Growth of all-chemical high critical current YBa₂Cu₃O_7−δ thick films and coated conductors

For the deposition of thin and thick layers of YBCO, we have used 5 × 5 mm LAO or ^CSDCZO/YSZ single crystals and ^CSDCZO/^ABADYSZ/SS metallic substrates (table 1). The ^ABADYSZ/SS metallic substrates have been provided by Bruker HTS, Germany [17]. The tapes are 4 mm wide and 0.1 mm thick. On top of these substrates, ^CSDCZO cap layers were deposited by spin coating, although they can also be produced by IJP [44, 45, 64]. These CZO cap layers were grown at 900 °C during one hour in air, they are ∼25 nm thick and atomic force microscopy (AFM) analysis indicates that they have a rms (root mean square) roughness of ∼1.6 nm [61, 65]. In the case of single crystals, an rms roughness of ∼0.9 nm was systematically achieved.

The low-fluorine metalorganic precursor solution for the deposition of thin YBCO layers is composed of two acetates (barium acetate, Sigma Aldrich, and copper acetate anhydrous, Alfa Aesar—Cymit Química) and one TFA (yttrium trifluoroacetate anhydrous, 99.99 + %, Sigma Aldrich). They are dissolved in methanol and propionic acid in a molar concentration of 1.5 M (sum of metals). The fluorine reduction of this precursor solution is ∼80%, as compared to the all-TFA route [39]. From this precursor solution, ∼200 nm thick YBCO layers are obtained by single deposition spin coating, at 6000 rpm for 2 min. The pyrolysis of this type of solution has been described in detail elsewhere [39]. For the thicker films (t ≥ 700 nm), single pass IJP or spin coating multideposition was used (although multiple pass IJP has also been demonstrated as feasible) (table 1). The precursor solution for IJP had to be modified to achieve suitable properties for IJP deposition [43, 44]. Therefore, the solvent was changed from methanol to butanol and the molar concentration of this low-fluorine precursor solution was decreased down to 1 M (sum of metals). Prior to the deposition by IJP, two additives were added to the YBCO solution: one additive was used to modify the interaction between the ink and the substrate, which was provided by KAO Chimigraf Company (Rubi, Spain). The other, only used when indicated, was 5 molar% of Ag-TFA (99.99%, Aldrich) (table 2). As we will comment later, Ag salts decrease the supersaturation of reaction (1) and so, allow us to decrease the nucleation temperatures of epitaxial YBCO. In the thick YBCO layers, the pyrolysis process was performed at 500 °C, due to the higher decomposition temperature of the first additive. After the pyrolysis process, the YBCO thin and thick layers were heated at 25 °C min⁻¹ up to the nucleating temperature, where they were subjected to the growth treatment.

Growth of YBCO thin layers deposited on LAO single crystals was performed at 820 °C for 180 min, in an atmosphere of N₂ and 200 ppm of O₂ (PO₂ = 2 × 10⁻⁴ bar), using a water partial pressure of P_H2O = 23 mbar and a total pressure of 1 bar (table 2). The oxygenation of these samples was performed at 450 °C–550 °C in a dry oxygen atmosphere for 3.5 h. For the thick YBCO layers deposited on LAO and ^CSDCZO/YSZ single crystals, the crystallization was performed at 770 °C–780 °C for 180 min in a mixture of N₂ and 200 ppm of O₂ (PO₂ = 2 × 10⁻⁴ bar) and a higher water partial pressure (P_H2O = 200 mbar) (table 2).

For the crystallization of thin and thick YBCO CCs, we used Ag salts added to the precursor solution to lower the YBCO c-axis nucleation temperature and thus, minimize the chemical reactivity with the CZO cap layers. The crystallization process, in this case, is a two-step thermal treatment: at 700 °C–730 °C we perform the YBCO nucleation for about 60 min and then we rise the temperature up to 770 °C–800 °C, where we keep the sample for 135 min until it is completely grown. The atmosphere inside the tubular furnace during the crystallization step is a humid (P_H2O = 23 mbar and 200 mbar for thin and thick films, respectively) mixture of N₂ and 200 ppm of O₂. The oxygenation conditions were the same as described above.

X-ray diffraction (XRD) was conducted on the fully converted YBCO samples to analyze the crystalline phases that are present and the quality of the texture. For this, a Bruker AXS GADDS diffractometer equipped with a 2D detector was used. For the surface morphology studies of the YBCO layers, an FEI Quanta 200 FEG scanning electron microscope was used. The landscape of the structural defects of these layers was studied by transmission electron microscopy (TEM), using an FEI Titan 60–300 microscope equipped with an X-FEG gun, a CETCOR probe corrector and a Gatan TRIDIEM 866 ERS energy filter operated in scanning TEM (STEM) mode at 300 kV. The superconducting properties were measured inductively, using a superconducting quantum interference device (SQUID) magnetometer (Quantum Design, San Diego, CA). The critical current density (J_c) was calculated using the Bean critical state model of a thin disk [66].

Besides XRD, the CZO cap layers were also subjected to a morphological study using AFM. The equipment used for this purpose was an Agilent 5100 with silicon tips and the topography of the samples was obtained using a tapping mode. The images were processed with a Mountains Map Premium 7.2 software. Besides the quantification of the roughness, this software permits a quantification of the flat area fraction of the surface of these CZO cap layers. To do this, we have to first define a threshold height value above which the surface is not considered flat. In our case, we have used a 1.5 nm threshold value, which represents one and a half fluorite unit cell height. Then, using the threshold, value a binary mask is applied to the whole AFM image. The result is a two color image, where the flat areas are represented by the bright sections [65]. For the CZO cap layers used in this work, flat area fractions between 62% (in the ^CSDCZO/YSZ/SS metallic tapes) and 90% (in the ^CSDCZO/YSZ single crystals) have been quantified.

In this section, we will report an analysis of the growth process and the superconducting properties of thin and thick YBCO films grown on LAO substrates as a model system. The processing conditions to grow YBCO thin films from low-fluorine precursors on LAO substrates (<300 nm) with optimized superconducting properties have been previously investigated and they are described in the experimental section (tables 1 and 2). In figure 3(a) we present a typical 2D-XRD pattern of a 200 nm thick YBCO film grown on LAO single crystals and in figure 3(b) the corresponding integrated θ-2θ spectrum is shown.

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As it is observed, the YBCO thin layers are highly epitaxial, i.e. no random or a–b nucleation and no intermediate or unreacted phases are present. The lack of a-b grains is also confirmed by the SEM image (figure 3(c)) that reveals a homogeneous surface with low porosity of the YBCO layer. These thin YBCO layers exhibit reproducible excellent superconducting performances: J_c(77 K)^sf = 5 MA cm⁻², T_c = 91.2 K and ΔT_c = 1.8 K (figure 4(a)). Also, they exhibit a very homogeneous microstructure showing only a few YBa₂Cu₄O₈ (Y248) stacking faults (figures 4(b) and (c)), whose atomic structure and distribution has been recently described in detail [67–69]. This solution becomes, therefore, an excellent candidate for the multideposition of thick YBCO layers by the spin coating technique.

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We prepared thick YBCO layers, both by multideposition and by single deposition IJP, as described in the experimental section (table 1). Optical microscopy images recorded after the pyrolysis showed, in both cases, film homogeneities similar to that of thin films. However, when these thick films are grown following the same processing conditions of high quality thin films, a strongly perturbed texture is achieved (see the XRD patterns and SEM image in figures 3(d) to (f)), indicating that some random nucleation occurs under these processing conditions. This means that the degree of supersaturation was excessively high for films with increased thickness. It is clear, therefore, that to achieve high quality YBCO thick films the processing conditions need to be deeply re-examined.

As we have explained in section 1, in the thin film limit, the nucleation process of epitaxial YBCO films can be described on the basis of the formation of independent nuclei stabilized when its energy overcomes the energy barrier described in equation (3).

On the other hand, the theoretical analysis of the heterogeneous nucleation process in epitaxial YBCO thick films based on the BaF₂ approach (equation (1)) was first performed by Solovyov et al [22, 51]. These analyses clarified that consideration of the thick limit requires extending the schema of the classical nucleation theory exemplified by equation (3) to the consideration of how the HF gas flow balance induces collective effects in the nucleation events. While in the thin film limit, it could be safely assumed that there was no influence of the HF gas impedance towards the film surface, in the thick film approach this effect needs to be explicitly taken into account [51–53]. The main conclusion in the analysis of nucleation in thick films is that the supersaturation degree s (s = Δμ/kT) in a high surface energy site, i.e. a site with a retarded nucleation event, is influenced by the already existing stable nuclei, i.e. the nucleation process becomes a collective effect (figure 1). The supersaturation of the sites with delayed nucleation could then be correlated with several processing parameters:

$$\begin{eqnarray}&&s=\displaystyle \frac{{\rm{\Delta }}\mu }{kT}\propto \displaystyle \frac{dr}{P{(H2O)}^{1/2}{D}_{s}}\end{eqnarray} \tag{ 4 }$$

where d is the film thickness, r is the growth rate and D_s is the HF diffusion permeability through the film precursor. It is clear from equation (4) that the supersaturation degree s is directly influenced by the film thickness and so, we can understand that the processing conditions need to be modified to grow thick films. Based on the predicted dependence of supersaturation for the processing parameters, we decided to investigate new nucleation conditions, mainly by increasing the P(H₂O) values and reducing the growth rate r. The temperature dependence of YBCO growth rates r has been previously investigated by several authors [70–73] and it was found to follow the thermally activated behavior arising from the equilibrium constant K_e of reaction (1) [72]. For example, r was found to decrease by a factor of 2.5 when reducing the growth temperature from 820 °C to 770 °C. Concerning P(H₂O), our experimental set-up allowed us to increase it by a factor of 10. Overall, therefore, supersaturation s is predicted by equation (4) to be reduced by a factor of ∼7 in our modified nucleation experimental conditions of thick films (table 2). Although the growth experiments of YBCO thick films under these modified conditions showed a significant improvement concerning texture quality, the presence of a/b grains oriented perpendicular to the substrate was not fully eliminated and so, further improvement of the nucleation process was still required (figures 5(a)–(c)).

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To further enhance the texture quality of the YBCO thick films, we decided to implement in our growth process of thick films the use of an additional YBCO-Ag seed layer (50 nm thick) at the substrate interface. In previous investigations, it was shown that, due to the decreased peritectic temperature of the Ag-YBa₂Cu₃O_7−x system, the supersaturation degree is decreased under non-modified processing conditions, compared to the YBa₂Cu₃O_7−x composition [43, 74]. As a result, enlarged temperature windows for c-axis nucleation, reduced final YBCO film porosities and enhanced film planarities were achieved [46]. This strategy was found to be indeed successful, as it is concluded from the analysis of the results reported in figures 5(d)–(f) corresponding to a 1.2 μm thick film prepared by spin coating multideposition (five deposition repetitions with the first being a YBCO-Ag seed layer) (table 1). The film texture quality was clearly improved: no a/b nucleated grains were detected in the XRD pattern and only a few a/b axis nucleated grains could be observed in the SEM images. Only minor traces of Y₂O₃ are visible in both cases (using or without Ag) and so, we expect a near completion of the reaction (1).

Even if the new supersaturation conditions lead to full c-axis nucleation and growth, we could still go further and investigate whether there is any influence of the solution deposition methodology on the final film quality. For example, it is known that during the pyrolysis there is a tendency to segregate a CuO layer at the film surface and thus, when using a multideposition approach to prepare thick films we are facing the problem of compositional inhomogeneity [23, 31, 56]. This is clearly visible in figures 6(a) and (b), where we display a focused ion beam (FIB) cross section image of a film with five pyrolyzed layers (final film thickness ∼1 μm) and a cross section TEM image of the same film quenched after heating it up to the growth temperature, respectively.

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As we can see from the FIB image, between each deposited and pyrolyzed layer some CuO interlayers are formed. Therefore, the high temperature processing conditions chosen to grow these films need to promote Cu atomic diffusion out of these interlayers to be able to form YBCO. Eventually, this diffusion process may lead to the formation of some layered porosity associated with the phase transformations and shrinkage occurring during the heating ramp (layered pores in figure 6(b)) [23, 56, 57]. This effect can be minimized, however, by following different strategies. For example, if the layer thickness of each deposition is small the Cu diffusion length is large enough to overcome the initial inhomogeneous atomic distribution. In addition, through careful intermediate high temperature annealing treatments, the films can be compacted and homogenized [26, 27]. Therefore, achieving optimized nucleation and growth conditions of thick YBCO films involves careful processing requirements, to minimize the effect of the undesired CuO interlayers and the associated porosity that, eventually, may be created after the film growth.

An alternative approach to minimize this source of film inhomogeneity is to perform a single pass solution deposition and pyrolysis step for the thick films. This is, in principle, possible by using IJP, where the total amount of salts deposited can be controlled electronically (table 1). The use of this deposition methodology for CSD YBCO thin films has been previously described [43–45], while its extension to thick films and the corresponding optimization of the pyrolysis process will be presented elsewhere. Applying this deposition method, we were able to achieve homogeneous films, from one single deposition, up to 1 μm thick onto LAO single crystals, without any CuO interlayer phase separation after the pyrolysis process and so with reduced film porosity after the heating stage (figures 6(c) and (d)). The FIB image confirms that the microstructural and compositional homogeneity is indeed preserved following IJP and a single pyrolysis process. Actually, the difficulties with keeping thick homogeneous films in a single deposition are associated with both the solution deposition and drying steps and the pyrolysis treatment, where buckling or cracking should be avoided [75]. Using the IJP technique and applying the new supersaturation nucleation and growth condition for c-axis nucleation in thick YBCO films we were able to grow YBCO layers with thicknesses in the range of 0.6–1 μm having high texture qualities and good superconducting performances. In figures 7(a)–(c) the 2D-XRD patterns of a 600 nm thick YBCO layer on LAO single crystals and the corresponding SEM image are shown. We see that at this film thickness the YBCO layers are highly epitaxial. When the YBCO film thickness is increased to ∼1 μm the texture begins to degrade and a small amount of random grains is observed (figures 7(d)–(f)). Besides the texture degradation, also confirmed by some broadening of the (00l) XRD Bragg peaks, we can also detect the presence of some unreacted crystalline phases (in this case Y₂O₃). This indicates that the supersaturation conditions for c-axis nucleation at these high film thicknesses still requires some improvement, even if no a–b nucleation is detected in the XRD patterns and SEM images (figures 7(d)–(f)).

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Our success in achieving films with a thickness in the range of ∼0.6–1 μm approaches the best results registered so far in other functional CSD films [76]. In figures 8(a)–(c) we present TEM cross section images and energy dispersive x-ray microscopy (EDX) compositional analysis of a 1 μm thick single deposited IJP layer after the growth process. The TEM image shows that the porosity has been drastically reduced, as compared to the thick YBCO layers deposited by the multideposition method and grown using a conventional high temperature annealing process. The TEM and EDX images in figures 8(a)–(c) show that some secondary phases persist (BaO, BaCuO_x). This defect structure, however, does not show any sign of the layered porosity observed in the multideposition process (figure 6(b)), thus suggesting that further optimization of the nucleation and growth processes of IJP films should allow further improvement of the microstructural homogeneity.

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An overall view of the superconducting performances of the YBCO layers with different thicknesses deposited and grown on LAO single crystals is reached looking at the thickness dependence of J_c and I_c at 77 K and self-field (figure 9). Thin films in the range of 200–300 nm display J_c(77 K)^sf values around 4–5 MA cm⁻² and I_c(77 K)^sf ∼ 120 A/cm-w. Thick films prepared using multideposition by spin coating or single pass IJP, with thicknesses in the range 0.6–1.2 μm, preserve fairly good critical current densities: J_c(77 K)^sf = 2–2.7 MA cm⁻² and I_c(77 K)^sf = 240 A/cm-w is the highest value registered on LAO substrates, corresponding to a 1 μm thick film. These I_c(77 K)^sf values were obtained using inductive SQUID magnetization measurements. We should note that these values of the critical currents are close to those achieved by other vacuum-based film deposition techniques [14, 15, 18, 20, 21]. Finally, figure 10 displays the measured temperature dependence at self-field ${I}_{{\rm{c}}}{(T)}^{sf}$ for several films, which summarizes the success in developing thick YBCO films on LAO substrates. A detailed analysis of the in-field, temperature dependence and orientation of J_c(H,T,θ) is an issue related to the vortex pinning landscape of CSD YBCO films. This topic goes beyond the scope of the present work, but it has been recently analyzed in detail in relationship with the development of CSD nanocomposite YBCO films [77, 78].

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As we have previously mentioned, when growing YBCO films on CeO₂-related cap layers there is an interfacial reactivity issue which needs to be properly handled (figure 2) [62, 63]. This reactivity takes place between the BaF₂ precursor of equation (1) and the CZO cap layer and it leads to the formation of BaCeO₃ (BCO). The formation of this perovskite and the nucleation of YBCO on CZO are two competitive processes. If YBCO nuclei appear first, then we have chances to induce c-axis nucleation on top of the CZO cap layers. The formation of an epitaxial BCO layer will take place later underneath the YBCO layer and will not affect the texture of the YBCO in any way [61]. However, if BCO is formed first at the interface a significant amount of nucleation with random orientation occurs for both the BCO and YBCO layers and so the superconducting properties of the film are degraded.

Previous works on YBCO thin film growth on ^CSDCZO/YSZ single crystals showed that a key parameter to achieve high J_c(77 K)^sf values is the planarity of the CZO cap layers [61]. High performance was achieved with CSD-grown YBCO thin films (thickness ∼300 nm): J_c(77 K)^sf = 4–5 MA cm⁻² [61, 65]. These optimal performances are achieved when the surface planarity of the CZO cap layer is above ∼70%. The reactivity issue can then be properly handled and the supersaturation conditions can be selected to nucleate c-axis oriented YBCO islands. Typically, the growth conditions for these YBCO thin films were P(H₂O) = 22 mbar and annealing temperature T ∼ 770 °C–780 °C [61, 65] (table 2). When Ag-TFA additives are used in the solution, the nucleation temperature can be reduced down to ∼700 °C (table 2).

Now, our objective here is to use the knowledge generated when growing YBCO thick films (∼1 μm) on LAO substrates (section 3.1) to adapt the process to ^CSDCZO/YSZ single crystal substrates, as a model system for 'all-chemical' CCs. The strategy was to mimic, as much as possible, the processing conditions used for LAO and to introduce fine adjustments when necessary. To ensure that the crystallization window of YBCO remained as wide as possible, we included the Ag-TFA additive, either in a YBCO seed layer or in all the YBCO layers of a multideposition process. We focused our nucleation and growth analysis on YBCO films with a final thickness of ∼1 μm obtained after five sequential depositions by spin coating with the corresponding pyrolysis processes, as described in the experimental details in section 2 (table 2). Our previous nucleation analysis of YBCO thin films on ^CSDCZO/YSZ substrates determined that interfacial reactivity was low enough at T < 780 °C. We therefore kept this temperature while the low supersaturation s (equation 4) was kept by increasing P(H₂O) (table 2), i.e. P(H₂O) = 200 mbar.

Figures 11(a)–(c) show a GADDS XRD pattern, the corresponding integrated XRD pattern and an SEM morphological image, respectively, of a 1.050 nm thick YBCO film grown on top of a ^CSDCZO/YSZ single crystal. The XRD patterns show that the YBCO film is highly epitaxial while the CZO cap layer has been fully converted to epitaxial BCO. As it is observed, the microstructural quality of the film is fairly good and with just some small pores observed in the SEM image. Figures 12(a)–(e) show several cross section TEM images of such YBCO thick films, grown under the above-mentioned conditions. The low resolution cross section TEM image (figure 12(a)) displays a lower residual porosity than the YBCO thick films grown on LAO substrates (figure 8(a)), while a few secondary non-superconducting phases (BaCuO₂, Y₂O₃, Y₂Cu₂O₅) still remain (figures 12(b)–(d)). Figures 12(b)–(e) also show that the concentration of Y248 stacking faults is similar to that observed in YBCO films grown on LAO single crystals (figures 4(b)–(c)). We suggest that the nucleation rate of YBCO films is higher when using CZO cap layers owing to its lower lattice mismatch (ε = −0.52%) than with LAO (ε = −1.56%) because under these unstrained conditions the nucleation barrier (equation (3)) usually decreases [59]. When the nucleation rate is increased, merging of individually nucleated YBCO grains is easier because the density of the nuclei is larger and thus the required atomic diffusion lengths are shorter. As a consequence, the microstructures tend to be more homogeneous. This is an issue that has already been previously discussed by other authors, who suggested that secondary phases could be accumulated at the merged YBCO grain boundaries [79].

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Very competitive superconducting performances have been achieved so far in these ^CSDCZO/YSZ thick films (∼1 μm) (see figure 9): J_c(77 K)^sf = 3.7 MA cm⁻² and I_c(77 K)^sf = 390 A/cm-w or J_c(65 K)^sf = 7.2 MA cm⁻² and I_c(65 K)^sf = 750 A/cm-w (figure 10). These values are very close to the best performances achieved so far in vacuum-deposited YBCO films with similar film thicknesses [14, 15, 18, 20, 21]. Therefore, the competitiveness of the 'all-chemical' CC approach is confirmed. As we mentioned before, TEM images show that some residual porosity and secondary phases still remain. This suggests that further improvement of the microstructural and superconducting characteristics of the all-chemical ^CSDYBCO/^CSDCZO/YSZ architecture should be achieved through optimized processing conditions of the CZO cap layer and the YBCO layer.

The 'all-chemical' CC architecture we selected to develop is ^CSDYBCO/^CSDCZO/^ABADYSZ/SS (figure 2). The SS metallic substrates with YSZ deposited by ABAD on top were provided by Bruker HTS from Germany and are described in the introduction of this article.

The ^CSDCZO cap layers were deposited using spin coating, although IJP can also be used for long length CC production [43, 44] and grown according to the thermal treatment described in the experimental section (table 2). In order for these cap layers to be suitable, they have to comply with some requirements: they have to be epitaxial, they have to exhibit a low roughness and a high degree of surface flatness. The influence of the CZO cap layer flatness on YBCO film growth was previously investigated and it was concluded that the critical current density ${{J}_{{\rm{c}}}}^{sf}$ and the film texture quality is improved when the film flatness is enhanced [61, 65]. Therefore, we performed a complete analysis of all these features (epitaxy, roughness and flatness) in our CZO cap layers [44, 61]. The CZO cap layers exhibit a high quality epitaxy as only (00l) Bragg peaks are observed. The AFM measurements reveal rather low values of roughness, ∼1.6 nm for an area of 5 × 5 μm² and ∼1.3 nm for an area of 1 × 1 μm². For the same areas, we notice that the surface flatness of the CZO layer (measured as the percentage of the area with heights differing by less than 1.5 nm) is higher than 60% [43, 80].

The reactivity issue discussed in the previous section meant that the growth conditions of thick YBCO films on ^CSDCZO/YSZ single crystal substrates needed to be revised when growth is performed on metallic substrates. This is because the enhanced surface roughness of the CZO cap layers increases the interfacial chemical reactivity and, very likely, it also modifies the heterogeneous nucleation barriers ΔG* for YBCO islands and thus, their nucleation rate.

To comply with all the requirements for obtaining high critical current CCs we first investigated the temperature where the interfacial chemical reactivity is slow enough in YBCO thin films (∼200 nm). To promote c-axis nucleation we included YBCO-Ag seed layers in our architecture (figure 2). Two annealing steps at different temperatures were used, one at 700 °C to nucleate the film and another at 770 °C, to complete the film growth. The two-step annealing process becomes necessary to minimize the residual porosity in the YBCO films [27, 46]. As is shown in the XRD patterns of figures 13(a) and (b), a fairly good YBCO film texture was achieved, i.e. only sharp single (00 l) Bragg peaks are observed in the GADDS XRD pattern. There is no evidence of a-b or random YBCO nucleation, an issue also confirmed in the SEM images (figure 13(c)). The typical values of the texture quality determined by XRD are Δω = 2.5° and Δϕ = 6–7° [45]. Once again, the CZO cap layer has been fully transformed to epitaxial BCO with two orientations perpendicular to the substrate ({200} and {110}). These YBCO CCs exhibit high critical temperature, T_c = 91 K, and fairly high critical current density, J_c(77 K)^sf = 1 MA cm⁻², a value usually limited by granularity effects in these CCs [43, 80] and similar or slightly lower than those usually displayed by other vacuum deposition techniques [18, 20]. These nucleation and growth conditions were found to be robust, and allowed us to achieve a high reproducibility in the film qualities and the corresponding superconducting performance. We should stress that these results confirm that the use of YBCO-Ag seed layers allows developing c-axis nucleation at sufficient low temperatures.

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The next step was to combine the knowledge generated to grow thick YBCO layers on LAO and ^CSDCZO/YSZ single crystal substrates and the knowledge about the kinetics of interfacial reactivity forming BCO layers, to achieve thick (∼1 μm) YBCO CCs with the 'all-chemical' architecture on ^ABADYSZ/SS substrates (figure 2). Owing to the advantages that the IJP deposition method provides for scaling-up the CC production process and grounded in the good results reported in section 3.1, we decided to deposit thick YBCO CCs using this deposition method. A single pass IJP deposition was used leading to a final YBCO film thickness of ∼1 μm on top of ^CSDCZO/YSZ/SS substrates. As we have described in sections 3.1 and 3.2 when the thickness of the YBCO layers is increased the supersaturation degree s tends to increase (see equation (4)) and so the nucleating conditions need to be modified. At the same time, however, we need to take into account the modified kinetics of the interfacial reactivity reaction, when the processing conditions of thick films are used. Therefore, we need to handle simultaneously, both the control of the supersaturation degree s and the interfacial reactivity.

The increase in the YBCO film thickness by a factor of ∼5 has a strong influence on the percentage of c-axis nucleation achieved in our CC architecture. As in the case of single crystalline substrates, the supersaturation conditions used to grow thin films are again no longer valid here. We adjusted the supersaturation degree s (see equation (4)), following the same approach used in single crystals, i.e. by increasing the water pressure to P(H₂O) = 200 mbar and by decreasing the nucleation temperature which directly influences the growth rate r. To disclose the optimum nucleation temperature in our CCs we performed short isothermal treatments followed by a quenching process to mainly analyze the nucleation stage.

The as-obtained samples were analyzed by XRD. Typical XRD patterns of such quenched films are presented in figures 14(a) and (b). It can be seen that at 700 °C YBCO nucleates randomly while a large fraction of unreacted crystalline phases is still present after 30 min of annealing at this temperature (figure 14(a)). We also noticed that no BCO has formed yet, i.e. at this temperature and under the selected processing conditions no BCO formation is detected. This is also confirmed by the presence of an intense (200) Bragg peak of the CZO film, which shows that no cap layer reactivity occurred. So, at this temperature the reactivity issue is controlled but we have random nucleation of YBCO due to an excessive supersaturation s. It is necessary, therefore, to increase the annealing temperature to optimize the nucleation process. If we look to the XRD pattern of a film quenched after annealing during 30 min at 730 °C (figure 14(b)), we can see that the fraction of randomly oriented YBCO grains ((103) YBCO Bragg peak) has drastically diminished while the unreacted precursor crystalline phases are no longer present. Taking into account all this information, we designed a double step annealing treatment for growing thick YBCO coated conductors. The first step was the nucleation stage at 730 °C (60 min) followed by the growth step at 770 °C–800 °C (135 min) (table 2).

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Figures 15(a)–(c) show XRD patterns and a morphological image (SEM) of a fully grown 1 μm thick YBCO CC based on the new designed nucleation and growth treatment (table 2). Even though a small amount of randomly oriented YBCO grains are still present ((103) YBCO Bragg peak), it is noticeable that c-axis nucleation and growth is strongly favored under these new supersaturation conditions. Again, no a-b grains seem to be present, neither in the XRD patterns nor in the SEM image. These thick YBCO superconductors exhibit fairly good superconducting properties. For a single pass IJP deposition (∼1 μm) a YBCO film was achieved exhibiting a critical temperature T_c = 90 K, a sharp transition with a width of ΔT_c = 6.5 K (see figure 16(a)), and J_c(77 K)^sf = 1 MA cm⁻² which corresponds to I_c(77 K)^sf = 100 A/cm-w (see figure 9). STEM cross section images of a 1 μm thick YBCO CC (figures 16(b)–(d) show that some porosity and secondary phases still remain and so our analysis is valid as a 'proof-of-principle' for how to grow 'all-chemical' CCs but there is still further room for improvement for both the film texture and the final microstructure. Figures 17(a)–(c) show a STEM cross section and EDX analysis of the interface in this CC where the formation of a BCO layer is also confirmed. As we have demonstrated in section 3.2, optimizing the nucleation and growth conditions on ^CSDCZO/YSZ single crystals, where the cap layer planarization is higher, allowed us to reach much larger critical currents, i.e. I_c(77 K)^sf = 390 A/cm-w. The refinement of the YBCO nucleation and growth conditions, as well as an increase in the CZO cap layer planarization, should allow us to further minimize the concentration of misoriented YBCO grains and decrease the residual porosity. Consequently, the highest performance achieved with single crystalline substrates would be within reach.

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Finally, we would also like to stress that all the processes reported here to prepare 'all-chemical' CCs are essentially valid for preparation of CSD nanocomposite films and CCs. This enhances the performance in an extended range of magnetic fields and temperatures by introducing preformed nanoparticles in the initial metalorganic solutions [64]. This is an issue, however, that will be reported elsewhere in the future.