Microstructure, pinning properties, and aging of CSD-grown SmBa 2 Cu 3 O 7 − δ films with and without BaHfO 3 nanoparticles

In order to improve the electrical transport properties of RE Ba 2 Cu 3 O 7 − δ nanocomposite films, SmBa 2 Cu 3 O 7 − δ films with and without BaHfO 3 nanoparticles were grown by chemical solution deposition, and their microstructural and transport properties were investigated in a detailed study using transmission electron microscopy and transport measurements in magnetic fields up to 24 T. The optimization process of the crystallization step (temperature and oxygen partial pressure) as well as an aging effect, which is due to the release of trapped fluorine, are described. Critical temperature and critical current densities surprisingly improve initially during the aging. Due to the complex microstructure, the additional BaHfO 3 nanoparticles have only a positive effect at low magnetic fields for our samples.


Introduction
Due to their exceptional performance not only at liquidnitrogen temperature but also in high magnetic fields at low temperatures or medium field and temperature ranges, REBa 2 Cu 3 O 7−δ (RE: rare earth, REBCO) phases are the compounds of choice for advanced high-temperature superconducting wires and tapes. These are the so-called coated conductors, where the functional material REBCO is deposited epitaxially on long metallic tapes with certain buffer layer architectures [1,2].
One of the REBCO compounds that gained interest in recent years is Sm 1+x Ba 2−x Cu 3 O 7−δ with superconducting transition temperature (T c ) as high as 95 K [3]. Due to the relatively large ion size of Sm 3+ , it is prone to Sm-Ba disorder and off-stoichiometries degrading T c and the transport properties. Therefore, the phase formation is somewhat challenging, yet not as severe as for LaBCO or NdBCO. Nevertheless, SmBCO films have been successfully prepared not only by pulsed laser deposition (PLD) [4,5], but also by metalorganic chemical vapour deposition (MOCVD) [6], laser CVD [7], sputtering [8], co-evaporation [9], liquid-phase epitaxy [10], and chemical solution deposition (CSD). The latter has not only been shown in the usual trifluoroacetate metalorganic decomposition (TFA-MOD) [11,12], but also with extremely-low-fluorine [13] and fluorine-free solutions [14]. In fact, SmBCO has already been used as functional material for coated conductor research and production [15,16].
In order to improve and optimize the current carrying capabilities of REBCO films and coated conductors for a certain application with a specific range in temperature and magnetic field, artificial pinning centres have to be introduced [2]. Besides mixing RE elements [17], transition metal doping [11], adjusting the stacking fault (SF) density [18] (here YBCO), multilayering and adjusting the Sm-Ba ratio [19], surface decoration [20], and irradiation [21], the best-investigated method is probably the insertion of non-superconducting nanoparticles (for reviews see [22][23][24]), such as perovskites or occasionally pyrochlores [25]. For SmBCO, the former has been intensively investigated by Yoshida et al on PLDgrown films [26]. Due to the columnar growth mode during PLD, perovskites usually self-organize in more or less c-axisoriented nanocolumns [27], leading to huge c-axis peaks in the angular dependence of the critical current density J c [26], distinct matching effects [28] as well as a usual slight T c decrease [29].
For CSD-grown films, the addition of perovskite nanoparticles has mainly been investigated for YBCO, e.g. [30][31][32], as well as GdBCO [30,33] and (Gd,Y)BCO, e.g. [34,35]. Recently, we started to systematically extend these investigations to other REBCO compounds [12]. According to this study, the most suitable substrate and the optimum deposition conditions depend on both the RE and the presence of the perovskite, here BaHfO 3 (BHO), often also BaZrO 3 (BZO).
Here, we investigate the influence of BHO nanoparticles on microstructure and current carrying capability of optimized SmBCO films grown by full-TFA MOD, which has not been studied before.

Sample preparation
The TFA-MOD precursor solutions were prepared following the recipe of [12,36]: the acetates of Sm, Ba, and Cu (>99.99%, Alfa Aesar) of ratio 1:2:3 were dissolved in water, mixed with surplus of trifluoroacetic acid (99.5+%, Alfa Aesar), and stirred until complete dissolution of the metalorganic salts at room temperature. These solutions were concentrated with a rotary evaporator, re-diluted in ultra-dry methanol (>99.9%, H 2 O < 50 ppm, Carl Roth) and filled up to the final concentration of the RE of 0.25 mol l −1 . For the nanocomposite solutions, hafnium(IV)-2,4-pentanedionate (Hf(acac) 4 , 97+%, Alfa Aesar) and an according molar amount of barium acetate, dissolved in water, were added to the solutions for nominally 12 mol% of BHO in the films. For BZO as nanoparticles, it had been reported that the amount should be limited to 10 mol% in order to avoid nanoparticle coarsening [37,38]. For BHO, our earlier (unpublished) experiments with several REBCO phases suggest a quite similar amount of 12 mol% as optimum because the coarsening is still not excessive but the pinning improvement considerable. Afterwards, minute amounts of acetylacetone (0.6 acac/RE ion or 1.5 vol% in final solution) were added to protect them from impurities such as water. The solutions were then filtered through polytetrafluoroethylene with 0.2 µm pore size and spin-coated with 6000 rpm for 30 s on cleaned 10 mm × 10 mm (001)-oriented single crystals of LaAlO 3 (LAO) resulting in a film thickness of around 220 nm. The films were heat-treated (for details see [36]) with a crystallization oxygen partial pressure p O2 of 50 or 150 ppm in nitrogen at 1 atm total pressure (dew point T dew = 19 • C) and crystallization temperatures between 780 • C and 860 • C. Afterwards, the sample cooled down in dry nitrogen with the same p O2 and gas flux as used during the film growth to 450 • C for the final oxygenation in 1 bar oxygen.

Structural characterization
Structural features of the films were measured by x-ray diffraction (XRD; D8 Discover, Bruker, Cu-Kα radiation) and scanning electron microscopy (SH-5000P, Hirox, tungsten cathode, SE-detector, 10 kV acceleration voltage). Layer thickness and surface topography were determined by atomic force microscopy (Dimension Edge, Bruker) in tapping mode on bridges used for electrical transport measurements, and occasionally cross-checked by transmission electron microscope (TEM).
Cross-section samples for scanning TEM (STEM) were prepared by an in-situ lift-out technique [39] in a FEI Strata 400S focused-ion beam (FIB)/scanning electron microscope system. High-and low-angle annular dark-field (HAADF/LAADF)-STEM images were taken with a FEI Titan 3 80-300 TEM operated at 300 kV. The HAADF-STEM image intensity is sensitive to the average atomic number Z, where a larger Z results in higher HAADF-STEM intensity for a sample with reasonably constant TEM-sample thickness. In contrast, LAADF-STEM images also show diffraction contrast and reveal more clearly crystalline defects and lattice strain. Chemical analysis with energy-dispersive x-ray spectroscopy (EDXS) was carried out with a FEI Tecnai Osiris microscope operated at 200 kV with ChemiSTEM technology [40]. The EDXS mapping data set was denoised with principal component analysis [41] before extracting the elemental signals with a peak-fitting routine using the HyperSpy Python-package [42,43].

Electromagnetic measurements
The films were scanned for their superconducting properties by inductive techniques: the transition temperature T c (defined as 10%, 50% and 90% of the normal-state value) by a selfmade mutual-inductance device, and the self-field critical current density at 77 K in a Cryoscan (THEVA, 50 µV criterion). The ∼50 µm wide and 1 mm long bridge structures for 4-point transport measurements on selected samples were photolithographically prepared with an image-reversal photoresist (AZ5214E, Microchemicals) plus wet-chemical etching in an aqueous 0.6 wt% HNO 3 solution. For lowering the contact resistances, Au pads were prepared by PLD. The electrical contacts were provided by Cu wires attached with silver paint.
Field, temperature, and orientation dependence of J c in magnetic fields up to 14 T were determined via the measurement of voltage-current (V-I) characteristics in a Physical Properties Measurement System (Quantum Design) with rotator stage and a self-made LabVIEW measurement software controlling a Nanovoltmeter 2182 A and a SourceMeter 2460 by Keithley. J c was determined with an electrical field criterion E c of 1 µV cm −1 . The N value (V ∼ I N ) was determined by fits of log(V) vs log(I) in 1-2 decades above E c .
The upper critical field B c2 and the irreversibility field B irr were determined via temperature dependent resistivity measurements, ρ(T), at constant applied fields B||c and a bias current of 100 µA. The criteria were 75% of normal state resistivity above T c and ρ irr = E c /100 A cm −2 , respectively. T c and T irr , i.e. zero-resistivity T c , were determined accordingly. For evaluating B c2 additionally, extrapolation to zero of the logarithmic derivative (dlnR/dT) −1 was used in order to be more criterion-independent, figure S1. The activation energy U 0 of thermally activated flux motion was determined by linear fits of the Arrhenius plots ln(R) vs. 1/T in the region of highest linearity.
The two top-performing samples of the pristine and the nanocomposite series were analysed further in d.c. fields up to 24 T in the 25 T cryogen-free superconducting magnet at IMR Tohoku University. For measuring the angular dependence of J c , the magnetic field B was applied in maximum Lorentz force configuration at different angles θ measured from the c-axis.

Microstructural properties
The optimized films with respect to their superconducting properties show a rather smooth and dense surface morphology, figure 1, with occasional pinholes (left image) and superficial grooves in nanocomposite films (right image). The pristine samples are in general slightly smoother, yet contain some amount of needle-like a-axis oriented grains. Lower crystallization temperatures (<800 • C) lead to an increased density and size of these a-axis grains and other needle-like outgrowths, especially for the nanocomposites and/or 150 ppm O 2 . An increase of a-axis grains with oxygen partial pressure in TFA-MOD SmBCO films was also found by Mitani et al [44]. At higher crystallization temperatures (>830 • C), round precipitates of ∼1 µm diameter (most likely CuO) are visible for the pristine films, especially at 50 ppm O 2 . The nanocomposites do not show these precipitates but occasionally contain very large foreign phases of ∼10 µm at even higher temperature (>840 • C).
The XRD θ-2θ scans, figure 2, show the expected pattern of c-axis-oriented films with only REBCO(00l) reflections besides reflections of the substrate and occasional foreign phases. The above-mentioned a-axis grains (cf the 200 reflection of SmBCO) are visible for the as-grown films next to the SmBCO 006 peak. Re-measuring the samples 1.5 years after deposition reveals an interesting aging effect by the appearance of extra peaks mainly between 20 • and 40 • , especially CuO, Sm 2 Cu 2 O 5 , and F-containing phases of the system BaF 2 -SmF 3 [45]. This is surprising. Apparently, the fluorine is not fully taken out of the film during our growth process and is either trapped in small voids as OF phase [46] (peak at 24.9 • for the as-deposited films) or presumably even partly entering the SmBCO crystal structure forming Sm 1+x Ba 2−x Cu 3 F 2y O 6.5−y+δ as described in literature for the YBCO pendant [47,48], which may release the fluorine over time. This released fluorine reacts with Sm and Ba species which leads to the appearance of other foreign phases for stoichiometric reasons. Due to possible strain and composition variations, a final identification of all phases is not possible from XRD alone. Also, neither Sm 2 Cu 2 O 5 nor the Ba-Sm-F phases are available in the Inorganic Crystal Structure Database (ICSD). However, both have been extensively described for the Y case, e.g. [49], and were here inferred from Y 2 Cu 2 O 5 and Ba-Y-F, respectively. Smith et al recently mentioned the possibility of the RE 2 Cu 2 O 5 phase for GdBCO coated conductors [50], although its stability is assumed to decrease with increasing RE ion radius and had not been reported for larger lanthanoids than Tb [51]. Noteworthy, Y 2 Cu 2 O 5 has been tried as artificial pinning centres recently in bulk YBCO samples, although with considerable T c reduction [52]. Nevertheless, also BaCuO 2 (which is in equilibrium with REBCO and RE 2 Cu 2 O 5 at high temperatures [53]) has been described by Park et al [54] as a peak triplet around 30 • for GdBCO coated conductors fabricated by reactive coevaporation. Wang et al [55] show Sm 2 CuO 4 and CuO in MOCVD-grown SmBCO films with peaks near 30 • , and Naik et al [56] identify a peak near 36 • with CuO for infiltrationgrown YBCO bulk samples. This aging leads to an increase in T c by ∼0.5 K and of the self-field J c by a factor of up to 1.5 within 2 years (supplement S2) even though the intensity of the SmBCO peaks is decreasing. The newly developed foreign phases most likely contribute to the flux pinning.
The microstructure of one of the nanocomposite films (150 ppm, 830 • C) was investigated in more detail by STEM.   figure 3). A few Sm-rich and Ba-depleted precipitates are visible at the substrate-film interface (marked by arrows). One of these precipitates was identified to consist of the Sm 2 CuO 4 phase as shown below.
The LAADF-STEM overview image of the film in figure 4(a) shows a region with BHO nanoparticles at the surface, in the film, and at the film/substrate interface marked by red, green, and blue lines, respectively. The film has a thickness of about 260 nm. However, a thinner region (marked by dotted white lines) is observed, where BHO particles are absent in the film but seem to float on the surface. Such a thin region was only observed once in the investigated TEM sample. An average equivalent diameter µ a of 20.5 nm with a standard deviation σ of 7.3 nm (figure 4(d)) was determined from all BHO nanoparticles in figure 4(a). A Sm-rich precipitate at the film/substrate interface with slightly increased intensity compared to the film is visible at higher magnification (figure 4(c)). Extended SFs, i.e. extra CuO planes, appear as dark horizontal lines (marked by yellow dashed arrows in figure 4(b)) at higher magnification. A particularly high SF density is found in the vicinity of the film/substrate interface, near BHO particles, and close to the surface, which was also observed for other REBCO films [12]. In addition, short SFs of only a few nm in length are distributed throughout the film (e.g. in the region marked by a dashed circle in figure 4(b)), which cause a local buckling of the surrounding SmBCO unit cells. We suspect that the short SFs in the shown TEM samples were induced during TEM-sample preparation as suggested by Zhu et al [57] or by TEM-sample degradation in air [58,59]. The TEM samples in this work were investigated after FIB preparation and exposure to ambient air of about 20 min during transfer between the FIB and TEM instruments. At the film/substrate interface, bright structures are visible in LAADF-STEM images (cf black arrows in figure 4(b)) which are most likely induced by strain fields [60] from steps at the film/substrate interface or dislocations.
High-resolution HAADF-STEM imaging of the film near the interface reveals an in-plane shift by a/2 (or b/2) for the first unit cell of SmBCO (inset in figure 5(a)), whereas cube-on-cube growth without such an in-plane shift was also reported for REBCO growth on LAO [57]. The HAADF-STEM image also shows the aforementioned high density of short SFs in the film (examples marked by circles in figure 5(a) Note that the Cu Lα x-ray transition was used here to extract the elemental map of Cu, which is less affected by spurious Cu x-rays than the Cu Kα x-ray transition. The high-energy Cu Kα x-rays (∼8.05 keV) can more easily escape and penetrate the Cu TEM grid, whereas the lowenergy Cu Lα x-rays (0.93 keV) mainly originate from the TEM sample.

Sample optimization
From experience, we regarded the crystallization temperature T crys and the oxygen partial pressure p O2 during crystallization as the most significant parameters. Consequently, while keeping the water vapour, i.e. the dew point T dew constant at 19 • C, we varied T crys at two oxygen partial pressures, which were promising from former studies [12] taking into account that the optimum partial pressure decreases with increasing RE ion size [17,62]. The lower oxygen partial pressure of 50 ppm yielded higher critical temperatures and self-field critical current densities compared to 150 ppm, figure 6. The optimum crystallization temperature, especially with respect to J c , lies at 820 ± 10 • C for nanocomposites at both partial pressures as well as the pristine films at 150 ppm, whereas 50 ppm O 2 allows even lower T crys < 800 • C for the pristine films. Crystallization temperatures below 780 • C were not reasonable anymore because of increasing porosity and formation of misorientations. Also for lower oxygen partial pressures, the system seems to be not stable anymore. For extreme-lowfluorine CSD-grown SmBCO films for example, 20 ppm yielded lower J c values than 50 ppm [13].
In this optimized parameter range, the resistive transition width, the zero-resistivity T c value, and the temperature dependence of the irreversibility field B irr are very similar, except for the highest T crys values at 50 ppm and the nanocomposites at 150 ppm, which show lower irreversibility lines B irr (T), figure 7.
The pinning force densities F p (B), measured 1.5 years after deposition, show some variability for the optimized parameter range, see figure 8 for the 65 K data as an example. Where the pristine samples have a maximum pinning force density F pmax of 10 ± 2 GN m −3 , the nanocomposites show slightly better values of 14 ± 3 GN m −3 . This trend is also visible for other measurement temperatures. This increase due to the BHO nanoparticles of ∼30% (best samples) compares well with earlier results on GdBCO + BHO [30], however are well below values for recent GdBCO, as well as GdYBCO, and YBCO systems, see table 1. So, the achievable increase in F pmax due to randomly oriented nanoparticles seems to depend in part also on RE ion size (most likely via the respective microstructures) besides nanoparticle size [34] and content. However, one has to take into account the much larger advance and data availability for the other systems.

Resistive transition, activation energy and phase diagram
The resistive transitions of pristine and BHO-added SmBCO films are remarkably similar, figure 9. Both kinds of samples, if optimized, show a very small zero-field transition width of ∼1 K and comparable normal state resistances of 120 Ω at 100 K (i.e. ∼130 µΩ cm). Via the Arrhenius plots of these data, the activation energy U 0 = −dlnR/d(kT) −1 of thermally activated flux motion was determined, figure 10. The BHO-SmBCO nanocomposite shows slightly larger values especially for low fields (1 T values: 4.6 10 4 K vs. 5.4 10 4 K), and both samples show roughly a square root dependence in low fields (characteristic for plastic pinning) and a linear dependence at higher fields (characteristic for collective pinning). A linearity analysis between the residual resistance near T c and U 0 yields an estimate for T c , which is 93.9 K for both samples within error bars. This value is consistent with two other criteria: maximum slope (93.5 K) and extrapolation of the logarithmic derivative (dlnR/dT) −1 to zero (94.0 K), which is related to the excess conductivity analysis (not followed here), see also supplement. This corresponds to the aforementioned 75% onset-T c resistivity criterion for B c2 .
At higher fields, both samples show a characteristic kink in their transitions (arrow in figure 9). Since this kink is not observed at low fields, inhomogeneities and grain boundary effects can be excluded as explanations. It rather characterizes the beginning of a transition region between the thermally activated flux motion and the fluctuation region near T c with strong excess conductivity. The transition region itself shows a rather linear behaviour, characteristic for flux flow. This kink behaviour has been observed in different REBCO films, e.g. [66][67][68][69], for a theoretical description see e.g. [69,70]. Figure 11 shows the B-T phase diagram of the BHO-SmBCO nanocomposite for B||c (the one for pristine SmBCO is very similar). B c2 (T) was linearly fitted with  slope dB c2 /dT| Tc = 1.50 ± 0.02 and T c = 93.9 ± 0.1 K. The werthamer-helfand-hohenberg (WHH) estimate [71] of the zero-temperature orbital upper critical field B c2 (0) = 0.69dB c2 /dT| Tc T c yields 97 ± 1.5 T and consequently a coherence length ξ ab (0) = 1.84 ± 0.01 nm. The irreversibility field could well be fitted with a power law B irr (T) = B irr (0) (1 T/T 0 ) q , with B irr (0) = 76 ± 1.6 T, T 0 = 93.1 ± 0.2 K, and an exponent q = 1.32 ± 0.02. The 95% confidence intervals for  a Calculated from 12 vol% in [34]. b Calculated from 1.5 at% in [63].   both fits are roughly the line thickness in figure 11. Within the error bars, q is 4/3, which is characteristic for a glass-liquid transition. This transition, however, could not be determined unambiguously from the present data set. B irr (0) being much smaller than B c2 (0) can have two origins: either B c2 is not orbitally limited at low T and/or B irr (T) shows an upturn at low T due to a possible extra pinning component. High-field measurements may clarify this.

Field and angular dependence of Jc
The addition of BHO nanoparticles to the SmBCO films enhances the pinning force densities and hence the critical current densities mainly at low fields with a crossover at low temperatures around 10 T, figure 12. The self-field critical current The medium-field range of increased J c and F p by BHO addition is also clearly distinguishable by an increase in N value, though in somewhat smaller ranges at high temperatures, figure 12. The N value, i.e. the exponent of power law E(J) curves E ∼ J N , is in first approximation related to the flux creep rate S ∼ = 1 N−1 . In lowest fields however, the BHO nanocomposites show somewhat lower N values for T ⩾ 50 K, which may be due to an increased inhomogeneity, i.e. in local J c variance. For medium fields and temperatures (T ⩽ 65 K), the N value seems not to depend on temperature for both samples. For the lower temperatures, and especially for the nanocomposites, a region of nearly field-independent N value (i.e. creep rate) is observed before the N value tends towards 1 for fields going to B irr .
The pinning force density curves F p (B) clearly show the above-mentioned crossover at around 10 T (at low T). This indicates a reduction in the irreversibility field by BHO addition and an already high pinning efficiency of the pristine films' microstructure. The pinning force curves of the pristine samples can be approximated by a Dew-Hughes function F p (B) ∼ B p (1 − B/B irr ) q with (p,q) ∼ (0.5,2), which suggests surface pinning or random pinning at sufficiently small defects. The parameter q is actually often slightly larger than 2 (close to 3) due to the stronger flux creep effects near B irr . The F p curves of the BHO nanocomposites, more than for the pristine films, however cannot be described well by a single Dew-Hughes function with physically reasonable combinations of B irr and q. This is due to an extra low-field component of the extended BHO nanoparticles at around 1 T. A similar behaviour has been observed recently on all-CSD-grown YBCO nanocomposite coated conductors with BZO and BHO nanoparticles [72].
The angular dependence of the critical current density and of the corresponding N value, figure 13, with respect to field orientation in fields up to 14 T of the pristine and the BHO nanocomposite film are very similar. A strong ab-peak in J c (θ) due to electronic anisotropy (random pinning and intrinsic pinning at low temperatures) as well as correlated pinning at planar defects (SFs) is clearly visible. Close to the irreversibility line, i.e. at sufficiently high fields and temperatures, a small c-axis peak is appearing; most likely due to extended planar (anti-phase boundaries, twin boundaries, domain boundaries) or linear defects (dislocation networks) parallel to the c-axis. At medium fields and low temperatures-more for the pristine sample-J c (θ) shows a dip near B||c. This also can be attributed to planar defects, namely vortex channelling effects at twin boundaries [73]. The BHO particles hence reduce this channelling effect either directly by increased pinning or indirectly by leading to a reduced density of extended twin boundaries. By BHO nanoparticles being effective in a wide angular range around B||c, cf figure 12, the J c anisotropy is decreased by BHO incorporation at all temperatures. Figure 14 compares the J c anisotropy at 24 T for several temperatures between the pristine and the BHO-nanocomposite film. Above 30 K, both films behave surprisingly similar.
Further information about the present pinning species and mechanisms are gained by the study of the N value anisotropy. First of all, it usually enhances features due to its averaging effect, see e.g. the extra pinning component of additional medium angle peaks in N(θ) (e.g. peak around 135 • at 65 K, 3 T), which are not or barely visible in J c (θ) and due to the sparsely distributed nanoparticles larger than ξ. In these regions, N scales well with J c , since both depend on the pinning potential U 0 , and the pinning behaviour is dominated by the elastic properties of the flux lines. In contrast, at sufficiently low temperatures, N(θ) shows a strong dip whereas J c (θ) has a maximum. Here, N and J c are anticorrelated and the creep rate is not determined by the pinning potential anymore but influenced by a second creep mechanism. In this region, the flux lines build staircase patterns of segments strongly pinned (trapped) within the ab-planes/ ab-planar defects and weakly pinned segments with oblique angle to the ab-planes. The easy movement of the weakly pinned segments leads to a net hopping behaviour of the trapped segments from one planar pinning centre to the next, increasing the creep rate and hence decreasing the N value. At the lowest temperatures (see 30 K) and even smaller angles between B and the ab-planes, the whole flux lines will be trapped in the planar pinning centres, which hinders the flux lines from the net hopping movement, thus decreasing the creep rate, and hence increasing the N value. As is apparent in figure 13, especially at low T, the addition of BHO nanoparticles increases the widths of these regions. This is due to the shortening of the trapped segments due to the direct interaction with the nanoparticles and/or a shortening of the planar defects. Interestingly, the trapping angle seems rather field-independent for the pristine sample at medium temperatures, while a B −3/4 dependence would be expected by theory. This had been observed also for NdFeAs(O,F) films and was explained by the vicinity of the 2D-3D transition [74]. Finally, we would like to point the attention to the slight c-axis peak in N(θ) at 14 and 24 T, 30 K for the pristine sample, indicating a

Conclusion
The growth of SmBa 2 Cu 3 O 7−δ films with and without BHO nanoparticles by CSD was investigated and optimized regarding the crystallization step. Crystallization temperatures around 810 • C and a low oxygen partial pressure of 50 ppm yield the best results regarding T c and self-field J c . With XRD as well as a detailed TEM study, a complex microstructure with several foreign phases, such as CuO, Sm 2 CuO 4 , and Sm 2 Cu 2 O 5 , was detected. Due to the release of trapped fluorine (either within the crystal structure or as OF phase in voids), the samples undergo an aging process with distinct changes in microstructure and-at least initially-an improvement in T c and J c . The transport properties, measured in fields up to 24 T revealed a positive effect of the BHO nanoparticles only in small fields up to ∼10 T, which is due to the already rich pinning landscape in the pristine films. The exact microstructural reasons for the aging process as well as an optimization of nanoparticle insertion for pinning improvement may be tasks of future studies.

Data availability statement
TEM images and data analysis for figures 3 and 4 are available at Zenodo [75].
The data that support the findings of this study are available upon reasonable request from the authors.