³⁺), 872.09 eV (Ni 2p_1/2, Ni²⁺), 856.17 eV (Ni 2p_3/2, Ni³⁺), and 854.5 eV (Ni 2p_3/2, Ni²⁺) [58]. The spin−orbit splitting energies for Nd 3d_5/2-Nd3d_3/2 and Ni 2p_3/2-2p_1/2 were \mathrm{\Delta }{E}_{\mathrm{s}\mathrm{o}}\left(\mathrm{N}\mathrm{d}\right)=22.3\mathrm{e}\mathrm{V}\pm 0.1\mathrm{e}\mathrm{V} and \mathrm{\Delta }{E}_{\mathrm{s}\mathrm{o}}\left(\mathrm{N}\mathrm{i}\right)=17.6\mathrm{e}\mathrm{V}\pm 0.1\mathrm{e}\mathrm{V}, respectively. The Sr 3d spectrum was attributed to signals from both the surface and the bulk, with binding energies of 135.2 eV (Sr 3d_3/2, surface), 133.57 eV (Sr 3d_5/2, surface), 133.2 eV (Sr 3d_3/2, lattice) and 131.61 eV (Sr 3d_3/2, lattice) [59].

For the satellite peaks in the Nd 3d and Ni 2p spectra, the energy differences and area ratios relative to the main peaks were consistent. In the Nd 3d_3/2 spectrum, satellite peaks were located at 1008.73 eV, 999.68 eV, and 993.83 eV. For Nd 3d_5/2, satellite peaks appeared at 981.7 eV and 971.75 eV. Additionally, satellite peaks were fitted at 879.61 eV for Ni 2p_1/2 and 861.5 eV for Ni 2p_3/2.

The element mole ratio is caculated by the ratio of the main peak area to the sensitivity factor, with the sensitivity factors referenced from the values compiled by Wagner et al. [60].

2.3 Transport measurement and X-ray absorption spectra

The transport behaviors were measured contacts using a Quantum Design Physical Property Measurement System (PPMS-9T) in a van-der-Pauw geometry with Al wire bonded. X-ray absorption spectroscopy (XAS) were carried out at the MCD-A and MCD-B beamlines (Soochow Beamline for Energy Materials) at the Hefei National Synchrotron Radiation Laboratory (NSRL). The Ni L_2,3-edge and O-K-edge were measured along the direction of normal incidence in the Total Electron Yield (TEY) mode.

3 Results and discussion

Firstly, a mixed-phase target with a nominal stoichiometry Nd:Sr:Ni = 0.8:0.2:1 was used for the growth of Nd_0.8Sr_0.2NiO₃ and optimised using high-fluence growth conditions [27]. While this approach successfully yields superconducting thin films, it is hampered by low reproducibility and challenges in controlling nickel kinetics during growth, primarily due to the substantial entropy introduced by Sr doping. Fig.1(a) illustrates the two possible nickel dynamics of elemental Ni during pulsed laser deposition of Nd_0.8Sr_0.2NiO₃ thin films. (i) Since the target synthesised using the solid-phase reaction is a mixed-phase target, some of the nickel paticles generated during the laser sputtering process may adhere to the surface of this mixed-phase target, leading to a lack of nickel stoichiometry in the film over a long period of time during deposition [28]. This off-stoichiometry problem can be mitigated by using a single-phase target or polishing the target. However, the success rate of synthesising superconducting infinite-layer nickelate films remains low, despite the fact that we have polished the target after growth of each film. (ii) Upon deposition of nickelate oxides onto a heated SrTiO₃ substrate, some nickel oxide particles may diffuse on the surface of the films. The diffusion can be visualized through the surface morphology of as-grown film characterized by atomic force microscopy (AFM), as shown in Fig.1(b) and (c). As can be seen from the surface morphology, the films have relatively flat atomic-level steps, which proves that our films are grown with good quality. However, there are a large number of 5 nm-high particles diffusely distributed on the surface, which should be the NiO_x particles precipitated during the high temperature process of growth [34, 35]. The cumulative effect of nickel loss through these two mechanisms significantly impacts the cation stoichiometry of the parent Nd_0.8Sr_0.2NiO₃ films epitaxially grown on the substrate, consequently reducing the reproducibility of superconducting films.

To compensate for the lack of cation stoichiometry due to nickel dynamics during laser pulse deposition growth, we implemented a strategy of controlled nickel overdoping in the target. Specifically, we prepared four mixed-phase targets with varying stoichiometric ratios, designated as Ni-1.0, Ni-1.05, Ni-1.1, and Ni-1.15, corresponding to nickel excesses of 0%, 5%, 10%, and 15%, respectively, for the synthesis of parent films. Fig.2(a) shows the polycrystalline X-ray diffraction (XRD) patterns of these four targets. Under different nickel compensation conditions, the target composition is a mixed phase of tetragonal Nd_1.6Sr_0.4NiO₄ (JCPDS 80-2324) and cubic NiO (JCPDS 73-1523), which is consistent with the structure of the target prepared by the same method before [24]. This shows that the structure of the target powders is not affected by the doping of excessive nickel. The nickel stoichiometry in the target powder was further quantitatively assessed by energy dispersive spectrometer (EDS) (see Supplementary Information, Fig. S2), and the actual versus nominal values of Ni/(Nd+Sr) are shown in Fig.2(b). For the target, the stoichiometry is essentially the same as the content prior to the blending reaction, indicating that there is no loss of nickel or other cations during the solid-phase reaction of the target.

The structure of as-grown Nd_0.8Sr_0.2Ni_xO₃ films were characterized using X-ray diffraction (using Cu-K_α radiation), as shown in Fig.3(a). A systematic shift of the (001) and (002) 2θ peaks to higher angles is observed with increasing nickel content in the precursor target, reflecting a corresponding reduction in the c-axis lattice constants. Notably, when the nickel stoichiometry exceeds 1.05 (Ni ≥ 1.05), the (002) peak position shifts beyond 48°, a critical threshold of parent films for achieving superconducting properties after subsequent reduction [27]. The intensity of the (001) diffraction peak serves as an additional indicator of film quality. Films with excessive secondary phases exhibit a significantly weakened (001) peak, whereas those with a pure perovskite phase display a pronounced (001) peak intensity. For the aforementioned parent phase films, the (001) peaks are all clearly visible. However, for Ni-1.0 sample, the (002) diffraction peak value is 47.78° (c ≈ 3.805 Å), and the (001) diffraction peak nearly vanishes, indicating the formation of secondary phase Nd₄Ni₃O₁₀ [27, 36, 37]. With the gradual increase in nickel compensation content, the parent films exhibit a smaller c-axis lattice parameter and a more pronounced (001) peak. These observations underscore the importance of precise nickel stoichiometry control in achieving the desired structural properties in these materials.

Fig.3(b) shows the structure of as-reduced Nd_0.8Sr_0.2Ni_xO₂ films, denoted as “R-Ni-x”. Only R-Ni-1.1 and R-Ni-1.15 samples exhibit a well-defined infinite-layer structure, while R-Ni-1.0 and R-Ni-1.05 samples failed to achieve this structural configuration. It is important to note that even though the (002) diffraction peak of the Ni-1.05 sample is 48.15° (higher than the threshold 48°) and high quality (clear Laue oscillations), the peak position of R-Ni-1.05 shifts only rightward to 50.4° (c ≈ 3.619 Å ) after a full reduction process, similar to that of the previously reported LaNiO_2.5 [5]. In contrast, the Ni-1.0 sample, which initially contains a secondary phase structure, shows a downward shift in the 2θ value after reduction, accompanied by a larger out-of-plane lattice parameter (c ≈ 3.842 Å) compared to its parent film. Fig.3 summarizes the impact of varying nickel contents in the target on the lattice parameters of both as-grown and as-reduced films. While excessive nickel stoichiometry in the target facilitates the growth of parent-phase films with reduced c-axis lattice constants, it is only when the nickel compensation content exceeds a critical threshold, (i.e., Ni > 1.1), that high-quality infinite-layer films can be successfully obtained through the reduction process. Fig.3(d) presents the distribution of thin-film synthesis outcomes using nickel stoichiometric targets with varying compositions. The scatter plot of c-axis lattice constants aligns with the trends observed in Fig.3(c). Notably, targets with 15% nickel excess (Ni-1.15) demonstrated substantially higher success rates in achieving the infinite-layer phase compared to stoichiometric (Ni-1.0) targets. The corresponding bar chart quantifies the superconducting film yield. Analysis of comparable sample sizes reveals that 15% nickel doping significantly enhances the reproducibility of superconducting films. For thin film samples prepared from Ni-1.0 stoichiometric target, due to the influence of growth kinetics, most results tend toward nickel-deficient secondary phases, with superconducting films only obtained in a few cases.

The electrical transport behavior of different films obtained with different nickel overdoping was further analyzed, respectively. As shown in Fig.4(a), all as-grown Nd_0.8Sr_0.2Ni_xO₃ films exhibit metallic with few upturns. The temperature-dependent resistivity (ρ(T)) of this metallic state can be fitted using a power law:

(1){\rho =\rho }_0+A{T}^n,

where A represents a coefficient associated with the strength of electron scattering, and n denotes the apparent power-law exponent. For normal metals with weak electron interactions (Fermi liquids), n = 2. However, for strongly correlated systems, the introduction of multiple ordering covariates (e.g., pressure, magnetic field, or doping) suppresses some of this ordering, rendering the transport properties no longer classical Fermi-liquid behavior [38]. This phenomenon is observed in systems like doped copper-based oxides and heavy fermionic systems [39, 40], where n = 1 or other values are not equal to 2. The ρ(T) curve of as-grown Ni-x films were fitted using the above equation, with representative fitting results illustrated in Fig.4(b). Equation (1) provides an excellent fit to the ρ(T) curves for temperatures above 30 K. The fitted parameters n and A are summarized as functions of nickel content (x) in Fig.4(c). All films exhibited a non-Fermi liquid behavior with n = 1.15 − 1.35. In addition, a clear evolution of n and coefficient A with increasing x is observed. As x increases, n progressively decreases and approaches 1, while the coefficient A steadily increases. Similar trends were also observed in NdNiO₃ films, which were attributed to structural changes due to stresses and defects [29].

Fig.4(d) shows the temperature-dependent resistivity of as-reduced Nd_0.8Sr_0.2Ni_xO₂ films. For samples R-Ni-1.0 and R-Ni-1.05, which failed to stabilize an infinite-layer structure, both exhibit insulating behavior across the entire temperature region. This insulating nature arises from the significant disorder introduced by uncontrollable elemental dynamics and oxygen vacancy formation during both the growth and reduction processes. To describe the electrical transport behavior of such disordered systems, the variable range hopping (VRH) model serves as a classical framework. At low temperatures, the resistivity satisfies [41]:

(2)\rho \left(T\right)={\rho }_0\exp {\left(T_0/T\right)}^{\frac{1}{d+1}},

where T₀ depends on the density of localized states at the Fermi level and the falloff rate of the wave functions associated with these states, and d represents the dimension of this noninteracting system [42]. Fig.4(e) shows ln ρ versus T^{-1/4} curves for the insulating R-Ni-1.0 (Below 300 K) and R-Ni-1.05 (Below 90 K) samples. A good agreement linear fit can be observed, indicating the 3D VRH model.

For infinite-layer R-Ni-1.15, it shows superconductivity with a T_c,onset of 10 K and a T_c,offset of 5 K [Fig.4(f)], consistent with previously reported results [6]. However, this temperature is moderate among the values reported for the same system. It can be more effectively increased by further improving the quality of the infinite-layer structure. This can be achieved by adding a capping layer to the film to protect the surface of the infinite-layer and using substrates with a more compatible lattice, such as (LaAlO₃)_0.3(Sr₂TaAlO₆)_0.7 [24]. In addition, co-doping technology is also an effective method for optimizing the infinite-layer structure and increasing the T_c value [12, 43]. For another infinite-layer sample R-Ni-1.10 with a nickel content between the above samples, it shows a distinctive “U” type ρ(T) curve. Concretely, it exhibits metallic behavior in the temperature range of 200–300 K, crosses over to insulating behavior below 200 K, and displays a slight resistivity downturn at 4.5 K. This resistivity−temperature dependence appears to lie intermediate between the behaviors of R-Ni-1.05 and R-Ni-1.15. Considering that inadequate reduction may also contribute to this transport behavior [44, 45], we made subsequent measurements after continued reduction of this sample, which still exhibited similar “U” type curve but insulating properties below 4.5 K (See Supplementary Information, Fig. S3).

Taken together, these results suggest that employing a Ni-overdoped target during pulsed laser deposition facilitates the synthesis of infinite-layer nickelates with superconducting properties. To further investigate the impact of target nickel compensation on cation stoichiometry in epitaxial films, we quantified the relative stoichiometric ratios of Nd_0.8Sr_0.2Ni_xO₃ films using X-ray photoelectron spectroscopy (XPS). The wide range of XPS spectrum of different parent samples are presented in Fig. S4, revealing the elemental composition and their corresponding core lines. For rare earth elements, due to the simultaneous effect of solid-state hybridization and intra-atomic electrostatic coupling between 3d hole and outer unpaired 4f electrons, the shape of 3d state lines will become complex [46−48]. Additionally, the shake resulting from core−hole screening effects in the final state also complex the lines shape [49−51]. These factors collectively lead to the splitting of the peaks into a main peak and multiple satellite peaks. The Nd 3d, Ni 2p and Sr 3d XPS spectra are presented in Fig.5(a)−(c) and Fig. S5, the peak shape was fitted using Lorentzian and Gaussian function, fitting details are presented in experiment section. The nickel element in the film exhibits a mixed valence of +2 and +3, consistent with observations in NdNiO₃ films [52]. The relative molar ratios of the elements were quantitatively analyzed based on the areas of the main peak regions (A−F).

We extracted the Ni/Nd and Nd/Sr molar content ratios as functions of the nickel content in the target, as illustrated in Fig.5(d). For the nominal Nd_0.8Sr_0.2NiO₃, the expected Nd/Sr and Ni/Nd ratios are 4 and 1.25, respectively. In the case of thin films grown using targets with varying Ni contents, the Nd/Sr ratio remains remarkably close to the ideal value, deviating by less than 2%. In contrast, the Ni/Nd ratios consistently fall below the nominal values. As the nickel content in the target increases, the Ni/Nd ratio of as-grown films rises monotonically, gradually approaching the nominal value. This trend suggests that excess nickel doping in the target effectively compensates for nickel loss during the growth process, ultimately resulting in as-grown films that more closely align with the intended composition.

The electronic structure of the films was further measured using x-ray absorption spectroscopy (XAS) at the Ni-L_2,3 and O-K edges, as shown in Fig.6. Fig.6(a) presents the Ni-L_2,3 XAS of the parent Nd_0.8Sr_0.2Ni_xO₃ films, matching well with previously observed in NdNiO₃ [53]. There are clear separated peaks at Ni-L₃ edge, which is associated with the 2p⁶3d⁸-2p⁵3d⁹ (~852.4 eV) and 2p⁶3d⁸Lⁿ-2p⁵3d⁹Lⁿ (~853.8 eV) multiplet transition [54, 55]. After the reduction process, the Ni-L_2,3 XAS peaks shift towards lower energy (~852.2 eV), indicating the lower valence of Ni in the as-reduced films. At approximately 854 eV, a weak shoulder peak can still be observed, proving that the above-mentioned transition process still exists in the as-reduced films [56]. For the negative charge-transfer insulator perovskite nickelates, due to the distribution of breathing mode distortions of nickel in different electronic states in this structure, the 2p electrons in the oxygen ligands of ⟨{\mathrm{N}\mathrm{i}\mathrm{O}}_6⟩ octahedra spontaneously transfer to the 3d orbitals of the central nickel cation and leave holes in the oxygen ligand, which known as the self-doping effect [54, 55, 57]. This results in the presence of a pre-edge (~528 eV) in the O-K edge, as shown in Fig.6(c). After reduction by CaH₂, the condition will change. The infinite-layer nickelate will transform into the Mott−Hubbard regime, causing the pre-edge to disappear [Fig.6(d)], just as the observation in the parent infinite-layer (La, Nd)NiO₂ [55]. However, a new weak pre-edge peak appears at 530 eV, indicating the presence of oxygen ligand hole in the Nd_0.8Sr_0.2Ni_xO₂ [56].

Remarkably, the XAS peak shapes and compositions remain largely consistent across films with varying nickel contents, both before and after reduction, despite differences in their transport behavior. This indicates that the electronic structure is robust to changes in nickel stoichiometry, even as the films exhibit distinct electrical properties. Although significant stoichiometric differences exist in the targets, the primary electronic structure of the thin films remains consistent with that of the Nd_0.8Sr_0.2NiO₃ or Nd_0.8Sr_0.2NiO₂ phase. This occurs because the intentionally introduced nickel excess in the targets primarily suppresses secondary phase formation during film growth and compensates for minor surface nickel segregation. Crucially, these secondary phases and surface nickel oxide particles constitute non-dominant components, with their abundance substantially lower than the Nd_0.8Sr_0.2NiO₃ or Nd_0.8Sr_0.2NiO₂ phase. The electronic structure variations induced by such trace constituents likely fall below the detection limit of XAS.

4 Conclusion

In summary, precise stoichiometric control is as crucial as growth condition optimization for synthesizing infinite-layer nickelate films. Kinetic factors, such as elemental segregation and preferential precipitation during pulsed laser deposition, often lead to nickel deficiency in as-deposited parent-phase films. These deficiencies promote the formation of secondary phases in the parent films and impede the stabilization of infinite-layer structures, even after thorough hydrogenation reduction. By introducing 15% excess nickel doping into the target, we effectively compensate for nickel stoichiometry loss and facilitate the formation of the perovskite phase, enabling the production of corresponding superconducting thin films after reduction. And the structural and electrical properties of nickelate films, both before and after reduction, are intimately correlated to the nickel compensation content. Increasing nickel compensation systematically reduces secondary phase content in parent films and enhances the likelihood of forming phase-pure infinite-layer structures with superconductivity after reduction. XPS quantitative analysis reveals that, with this nickel compensation strategy, the Ni/Nd ratios in the parent-phase films closely align with the ideal values. Furthermore, this approach does not alter the electronic structure of either the as-grown or reduced films, demonstrating that its primary influence lies in finely adjusting the stoichiometric ratios and structural properties of the nickelate films. Our findings establish a stable synthesis protocol for synthesizing superconducting infinite-layer nickelate films and significantly alleviating the challenges inherent to their preperation.

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