High critical current solution derived YBa₂Cu₃O₇ films grown on sapphire

A low fluorine YBCO precursor solution was synthesized using Ba and Cu acetates and Y trifluoroacetate, all precursor salts were acquired from Sigma-Aldrich. First, the Ba acetate is dissolved in methanol and propionic acid (70 vol%:30 vol% = CH₃OH:CH₃CH₂COOH) under continuous stirring. After the barium (Ba) precursor is completely dissolved, copper (Cu) acetate is added to the solution. Once the Cu acetate dissolution is complete we proceed with the addition of yttrium trifluoroacetate and leave it under stirring overnight. The precursor metalorganic solution is subjected to rotoevaporation at a temperature of 75 °C until a blue powder is obtained. This powder is then re-dissolved in methanol in order to achieve a concentration of 1.5 M (sum of metals). The stoichiometry of the metal ions is Y:Ba:Cu = 1:1.5:3. The reduction in fluorine content is about 80% with respect to an all TFA YBCO precursor solution [30] and a reduction of Ba content is used because it allows a considerable reduction of the YBCO film growth temperature [31, 32]. Once the solution is filtered it is stored in close vials, under dry atmosphere until deposition.

The YBCO layers were deposited and grown onto CZO buffered sapphire substrates of 5 × 5 mm. The unpolished and BPS substrates were prepared as ribbons in lengths up to 1 m by the EFG technique [29] and they were R-cut. These ribbons were then coated with a YSZ layer (∼35 nm) by sputtering [2]. The CZO buffer layers were deposited by CSD using spin coating at 6000 rpm for 2 min onto the ^sputtYSZ/sapphire substrates with sizes up to 10 × 50 mm. The use of CZO buffer layers grown by CSD has been shown to lead to highly flattened layers and so they were selected for this multilayer architecture [33]. The details on CZO precursor solution synthesis are reported elsewhere [34]. In order to achieve epitaxial CZO buffer layers the as-deposited layers were subjected to a growth treatment at 900 °C for 1 h in static air atmosphere. The thickness of the grown CZO buffer layers was between 15 and 25 nm.

Onto the ^CSDCZO/^sputtYSZ/sapphire architecture, the Ba-deficient YBCO precursor solution is deposited by spin coating at 6000 rpm during 2 min. The deposition of the YBCO precursor solution is performed under dry N₂ atmosphere in order to reduce as much as possible the humidity during the deposition process. The as deposited layers are dried afterwards on a hot plate at 70 °C for 5 min.

To achieve the YBCO orthorhombic crystalline structure, the deposited layers are subjected to two thermal treatments. The first thermal treatment (pyrolysis) is performed at 310 °C in humid O₂ atmosphere [30]. The pyrolyzed films are then grown in a tubular furnace at 750 °C, under a humid atmosphere of N₂ + 200 ppm O₂ (P_O2 = 0.2 mbar and P_H2O = 23 mbar). The oxygenation of the YBCO layers was performed between 450 °C and 350 °C for more than 6.5 h using dry O₂ atmosphere.

The texture quality of the YBCO layers has been investigated by x-ray diffraction using a Bruker D8 Advance GADDS diffractometer equipped with a 2D detector, while a Bruker D8—discover was employed for the reciprocal space maps (RSM) measurements. The surface morphology of the YBCO layers has been analyzed via scanning electron microscopy performed on FEI Quanta 200 FEG microscope.

Both the sapphire substrates with sputtered YSZ and the CZO buffer layers have been characterized by means of x-ray diffraction and atomic force microscopy (AFM) to ensure the texture quality of the crystalline phases and the surface quality in terms of roughness and planarity [33]. The microscope used for this purpose was an Agilent 5100 equipped with silicon tips. Tapping mode was used to obtain the topographic images that where processed using Mountains Map Premium 8.0 software. This software allows not only the quantification of roughness, but also of the flat area fraction of the analyzed surface (planarity) [21]. The threshold value above which a surface is no longer considered flat, in this case, is 1.5 nm.

The microstructural investigation was performed using cross-sectional intermediate magnification and high-resolution transmission electron microscopy (TEM and HR-TEM). FEI Helios Nanolab 650 dual beam focused ion beam system equipped with an omniprobe manipulator was used to prepare cross-sectional TEM specimen of YBCO thin film grown on ^CSDCZO/^sputtYSZ/BPS. TEM and HR-TEM images were obtained using FEI Tecnai G2 F20 S-TWIN HR(S)TEM microscope operated at 200 kV. Gatan digital micrograph software (Gatan, Inc., Pleasanton, CA, USA) was used for image analysis and to calculate fast Fourier transform (FFT) diffraction patterns from HR-TEM images.

The performances of the YBCO layers have been investigated using a superconducting quantum interference device (SQUID) magnetometer (Quantum Design, San Diego, CA) equipped with a superconducting magnet of 7 T. The critical current density (J_C) was calculated using the Bean critical state model of a thin disk [35, 36].

The sapphire substrates from ribbons had different surface qualities, from high quality polishing (epi-polished) to unpolished. The unpolished sapphire substrates from EFG ribbons were characterized by AFM which indicated that rms roughnesses as high as ∼60 nm exist (figures 1(a) and (b)). As observed from the line profiles, outgrowths higher than 150 nm are present.

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On these unpolished substrates, the deposition of the CZO metalorganic precursor solution became inhomogeneous and achieving a continuous layer is very difficult. The growth of CZO buffer layers on top of these surfaces always led to the formation of some (111) CZO polycrystalline grains. Therefore, we conclude that the surface quality of unpolished sapphire substrates is not suitable for the growth of epitaxial ^CSDCZO buffer layers. Much smoother and therefore, plane surfaces are needed in order to achieve good quality and epitaxial CZO buffer layers. In order to improve the surface quality of the unpolished sapphire substrates, high temperature annealing treatments were performed expecting to induce planarization [33].

The unpolished sapphire substrates were thermally treated at 1100 °C and 1200 °C for 6 h and 2 h, respectively, in air atmosphere. As it can be seen in figures 1(c) and (d), the surface planarity has been greatly enhanced and the unit cell steps are visible. For comparison purposes, a line scan with a 3 µm length for an untreated sapphire substrate was introduced in figure 1(b). As it can be seen the outgrowths decrease from 150 nm to 4 and 5 nm in height. The unit step cells are detailed in figure 2. These plateaus are 600–700 nm long and exhibit a roughness of 1 nm or lower.

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On these improved sapphire surfaces, the epitaxial growth of CZO buffer layers becomes feasible, but their surface quality is still poorly reproducible due to the inhomogeneity of unpolished sapphire substrates. Even though the planarity of the substrates surfaces has been strongly enhanced, there are still outgrowths present that can't be removed even with high temperature annealing. It has become quite clear then, that some type of polishing process is needed in order to achieve high quality surfaces, and so, comply with the requirements for a proper epitaxial growth of CZO buffer layers. However, a high quality polishing process is too expensive for the long length sapphire substrates (several meters) which are required for a SCFCL. Therefore, an intermediate quality polishing process, consisting of polishing the sapphire substrates just enough that the inhomogenities cannot be seen with an optical microscope, was tested. The surface of these BPS substrates is drastically improved as compared to unpolished ones, the planarity is greatly enhanced and the cost is extremely reduced, making them suitable for scalability. As it can be seen in figures 3(a) and (b), the remaining outgrowths are not higher than 5 nm, large flat areas are present and the planarity is higher than 70%. These surfaces are then apt for the deposition of highly epitaxial CZO buffer layers [33].

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In the case of BPS substrates, epitaxial ^CSDCZO buffer layers could be easily achieved (figure 4(a)). The planarity percentage (figure 4(b)) is higher than 65%, complying with the requirements for the growth of epitaxial superconducting ^CSDYBCO layers (>60%) [32, 33]. The reproducibility in achieving this high planarity is higher than 75%.

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Achieving epitaxial and flat surfaces for ^CSDCZO buffer layers enable a higher control of the reactivity between CeO₂ and the intermediate BaF₂ phase [21]. This reactivity leads to the formation of BaCeO₃ at the interface which has a high lattice mismatch with YBCO. If this reactivity is not properly controlled, it will lead to the nucleation of misoriented YBCO grains which degrade the superconducting properties. In the case of CeO₂/epi-polished sapphire substrates, as it can be observed in figure 4(c), when the CeO₂ buffer layer exhibits a high planarity the reactivity is avoided in 200 nm thick YBCO films and it is controlled for higher film thicknesses. For those achievements, the use of Ba-deficient YBCO precursor solutions was very beneficial, since this type of metalorganic precursor solution leads to high performance YBCO layers, at lower growth temperatures than stoichiometric YBCO [37], thus allowing the nucleation of YBCO without the formation of BaCeO₃ crystalline phase.

The next step is the deposition and growth of thick CSD YBCO layers on BPS. In previous works [21, 22], we have demonstrated 1 µm YBCO growth on single crystalline substrates (LaAlO₃, SrTiO₃, CeO₂/YSZ). Being aware of the ease to generate microcracks during YBCO growth beyond a critical thickness d_c and its dependence with microstructural features [13, 18, 38], we investigated whether microcracks are generated in substrates with different degrees of polishing. With this purpose, different thicknesses of YBCO layers have been grown on unpolished, epi-polished and BPS substrates. As shown in figures 5(a)–(c), independently of the polishing quality of the substrates, microcracks parallel to the {100} planes are formed. We also observe that the microcracks density increases with the thickness of the YBCO. As earlier mentioned, the calculated critical thickness of YBCO layers on sapphire, above which the layer will present microcracks was expected at about 300 nm. In particular, CSD-TFA growth has never surpassed YBCO thicknesses of 250 nm on sapphire without microcrack formation in the YBCO layer. In order to analyze in which stage the microcracks are formed, we have investigated the different steps of the growth process of 600 nm CSD YBCO films deposited onto ^CSDCZO/^sputtYSZ/BPS substrates. The as-deposited 600 nm YBCO films are first subjected to the growth process without going through the oxygenation step. The surface morphology of the tetragonal YBCO films are then investigated by scanning electron microscopy (figures 6(a)–(e)).

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These layers exhibit a low amount of porosity and surprisingly no presence of microcracks is revealed. So, even though the ^CSDYBCO layers thickness surpasses the critical thickness by a factor of 2, no microcracks are generated on the YBCO layer grown on BPS, confirming that they do not arise from the tetragonal phase. These YBCO layers are then subjected to a separate oxygenation process at 450 °C for 3.5 h (our standard oxygenation process). After this step, we indeed detect the presence of microcracks in the orthorhombic phase YBCO layers (figure 6(f)). Therefore, we conclude that the appearance of microcracks occurs during the transition from the tetragonal to the orthorhombic phase as a result of the change in in-plane lattice parameters.

With the aim of minimizing the risk of microcrack formation, we performed the oxygenation of the YBCO layers at a lower temperature (350 °C) to slow down oxygen diffusion and strain generation [39]. Longer periods of time were then required to achieve full oxygenation (more than 6 h). This approach was proved to be successful for layers up to 400 nm in thickness. As it is seen in figures 7(a) and (b), up to these thicknesses, we are able to achieve highly epitaxial and microcrack free YBCO layers. Keep in mind that 400 nm is higher than the calculated critical thickness (∼300 nm) of YBCO layers on sapphire substrates.

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Even though we distinguish the presence of BaCeO₃, (figure 7(a)) the reactivity between BaF₂ and CeO₂ did not have any effect on the YBCO texture. The 400 nm ^CSDYBCO layers exhibit a fairly low porosity (figure 7(b)), a self-field J_C at 65 K of 3.7 MA cm⁻² and at 77 K of 1.9 MA cm⁻², as well as a sharp superconducting transition with an onset critical temperature T_C = 90 K (figures 7(c) and (d)). The self-field total critical currents at 65 K and 77 K are then I_C ∼ 150 A cm⁻¹ w⁻¹ and I_C ∼ 80 A cm⁻¹ w⁻¹, respectively. These values are close to those achieved in YBCO films grown on epi-polished sapphire by vacuum deposition techniques or CSD [1, 16, 40, 41], and by CSD on CZO/YSZ-single crystal substrates for similar thicknesses [21].

Also, in figure 7(d) we show a moderate decrease of self-field J_c(T) in 400 nm thick films grown on BPS as compared to 200 nm ones on epi-polished sapphire substrates. The I_c values are already fairly promising in view of using these elements for SFCL applications [1], thus further understanding of the microstructure and strain generation mechanism is a very relevant issue to pave the way to increase the YBCO film thickness and prevent the generation of microcracks. We have, therefore, first investigated the epitaxial growth of YBCO on ^CSDCZO/^sputtYSZ/BPS, through cross-sectional TEM observations (figure 8).

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Figure 8(a) shows an intermediate magnification cross-sectional TEM image of the combined stack comprised of YBCO grown on 15 nm of CZO buffer layer/35 nm of YSZ on BPS substrate, as indicated. We can observe in HR-TEM image (figure 8(b)) that the CZO/YSZ interface is flat and free of secondary phases or other reaction, with clear boundaries between the CZO buffer layer and YSZ. FFT pattern in the inset in figure 8(b) shows that the buffer layer of CZO grown onto YSZ/BPS is epitaxial. HR-TEM image in figure 8(c) combined with the FFT pattern of YBCO/CZO interface (inset in figure 8(c)) reveals the epitaxial c-axis growth of YBCO film on ^CSDCZO/^sputtYSZ/BPS, coinciding with the results observed in XRD data (figure 4).

We have then investigated the strain generation mechanism in the multilayer ^CSDYBCO/^CSDCZO/^sputtYSZ/BPS architecture to minimize the risk of microcrack formation when the film thicknesses are increased. To acquire this knowledge, we analyzed the cell parameter evolution at different steps of the growth process of all the layers in a 200 nm film of ^CSDYBCO/^CSDCZO/^sputtYSZ/BPS through RSM.

RSM were first performed for the ^CSDCZO/^sputtYSZ/BPS system and the results are shown in figure 9. For the sputtered YSZ layers and for the ^CSDCZO buffer layers the calculated values for the in-plane cell parameters $a$ were found to be ${a_\parallel }$ = 0.5067 nm and ${a_\parallel }$ = 0.5296 nm, respectively. When grown on sapphire substrates both of these layers become compressive (${\varepsilon _a}$ _(YSZ) = −1.03% and ${\varepsilon _a}$ _(CZO) = −1.65%, where ${\varepsilon _a} = \,\frac{{{a_f} - {a_b}}}{{{a_b}}}*100\,\left( \% \right)$ is the lattice mismatch, a_f is the film in-plane parameter and a_b is the bulk parameter) with respect to the bulk values (a_bulk = 0.512 nm for YSZ and a_bulk = 0.5385 nm for CZO). The in-plane lattice misfit of CZO versus YSZ turns out to be compressive with ${\varepsilon _a}$ = −4.8%.

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When YBCO is grown on top of this buffered sapphire substrate the cell parameters for both YSZ and CZO are modified. This has been previously observed as well by Nie et al [38]. After the oxygenation of YBCO, the cell parameters of YSZ and CZO become $a_{{\text{exp}}}^{{\text{YSZ}}}$ = 0.5106 nm and $a_{{\text{exp}}}^{{\text{CZO}}}$ = 0.5429 nm, respectively. The compressive strain for YSZ becomes smaller (${\varepsilon _a}$ _(YSZ) = −0.27%), but the growth of YBCO appears to induce a tensile strain to the CZO buffer layer (${\varepsilon _a}$ _(CZO) = +0.81%), thus the misfit of CZO versus YSZ increases to ${\varepsilon _a}$ = −6.4%. The RSM for the ^CSDCZO/^sputtYSZ/BPS films after the growth and oxygenation of ^CSDYBCO is presented in figure 10(a), along with the RSM of the (308) and (038) reflection of YBCO (figure 10(b)).

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Calculating the strain of YBCO with respect to the CZO buffer layer grown on BPS substrates we found that YBCO suffers a tensile strain in the $a$ direction (${\varepsilon _a}$ = +0.28%) and a compressive strain in the $b$ direction (${\varepsilon _b}$ = −1.3%). These values were calculated with the $d_{{\text{exp}}}^{\,{\text{CZO}}}$ = 0.3838 nm $\left(\,d\, = \,\frac{a}{{\sqrt 2 }}\,\right)$ because the YBCO lattice is rotated 45° with respect to that of CZO.

In the case of non-oxygenated YBCO, the strain in both directions $(a,\,b)\,$ is compressive with ${\varepsilon _a}$ = −0.57% and ${\varepsilon _b}$ = −2.13% (figure 11). Non-oxygenated YBCO (^NOYBCO) seems then to be compressive with respect to the CZO buffer layers grown on BPS substrates. Therefore, no microcracks are generated during growth and cooling process under an oxygen pressure of 0.2 mbar which keeps a small orthorhombic distortion.

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When the transition from the non-oxygenated phase to orthorhombic phase takes place, the oxygen doping leads to an increase of the $b$ parameter generating a compressive strain along $b$-direction, but to a decrease of $a$ lattice parameter which creates a tensile strain in the orthorhombic $a$-direction. So, when oxygenating ^CSDYBCO/^CSDCZO/^sputtYSZ/BPS, YBCO is subjected to a tensile stress with respect to the CZO buffer layer in the in-plane $a$-direction. Consequently, the microcracks are formed somewhere parallel to (100) crystallographic planes, although due to the twinned structure of YBCO, microcracks along the two in-plane orthogonal directions are observed (figure 5). As we have mentioned before, we could avoid the microcrack formation in films with thicknesses up to 400 nm by lowering the oxygenation temperature, which smoothens the kinetics of tensile strain generation. These results suggest that for higher film thicknesses further engineering of the interlayer strain generation is required in order to minimize the tensile stress in the $a$ direction.

In order to confirm our conclusion, we have performed RSM on a similar architecture without sapphire: ^CSDYBCO/^CSDCZO/YSZ single crystal. We have found that in the case of this architecture, YBCO is compressive with respect to the CZO in both in-plane directions, in the non-oxygenated and fully oxygenated orthorhombic structure. In the case of the orthorhombic YBCO, the compressive strain along $a$ and $b$ directions were ${\varepsilon _a}$ = −0.23% and ${\varepsilon _b}$ = −2.07%, respectively. This is the reason that we do not observe microcracks when growing YBCO onto this architecture, even when 1 µm of thickness is achieved [21].