Superconductive REBCO Thin Films and Their Nanocomposites: The Role of Rare-Earth Oxides in Promoting Sustainable Energy

The lossless transmission of direct electrical currents in superconductors is very often regarded as an “energy superhighway” with greatly enhanced efﬁciency. With the discovery of high temperature superconductors (HTS) in the late eighties, the prospect of using these materials in efﬁcient and advanced technological applications became very prominent. The elevated operating temperatures as compared to low temperature superconductors (LTS), relaxing cooling requirements, and the gradual development of facile synthesis processes raised hopes for a broad breakthrough of superconductor technology. The impact of superconductor technology on the economy and energy sectors is predicted to be huge if these are utilized on a large scale. The development of superconducting tapes with high critical current density ( J c ) is crucial for their use in transmission cables. Many countries these days are running projects to develop wires from these HTS materials and simultaneously ﬁeld trials are being conducted to assess the feasibility of this technology. These HTS wires can carry electrical currents more than 100 times larger than their conventional counterparts with minimal loss of energy. The increased efﬁciency of HTS electric power products may result in greatly reduced carbon emissions compared to those resulting from using the conventional alternatives. In order to use the thin ﬁlms of YBa 2 Cu 3 O 7 − δ (YBCO) and RE BCO [ RE (rare-earth) = Sm, Gd, Eu etc.], members of the HTS family, for future technological applications, the enhancement of J c over wide range of temperatures and applied magnetic ﬁelds is highly desired. The enhancement of J c of YBCO and RE BCO ﬁlms has been successfully demonstrated by employing different techniques which include doping by rare-earth atoms, incorporating nanoscale secondary phase inclusions into the RE BCO ﬁlm matrix, decoration of the substrate surface etc. which generate artiﬁcial pinning centers (APCs). In this review, the development of the materials engineering aspect that has been conducted over the last two decades to improve the current carrying capability of HTS thin ﬁlms is presented. The effect of controlled incorporation of APCs through various methods and techniques on the superconducting properties of YBCO and RE BCO thin ﬁlms is presented, heading toward superior performance of such superconducting thin ﬁlms.


INTRODUCTION
Superconductivity is a phenomenon in which the material in consideration has virtually no resistance to direct electric current. This phenomenon takes place below a certain characteristic temperature, known as critical temperature, or transition temperature (T c ), which is different for different superconductors. Ever since its discovery by Onnes [1], it has been a subject of great interest because of the unique properties it exhibits. In addition to transport of electrical current without any resistance, the superconducting state is also characterized by perfect diamagnetism (χ m = −1). The magnetic flux is expelled from the interior of the superconductor at T < T c . The screening currents are induced in the surface layer of the superconductor which generates a flux density opposite to that of the applied magnetic field. This phenomenon is known as Meissner effect [2].

Types of Superconductors
On the basis of their magnetic properties, superconductors have been broadly classified into two classes: (a) Type I Superconductor: Type I superconductors expel the magnetic field completely for applied magnetic field (H) smaller than the critical magnetic field (H c ). Above H c , the material abruptly goes back into the normal state. The variation of magnetization (M) with respect to H for type I and type II superconductors is presented in Figure 1A.
Most of the elemental superconductors (e.g., Pb, Hg, Sn, etc.) exhibit type I superconductivity except Nb, V, and Tc [3]. In addition, type I superconductivity is also exhibited by an alloy, TaSi 2 [4] and a compound, SiC (with heavy doping of B) [5]. The T c s for type I superconductors are normally lower (< 10 K). The values of H c for type I superconductors are in the range of 5-200 mT [3]. Because of their low T c and H c values, type I superconductors are of limited use. (b) Type II Superconductor: Type II superconductors have two critical magnetic fields: H c1 and H c2 as depicted in Figure 1A. Below H c1 , the superconductor remains in the Meissner state, completely expelling the magnetic flux from its interior. For H > H c1 but smaller than H c2 , the magnetic flux starts penetrating the sample in the form of discrete bundles termed "flux lines" and the sample goes into the mixed state (or vortex state). When H becomes > H c2 , the superconductor comes into the normal state.
In the vortex state, the specimen comprises of alternating normal and superconducting regions. The normal cores are surrounded by superconducting regions allowing some magnetic field penetration. Type II superconductivity is exhibited by metallic compounds, alloys, and complex oxide ceramics. Type II superconductors have much higher critical magnetic fields as compared to type I superconductors. Also, Type II superconductors can carry much larger current densities while remaining in the superconducting state. Due to the above advantages, type II superconductors have greater potential for practical applications.

Flux Pinning in a Superconductor
Apart from the critical temperature (T c ), the critical current density (J c ) is the most relevant parameter of a superconductor which is directly related to its possible use in practical applications [6][7][8]. J c in a superconductor is determined by the critical temperature, electronic structure, and the flux pinning mechanism governed by the microscopic defects that are generated naturally or artificially during the growth of the superconducting films. The upper limit of the J c is determined by splitting of the paired electrons that carry the supercurrent (the so-called Cooper pairs), and Ginzburg-Landau gave an equation to estimate the depairing current density (current density at which splitting of Cooper pairs takes place) J d [9]: where Φ 0 , µ 0 , λ, and ξ are the flux quantum, permeability in vacuum, the London penetration depth and the Pippard coherence length, respectively. J c cannot exceed this value even by the optimum vortex pinning [10]. For YBa 2 Cu 3 O 7−δ (YBCO) and REBCO [RE (rare-earth) = Sm, Gd, Eu etc.] films, J d at 77 K and zero applied field is estimated to be 40-50 MAcm −2 [11]. This value is quite large but the observed values of J c are limited to about 10% of J d . Thus, there is large room for sufficient enhancement of J c by employing different methodologies. Figure 1B shows the H-T phase diagrams for type I and type II superconductors. In the mixed state (H c1 ≤H≤H c2 ) of a type II superconductor, the magnetic flux penetrates into the superconducting specimen in the form of small "tubes" called vortices. These vortices have a normal-conducting core and are surrounded by circulating supercurrents generating a magnetic flux quantum Φ 0 = h/2e (h being Planck's constant and e being electronic charge). The circular currents make any two vortices repel each other forming an ordered hexagonal lattice called the Abrikosov vortex lattice. If an electrical current is passed through a superconductor in its mixed state, the vortices would experience a Lorentz force, whose density is given by F L = J × µ 0 H. Due to the influence of this Lorentz force, the vortices start to move in a direction perpendicular to the directions of both the transport current and the applied magnetic field. A schematic representation of this situation is presented in Figure 2. There are, however, some kinds of defects or impurities in superconducting samples, such as dislocations, voids, grain boundaries, etc. which act as pinning centers for vortices and the magnetic flux gets trapped. The force which resists the motion of the vortices under the influence of the Lorentz force is called the pinning force, whose density is termed as pinning forced density (F p ). The flux lines remain stationary, as long as F p is > F L . When F L exceeds F p , vortices start moving across the superconductor. If the vortices move with a velocity v, an electric field E = µ 0 H × v would be generated. Since, both the current and the generated electric field would be parallel, a finite power would be dissipated in the system and the superconductor would lose its ability to sustain dissipation-free current flow.
The penetration of magnetic flux in a type II superconductor is gradual over a wide range of applied magnetic field. However,  the presence of lattice defects prevents the easy entry or exit of the flux lines and the magnetization is irreversible. The presence of lattice defects modifies the vortex structure; vortices may be "pinned" down at the defect sites and no longer free to move. From the energetics point of view, the defect sites are surrounded by an energy barrier which the pinned vortex must climb before it can move. The Lorentz force effectively lowers this barrier and the critical current density of a specimen would reach when the pinning force is balanced by the Lorentz force. At finite temperatures, however, thermal activation also lowers the effective barrier height and gives rise to a strongly temperature dependent critical current.

YBa 2 Cu 3 O 7−δ : A HIGH TEMPERATURE SUPERCONDUCTOR
YBCO belongs to the class of type II superconductors, which are known for their ability to maintain the superconducting properties even at higher applied magnetic fields. YBCO was the first superconductor which was shown to have T c above the liquid nitrogen temperature (77 K) [12]. The discovery of YBCO fuelled lot of research activities in the field of superconductivity and most of the earlier works were focused on discovering the superconductor which could have an even higher T c . Subsequently several other superconductors such as Bi 2 Sr 2 Ca 2 Cu 3 O 10+δ [13], Tl 2 Ba 2 Ca 2 Cu 3 O 10 [14], and HgBa 2 Ca 2 Cu 3 O 8+δ [15] have been discovered which were found to have T c s above 77 K. However, among all other cuprate superconductors, YBCO has several advantages over others which include many facile synthesis routes for thin film fabrication and much larger irreversibility field (H irr ).
H irr of a superconductor is another very important parameter which is defined as the magnetic field at which the resistivity value due to flux motion is significant and the pinning strength goes to zero [16,17]. The irreversibility line divides the H-T phase diagram of a superconductor into two portions: reversible and irreversible (or hysteretic). The reversible portion of this phase diagram refers to the vortex-liquid state in which vortex movements due to thermal fluctuation are so high that the ordered vortex lattice state (referred by irreversible portion of the phase diagram) is destroyed [18,19]. Thus, the irreversibility line or the melting line demarcates the solidliquid phases of the vortex matter and it is highly desired to shift this line toward higher H-T regime by artificial pinning center (APC) technology [6,8,20,21]. Figure 3 shows the comparison of the irreversibility lines of various superconductors which include superconducting alloys, metallic compounds, MgB 2 , and high temperature superconductors (HTS). The coated conductor technology employing APC incorporation in superconducting REBCO matrix is set to usher new frontiers of superconductivity applications between a wide temperature range of 5-77 K. Following subsections present different properties of YBCO superconductor.

Crystal Structure of YBCO
The crystal structure of YBCO is a complex variation of the perovskite structure [22], which is shown in Figure 4. As shown in the figure, the YBCO unit cell consists of an YCuO 3 cube with adjacent BaCuO 3 cubes above and below, but with some oxygen sites not occupied. The oxygen sites on the same horizontal plane as the Y atom are never occupied, which causes the oxygen atoms to move slightly toward the Y atom. The orthorhombic phase of YBa 2 Cu 3 O 7−δ has lattice parameters, a = 0.382 nm, b = 0.388 nm, and c = 1.168 nm when δ is very small. The oxygen content in YBCO determines its crystal structure and the hole concentration in CuO 2 planes. For δ = 1, the compound (YBa 2 Cu 3 O 6 ) has the tetragonal structure and it is an insulator. Increasing the oxygen content up to δ = 0.4, the compound undergoes a phase transition from tetragonal to orthorhombic and the Y-Ba-Cu-O system becomes superconducting. T c approaches its maximum value of 92 K for δ ≈ 0.06 [23] which is ascribed to the optimum hole doping. For δ < 0.06, T c is found to decrease which is attributed to the overdoped state of the phase in which the concentration of the holes in CuO 2 planes exceed the optimum limit. The formation of tetragonal phase is observed in the temperature range of 700-900 • C and the orthorhombic phase is formed when the tetragonal phase is slowly cooled in an oxygen atmosphere at ≈ 550 • C. The transition from tetragonal to orthorhombic phase creates a large number of different twin domains because of the release of stress in the material. In the tetragonal phase, the oxygen atoms randomly occupy about half of their respective sites in the basal planes whereas they are ordered along the b-direction into Cu-O chains in the orthorhombic phase. This creates oxygen vacancies along the a-direction in the orthorhombic phase which subsequently leads to slight compression of the unit cell so that a < b. The contribution to superconductivity comes both from the CuO 2 planes and CuO chains in the orthorhombic phase.
As shown in Figure 4, the crystal structure of YBCO is highly anisotropic. This anisotropy is observed in other superconducting properties also, such as energy gap ( ), coherence length (ξ ), and penetration depth (λ). The electrical conduction in YBCO, like other HTS cuprate superconductors, is also highly anisotropic with conductivity along the abplane being much higher than along the c-axis [24,25]. The transport of electrical currents is conducted by holes induced in the oxygen sites of the CuO 2 planes. The oxygen occupancy at the chain site influences the carrier density in the CuO 2 planes and subsequently the macroscopic electronic properties of YBCO superconductor. The anisotropic character of YBCO is observed in its characteristics both in normal and superconductive states. The properties along the ab-plane, parallel to the layers, are very different compared to along the c-axis, which is normal to the layers. Apart from other superconducting properties, J c also exhibits strong anisotropy being higher when the applied magnetic field is oriented along the ab-plane and minimum when the applied magnetic field is oriented along the c-axis of the superconducting sample.

The Evolution of Critical Current Density of YBCO and REBCO Superconductors Over Time
Superconductivity in YBCO was discovered in its polycrystalline bulk sample. Although it has T c ∼ 92 K and relatively large H irr (∼ 7 T at 77 K) in its bulk polycrystalline form, the values of J c were observed to be ∼ 10 2 A/cm 2 (77 K, self-field) [26,27], which is not high enough as compared to metallic copper. Such a low value of J c was attributed to low coherence length of YBCO, which results in limited percolation of electrical current across grain boundaries [28]. The alignment of the grains in bulk YBCO and REBCO materials was improved by melt-texturing-growth (MTG) technique [29]. This new approach resulted in much higher J c ∼ 10 4 A/cm 2 (77 K, self-field), two orders of magnitude larger than for polycrystalline bulk samples of YBCO. The recent development of infiltration growth technique has resulted in superior J c performance (∼ 10 5 A/cm 2 ; 77 K, self-field) of YBCO and REBCO melt-textured bulk samples [30][31][32][33]. In addition, the MTG REBCO superconductors exhibited superior capacities to trap very high magnetic fields (∼16-17 T at 25-30 K) [34,35] which could be very useful for permanent magnet applications.
Within a couple of years of its discovery in the polycrystalline bulk sample, it became possible to make thin films of YBCO on single crystal substrates such as SrTiO 3 , Al 2 O 3 , MgO, etc. [36]. Highly c-axis oriented YBCO films on single crystals exhibited J c > 10 6 A/cm 2 (77 K, self-field) which was again two orders of magnitude larger than the MTG YBCO samples. In a recent work, highly oriented YBCO thin films on different single crystal substrates were deposited and the role of the substrate material and thermal contact during the deposition are carefully investigated [37]. Thus, with the combined progress in the material and its processing, the current carrying capability of REBCO superconductors improved significantly. Highly c-axis oriented YBCO films deposited on single-crystal substrates by different techniques such as pulsed laser deposition (PLD) [38], chemical solution deposition (CSD) [39] and metal organic chemical vapor deposition (MOCVD) [40] exhibited high J c of 1-5 MA/cm 2 at 77 K, self-field [41,42]. The high value of J c was attributed to various crystal defects which will be discussed in the next section. The distribution and density of such crystal defects, however, is very difficult to control, which is very much needed for in-field enhancement of J c . The challenge of improving J c of YBCO thin films under applied magnetic field remained unresolved for almost a decade.
After successful demonstration of high critical current in epitaxial YBCO films on single crystal substrates, it was desired to develop the technique of making these epitaxial films on flexible metal tapes for technological applications. Subsequently, the development of second-generation wires and tapes or so called coated conductors was carried out by ensuring both in-plane and out-of-plane texturing of REBCO superconducting layers on buffered metal substrates in a biaxial alignment [43]. In order to address this issue, two approaches were employed: ion beamassisted deposition (IBAD) [44] and rolling-assisted biaxially textured substrates (RABiTS) [45]. In IBAD technique, biaxially oriented buffer layer on polycrystalline metallic substrate is deposited before the deposition of superconducting film. Although IBAD process provides excellent in-plane texture; it involves high equipment costs and is a time-consuming process. Another approach which has been used for the texturing of metal tapes is inclined substrate deposition (ISD) in which textured film of MgO is deposited by electron beam evaporation [46,47]. The ISD method provides higher deposition rates as compared to IBAD and is independent of recrystallization properties of the metallic substrates. In RABiTS, biaxial texturing is carried out through cold rolling and recrystallization of metallic substrate (Ni). By continuous development of the fabrication methods involving IBAD, ISD, and RABiTS, it was possible to get high self-field J c in the coated conductors as well.

VORTEX PINNING IN YBCO AND REBCO THIN FILMS: NATURAL PINNING CENTERS
The defects which are naturally generated during the growth of a superconducting YBCO and REBCO thin films can act as pinning centers are as follows [8]: (a) Point defects: Point defects are a kind of imperfection which occur when the crystal structure is disrupted on atomic scale. It can be due to impurities, vacancies or interstitials (missing or extra atoms). When YBCO is doped with rareearth elements like Nd, Sm, and Eu [48,49], yttrium and barium atoms sometimes exchange their places which also results in point defects. (b) Voids: During the deposition of the thin film, the vapors of the material are unable to fill the valleys because of the shadowing effect or its reduced mobility. This causes surface roughness in the thin film. When deposition and crystallization of the film is done in separate steps (ex-situ process), pores and voids may result during the crystallization step due to volumetric changes. (c) Misfit dislocations: Due to the lattice mismatch between REBCO and the substrate on which it is deposited, strain is developed. But when the strain is more than a limit, dislocations are generated in the form of missing (or extra) half-plane of atoms to relieve the stress. phase of YBCO such as the copper rich 124 phase leads to the generation of planar defects. The layered structure of YBCO is sometimes disrupted by an extra or missing layer of, for example, CuO 2 planes, which results in another kind of planar defect called the stacking fault [50]. (f) Grain boundaries: Grain boundary is one of the most common crystal defects in thin films which separates regions of different orientation of the crystals. The non-uniform growth of the solids during crystallization process results in the formation of grain boundaries. If the angle between the adjacent grain boundaries is small, it is effective for pinning the vortices [51,52]. However, if the angle between grain boundaries is larger, it results in weak coupling between the adjacent grains and J c decreases rapidly. In all REBCO films, particularly those which are deposited on coated-conductor substrates that inherently have imperfect crystalline structures, the alignment of a-and b-axes deviates. At the boundaries, where misaligned grains meet, atomic order is disrupted resulting in strain and dislocations which provide pinning to the vortices. (g) Twin boundaries: The crystal structure of REBCO materials, in their superconducting state, is orthorhombic. During the deposition of REBCO films on crystalline substrates, when c-axis of the material grows perpendicular to the substrate, domains with perpendicular a-and b-axes orientations form. The boundaries where these domains meet are called twin boundaries, which are also a kind of planar defect [53,54]. (h) Antiphase boundaries: YBCO has a layered structure in which appropriately oxygenated yttrium, barium and copper layers are arranged in a particular order. During the process of growth of the film, domains coalesce with the layers matching yttrium-to-yttrium, barium-to-barium, and copper-to-copper. Sometimes, however, this sequence is disrupted and a boundary between imperfectly matched domains is generated which is called an antiphase boundary [55]. (i) Threading dislocations: During the growth of the film, the dislocations between misoriented grains may also grow simultaneously and run entirely through the film thickness. Sometimes, dislocations form between the growth islands due to lattice mismatch between the substrate and REBCO phase [56,57].

VORTEX PINNING IN YBCO AND REBCO THIN FILMS: ARTIFICIAL PINNING CENTERS (APCs)
It has been discussed earlier that the immobilization of vortices is required to achieve high J c in the presence of large magnetic field. Due to the presence of naturally occurring defects which are described in the previous section, the vortices are pinned in REBCO films. However, the pinning efficiency of these naturally occurring defects against thermal fluctuations is not sufficient to sustain necessary level of J c at high applied magnetic fields [58,59]. It has been, therefore, a subject of great interest to improve the J c values of YBCO and REBCO thin films by introducing additional defects into the superconducting matrix. There are many methods which have been applied to introduce APCs into YBCO and REBCO superconductors. These can be classified into three main categories: Doping of Rare-Earth Elements (Addition and/or Substitution) Different rare-earths in REBCO have different ionic-radii and the effect of ionic-radii on the T c of REBCO superconductors has been reported in an earlier study [60]. It has been observed that the T c varies linearly with the ionic-radius of RE ions which was attributed to the strain-induced charge redistribution between the charge reservoir (CuO-chains) and the CuO 2 planes. The doping of Ho in YBCO melt-textured samples was conducted and significant enhancement in the in-field J c was observed in the Ho-doped samples [61]. In order to improve the vortex pinning properties of YBCO superconducting thin films, several rare-earth elements such as Sm, Eu, Nd have been doped in place of Y with different molar cationic ratio [62]. Figure 5 shows the enhanced in-field J c by up to a factor of 3 as observed in Y 2/3 Sm 1/3 Ba 2 Cu 3 O 7−δ films deposited by PLD technique on single crystal substrates. The enhanced infield J c in the doped films has been attributed to the additional random defects and tilted linear defects present in the doped sample. However, the enhanced field-dependent J c behavior of Y x RE 1−x BCO compounds as compared to that of YBCO is not very clearly understood.
According to some reports [63][64][65][66], in Y-Ba-Cu-O compound, the Y atom is completely replaced by another rare-earth atom or combination of 2 or more rare-earth atoms, which resulted in improved vortex pinning. In order to see if there is additional enhancement resulting from the strain induced by lattice mismatch when mixtures of rare-earth elements were used instead of a single rare-earth element, various combinations were reported which include (Gd 0.8 Er 0.2 ) [65] and (Nd 1/3 Gd 1/3 Eu 1/3 ) [66]. However, in all cases, the enhancement was not remarkable except for the case when the defects were present randomly and not correlated.
Several other elements (Tb, Ce, Pr, Nd, La, Co, Dy, and Eu) have been attempted for substitution at the Y site of YBCO and RE site of REBCO [49,[67][68][69][70] films. These substitutions resulted in enhanced J c and F p values over a broad range of applied magnetic fields, which was attributed to increased density of nanoprecipitates of these substituents in doped REBCO films as compared to pristine REBCO film which subsequently resulted in stress field due to lattice mismatch between the phases in the resulting REBCO films.

Nanoscale Secondary Phase Inclusions/REBCO Based Nanocomposites
Another way of introducing APCs into REBCO superconductors is the incorporation of nanoscale non-superconducting secondary phase materials into the superconducting matrix. There are several reports in the recent past in which the vortex pinning properties of REBCO superconductors has been improved by intentionally adding non-superconducting secondary phase nanoparticles. Incorporation of different secondary phases such as Al 2 O 3 [71], TiO 2 [72], WO 3 [73], BaTiO 3 [74], BaZrO 3 [75,76] etc. into YBCO polycrystalline bulk samples have been reported in the past which resulted in the enhancement of in-field J c of YBCO bulk samples. The melttextured bulk samples of REBCO superconductors consisting of precipitates of secondary phases such as (RE) 2 BaCuO 5 (RE-211) [77,78], NbO 3 , MoO 3 [79] etc. have also exhibited superior vortex pinning properties resulting in enhanced in-field J c and H irr values.
The incorporation of BaZrO 3 (BZO) nanostructures into YBCO thin films was reported in an earlier study [80] in which a composite (premixed YBCO:BZO) ablation target was used in the PLD technique. It was found that although BZO nanoinclusions were distributed randomly throughout the YBCO matrix, it produced a significant c-axis correlated enhancement of J c . In another report [81], yttria-stabilized zirconia (YSZ) was added to the YBCO target, which led to the formation of BZO nanostructures in the as deposited YBCO thin film, presumably leaving a Ba-deficient YBCO film matrix. In this composite film, a self-assembly of vertical arrays of BZO phase was observed. These vertical arrays of self-assembled BZO phase were supposed to arise from the preferential nucleation of the impurity islands in the strain field above the impurity particles [82]. These selfassembled vertical arrays of BZO phase resulted in strong pinning of vortices especially when the applied magnetic field is along the c-axis. Goyal et al. [83] have also reported a strong c-axis correlated pinning enhancement in YBCO:BZO nanocomposite thin films on RABiTS.
The incorporation of BZO nanoinclusions into YBCO thin films has been conducted in many other studies which include study of temperature dependence of J c for YBCO and YBCO:BZO thin films [84] and field dependence of the optimum concentration of BZO nanoinclusions for effective pinning of vortices [85]. Figures 6A-C shows the cross-sectional and planar views of the microstructure of YBCO:BZO thin films, deposited on IBAD-MgO templates, with varying concentrations of BZO. The variation of the average spacing between the BZO nanorods and the corresponding matching field (B ϕ ) with respect to BZO concentration is also shown in Figure 6D. As shown in Figure 7, lower concentrations of BZO nanoinclusions are more effective at 77 K, exhibiting higher F pmax values. However, at 65 K and higher magnetic field, the performance of the YBCO:BZO films with higher concentration of BZO nanoinclusions is enhanced as compared to YBCO:BZO films with low BZO content.
The CSD method has also been employed to prepare YBCO:BZO nanocomposite thin films which exhibited strongly enhanced vortex pinning [86]. Unlike the previous results in which films were deposited using PLD technique, the solution process was found to produce non-epitaxial secondary phase particles in the film matrix surrounded by many crystalline defects. Figure 8 shows the microstructure of the YBCO:BZO nanocomposite thin film using transmission electron microscopy (TEM) in which the separate phase of BZO nanoparticles and crystalline defects surrounding them can be clearly observed. The formation of BZO nanoparticles inside YBCO matrix resulted in isotropic pinning characteristics of these films [87,88]. The significant improvement in the value of J c and F p of the YBCO:BZO nanocomposite thin films can be seen in Figure 9.
By adopting a different approach in PLD technique, Haugan et al. [89] have successfully incorporated a nonsuperconducting phase, Y 2 BaCuO 5 (Y211), into YBCO thin film in a controlled manner. Two different targets, one of YBCO and another of Y211, were used and by depositing thin YBCO layers and discontinuous Y211 layers alternately, a pancake-like array of precipitates were formed in the resulting film matrix. In this case, enhanced vortex pinning properties was found for both the field orientations: when H was parallel to the c-axis and also for H parallel to the ab-plane [90]. Very recently it was demonstrated that these Y211 nanoparticles are effective not only in increasing the in-field J c but also in reducing the critical current anisotropy [91].
Apart from BZO and Y211, the nanostructures of other materials such as Y 2 O 3 (YO) [92], BaSnO 3 (BSO) [93][94][95], and BaTiO 3 (BTO) [96] have also been successfully incorporated into YBCO thin films using the PLD technique. In all these cases, the enhancement in the value of J c was more prominent at higher applied magnetic field. In one of the reports [95], the F pmax value for YBCO:BSO nanocomposite film turned out to be 28.3 GNm −3 , a record of that time, which reflects the excellent in-field performance of J c . The incorporation of a double-perovskite material, YBa 2 NbO 6 (YBNO), was also reported and it was found that YBNO phase grows in the form of columnar nanostructures [97]. In another report, the YBNO nanocolumns were introduced into YBCO thin films by surface modified target method in which YBNO concentration inside YBCO thin film was controlled by controlling the target rotation speed [98]. YBNO nanocolumns were observed to be very effective in enhancing the in-field J c of YBCO thin films.  Rare-earth tantalates (RE 3 TaO 7 , REBa 2 TaO 6 ) also turned out to be excellent secondary phase materials which were incorporated inside REBCO superconducting matrix for enhancing in-field J c of REBCO thin films [99,100]. It has been suggested that the lattice mismatch is appropriate for superior vortex pining if it is in the range of 5-12% [100]. The nanoinclusions of the rare-earth tantalate and niobate have simultaneously been incorporated into YBCO thin films which significantly enhanced the J c of YBCO films and improved the angular anisotropy of J c [101]. Extremely high F pmax values have been observed in YBCO films with YBa 2 (Nb/Ta)O 6 (YBNTO) nanoinclusions [102]. Figure 10 shows the microstructure of the YBCO film with YBNTO nanoinclusions. Apart from YBNTO nanocolumns, other pinning centers such as Y 2 O 3 nanoparticles and planar defects can also be observed in this figure. Figure 11 shows the variation of J c and F p , at 77 K and 30 K, for YBCO films with YBNTO nanoinclusions deposited at different frequencies in the PLD system. At 77 K, F pmax exceeded 25 GNm −3 and at 30 K, it exceeded 300 GNm −3 which can be seen as excellent in-field J c performance of these YBCO nanocomposite thin films.
More recently BaHfO 3 (BHO) has emerged out as a very promising secondary phase material whose nanoinclusions in the form of columnar or spherical structures inside REBCO matrix resulted in significantly enhanced in-field J c of REBCO thin films deposited on both single crystals and metallic tapes [103][104][105][106][107][108][109][110]. By adopting low-temperature growth technique in PLD, SmBCO:BHO films have been demonstrated to exhibit very high pinning force density (F pmax = 28 GN/m 3 ) at 77 K for H parallel to the c-axis [107]. Even on metallic tapes, GdBCO:BHO nanocomposite films exhibited large values of pinning force density (F pmax = 23.5 GN/m 3 ) and irreversibility field (µ 0 H irr = 15.8 T) at 77 K for H parallel to the c-axis [108]. The CSD approach has also been employed to incorporate BHO nanoparticles into YBCO [109] and GdBCO [110] thin films, which resulted in the enhanced in-field J c values of the nanocomposite thin films.
By adopting a different synthesis route called MOCVD, researchers at University of Houston in USA have demonstrated that the incorporation of BZO in high volume fractions results in outstanding in-field J c performance of REBCO films which are relatively thicker (∼ 1-2 µm) [111][112][113][114][115][116]. The F pmax value of one of these composite films reached a record value of 1.7 TN/m 3 at 4.2 K together with extremely high µ 0 H irr ∼ 14.8 T (at 77 K), a value that is much higher than the µ 0 H irr ∼ 11 T of NbTi superconductor at 4.2 K [111]. Figure 12 shows the cross-sectional TEM image of a heavily doped (Gd,Y)BCO superconducting film on IBAD substrate prepared by MOCVD technique in which c-axis oriented self-assembled BZO nanocolumns can be clearly observed. Figure 13 shows the variation of J c and F p with respect to applied magnetic field for the films with different concentration of Zr inclusions (in mol %). It can be observed that in heavily doped film, F pmax exceed the value of 1 TNm −3 which is excellent in terms of critical current performance.
The so-called quasi-multilayer (multilayers in which layer/s of one phase is incomplete) approach has also been employed for improving vortex pinning in YBCO thin films. The quasimultilayers of YBCO with different materials such as Y 2 O 3 [117,118], BZO [119], and transition metals such as Ir [120] and Ti, Zr, Hf [121] and prepared by PLD technique have been investigated. The transition metals when incorporated into YBCO thin films, through this approach, have been observed to form BaMO 3 (M = Ti, Zr, Hf, Ir) phase. Not only enhancement in the infield J c of YBCO based quasi-multilayers was observed, but the irreversibility line was also observed to shift toward higher H-T regime [119].
The materials, whose precipitates or nanoparticles have been mentioned so far, for their use as secondary phase inclusions in the YBCO/REBCO thin films, are mostly insulating and nonmagnetic in nature. However, there have been successful attempts of using ferromagnetic secondary phase inclusions as a source of APCs in YBCO thin films. A thin layer of Fe 2 O 3 has been used either as a cap layer on YBCO thin film or as a buffer layer on STO substrate before YBCO film deposition [122,123]. It has been observed that the samples consisting of Fe 2 O 3 cap layer have exhibited significant enhancement in the J c (both self-field and in-field) of YBCO thin films. In addition, other ferromagnetic nanoinclusions such as of YFe 3 O 4 [124] and CoFe 2 O 4 [125] have also been successfully used to enhance the vortex pinning properties of YBCO thin films. The incorporation of Fe into the YBCO matrix at sufficiently low concentrations was effective for vortex pinning and the so-called "poisoning effect" was observed as its concentration is increased [126]. The enhanced vortex pinning properties of the YBCO films with ferromagnetic nanoinclusions have been discussed in terms of Lorentz force reduction pinning mechanism which is effective for low field regime (up to ∼ 1 T) [127]. It has been suggested that in the low field regime, the magnetic flux is effectively reduced from the vortices which subsequently reduces the Lorentz force on the vortex lattice.

Incorporation of Hybrid APCs Into REBCO Thin Films
On the basis of numerous experimental results, it became apparent that some of the secondary phase nanoinclusions tend to self-assemble in the form of nanocolumns along the c-axis direction, while others become nanoparticles with different sizes. The self-assembled columnar nanostructures of perovskite and double-perovskite materials such as BSO, BZO, BSO, YBNO etc. are very effective in enhancing J c particularly when the applied magnetic field is normal to the film surface. However, at higher temperatures, the vortices tend to form double kink structures under the influence of thermal excitations. Even if the matrix contains c-axis correlated defects (nanocolumns), the unpinned vortex segments can move due to the acting Lorentz force. Moreover, as applied magnetic field direction changes from normal to the film surface toward ab-plane, the columnar nanostructures start losing the vortex segments gradually and at enough inclination these vortices become free from the columnar pinning sites. Such a situation is represented in Figure 14. It was, therefore, considered to incorporate mixed pinning landscapes in the REBCO films so that vortices may remain pinned even if the applied magnetic field is tilted from the c-axis of the films. This would result in increased J c in the intermediate angular regime (between ab-plane and c-axis), and the angular anisotropy of the critical current density would thus be reduced.
The use of two different kinds of pinning landscapes was reported by Mele et al. [128] in which BZO nanocolumns and YO nanoparticles were simultaneously incorporated into YBCO film deposited by PLD technique. Although, J c was increased slightly in the intermediate angular regime, but it was found to decrease along the c-axis of the film as compared to BZO added YBCO film (columnar pins only). Later, the combinations of nanocolumns and nanoparticles of different materials have also been reported where the in-field J c was enhanced and anisotropy of J c of YBCO thin films was sufficiently reduced [129][130][131][132]. The use of BSO nanocolumns together with YO nanoparticles was attempted on metallic tape as well and significant enhancement of in-field J c and reduction in J c anisotropy was reported [133]. The combination of BSO nanocolumns and Y 2 O 3 (YO) nanoparticles as hybrid APCs was very effective in controlling the angular anisotropy of the J c of YBCO thin films [130]. The cross-sectional TEM images of YBCO+BSO3% and YBCO+BSO3%+YO nanocomposite films are shown in Figure 15 in which formation of only columnar nanostructures in YBCO+BSO thin film and that of both columnar and spherical nanostructures in YBCO+BSO+YO thin film can be clearly observed. Figure 16 shows the angular variation of J c measured at 77 K, 1 T and 65 K, 3 T for pristine YBCO, YBCO+BSO, and YBCO+BSO+YO thin films. The superior in-field J c performance with reduced anisotropy for YBCO thin films consisting of hybrid APCs can be clearly observed in Figure 16.
Successful incorporations of different combinations of nanocolumns and nanoparticles into YBCO thin films were later reported in which BSO nanocolumns and Y211 nanoparticles [132], BZO nanocolumns and BaCeO 3 nanoparticles [134], and BHO nanocolumns and YO nanoparticles [135] were simultaneously incorporated into YBCO thin films, which resulted in superior in-field J c performance with reduced anisotropy. Hybrid APCs were also incorporated in the form of segmented BSO nanocolumns [136] and segmented nanocolumns with YO nanoparticles [137].
From an energetics point of view, a pinned vortex line is more stable than a free vortex in a superconductor. For a nonsuperconducting particle, the energy gain by a vortex line per unit of interaction volume equals to the condensation energy in the core region. The total pinning energy for a vortex line is, therefore, given by [138]: where, B c (T) = Φ 0 2 √ 2π ξ ab (T)λ ab (T) is the thermodynamic critical field, Φ 0 is the flux quantum, ξ ab and λ ab are the coherence length and the penetration depth along the ab-plane and d the size of the non-superconducting inclusion. ε (θ ) = (cos 2 θ + γ −2 sin 2 θ ) where, γ (= 5-7) is the electronic mass anisotropy parameter. When the vortex line is pinned at the normal core, energy has to be spent to move it away from the core. The force needed to eject a vortex line segment is called the elementary pinning force and is related to the condensation energy as follows: Considering each vortex is individually pinned by a nonsuperconducting core region, and the elastic energies are considerably smaller than the pinning energies, the global pinning force per unit volume may be given by directly summing the elementary pinning forces over all the pinning sites [138].
where N is the density of non-superconducting precipitates, and d p and a 0 are the average separation among the pinning centers and flux lines, respectively. According to this equation, the global pinning force is directly proportional to the concentration of pinning centers. It is, however, to be noted that the nature of Equation (4) is simplistic in the sense that it assumes that all the pinning centers are of equal strength. The density of APCs, however, has an upper limit also for being efficient enough for vortex pinning before the superconductive matrix is severely degraded. Gurevich [139] and other researchers [140,141] considered nanoparticles of materials such as YO or Y211 as strong pinning centers where  parts of the vortex segments can be pinned strongly. If a vortex segment is pinned by two non-superconducting nanoparticles separated by a distance d, then the expression for the upper limit of critical current density, according to their model, is given by: where the terms have their usual meanings. As per this equation, the J c value is proportional to the term 1 d l n d ξ c and for smaller separation among the nanoparticles (higher the concentration), higher J c is anticipated. However, too much incorporation of such non-superconducting nanoparticles may result in the degradation of the superconducting matrix and hinder the transport of electrical current leading to low J c [142].

Factors Determining the Geometry/Morphology of Nanoscale Inclusions
In order to improve vortex pinning in REBCO superconductive films, the inclusion of nanoscale secondary phases remains the most extensively studied among all the techniques. These nanoscale inclusions are considered as strong pinning centers and are very effective depending upon their density and geometry/morphology within the superconducting matrix [8].
It is, however, very interesting to observe that these secondary phase nanoinclusions have different morphologies within the REBCO matrix and accordingly they are efficient for different range of H orientations. While columnar nanostructures are effective for H oriented along the c-axis (exhibiting a strong J c peak for H//c-axis), spherical nanoparticles are observed to be effective for larger angular range of H. It, therefore, becomes imperative to understand the physical origin of varying morphology of these nanoscale inclusions. One of the fundamental parameters which decides the geometry of these nanoscale structures within REBCO thin films, is the lattice mismatch between the two phases. Apart from the lattice mismatch, the surface diffusion of adatoms also play a crucial role in determining the interface between the two phases, which may be coherent, semi-coherent or non-coherent [86]. While the surface energies of coherent and semi-coherent interfaces are much lower, that of the non-coherent interfaces are much higher which results in the observation of many phenomena such as coarsening of the grains in polycrystalline films [143] and faceting of precipitates and grain boundaries [144].
In recent years, Wu et al. in a series of reports [11,[145][146][147][148][149], have attempted to explain the morphology of the nanoinclusions on the basis of elastic strain model. They have found that the APC morphology is determined by the combined effect of the lattice mismatch and elastic properties of REBCO and secondary phase materials. Based on their calculations, they defined a phase boundary separating two regions: c-axis aligned columnar nanostructures are preferred energetically on one side of the boundary and not preferred on the other side of the boundary. Their calculations were in good agreement with experimentally obtained results on well-aligned nanocolumns of materials such as BZO [83,150], BSO [94], BHO [103,151,152], and YBNO [97] and spherical nanoparticles of other materials such as Y211 [153] and YO [154,155] which are on the other side of the calculated phase boundary.

Substrate Surface Modifications
Substrate surface modification is one of the earliest methods used to introduce APCs into superconducting thin film to improve its vortex pinning characteristics. This method provides a means for generating nanostructured APCs by decoration of the substrate surface by non-superconducting secondary phase nanoparticles which generates interfacial defects between the phases. Before the deposition of superconducting thin film, the decoration of the substrate surface is accomplished by growing nanoparticles of various species such as metals [156,157] or oxides [158][159][160][161][162][163][164] on the substrate. The substrate surface is modified by processing an oxide layer in such a way that nanoscale outgrowth develops naturally on the deposition surface. Due to the presence of these nanoparticles at the substrate/film interface, the lattice planes of YBCO are distorted or buckled above the nanoparticles, resulting in low-angle grain boundaries or dislocations that may extend through the entire thickness of the thin film. Matsumoto et al. [159] have reported that the presence of YO nanoparticles at the substrate/film interface produced c-axiscorrelated pinning which was reflected in the enhancement of J c when H was parallel to the c-axis. In another report by Aytug et al. [157], where STO is treated with Ir nanoparticles, the enhancement in J c was observed. The authors have shown that the YBCO planes above the grown Ir nanoparticles are buckled and give rise to random pinning. Random pinning, generally, is caused by the homogeneous distribution of the point defects throughout the film volume. Since in this case the film thickness is 100-200 nm, the presence of strain fields through the entire thickness is supposed to create the appearance of a volumetric particle distribution. The difference in the pinning mechanism due to YO and Ir nanoparticles can be understood in terms of their chemical reaction with the YBCO phase. In the case when Ir is present at STO/YBCO interface, these Ir nanoparticles might have reacted partially with YBCO and the volumetric change that could have taken place might have provided an alternate way to relieve the strain thereby reducing the driving force for dislocation formation and eliminating correlated pinning enhancement. But, YO nanoparticles are chemically inert with YBCO, and therefore, these are intact at the substrate/film interface. Later, ferromagnetic La 0.7 Sr 0.3 MnO 3 (LSMO) nanoparticles were also successfully grown on the  [165], with the permission from AIP Publishing. FIGURE 18 | Variation of critical current density with applied magnetic field as measured at (A) 60 K and (B) 77 K for YBCO films on the undecorated and LSMO decorated STO substrates. Variation of the corresponding pinning force density with applied magnetic field is shown in the insets of (A,B). Reproduced from Jha et al. [165], with the permission from AIP Publishing.
STO surface by PLD technique before the growth of YBCO thin film [165]. Figure 17 shows the atomic force microscope (AFM) images of undecorated and LSMO nanoparticle decorated STO substrates onto which YBCO thin films were deposited. The decorated sample exhibited enhanced in-field J c than the undecorated sample which can be observed in Figure 18. It was suggested in this report that the presence of LSMO nanoparticles at the YBCO/STO interface could generate structural defects such as threading dislocations which would result in c-axis vortex pinning in the resulting YBCO film. The enhanced J c values for the YBCO film on decorated substrate can also be understood in terms of Lorentz force reduction pinning mechanism [127]. In such a scenario, the Lorentz force splits between the vortices and the magnetic pinning sites which results in higher current density before the vortices start depinning from the defects. The metal organic decomposition method has also been employed for decoration of the substrate surface for improving the vortex pinning properties of YBCO thin films [166][167][168].

CONCLUSIONS
REBCO superconductors are the most promising HTS for their high T c and J c values. However, for their use in several applications in the practical range of temperature and applied magnetic field, pinning of vortices is an essential requirement. It has been shown that the defects that are generated naturally during the growth of the sample (bulk or thin film) are not sufficient to pin the vortices at elevated temperatures and applied magnetic fields. That is how the need of introducing artificial pinning centers arose and extensive work has been carried out to address this issue.
In recent years, a variety of methods have been reported to intentionally introduce nanostructured defects into the REBCO superconducting samples which include doping of rare-earth elements, addition of secondary phase nanostructures and modification of the substrate surface. The different methods for introducing artificial pinning centers lead to different and interesting underlying physics. The dominant pinning mechanism needs to be investigated which governs the role of artificial pinning centers in these superconducting samples.
The various interesting issues and future potential applications based on REBCO superconducting films have provoked researchers to continue activities in this field of research. The generation of nanoscale APCs with desired density, geometry and orientation is very much needed for vortex pinning in REBCO nanocomposite films deposited on single crystals and metallic tapes. Precise control of these parameters is very promising for optimal pinning of the vortices as required to push the limits of critical current performance.

AUTHOR CONTRIBUTIONS
AJ conceived the idea of highlighting the role of rare-earth oxides in promoting sustainable energies. AJ and KM discussed the development in this area over the years and contributed to the manuscript.