Peritectic melting of thin films, superheating and applications in growth of REBCO superconductors

Progress in Materials Science 68 (2015) 97–159

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Peritectic melting of thin films, superheating and applications in growth of REBCO superconductors Yuanyuan Chen a, Xiangxiang Cui a,⇑, Xin Yao a,b,⇑ a State Key Lab for Metal Matrix Composites, Key Lab of Artificial Structures & Quantum Control (Ministry of Education), Dept. of Physics and Astronomy, Shanghai Jiao Tong University, 800 Dongchuan Road, Shanghai 200240, China b Collaborative Innovation Center of Advanced Microstructures, Nanjing 210093, China

a r t i c l e

i n f o

Article history: Received 18 April 2014 Received in revised form 12 September 2014 Accepted 13 September 2014 Available online 19 October 2014 Keywords: Peritectic melting Thin film Superheating phenomenon REBCO superconductor Phase transition Seed materials Crystal growth

a b s t r a c t Superheating of solids, an unconventional phenomenon in nature, can be achieved by suppressing the heterogeneous nucleation of melt at defect sites, such as free surfaces and internal grain boundaries. In recent years, experimental evidences have clearly proved that the YBCO (Y123) thin film with a free surface possesses a superheating capacity, which is mainly attributed to the film/substrate structures, distinctively consisting with low-energy surface and semi-coherent interface. Like most functional oxides, YBCO (denoted as a phase) is characterized by a peritectic melting: a ? b + liq. Its superheating behavior certainly relates to this peritectic reaction. Furthermore, REBCO (RE123, RE: rare earth elements) thin films with high thermal stability have been successfully employed as seed materials in inducing the growth of REBCO materials, such as thick film, single crystal and single domain bulk. Therefore, this superheating property of thin films is of great importance in both scientific study and practical application. In this paper, the up-to-date researches covering on the superheating phenomenon of the a phase film, its mechanism and applications in growth of REBCO superconductors are reviewed, which is supposed to be valid for more thin films of functional oxides that have the same nature as the YBCO film/substrate. Ó 2014 Elsevier Ltd. All rights reserved.

⇑ Corresponding authors at: State Key Lab for Metal Matrix Composites, Key Lab of Artificial Structures & Quantum Control (Ministry of Education), Dept. of Physics and Astronomy, Shanghai Jiao Tong University, 800 Dongchuan Road, Shanghai 200240, China (X. Yao). E-mail address: [email protected] (X. Yao). http://dx.doi.org/10.1016/j.pmatsci.2014.09.001 0079-6425/Ó 2014 Elsevier Ltd. All rights reserved.

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Contents 1. 2.

3.

4.

5.

Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 99 First two reports on superheating phenomenon of YBCO thin films . . . . . . . . . . . . . . . . . . . . . . . . . . . 100 2.1. YBCO-film-seeded NdBCO-growth by liquid-phase-epitaxy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 101 2.2. In-situ observed melting mode of YBCO thin film and its superheating nature . . . . . . . . . . . . . 103 Factors affecting thermal stability of thin films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 106 3.1. Film/substrate structure__primarily relating to melting nucleation . . . . . . . . . . . . . . . . . . . . . . 106 3.1.1. Film orientations . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 106 3.1.2. Substrate materials. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 117 3.2. Phase diagram nature__primarily relating to melting growth . . . . . . . . . . . . . . . . . . . . . . . . . . 121 3.2.1. Superheating related supersaturation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 122 3.2.2. Temperature coefficient of solubility . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 124 3.3. Correlations between film/substrate structure and phase diagram nature . . . . . . . . . . . . . . . . 128 Exploits in fundamental study and engineering application of thin films with high thermal stability 131 4.1. New field for basic research: unconventional phase transition in high-superheating state. . . 131 4.2. As seed materials: effectively inducing the growth of high performance REBCO crystals . . . . 135 4.2.1. Top-seeded melt-growth and progresses of seed materials searching . . . . . . . . . . . . . 135 4.2.2. Demonstrations of superheating capacity: thin-film as homo-seed in REBCO cold-seeding process . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 138 4.2.3. YBCO-buffered NdBCO film with enhanced thermal stability for growing higher processing-temperature required REBCO crystals . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 140 4.2.4. Long tolerability of film-seed with melt and two-layer batch growth of YBCO . . . . . . 143 4.2.5. Thermal stability of film-seed in low supersaturated melt for growing YBCO crystals 147 4.2.6. Highly oriented a/b-direction and a–b plane of film-seed in multi-seed process for controlling grain boundary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 149 4.2.7. Large-sized film-seed inducing the growth of REBCO crystal with large c-growth sector 152 Conclusion and prospects . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 153 Acknowledgments . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 155 References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 155

Nomenclature

English symbols a/a GSB a/a growth sector boundary AFM atomic force microscope a-GS a-growth sector BCO Ba–Cu–O liquid CE the equilibrium concentration of the RE element in the Ba–Cu–O liquid c-GS c-growth sector CL the RE element concentration in the liquid CL123 the RE element concentration in the liquid at the front of the interface of RE123 CL211 the RE element concentration in the liquid at the front of the interface of RE211 CS123 the RE element concentration of the solid RE123 CS211 the RE element concentration of the solid RE211 EDS energy-dispersive spectrometer EPMA electron probe microscopic analysis HTOM high temperature optical microscopy Jc critical current density JIn the density of RE element outflow from RE123 to the Ba–Cu–O liquid

Y. Chen et al. / Progress in Materials Science 68 (2015) 97–159

JOut LAO LPE LREBCO LSS MG MSMG NM P(O2) PCS PLD RE123 R211 grow

the density of solute flow from the liquid to the RE211 growth front LaAlO3 liquid phase epitaxy LRE123, LRE = light rare earth element, e.g., Nd, Sm, Gd large-sized seed melt-growth multi-seeded melt-growth non-melting oxygen partial pressure persistent current switch pulsed laser deposition REBa2Cu3Oy, RE = rare earth element the growth rate of RE211

R123 melt

the melting rate of RE123

SEM SIG STO sub. Tc t.c.s. TMG up

scanning electron microscope seeded-infiltration and growth SrTiO3 substrate superconducting critical transition temperature temperature coefficient of solubility the highest temperature, above which the film decomposes and fails to induce RE123 by melt growth the temperature at which films decompose completely the equilibrium melting point a maximum processing temperature employed in TSMG which is higher than Tp the peritectic temperature top-seeded melt-growth top-seeded solution-growth the melting starting temperature the up-limit temperature at which melting catastrophe (homogeneous nucleation) takes place at the RE123/sub. interface Y2O3 YBCO-seeded NdBCO-growth two-dimensional

Tend Tm Tmax Tp TSMG TSSG Tstart Tup Y200 YSNG 2D

99

Other symbols r supersaturation r-DTs curve supersaturation versus the degree of superheating

DTs DTsc DTM

degree of superheating capability of superheating temperature difference between TMG up and Tp

DC 123 melt

the difference of the solute concentration between the RE123 solid and the liquid

DC 211 grow

the difference of the solute concentration between the RE211 solid and the liquid

d

the ratio of concentration differences (CS211–CL211)/(CS123–CL123)

1. Introduction As an intrinsic property of solid material, the thermal stability is of critical concerns for fundamental studies and practical applications. Cahn [1] have predicted a situation of a so-called superheating or overheating phenomenon. It is the phenomenon that when melting at surface of a solid was

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suppressed, one may remain in its metastable solid phase above its equilibrium melting point (Tm). Investigations focused on this unconventional phenomenon can date back over decades. Since Däeges et al. first reported a superheating phenomenon in Ag particles coated by Au in 1986 [2], it became increasingly imperative to understand the mechanism of thermal stability against melting, i.e., superheating property. However, dissimilar to supercooling, superheating is extremely difficult to be achieved in the real nature, since melting always initiate from the defect sites such as free surfaces owing to their vanishing barrier for melting. Therefore, up to date, most strategies have been applied to suppress melting from surface or interface to observe a superheating phenomenon. In recent years, numerous researchers have reported this fascinating phenomenon in varied systems. Based on their intrinsic structures, three major superheating modes can be identified among the reported studies. The first one is superheating of confined-particle structure. To be specific, when the nano-sized particles were epitaxially embedded in the materials with high melting points [2–9], the particles can survive above their melting point with semi-coherent interfaces and no exposed surface. In that case, the melting initiates inside the bulk interior. In other words, the heterogeneous nucleation of melting at the surface and interface can be effectively suppressed in the particle–matrix system [10]. The second one is superheating of sandwich-confined film construction [11–14] such as Al/Pb/Al. Owing to the semi-coherent interfaces, the growth of molten Pb liquid is retarded due to a positive interfacial energy change of cPb(l)/Al  cPb(s)/Al, resulting in a superheated state in the Pb thin film. The third one is superheating of Pb single crystals with the so-called non-melting (NM) surfaces of crystalline planes (1 1 1), which are characterized by low surface energy. It was reported that this kind of crystal can be superheated for several degrees [15]. Melting nucleation can be effectively delayed due to the higher nucleation barrier on all exposed surfaces. As low-dimensional materials, their thermal stability is a crucial issue for industrial applications. For instance, achieving a superheating state of solid may provide novel approaches for studying phase transformation and stabilization phenomenon. However, it is practically difficult to prepare the confined or embedded structures. On the other hand, the superheating capability of the materials with non-melting free surfaces is not strong enough for technological utilization. More recently, REBCO7y (REBCO or RE123, RE = rare elements) thin films were observed to have a superheating property [16]. As a typical representative of functional oxides, REBCO materials possess a peritectic melting, thus the superheating property means they can endure a high temperature above their peritectic temperature (Tp). This superheated film material with free surface provides new perspectives for studying superheating origins. It is generally believed that this superheating phenomenon is mainly attributed to both low REBCO surface energy and low film/substrate interface energy, as we will discussed in this review. As typical functional oxides as well as high temperature superconductors, REBCO thin films with high thermal stability have wide applications in many fields. In the present paper, we review the major progresses on the superheating of REBCO oxide films including the discovery, various experimental investigations and analyses on this subject, to give a comprehensive summary on this topic. Emphasis is made on the mechanism and factors that affect the thermal stability of the REBCO film/substrate structure, including film orientation, substrate material, temperature coefficient of solubility and so on. In the following section, the extended research on the phase transformation behaviors in high superheating status is discussed. Finally, the achievements in technological applications of thin films as seed materials in inducing the growth of REBCO superconductor bulks are addressed in the last sections.

2. First two reports on superheating phenomenon of YBCO thin films As popular materials, functional oxides of thin films (denoted as a phase) mostly are characterized by a peritectic melting. That is, a solid phase decomposes into a second solid phase and a liquid at the peritectic temperature. By using liquid phase epitaxy (LPE) and high temperature optical microscopy (HTOM), the superheating phenomenon of REBCO (a typical representative of a phase) thin films and its mechanism have been reported and analyzed. The peritectic melting of REBCO oxide is: 2YBa2Cu3O7y (Y123, solid) ? Y2BaCuO5 (Y211, solid) + Ba3Cu5O9 (Ba–Cu–O, liquid).

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2.1. YBCO-film-seeded NdBCO-growth by liquid-phase-epitaxy Liquid-phase-epitaxy is one of the important avenues for preparing crystalline films. Compared with the other film-preparing techniques, it has several advantages. Firstly, high-quality crystals can be obtained since the film grows under near thermal equilibrium conditions. Secondly, the LPE process has a high growth rate, thus thick epitaxial-films can be obtained in high efficiency. Finally, it is a cost-saving method, since vacuum atmosphere is not necessary. Generally, this method is carried out on the basis of the concept of Czochralski growth, i.e., topseed solution growth (TSSG). In other words, it is a seed-required process. In that case, the pre-existing hetero-seed (the seed material is different from the growing film) is commonly used when the homoseed (the seed is the same material as growing film) is hard to gain. In the principle of the hetero-seed selection, firstly and most importantly, it should have a higher melting temperature than the processing temperature. Furthermore, the hetero-seed material should have a good lattice matching with and a similar type of crystal structure to LPE film materials. Remarkably, it was found that acting as hetero-seeds, YBCO (Tp about 1005 ± 5 °C) thin films were successfully used in inducing the growth of high Tp NdBCO thick films by LPE at the processing temperature of 1055 °C [17,18]. This YBCO seeded NdBCO growth (YSNG) demonstrates a completely different behavior that is certainly inconsistent with the essential prerequisite mentioned above, ‘‘the seed material should have a higher melting temperature than the processing temperature’’. This is the first report with respect to a superheating phenomenon on the 2-dimensional oxide film in seeding LPE growth. Note that these used YBCO films were deposited on MgO substrate by pulsed laser deposition (PLD), which possess a high c-axis orientation associated with a superior in-plane alignment, presenting a 0° relationship of h1 0 0iYBCO//h1 0 0iMgO. It can be deduced that this novel superheating phenomenon intrinsically relates to this film/substrate structure. Aiming at elucidating the mechanism of YSNG, the microstructural evolution at the initial stage was investigated. A vertically dipping experiment [19] was conducted with a continuous traveling mode as schematically illustrated in Fig. 1. The MgO substrate with an YBCO thin film was immersed vertically into the liquid at a rate of 1 mm/s. Once touching the liquid, the specimen was pulled up. The dipping experiment experienced approximately 3 s, and a film of about 3 mm in length on the dipped side was obtained. Due to the liquid wettability, the liquid climbed upon the substrate immediately at the beginning of the dipping. The boundary area between seed-film and as-grown film corresponds to nearly 0 s of growth, while the bottom region undergoes 3 s dipping time. Fig. 2 shows an optical micrograph of a vertically dipped LPE specimen, near the boundary between an undipped region on the upper side and a dipped region on the lower side. When the specimen was put into the liquid, the YBCO grains underwent heating from room temperature to 1055 °C, which is higher than the peritectic temperature of YBCO, and became thermodynamically unstable. As a result, Y2BaCuO5 (Y211) grains and solidified Ba–Cu–O (BCO) melts can be observed in the undipped region, which are the products of the YBCO decomposition. BCO covers Y211 due to its wettability. Needle-shaped Y211 grains with the c-axis in a long length direction [20] are well aligned on the MgO substrate due to good lattice matching. The preferential orientation presents a relation of 1 mm/sec Before touching

Dipping

Separating After separating YBCO-seed-film 0 sec 3 sec

0 sec

3 sec

Dipped region

Nd-Ba-Cu-O Solution Fig. 1. Schematic illustration of vertically dipping experiment.

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Fig. 2. An optical micrograph showing microstructures at the boundary between the Liquid-untouched region on the upper side and the Liquid-dipped region on the lower side.

h0 0 1iY211//h1 1 0iMgO, interpreted by a minor difference of the lattice constant between the Y211 caxis and the MgO h1 1 0i direction: 0.5659 nm and 0.5955 nm [21,22], respectively. Surprisingly, however, the undecomposed YBCO particles are clearly visible, and unexpectedly retained although they endured a high processing temperature, implying a superheating phenomenon. Most importantly, both YBCO and NdBCO are visible on each side of the boundary, showing the same orientation: (1 0 0)NdBCO//(1 0 0)YBCO//(1 0 0)MgO. In other words, it indicated the 0°-oriented NdBCO grain on the MgO substrate, which behaves same as the YBCO seed-grain on MgO. No independent YBCO grain is visible in the dissolution area due to its instability and the Y solubility in the Nd–Ba–Cu–O solution at 1055 °C. It can be deduced that YBCO grains could be responsible for the formation of NdBCO grains because both of them have the same crystal structure and most elements and bondings are similar [23]. However, the critical point is to probe a linkage between YBCO and NdBCO grains. The mapping analysis by electron probe microscopic analysis (EPMA) was used to determine the Y and Nd distribution near the boundary between dipped and undipped regions, as shown in Fig. 3. On the one hand, NdBCO grains can be recognized from the Nd mapping image in the dipped region. On the other hand, close to the boundary, we can clearly see that the Y element concentrates at the site of NdBCO grains from the Y distribution result (marked by the rectangle). Therefore it can be induced that these Y-segregated regions are previous YBCO grains, which act as hetero-seeds for the growth of NdBCO grains. This confirmed the existence of the Y element from the initially grown NdBCO grains, which provide direct evidence that high peritectic temperature NdBCO grains could grow from YBCO film-seeds. In short, the superheating phenomenon of YBCO thin films was found in YSNG process. On the basis of the investigation on the initial stage of YSNG, a comprehensive model was given. Since a biaxiallyoriented YBCO thin film is relatively easy to gain in comparison with that of other REBCO materials, the potential merit of the use of this film is obviously encouraging. Firstly, this so-called hetero-seeded growth can be extended to a wider application, suitable for all REBCO oxides and even other perovskite ones, such as ferroelectric materials. Note that against the dissolution the YBCO film even survived in the PbO solvent and acted as seed to induce PZNT LPE film at the growth temperature of 1050 °C [24]. This is of great significance in technological applications in various fields, such as HTS device, persistent current switch (PCS) device and coated conductor. So far, using YBCO/MgO

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Fig. 3. EPMA mapping images showing Y and Nd distributions at the boundary between the Liquid-untouched region on the upper side and the Liquid-dipped region on the lower side. The color gradient from white (top) to black (bottom) shows a reduced tendency for the element content Nd and Y [17].

Table 1 Using YBCO/MgO thin films as seed materials, various LPE thick films were successfully grown. Systems

Growth temperature (Tg)

NdBCO SmBCO Ni-NdBCO Yb-YBCO Ca-YBCO Zn-YBCO Pb[(Zn1/3Nb2/3)0.91Ti0.09]O3 (PZNT)

1057 °C [17,18] 1020–1057 °C [24] 1057 °C [25] 970–980 °C 957–980 °C [26,27] 966–984 °C 1050 °C [28]

thin-film-seed, the LPE thick-film growth have been successfully achieved in various oxide materials (see Table 1) [24–28]. Secondly, not only valid for the LPE growth, the principle of YSNG can presumably be utilized for other seed-required HTS processes, e.g., the single crystal growth, the bulk growth and the coated conductor process. Lastly, as a universal behavior, it can be generally used for all hetero-seeded growth of materials. It means more candidates can act as seeds for preparing crystals. 2.2. In-situ observed melting mode of YBCO thin film and its superheating nature In the previous section, it was discussed that an NdBCO thick film could grow on the YBCO thinfilm-deposited MgO substrate by LPE. The interesting point is that YBCO as a hetero-seed material has a peritectic temperature lower than the processing temperature of LPE-grown NdBCO thick films. This behavior is evidently opposite to what occurs in a conventional melting process. To clarify the possible origin of YSNG, and to extend the understanding of melting mechanism of the functional oxide materials, the melting process of the YBCO/MgO thin film employed in the YSNG method was in-situ observed by high temperature optical microscope with a heating stage. This method allows direct visualization of the peritectic decomposition and can provide real-time observations of the nucleation and growth of the melting process. The specimen was heated from room temperature to 1060 °C in air. Because of the poor thermal conduction of the substrates, the temperature was held at several fixed points for 3 min to avoid the real temperature of the films lagging behind the nominal setting. A series of micrographs are shown in Fig. 4, which illustrate an example of the observation of melting process of the YBCO thin film grown on the MgO substrate. Plenty of small black dots on the YBCO

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Fig. 4. An optical microscope showing microstructure of a YBCO thin film deposited on MgO at (a) room temperature; (b) 1010 °C for 3 min; (c) 1050 °C for 3 min; (d) 1060 °C for 3 min [16].

Fig. 5. X-ray pole figure of the Y211 crystal on the MgO substrate [16].

thin film can be observed from Fig. 4(a) at room temperature, which were either YBCO crystal defects or impurities. Then the sample was heated to 1010 °C, which was the peritectic temperature of YBCO, and held for 3 min. From Fig. 4(b) we can see the black dots gradually enlarged, implying Y211 crystals nucleated and grew in the defect regions [29,30]. Fig. 4(c) gives an optical micrograph of YBCO thin film when the temperature reached 1050 °C and was held for 3 min. It was remarkable that some needle-shaped Y211 gradually appeared from undersurface (see the dashed line in Fig. 4(c)), but only a small amount of YBCO thin film was melting at this temperature, although it is far higher than Tp of YBCO. The needle-shaped Y211 grains were well aligned on the MgO substrate except for a very few. Fig. 5 shows the pole figure analysis for the Y211 crystal. It can be clearly seen that h0 0 1i direction of Y211 grains oriented in a 0° relationship to the h1 1 0i direction of the MgO substrate, which is in good agreement with other experimental work [31,32]. The small lattice mismatch of Y211h0 0 1i//

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Fig. 6. Optical micrographs displaying the liquid movement from the Y123 melting front to the Y211 phase. Two framed areas present the melting evolution in the same site, while an arrow in the inserted picture on the upper right side points a shifted liquid from the Y123 melting front after holding for 10 s.

MgOh1 1 0i and Y211h1 0 0i//MgOh1 1 0i (0.0515 and 0.0995, according to the equation: lattice mismatch = 2(a1  a2)/(a1 + a2)) must be responsible for this epitaxial relation. Most importantly, the majority of the YBCO thin film was undecomposed although it endured a high processing temperature. Thus it proved again that the YBCO-thin-film possesses a superheating capacity. The morphology of the YBCO thin film after holding for 3 min at 1060 °C is shown in Fig. 4(d). We can observe a great many well aligned Y211 grains on the MgO substrate. The BCO liquid, as the product of the YBCO decomposition, covers on Y211 grains due to its high wettability to Y211 and relatively low wettability to the MgO substrate. Most importantly, the undecomposed YBCO particles were still clearly visible, and unexpectedly retained between some Y211 particles, especially when they are completely isolated. It can be concluded that some YBCO particles can even be superheated up to 1060 °C. According to experimental observations [33] and thermodynamic analysis [34], the melting temperature of freestanding two-dimensional (2D) thin film is lower compared to the bulk melting temperature, which is similar to the melting of nano-particles. However, unlike the superheating cases in confined systems, there is a free surface of YBCO thin film, thus the model of melting nucleation suppression does not work. To explain this novel superheating phenomenon, three factors have been mainly considered. Firstly and most importantly, the YBCO thin film is c-axis oriented, which correlates with a low surface energy. Consequently, the melting nucleation at the surface is effectively suppressed. Secondly, the semi-coherent interface between the MgO substrate and the YBCO thin film is a crucial factor. Because of good lattice match between the Y123 crystal and the MgO substrate (lattice fit  0.0952, based on the relationship of Y123h1 0 0i//MgOh1 0 0i), this interface is a low-energy interface. Therefore it is favorable to suppress the heterogeneous nucleation at the interface, which would result in the superheating of YBCO thin film. It can be reasonably deduced that the complete decomposition of YBCO thin film occurs when the sample is superheated to a certain extent, the bonding at the semi-coherent interface between the Y123 and the MgO substrate is destroyed. Thirdly, from the real-time observation of YBCO melting at a temperature of 1043 °C shown in Fig. 6, the BCO liquid, as marked at the melting front of Y123 grain, rolls away immediately from the melting interface of the Y123 grain and sweeps across the MgO surface to Y211 needles. This result suggests that the BCO liquid wets neither the parent phase of Y123 nor the MgO substrate. Generally, the melting continues as long as a driving force exists (even if it is very small). However, the nucleation

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of melting requires a sufficiently large driving force. Since liquid keeps moving away from the Y123 solid and cannot facilitate the growth of melting, a continuous driving force for nucleation of melting is required. All the above-mentioned factors may cause the delay of melting. In brief, apart from the YSNG result by LPE, the in-situ observation by HTOM confirms again the c-axis YBCO thin film with a MgO substrate can be substantially superheated above its Tp (about 50 °C). Unlike the previously reported superheating in metallic systems, the oxide superconducting materials are characterized by a peritectic melting. The origin of this superheating phenomenon is that the YBCO film has both low-energy surface and semi-coherent interface with MgO. This finding leads to a conclusion that superheating exists not only in metal but also in oxide materials. 3. Factors affecting thermal stability of thin films Following the rising interest in two-dimensional REBCO thin film materials, dependence of their thermal stabilities on several factors is reviewed in this section. 3.1. Film/substrate structure__primarily relating to melting nucleation It is well known that melting normally initiates from defects or grain boundaries. In addition, it was reported that while a thin film or particle was confined by a semi-coherent interface, resulting in no exposed surface, the melting nucleation can be suppressed, and then they can survive above its melting point. Thus the microstructures of the films and the intrinsic structure at the film/substrate interface should play a key role in heterogeneous nucleation of melting. From this respect, two effects are studied. One is effect of film orientations including out-of-plane and in-plane alignment, the other is effect of substrate material. 3.1.1. Film orientations I. Out-of-plane orientation Firstly, the out-of-plane orientation effect was investigated. Generally, during the growth of c-axis oriented REBCO films, a-axis oriented grains always appear, becoming a commonly-identified out-ofplane orientation. It has been demonstrated by Ichino et al. that a pure c-axis-oriented SmBCO thin film can be prepared by the usual PLD technique only in an extremely rigorous growth condition [35]. Therefore, it is necessary to study the thermal stability of the c-axis oriented REBCO thin films with various concentrations of a-axis oriented grains. Cheng et al. have performed an in-situ research on the out-of-plane effect on the thermal stability as well as melting behaviors of the thin film [36]. To investigate the out-of-plane effect, two types of SmBCO films deposited by the PLD technique on single-crystalline MgO substrates were used, labeled as film A and B, respectively [35]. Both films were dominated by 0°-oriented grains. According to the peak intensity in the XRD spectrum, the portion of

Fig. 7. SEM images of two SmBCO thin films: (a) film A and (b) film B [36].

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Fig. 8. Optical micrographs showing the melting process of the films A and B: (a–c) for film A at 1100 °C just achieved, at 1100 °C after 3 min, and at 1130 °C after 5 min, respectively. (e–f) for film B at 1041 °C just achieved, at 1050 °C just achieved, and at 1065 °C after 3 min, respectively [36].

the a-axis oriented grains in film A was established as 1%, while in film B 3.2%. Obvious difference in surface morphology between the two films can be seen from the scanning electron microscope (SEM) images in Fig. 7. Those needle-like objects on the flat area indicated by arrows are so-called a-axis-oriented SmBCO grains. The long axis of those grains runs along h1 0 0i or h0 1 0i, while the short axis is h0 0 1i-oriented [37,38]. We can observe that the portion of the a-axis-oriented grains in film B (Fig. 7(b)) is much higher than that in film A (Fig. 7(a)), while the size is much smaller. The HTOM was employed to observe the melting evolutions of the two films in air atmosphere. A series of micrographs in Fig. 8 illustrate the melting process of two thin films. Fig. 8(a) depicted film A at 1100 °C, higher than its Tp (1065 °C in air), remained almost unchanged, indicating a superheating property. Fig. 8(b) shows a section of the film held for 3 min at 1100 °C. Some needle-shaped Sm211 grains appeared, indicating a partial peritectic decomposition of SmBCO. Morphology of film A held for 5 min at 1130 °C, with proceeding peritectic melting, is shown in Fig. 8(c). Fig. 8(d–f) shows the evolution of film B on temperature rising from 1041 °C to 1065 °C. The SmBCO grains in film B completely decompose at 1075 °C. In short, the SmBCO decomposition in film A started at 1100 °C and finished at 1130 °C, while in film B it started at 1050 °C and ended at 1075 °C. The conclusion is that film A has a higher thermal stability than film B (see Table 2). After the melting process, the samples were taken out from the heating stage to observe and compare the final surface morphology (see Fig. 9). In film A the c-axes of the Sm211 crystals have a 45°

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Table 2 The differences in melting behavior of films A and B. The decomposition temperatures Tstart and Tend of the films indicate the temperatures at which the SmBCO thin films start and finish decomposition. SmBCO films

XRD (vol% of a-axis grains) (%)

Decomposition temperature Tstart (°C)

Decomposition temperature Tend (°C)

Preferential orientation of Sm211 (°)

Size of grains Sm211

Film A Film B

1 3.2

1100 1050

1130 1075

45 0, 90

Large Small

Fig. 9. A comparison of the surface morphology of SmBCO films after the melting process: (a) film A, (b) film B [36].

preferential orientation with respect to h1 0 0i of the MgO substrate, i.e., Sm211h0 0 1i//Sm123h1 1 0i// MgOh1 1 0i, denoted as 45° Sm211. In contrast, in film B, the orientation relations Sm211h0 0 1i// Sm123h1 0 0i//MgOh1 0 0i and Sm211h0 0 1i//Sm123h0 1 0i//MgOh0 1 0i are presented, denoted as 0° Sm211 and 90° Sm211, respectively. The size of the Sm211 crystals in film A is much bigger than that in film B. The distinctions in melting behaviors of the two films should be attributed to the difference in their microstructures. From the viewpoint of surface (or interface) energy, the discussion of the different decomposition modes will be focused on (i) thermal stability and (ii) orientation of decomposition product Sm211. As regard to thermal stability of the two films, film A, in which the number of a-axis-oriented grains is much less, can be superheated to a rather higher temperature. Similar result has also been obtained in the YBCO system [39]. The c-axis YBCO/MgO film with a few a-axis grains showed a higher thermal stability compared with the pure c-axis one. On the one hand, attention is paid to the difference in film microstructure. Fig. 10 shows the AFM results of these YBCO films, which reveals one of the major reasons for the distinction. As one can see, the film with a few a-axis grains has a relatively regular square shape, while the pure c-axis one presents a twisted shape. Taking lattice matching into consideration, the a-axis-oriented YBCO grain has a smaller misfit than the c-axis-oriented one. Therefore, the existence of a-axis-oriented YBCO grains may partially relieve the stress energy, resulting in the reduction of the total system energy. In brief, a small amount of a-axis-oriented grains can help to relax the interface stresses and let the film to endure superheating at a higher temperature. On the other hand, the different surface energy introduced by the a-axis-oriented grains should be taken into consideration, since it is commonly accepted that a solid starts melting at heterogeneous nucleation areas like the defect sites at interfaces or free surfaces. Furthermore, melting is more likely to take place at a surface with high free energy. Theoretical and experimental work on the REBCO system proved that the (0 0 1) surface has the lowest surface energy among the common three crystallographic layers (the indentation measurement results: 4.4 eV nm2 of (1 0 0), 6.8 eV nm2 of (0 1 0),

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Fig. 10. AFM images showing the microstructure of: (a) c-axis YBCO film with a few a-axis grains; (b) pure c-axis YBCO films [39].

Fig. 11. A schematic illustration of SmBCO crystal set on YBCO bulk in top-seeded melt growth [36].

3.5 eV nm2 of (0 01 ) surfaces, respectively) [40,41]. This means that decomposition of Sm123 should initiate at the less-stable (1 0 0) or (0 1 0) planes. A lower thermal stability of the (1 0 0) or (0 1 0) surface of the SmBCO crystal compared with the (0 0 1) surface was also verified by the following experiment. Due to the lack of pure a-axis-oriented film (extremely difficult to prepare) for comparison, the thermal stability of different crystalline planes in a SmBCO crystal was studied. Small SmBCO single crystal was placed on the top center of an YBCO pellet before melt processing. The upper surface of the SmBCO single crystal in the schematic illustration of Fig. 11 is (0 0 1) free surface, while the lateral view is (1 0 0) or (0 1 0) surface. In the heating procedure, the SmBCO crystal endured the highest temperature of 1045 °C for 2 h, lower than the Tp of SmBCO material. It was found that no reaction took place on the (0 0 1) surface. However, needle-shaped Sm211 grains appeared on the (1 0 0) or (0 1 0) faces with a relationship of Sm211h0 0 1i//Sm123h1 0 0i (see Fig. 12). Based on the peritectic melting reaction, the nucleation of Sm211 grains means that the Sm123 decomposition occurred. That is to say, the (1 0 0) or (0 1 0) surface possesses a lower thermal stability and the peritectic melting may initiate on those crystalline surfaces, even at a temperature lower than Tp of REBCO. Therefore, an increase in the proportion of a-axis-oriented grains can cause a rise in surface energy of the film, which weakens the thermal stability of the film. Since the introduction of a-axis-oriented grains may cause both positive and negative effects on thermal stability of c-axis-oriented films, it is necessary to clarify which effect is dominant. In that case, the locations of melting nucleation where the peritectic melting virtually took place, was taken into consideration. Particularly, a focus was paid on the orientation relationship of Sm211 with Sm123. According to the experimental results shown in Fig. 9, most of the Sm211 grains on film A have a 45° orientation on MgO, presenting a preferential epitaxial relation of Sm211h0 0 1i//MgOh1 0 0i [16].

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Fig. 12. Material melting observed on the lateral view of SmBCO single crystal [36].

It implies that the peritectic melting begins at the (0 0 1) film/substrate interface in film A. Dissimilarly, in peritectic melting of film B, the product phase of Sm211 shows a 0° or 90° orientation with MgO, which is consistent with that of the primary phase (a-axis grains in Fig. 7). It is an epitaxial relation (Sm211h0 0 1i//Sm123h1 0 0i) that was observed on the (1 0 0) or (0 1 0) face of Sm123 singe crystal as shown in Fig. 12. Thus, the directions of Sm211 grains in film B suggest a possibility of Sm211 nucleation on the surface of a-axis-oriented Sm123 grains. In such a case, the nucleation and growth of Sm211 grains would just proceed along the Sm123 (1 0 0) or (0 1 0) free surfaces, rather than on the film/substrate interface. Therefore, it can be deduced that there is an evolution of nucleation site change, from the interface to surface of Sm123/MgO films with increasing a-axis-oriented grains, which certainly relates to the thermal stability of films. Firstly, for a pure c-axis-oriented film, the peritectic melting of the Sm123 film is characterized by a normal mode: nucleating at the Sm123/MgO interface due to the misfit caused stress. Secondly, with a small amount of a-axis-oriented grains (film A), the positive effect, i.e., the relaxation of stress energy at the SmBCO/MgO interface is dominant. As a result, the peritectic melting delays to a higher temperature, and however, it still initiates at the interface between SmBCO (0 0 1) plane and MgO. Thirdly, with further increasing number of a-axis-oriented grains (film B), the negative effect of the enhanced surface energy becomes more important. In that case, the normal mode (the peritectic melting nucleates at interface) changes. The primary nucleation occurs at high energy (1 0 0) or (0 1 0) surface of a-oriented grains, at a temperature lower than Tp, demonstrating a low thermal stability. In short, the thermal stability of SmBCO films, the epitaxial relation between Sm211 and Sm123 and melting nucleation location were significantly influenced by a slight change in the proportion of a-axis-oriented grains. The out-of-plane orientation effect was verified from the view point of stress energy combining with surface energy. II. In-plane alignment In metals and its alloys, until now, many studies [42–44] have confirmed that there is a correlation between the intrinsic microstructure and their thermal stability. Generally, the intrinsic structure at the film/substrate interface strongly relates to the in-plane orientation. In this respect, the effect of in-plane alignment on the superheating capacity of REBCO thin films was studied. (1) Different melting modes Firstly, different melting modes were reported on basis of HTOM in-situ observation on two kinds of c-oriented YBCO films, with double (film C) and single (film D) aligned grains [45], respectively. In

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XRD pole figures as shown in Fig. 13, film C is characterized by an eightfold symmetry pattern indicating a mixture of both 0° (YBCOh1 0 0i//MgOh1 0 0i) and 45° (YBCOh1 0 0i//MgOh1 1 0i) grains, while film D features a fourfold symmetry one representing a single 0° aligned grains. The melting behavior of the two films showed great difference. Fig. 14(a and b) represents the initial melting stage of film C. There are two points that one should pay attention to. First, the number of Y211 crystals is large. Second, the grains with 0° and 45° orientation appeared simultaneously. For film D, firstly, Y211 crystals began to appear as several short dashed lines, see Fig. 14(c). Notice that the temperature was 1032 °C, higher than the melting point of Y123. Compared with film C, the amount of Y211 is small. Then an epitaxy growth of Y211 in the c-axis direction was evident while the Y211 phase gradually emerged as a solid straight line (1040 °C, Fig. 14(d)). Most Y211 crystals had a 45° orientation, which is known as the preferential growth orientation of Y211 on the MgO substrate as discussed before. From the viewpoint of interface energy, the different melting modes of the two films can be elucidated, as illustrated in Fig. 15. Due to the co-existence of 0° and 45° grains, film C possesses a weak epitaxial interface associating with additional defects and grain boundaries, which stores a large excess energy and constitutes an extra driving force for the melting. As a result, the nucleation and growth of Y211 grains were encouraged. This is supported by the large number of Y211 grains and the appearance of grains with non-preferred orientation. The nuclei grew along the interface quickly and broke the entire bonding between the film and the substrate (see Fig. 15(b)). As we observed, many Y211 appeared while the film surface kept uniform. As a result, film C melts as soon as the large area Y211 grains emerges, in other words, it cannot be superheated. As to film D, the semi-coherent interface played a critical role in understanding the melting behavior. It’s well known that the energy of coherent and semi-coherent interfaces is significantly lower than that of noncoherent ones [46]. At the same time, the nucleation of melting requires a sufficiently large driving force. Owing to the combination of these two points, Y123 decomposition occurred at a temperature higher than Tp of Y123. The small quantity of Y211 can be another proof of the adequately suppressed nucleation rate. Moreover, the semi-coherent interface barriers retard the growth of melt as well (see Fig. 15(a)). When melt nuclei grew along the interface, breaking the strong bonding between film and substrate demands enough extra energy. So the growth of Y211 is slow enough to be captured by bare human eyes under optical microscopy. Unlike film C, only 45° grains occupied the majority of Y211 population at the outset of melting in film D. Regarding the first report of thin film superheating by Zhang et al. [13], the nucleation of melting could not be prevented and the superheating was due to the suppression of ‘‘melting growth’’. Combining their results with Tang et al. [45], it can be deduced that a fine epitaxial confinement restrains both the melting nucleation and growth, while a partly epitaxial interface only delays the melting growth. The suppression of melting nucleation is of vital importance in respect that even

Fig. 13. XRD pole figure of two YBCO thin films: (a) film C; (b) film D [45].

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Fig. 14. Optical micrographs showing the initial stage of melting of two YBCO thin films: (a) film C at 970 °C; (b) film C at 978 °C; (c) film D at 1032 °C; (d) film D at 1040 °C [45].

Fig. 15. Schematic illustration of the melting process of YBCO/MgO thin film, (a) film D, and (b) film C. The red arrow indicates the melting growth direction.

partial melting may cause devices to be malfunctioning. In short, a poor in-plane alignment could be an excess energy source which encourages melting/decomposing at the interface; while a fine alignment, with lowered interface energy, could be an obstacle against melting/decomposition.

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(2) Different intrinsic structures Secondly, two kinds of c-axis oriented YBCO/MgO films with different four-symmetry in-plane orientations were in-situ investigated under 0.21 and 1 atm. oxygen partial pressures [47,48]. The films were denoted as film E (0° orientated grains) and film F (45° orientated grains), respectively, according to the pole figure results shown in Fig. 16. In pure oxygen atmosphere, dissimilar melting behaviors between these two kinds of YBCO thin films were observed during their melting process. Besides, after the heating process, the samples were taken out of the heating stage and the surface morphologies of the films were observed by the optical microscope. Fig. 17(a–c) shows the evolution of film E at several stages. When temperature reached 1060 °C, majority of the Y123 film decomposed (Fig. 17(b)). The morphology of film E at 1070 °C, with the proceeding peritectic reaction, is shown in Fig. 17(c). Obviously, the Y211 grains rapidly developed and quickly spread all over the surface. After the melting process, dense and small Y211 crystals can be observed from the optical micrograph (Fig. 18(a)), with preferential growth orientation on the MgO substrate (Y211h0 0 1i//MgOh1 1 0i), as shown in Fig. 18(b). With regard to film F, as shown in Fig. 17(f), the majority of the film at 1070 °C exhibited the same morphology as at the beginning, implying that an extensive melting behavior had not taken place yet. The YBCO grains of film F completely decomposed at 1095 °C (Tp of YBCO is about 1030 °C in O2). This means that a remarkable superheating exists in film F. Furthermore, it can be concluded that film F shows the better thermal stability compared with film E. Besides, it can be observed from Fig. 18(b) that various orientations coexist on the substrate of film F. In pure oxygen atmosphere, the melting modes of two films are similar to those in air. Furthermore, film E completely decomposed at 1060 °C, with the proceeding peritectic reaction. The YBCO grains of film F completely decomposed at 1020 °C. Obviously, both films presented superheating in air. However, it was notable that the thermal stability of film E (0° oriented grains) is higher than that of film F (45° oriented grains), which is in contrary to the situation in pure oxygen. The decomposition temperatures of the two films under both air and pure oxygen atmospheres are listed and compared in Table 3. In addition, previous works in the LPE growth also showed similar results. Under air atmosphere, a polycrystalline seed YBCO film with an eightfold symmetry resulted in a 0° LPE film, while it resulted in a 45° LPE film in pure oxygen atmosphere [49–51]. In other words, in pure oxygen, 45° oriented seed grains survived and prevailed in the preferential growth, while the 0° seed grains decomposed during this process. All the distinctions should be strongly correlated with both intrinsic and extrinsic conditions, namely, the structure and the melting/growing atmosphere, respectively.

Fig. 16. XRD pole figure of two YBCO thin films: (a) film E, (b) film F [47].

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Fig. 17. Optical micrographs showing the melting processes of film E and F in O2, respectively: (a and d) at the initial melting stage of about 1009 °C; (b and e) at 1060 °C; (c and f) at 1070 °C [48].

On the one hand, the differences of melting behaviors between film E and F in various aspects indicate their different melting modes. Similar to the crystallization process, in the melting process the solid may experience three states: stable state, metastable state and unstable state. According to the melting behaviors (see in Fig. 17), it is presumed that the film E is in the metastable state during the melting progress. Under this circumstance, Y211 epitaxially nucleated and grew on the substrate with a small driving force. This is supported by the well aligned Y211 crystals (h0 0 1iY211//h1 1 0iMgO). On the contrary, the film F underwent an unstable melting process, accordingly the Y211 crystals spontaneously nucleated with a large driving force. This presumption is supported by the large scale and irregular orientation of Y211 crystals on the MgO substrate. The differences of melting modes between film E and F could be attributed to the difference in their thermal stabilities.

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Fig. 18. A comparison of the surface morphology of YBCO films after the melting process: (a) film E, (b) film F [48].

Table 3 The differences in the melting behaviors and melting modes of films E and F grown in air and O2 atmosphere. The decomposition temperatures indicate the temperatures at which the YBCO thin films start and finish decomposition, respectively. Atmosphere

Film

Air

E(0°) F(45°)

Decomposition starting temperature (°C) 970 967

Decomposition ending temperature (°C) 1060 1020

O2

E(0°) F(45°)

1009 1007

1070 1095

On the other hand, the difference in the thermal stability ought to be strongly correlated to their intrinsic structures. The most important distinction between two films is the in-plane orientation, i.e., 0° oriented YBCO (film E) and 45° oriented YBCO (film F). Firstly, the geometrical factor should be taken into consideration. It is known that the lattice constants of YBCO vary with the oxygen content. The calculated lattice mismatch of the YBa2Cu3O7d grains (for different oxygen contents, d) with the MgO substrate is listed in Table 4. The data indicate that the variation of lattice constants caused by oxygen partial pressure has only a negligible effect on the lattice mismatch between the YBCO films and the MgO substrate. Note that the 45° in-plane orientation can keep the lattice matching degree less than 10% only under the matching of three YBCO units on two MgO units. However, under that mode, the interface between YBCO film and MgO substrate have to endure more dangling bonds or larger internal stress, which lead to the instability at the interface [19]. The coincidence of reciprocal lattice points calculation of REBCO/MgO system has been reported by Matsuda et al. [52]. The result shows that among the rotational disorientations on the h0 0 1i REBCO and h0 0 1i MgO, the 0° in-plane orientation gives the highest reciprocal point coincidence, while for the 45° in-plane orientation it is smaller. In other words, for YBCO/MgO system, the 0° in-plane texture is more stable than the 45° one in terms of the geometrical lattice matching. This conclusion is coincident with the experimental result that film E (0° oriented) is more stable than film F (45° oriented) in air during the heating processing. Furthermore, the terminal plane of the epitaxial film, which has close relationship with the energy at the REBCO/MgO interface, is also important to estimate the thermal stability of YBCO thin film. Matsuda et al. have reported the HRTEM result of 0° oriented YBCO film, which shows that 0° oriented YBCO grains are terminated by BaO layers [33]. On the contrary, the atomic configuration of film F which is 45° oriented film shows great difference from that of 0° film. Fig. 19(a) is the HRTEM image of the atomic configuration in the vicinity of the interface between film F and MgO substrate, and the inset is the diffraction pattern showing that the in-plane epitaxial orientation is 45°. The amplified part from the interface, as shown in Fig. 19(b), indicates that there are two unexpected rows of atoms between MgO and the BaO–CuO–BaO blocks of Y123. These arrays have been deduced as a Cu2O

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Table 4 The lattice mismatch of YBa2Cu3O7d with various oxygen contents and the MgO substrate for the 0° and 45° in-plane orientations. In order to achieve the best match, the number of units in the calculation was also varied. d

Lattice constants (Å) (measured data)

h1 0 0iY123//h1 0 0iMgO (0°) 1YBCO unit on 1MgO unit (%) (calculated data)

h1 0 0iY123//h1 1 0iMgO (45°) 3YBCO units on 2MgO units (%) (calculated data)

0.07

a = 3.8227 b = 3.8872

9.221 7.689

3.710 2.086

0.27

a = 3.8275 b = 3.8875

9.107 7.682

3.589 2.078

0.40

a = 3.8349 b = 3.8851

8.931 7.739

3.403 2.139

0.52

a = 3.8415 b = 3.8778

8.775 7.913

3.237 2.322

0.62

a = 3.8510 b = 3.8700

8.549 8.098

2.997 2.519

0.72

a = 3.8621 b = 3.8621

8.285 8.285

2.718 2.718

0.91

a = 3.86 b = 3.86

8.335 8.335

2.771 2.771

nanolayer [53], based on the following two reasons: (1) Afrosimov et al. have demonstrated that Cu2O possibly exists at the YBCO/MgO interface on the grounds of the middle-energy ion scattering study [54]. (2) The inserted atomic rows are shifted by about 0.21 nm along h1 0 0i MgO, which is consistent with the phenomena of Cu2O (0 0 1) thin-film growth on MgO substrate [55]. This HRTEM work revealed that the terminal layer of film B (45° oriented) is Cu2O, which does not belong to the YBCO lattice. It is the variation in the terminal plane of YBCO films that also takes responsibility for differences in the film thermal stability and further in the melting mode. In general, from the consideration of surface energy, the formation of a more stable interface corresponds to an energy cut requiring lower work. According to the calculation work of Granozio et al. [56], the minimum energy cut of YBCO along the h0 0 1i corresponds to the interface between the BaO and CuO2 layer under an oxygen deficient status. This result indicates that the BaO plane is the stable terminal plane under air atmosphere. Namely, film E with BaO bottom layer should be relatively stable during the heating process. On the other hand, acting as an intermediate layer, the Cu2O plane of film F could be epitaxially stable instead of thermodynamically stable. Therefore, its instability facilitates the YBCO decomposition at a lower temperature. This result indicates that the 0° oriented YBCO seed grains have a more stable interface than the 45° seed grains do, which is coincident with the experimental result that the thermal stability of film E (0° oriented) is better than that of film F (45° oriented) in air. However, under pure oxygen atmosphere, the thermal stability of the 45° oriented film F was considerably improved, getting even better than that of film E (0° oriented). From the chemical phase diagram, we see that the decomposition temperature of the oxide in a higher oxygen pressure is higher than that in lower oxygen pressure [57]. The thermal stability of Cu2O in various atmospheres (air and pure oxygen) by HTOM was also investigated. The structure of the Cu2O layer between the YBCO and the MgO substrate was simulated by smearing Cu2O powder on the MgO substrate surface. The Cu2O powder was heated in the same manner as the YBCO thin film. The melting temperature of the Cu2O powder in oxygen (1149 °C) is much higher than that in air (1076 °C), which is qualitatively consistent with the phase diagram analysis. Thus, it can be concluded that for film B, the most probable reason of its higher thermal stability in oxygen is the enhancing thermal stability of the Cu2O terminal nanolayer, which suppresses the nucleation of melting and shifts its occurrence to higher temperatures. In short, the in-plane orientation effect on film thermal stability can be concluded by two factors. One is the geometrical lattice matching. The other is the terminal plane at the film/substrate interface. More importantly, the thermal stability of a REBCO film may be significantly enhanced by inserting an

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intermediate buffer nanolayer. It is expected that the elucidated mechanisms of melting and thermal stability of thin films correlation with various in-plane orientations will be widely utilized in related experimental and theoretical works and will help to promote applications of 0° and 45° oriented YBCO thin films in electronics. 3.1.2. Substrate materials For two-dimension thin films, the substrate (sub.) materials significantly influence the intrinsic structure at the film/substrate interface, which can be regarded as another respect of the effect on the thermal stability. In addition, the mechanism of the substrate effect on the thermal stability of superconductor thin films is potentially universal and applicable for more thin film materials. To understand the substrate effect, study on thermal stability of YBCO thin films (c-axis dominated with fourfold symmetry) deposited on assorted substrates was performed in air, by means of in-situ observation [58]. The YBCO films (see Table 5 for details) were respectively deposited on MgO, SrTiO3 (STO) and LaAlO3 (LAO) single-crystalline substrates by varied fabrication processes. Since at present, these three substrates are the most popular substrates which are used to fabricate YBCO thin films. It is believed that they can represent the typical YBCO film/substrate interface constructions. It can be observed that films deposited on different substrates show different thermal stabilities as well as melting behaviors. Fig. 20(a–c) demonstrates the evolution of YBCO/LAO thin film at several stages. Some acicular crystallites of Y211 phase could be seen to epitaxially grow on the film in the initial stage of melting at 960 °C, with the orientation of Y211h0 0 1i//LAOh1 1 0i. The Ba–Cu–O liquid was observed to partially spread on the LAO substrate, which suggests the liquid slightly wets the substrate. Noticeably, not until the temperature reached 1105 °C, which is nearly 100 °C higher than Tp of YBCO, the Y123 film totally decomposed. The melting evolution strongly confirms a high superheating capacity of over 100 K in the YBCO/LAO film. For the YBCO/STO film, it started to melt at 970 °C (see Fig. 20(d)), and then completely decomposed when the temperature reached 1060 °C (see Fig. 20(f)). Y211 phase with multiple orientations simultaneously occurred, dominating along the h1 1 0i orientation of the STO substrate. Fig. 20(e) shows a better wettability of the STO substrate by the liquid

Fig. 19. The HRTEM image of the atomic configuration in the vicinity of the interface between the 45° oriented YBCO thin film and the MgO substrate. (a) The general view of the boundary region of the 45° oriented YBCO thin film with MgO substrate. The inset is the diffraction pattern of the 45° YBCO thin film, which determined the in-plane orientation of this film. (b) An enlarged image of the square area on figure (a), from which the Cu2O plane can be clearly observed [48].

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compared with LAO. As displayed in Fig. 20(g–i), the YBCO/MgO film fully decomposed at 1040 °C, and the c-axis Y211 grains show an epitaxial orientation relationship with the MgO substrate: Y211h0 0 1i//MgOh1 1 0i. In addition, the liquid characteristically does not wet MgO since most liquid appears around Y211 crystals or conglobately isolated on the MgO substrate, leading to an exposed substrate shown in Fig. 20(h). In brief, all these YBCO films can evidently survive at a high temperature above its Tp, suggesting a universal superheating phenomenon. However, the superheating property and the peritectic melting mode show great differences among these three films. It is noteworthy that the YBCO/LAO film has the highest superheating level. Besides, the wetting behavior of the Ba–Cu–O liquid on each substrate is obviously distinct from each other as well. Apparently, the film melting starts at the temperature (Tstart) when the Y211 phase nucleates, and ends up at the temperature (Tend) when the Y123 film completely decomposes. On the one hand, the Y211 phase preferentially nucleates at defect sites, due to their high free energy [29,30]. Therefore, Tstart is a pre-melting temperature, and a crystallinity-dependent factor related to microstructure characteristics. For thin films with good qualities, this pre-melting should have a negligible effect on its stability against a complete melting. On the other hand, a complete decomposition strongly correlates with the bonding energy between film and substrate, as well as the melting mode. Thus Tend represents the nature of the film/substrate construction and intrinsically indicates thermal stability of thin films. In order to gain reliable experimental results, more YBCO/substrates, prepared by different approaches, were used for in-situ observations. The values of Tstart and Tend of all samples are summarized in Fig. 21. With respect of Tstart, the discrepancy can be attributed to the different deposition methods. Notably, Tend values of the YBCO films on various substrates clearly display a descending order: YBCO/LAO > YBCO/STO > YBCO/MgO films. Remarkably, results of the YBCO/LAO films deposited by thermal co-evaporation process show the highest Tend value among other films at about 1120 °C. Such a superheating degree of over 100 K is exceedingly high in comparison with commonly reported results, which is certainly interesting for both fundamental study and practical application. Evidently, the thermal stabilities of YBCO films (represented by Tend) varied with different substrates. With respect to the substrate effect, three factors should be considered: involving in lattice misfit at film/substrate interface, substrate wettability by the liquid and elements doping from substrate. Firstly, for thin films feature a low energy surface and have a relatively higher nucleation barrier at the surface than at the interface for peritectic decomposition, normally the melting is supposed to initiate at the interface between film and substrate. Consequently, as a key factor which relates to the interface bonding and energy, the lattice mismatch between the substrate and the film is considered. Furthermore, for functional oxides, the phase transformation from solid to liquid is characterized by a peritectic melting (a ? b + liq., as mentioned at the beginning of this review). Therefore, the lattice match of substrate with both reactant phase a and product phase b of peritectic melting should be considered. (1) Basically, the large lattice mismatch of a/substrate will lead to dangling bonds, dislocations or distortion of the crystal lattice, associated with instability sites where the melting initiates [59,60]. In other words, the good lattice fit relationship of substrate with a phase intrinsically

Table 5 Details of various YBCO thin films: substrate material, deposition method, substrate lattice constant and crystal system. Film

Deposition method

Substrate lattice constant (nm)

Substrate crystal system

YBCO/MgO

Thermal co-evaporation Magnetron sputtering

0.4213

Cubic

YBCO/LAO

Thermal co-evaporation Pulsed laser deposition

0.3821

Cubic (T > 435 °C)

YBCO/STO

Magnetron sputtering Pulsed laser deposition

0.3905

Cubic

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Fig. 20. Morphologies of YBCO films captured during the melting process on: (a) LAO at 960 °C, (b) LAO at 1080 °C, (c) LAO at 1110 °C, (d) STO at 970 °C, (e) STO at 1040 °C, (f) STO at 1060 °C, (g) MgO at 982 °C, (h) MgO at 1020 °C, (i) MgO at 1040 °C [58].

corresponds to a strong interface bonding and low interface energy. It is well known that for complete melting of thin films, it is a result of the thermal dynamic instability of interface bonding. Thus the lattice fit between a phase and substrate thermodynamically decide the thermal stability (Tend) of the film. (2) With a good lattice match with substrate, the decomposition product phase b of peritectic melting requires a small driving force for nucleation, i.e., the nucleation barrier is lower. In such cases, a high growth rate can be obtained with a small driving force. Considering the peritectic melting reaction, a strong capacity of b epitaxial growth along the substrate is equivalent to a fast growth of a melting. Therefore the lattice-fit of b/substrate kinetically determines the rate of b formation or a melting. Accordingly, the superheating property of films can be enhanced by minimizing the lattice mismatch at a/substrate interface, and enlarging the one of b/substrate. In the case of YBCO thin films deposited on various substrates, the value of lattice mismatch was calculated by using the equation: lattice mismatch = 2(a1  a2)/(a1 + a2), where a1 is the lattice constant of Y123 or Y211, a2 denotes the lattice constant of substrate. In addition, the substrate lattice mismatch with Y123 was calculated according to the Y123(0 0 1)//substrate(1 0 0) relation because all selected films are c-axis oriented. The mismatch of Y211/substrate was calculated, based on the epitaxial relation of Y211h0 0 1i//substrateh1 1 0i and Y211h1 0 0i//substrateh1 1 0i, on the basis of

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experimentally results and Refs. [31,32], respectively. The calculated lattice mismatches are shown in Table 6. On the one hand, the good lattice fit of Y123 on the substrate has a positive contribution to suppression of the Y123 dissolving. Hence Tend should strongly correlate to the substrate bonding with Y123 at the interface. For further comparison and discussion, in Fig. 22 we summarized the correlation between film thermal stability and substrate lattice mismatch with the film. The both x-axes denote the lattice parameter of the substrates. The y-scale in the lower part represents the film/substrate lattice mismatch, and the fitting curve for peritectic decomposition reactant Y123 is derived from Table 6. Correspondingly, the y-scale in the upper part represents the Tend value of films, and a fitting curve is gained from the HTOM in-situ observation work. Obviously, according to the experimental results, with decreasing the lattice mismatch between substrate and peritectic melting reactant (Y123), the film thermal stability (Tend) becomes increasingly strong. It means that Tend possesses a pronounced dependence on the lattice mismatch of films with the substrate. Among the three kinds of films, the inferior lattice matching of the YBCO/MgO film structure facilitates nucleation of melting, rationally leading to its lowest degree of superheating. On the other hand, it seems that the lattice fit relationship of Y211/substrate has an insignificant effect on Tend. To interpret why, we regard the melting behavior of the YBCO films. It should be noticed that the melting growth of Y123 is discontinuous, since BCO liquid has a poor wettability with Y123 and substrate but a good wettability with Y211 phase. The growth of Y211 grains separates the films, resulting in isolated Y123 islands without any liquid surrounding (see Fig. 20). Therefore, the up limit temperature (the intrinsic thermal stability) for a complete melting of isolated Y123 phase significantly relies on the Y123/substrate related interface bonding, while the lattice fit of Y211/substrate is not a decisive factor. Secondly, the wetting behavior of liquid from film melting is considered to be an important factor. As is well known, in the case of metal and its alloy, liquid immediately covers the solid as long as the melting occurs, which facilitates the melting growth. However, in the melting mode of YBCO film structures in this work, the liquid from film melting correlates with the parent phase 123 and the substrate. In the present circumstance of YBCO films on varied substrates, though the liquid wetting behaviors with the Y123 phases should be the same, the wettability of each substrate by the liquid presents dissimilarity, showing a descending order: STO > LAO > MgO, as demonstrated in Fig. 20. This result agrees well with the reported observation from the dipping experiment [61]. In principle, as

Fig. 21. Summary of decomposing temperatures: Tstart (colored data at bottom) and Tend (black ones on the upside) of YBCO films on MgO, LAO, STO substrates [58].

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discussed in Section 2.2, the poor substrate wettability by the Ba–Cu–O liquid will give rise to a liquid migration from the melting front, which suppresses the melting growth since a continuous driving force for nucleation of melting is required, associating with a higher Tend. In contrast, the good wettability of the substrate by the liquid may lead to a contacting (unmoving) liquid with the melting front, which facilitates film dissolving or decomposition, leading to low thermal stability. Therefore, although the Y123 phase has almost the same lattice-fit relationships with LAO and STO, the YBCO/ STO film shows a poorer thermal stability than the YBCO/LAO film as shown in Fig. 21, due to the wetting behaviors. Nevertheless, the experimental results also show incoherent relation with the substrate wettability. Tend of the YBCO/MgO film was minimal compared with other films although MgO has the poorest wettability by the liquid. The distinction in thermal stability of various YBCO films cannot be well explained only by substrate wettability. In brief, the substrate wettability plays an important part in melting growth, but does not critically determine the superheating property of YBCO thin films. Finally, apart from main superheating related concepts discussed above, the substrate effect should be also associated with the potential doping from the substrate element. Since during the thin film fabrication process, films may have some impurities from the substrate because of the element diffusion. The doping from the substrate such as Mg and La elements may be advantageous to the superheating property of films [62,63], while with Al and Ti doping the peritectic temperature of YBCO seems to be unaffected [64]. In contrast, Tp of YBCO was identified to be slightly reduced in the Sr-doped YBCO system. According to the experimental results, the YBCO/LAO film has a better thermal stability compared with the YBCO/STO film, which seems in agreement with the element doping effect. However, the YBCO/STO film shows a higher degree of superheating than the YBCO/MgO film, which is contrary to the doping effect. Accordingly, it can be concluded that the element doping from the substrate may have a potential effect on the thermal stability of thin films, but not be a main factor. To sum up, the study on substrate effect provides experimental evidences of a universal superheating mode on the basis of YBCO/substrate construction. In particular, the YBCO/LAO thin film was firstly observed to have a conspicuous superheating degree over 100 K. More importantly, the research gives a novel, versatile criterion for comparing thermal stability of thin films which possess peritectic decomposition with a low-energy free surface.

3.2. Phase diagram nature__primarily relating to melting growth Normally, the melting evolution of the RE123/sub construction (rather than the RE123 oxide) should exhibit three stages during the heating process. (I) Melting nucleation (heterogeneous nucleation) occurs at the defect sites (known as pre-melting) at a temperature around Tp, which depends on the film quality. In this case, the construction of the RE123/sub interface is quite stable. (II) Melting growth proceeds with the increase of the temperature, associating with weakening the RE123/sub. interface bonding. In this stage, the liquid, a melting product, expands on the film, giving rise to the dissolution of the RE123 phase. (III) Melting catastrophe (homogeneous nucleation) takes place at the RE123/sub. interface when the film is heated to an up limit temperature (Tup). In other words, the interface bonding between the film and the substrate completely collapses caused by the thermal vibration. It should be pointed out that a complete melting (decomposition) of a RE123 film could happen either at the stage II or the stage III. Practically, the artificial RE123 films with perfect crystallinity Table 6 Lattice mismatch of Y123/substrates and Y211/substrates.

Y123h1 0 0i//sub.h1 0 0i Y123h0 1 0i//sub.h0 1 0i Y211h0 0 1i//sub.h1 1 0i 3Y211h1 0 0i//4sub.h1 1 0i

MgO

LAO

STO

0.0952 0.0828 0.0515 0.0995

0.0024 0.0153 0.0465 0.0066

0.0194 0.0064 0.0243 0.0237

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Fig. 22. Correlation between film thermal stability (Tend value in this work) and substrate lattice-mismatch with peritectic decomposition reactant (Y123/substrate) [58].

cannot be obtained, that is to say, defects cannot be avoided in the real cases. Therefore, kinetics during the stage II plays a key role in the complete melting of the RE123/sub construction. In other words, in most cases, the complete melting would like to take place at a temperature lower than Tup, which results from the kinetic factor (e.g., high supersaturation) at stage II rather than the complete collapse of interface bonding at stage III. In this section, in view of phase diagram, the temperature coefficient of solubility (t.c.s.) effect relating to both oxygen partial pressure and RE123 system on the melting kinetics is reviewed.

3.2.1. Superheating related supersaturation It should be pointed out that the melting growth of the RE123 thin films is in essence a process of RE element transfer from RE123 phase to RE211 phase. In other words, melting of a RE123 phase correlates with growth of a RE211 phase according to the peritectic reaction: RE123 (solid) ? RE211 (solid) + Ba–Cu–O (liquid). The RE element released from RE123 to Ba–Cu–O liquid provides the nutrition which sustains the growth of RE211. As is well known, at a temperature below Tp, RE211 liquidus extends its curve into the RE123 phase + L region (see Fig. 23), where the liquid is in a so-called supercooling status. For one given supercooling, it corresponds to the supersaturation of the RE solute in the liquid. It is the RE supersaturation that acts as a driving force maintaining a continuous growth of the RE123 phase. Similarly, above Tp, the RE123 liquidus extends its curve into the RE211 + L region (see Fig. 23). In contrast with the supercooling, this status is known as superheating. The RE element is supersaturated in the liquid as well. The supersaturation (r), which sustains the growth of the RE211 phase induced by the dissolving of RE123 phase, can be illustrated by the extended part of the RE123 liquidus. Obviously, the study of this correlation is of key importance for understanding the melting behavior during the melting growth procedure. On the other hand, taking account of the melting kinetics, the melting rate of RE123 obviously has great influence on its thermal stability. Therefore, a simplified model is given under an assumption of a

123

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Fig. 23. An illustration of the supersaturation for the growth of RE123 and RE211. Similar to the supercooling, the superheating can be represented by the extended part of the liquidus into RE211+L region [68].

quasi-equilibrium condition during the melting process, aiming at the study of the correlation between the melting rate of RE123 and the growth rate of RE211. It is well-known that the density of RE element outflow from RE123 to the Ba–Cu–O liquid (JIn) is proportional to the RE element concentration gradient at the border between the solid RE123 and the liquid [65,66]. Also, the density of solute flow from the liquid to the RE211 growth front (JOut) is proportional to the RE element concentration gradient at the interface of RE211/liquid. When the temperature is kept constant or has a minor change, we could suppose that the system is in a quasi-equilibrium status, which means the RE element concentration in the liquid is constant. Thus, RE element inflow (JIn) to the liquid and outflow (JOut) from the liquid should be equal to each other,

J In ¼ J Out

ð1Þ

Meanwhile, the density of the solute inflow (from the melting of RE123) can be described as 123 J In ¼ ðDC 123 melt ÞRmelt

ð2Þ

123 where DC 123 melt is the difference of the solute concentration between the RE123 solid and the liquid, Rmelt is the melting rate of RE123. Correspondingly, the RE element outflow, JOut, from the liquid follows the equation

211 J Out ¼ ðDC 211 grow ÞRgrow

DC 211 grow

ð3Þ R211 grow

where is the RE element concentration difference between the RE211 solid and the liquid, is the growth rate of RE211. Therefore, the correlation between the melting rate of RE123 and the growth rate of RE211 can be written as 211 123 211 R123 melt ¼ ðDC grow =DC melt ÞRgrow

ð4Þ

It means the melting rate of RE123 is proportional to the growth rate of RE211 and the ratio of con123 centration differences at the growth front and the melting border ðDC 211 grow =DC melt Þ. Furthermore, as is generally known, the growth rate of a crystal is proportional to the solute supersaturation. Thus, the growth rate of RE211 phase in the liquid can be described as

R211 grow / r

ð5Þ

Here, r can be calculated as

r ¼ ðC L  C E Þ

ð6Þ

where CL is the RE element concentration in the supersaturated liquid, CE is the equilibrium concentration of the RE element in Ba–Cu–O liquid. Therefore, the tendency of R211 grow can be deduced from r.

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Fig. 24. A schematic illustration of the solute concentration at the melting front and the growth front [68].

Fig. 24 illustrates the principle of solute diffusion during the melting process. The film partially melts above Tp, leading to the RE element concentration at the interface of RE123/liquid higher than that under non-superheating conditions. The RE element concentration at that status can be gained by the extended part of liquidus of RE123 in the RE211 + L region. Thus, DC 123 melt between the RE123 solid and the liquid can be calculated as CS123–CL123, where CS123 is the RE element concentration in the RE123 solid, CL123 is the RE element concentration in the Ba–Cu–O liquid, which is close to the melting front of S L S RE123. Also, DC 211 grow can be calculated as C211–C211, where C211 is the RE element concentration of the L solid RE211, C211 is the RE element concentration in the liquid in front of the interface of growing RE211. Hence, the melting rate of RE123 can be described as 211 R123 melt ¼ dRgrow

ð7Þ (CS211–CL211)/(CS123–CL123).

where d is the ratio of concentration differences Obviously, the melting rate of the RE123 phase increases with both increasing growth rate of RE211 and d. Furthermore, on the one hand, for a given superheating, the correlating supersaturation is in direct proportion to the temperature coefficient of solubility (t.c.s.) of RE element in the Ba–Cu–O liquid. That is to say, the growth rate of RE211 is proportional to the value of t.c.s. On the other hand, since the solute concentrations in the solid phases are constants, the temperature coefficient of solubility of RE element in Ba–Cu–O liquid is of key importance to determine the value of d. In a word, the t.c.s. effect plays a significant role in controlling the melting rate of RE123 phase. 3.2.2. Temperature coefficient of solubility I. Different REBCO systems (RE = Y, Nd) Aiming at studying how the effect of t.c.s. on the film thermal stability, we investigate different members in the RE123 family. Fig. 25 illustrates the solubility of yttrium and neodymium in the Ba–Cu–O liquid as a function of temperature in air atmosphere. The dots in the liquidus are plotted according to the results reported by Krauns et al. [67]. The extended part of the RE123 liquidus into the RE211 + L region is plotted using the classical Freundlich model, which can be expressed as y = axb, representing the RE element concentration at the melting front. According to the previous discussion in Section 3.2, the tendency of the growth rate of RE211 phase can be deduced from the level of supersaturation. Fig. 26 shows the RE element supersaturation level in liquid as a function of the superheating degree in the RE123 film, calculated from Eq. (6), indicating that r in the Nd–Ba–Cu– O system is always higher than that in the Y–Ba–Cu–O system. As indicated by Eq. (5), the growth rate

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Fig. 25. Solubility of RE elements in the Ba3Cu5Oy melt under air atmosphere. The extended part of the liquidus are plotted by using the classical Freundlich model, corresponding to the RE concentration at the melting front [68]. The others are plotted according to experimental results from Krauns et al. [67].

Fig. 26. The supersaturation (r) of the RE element in liquid as a function of superheating degree in the RE123 film.

of the RE211 phase is proportional to the supersaturation level of RE element, it can be concluded that Y211 the tendency of RE211 growth rate is: RNd211 grow > Rgrow . Furthermore, according to the Eq. (7), apart from the growth rate of RE211, melting rate of RE123 rises with increasing d as well. Fig. 27 shows the factor d as a function of the superheating degree (DTs) in the RE-Ba–Cu–O systems, calculated from Eqs. (4) and (7) for Y123 and Nd123, on the basis of solubility curves reported by Krauns et al. [67]. It is clear that with increasing the superheating degree, d greatly increases in the Nd–Ba–Cu–O system, while it presents a trivial upward tendency in the YBCO system. Thus considering both trends of the factor d and the RE211 growth rate as previously mentioned, a conclusion can be made that with a larger t.c.s. value, the film melting rate is much higher, particularly

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Fig. 27. d as a function of DTs showing that d for Nd123 is higher than that for Y123 [68].

Fig. 28. Top views of the YBCO and NdBCO bulks seeded by homo-seeds [68].

in a higher superheating status. Furthermore, a higher R123 melt facilitates a complete melting at stage II before the interface bonding collapse caused by the thermal vibration. That is to say, the film thermal stability should show an inverse proportion to the t.c.s. value of the system in melting kinetics point of view. For experimental work, by the melt-growth (MG) method, thermal stability of Y123 and Nd123 thin films deposited on MgO substrates were investigated [68]. Both films are c-axis oriented with fourfold symmetry. Several MG experiments were performed to determine the superheating capability (DTsc) of the thin films. Fig. 28 shows the top views of YBCO and NdBCO bulks, seeded by YBCO/MgO and NdBCO/MgO thin films, respectively. Table 7 lists TMGup and Tp for Y123 and Nd123, where TMGup is the highest temperature, above which the film decomposed and failed in inducing the bulk growth, Tp is the peritectic reaction temperature for RE123 in that furnace. Thus a large temperature difference (DTM) between TMGup and Tp corresponds to a high thermal stability of the film. It can be seen from Table 7 that both films show superheating property. However, the higher DTM indicates that the Y123 thin film is more stable than the Nd123 one, which is in agreement with the above discussions. In a word, after the heterogeneous nucleation at stage I, the following melting growth at stage II is controlled by the t.c.s. effect related melting rate. With a large t.c.s. value, a large melting rate of films is obtained in a low superheating status. As a result, the thermal stability of the film should decrease with increasing t.c.s. value of the RE element in the Ba–Cu–O liquid. In addition, the thermal stability

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Y. Chen et al. / Progress in Materials Science 68 (2015) 97–159 Table 7 The different thermal stability of the RE123 films investigated by melt growth. Melt growth Film

TMGup (°C)

Tp (°C)

DTM (°C)

Y123 Nd123

1033 1098

1000 1085

33 14

of the whole family of RE123 systems should be reasonably deduced by comparing their t.c.s. value in the Ba–Cu–O liquid. Furthermore, the tendency of the thermal stability of the RE123 thin films under various atmospheres different from air can also be inferred, which could be of key value for widespread applications and for fundamental understanding of the melting and superheating. II. Changed oxygen partial pressures (pO2 = 1, 21, 100%) As is well known, the oxygen partial pressure significantly influences the property of oxide ceramics. Firstly, the variation of atmospheres with different oxygen pressures diversifies Tp of RE123 [69]. Furthermore, the oxygen controlled melt growth process is generally utilized to suppress the RE/Ba substitute in LREBCO7y (LRE = Gd, Sm, Nd) system, which deteriorates the superconducting property [70]. In addition, the growth rate of REBCO bulks in an elevated oxygen pressure is higher than that in air [71]. With respect to the thermal stability of REBCO system, the oxygen affects the solubility of RE elements in the liquid and the t.c.s. value. To clarify that, the influence of various oxygen partial pressures on the thermal stability of Nd123/MgO thin film by HTOM was investigated [72]. Fig. 29 illustrates the solubility of neodymium in the Ba–Cu–O liquid as a function of temperature under different oxygen partial pressures (P(O2) = 1%, 21%, 100%). The liquidus curves are plotted according to experimental results [64]. Obviously, the t.c.s. value increases with increasing P(O2). Fig. 30 shows the Nd element supersaturation level calculated from Eq. (6) for various atmospheres. One can see that r is proportional to the t.c.s. value relating to various atmospheres. The factor d illustrated in Fig. 31 shows a monotonic tendency with respect to the oxygen partial pressure. Obviously, the liquidus slope increases as the P(O2) increases from 1% to 100%, which indicates that d increases as

Fig. 29. Liquidus of RE elements in the Ba3Cu5Oy melt under various atmospheres. The extended part of the liquidus (dash-spot ones) are plotted by using the classical Freundlich model, corresponding to the RE concentration at the melting front [72]. The others are plotted according to experimental results from Krauns et al. [67].

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Fig. 30. The supersaturation (r) of the Nd element driving the Nd422 growth under different atmospheres.

Fig. 31. d as a function of DTs showing that d for Nd123 increases with increasing oxygen partial pressure [72].

the t.c.s. value increases. Taken the Nd422 growth rate into consideration, a conclusion can be made that the melting rate of Nd123 thin film increases as the oxygen partial pressure increases. It can be deduced that the Nd123 thin film under pure oxygen atmosphere is the least stable one, while the one under 1% oxygen pressure should show the best temperature endurance. 3.3. Correlations between film/substrate structure and phase diagram nature In the previous sections, the dependence of thermal stabilities of RE123 films on the film interface structure and the melting kinetics were discussed, affecting the melting self-nucleation and melting growth, respectively. Which one is dominant? What is their potential correlation?

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To answer these questions, taking both factors into account, the superheating capabilities of different REBCO phases (Y123 and Sm123) on substrates MgO and LAO were studied [73]. Firstly, regarding to the substrates, MgO, SrTiO3 (STO) and LAO are most popular used in fabricating RE123 thin films. MgO and LAO substrates, which have the largest and smallest misfit with the RE123 lattice respectively, their effects on film thermal stability have the largest difference and are appropriate for comparative study. Secondly, REBCO phases including Y123, Sm123 and Nd123 were used in previous superheating study. The absence of Nd123 is attributed to the lack of high quality Nd123/LAO thin films in hand. The selected films are divided into two groups according to the two kinds of substrates. The orientations of the films were characterized by c-axis dominated out of plane and fourfold in-plane alignment. The representative photos about the evolutions of the RE123 phases melting and the RE211 phases formation were presented. The morphologies of Group A (RE123MgO) films at various heating temperatures are shown in Fig. 32. Tend of the melting reaction was 1070 °C for Y123 and 1118 °C for Sm123 films, respectively, as can be seen in Fig. 32(c and f). For comparison, the typical stages of melting behaviors of Group B (RE123LAO) films are illustrated in Fig. 33. For Y123 film, though Y211 appeared at temperature (990 °C) below Tp, the film experienced a rather slow melting process, and completely melted at a much higher temperature (1113 °C). The decomposition of Sm123/LAO also ended up at higher temperature (1125 °C) compared with Sm123/MgO film. As one can see from the capability of superheating values in Table 8, all films show high superheating characteristic. Firstly, deposited on the same substrates, Y123 films show higher DTsc values than the Sm123 films, as an evidence of the t.c.s. effect. Secondly, thermal stability of the RE123/LAO film in Group B is better than RE123/MgO film in Group A, indicating the universality of the substrate effect. Thirdly, more important and notable, it can be found that by changing substrate from MgO to LAO, DTsc is 43 °C for Y123, and 7 °C for Sm123, respectively, indicating the thermal stability of the Y123 film is much more sensitive to the substrate variation than that of Sm123. In other words, substratechange-caused enhancement in thermal stability differs greatly for different RE123 phases. Because of the similar lattice constant of Sm123 and Y123, they have close lattice misfit with the same substrate. The minor difference between them may not cause such large difference. Therefore, besides the substrates effect, the t.c.s. factor should be taken into consideration as well. Fig. 34 shows the solubility of Y and Sm in the Ba–Cu–O liquid at different temperatures in air atmosphere according

Fig. 32. Optical micrographs of the Group A films captured at various melting stages [73].

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Fig. 33. Optical micrographs of the Group B films captured at various melting stages [73].

Table 8 Different thermal stabilities of the a phase films on MgO and LAO substrates prepared by the PLD method. Film

Tend (°C)

DTsc = Tend  Tp (°C)

Group A

Y123/MgO Sm123/MgO

1070 1118

60 53

Group B

Y123/LAO Sm123/LAO

1113 1125

103 60

Fig. 34. Liquidus of RE elements in the BCO melt under air atmosphere. The extended liquid lines (the dashed lines) are plotted by using the classical Freundlich model, corresponding to the RE concentration at the melting front. a represents Y123 and Sm123, b represents Y211 and Sm211, respectively [73].

to the reported solubility work [67]. Based on Fig. 34, the supersaturation, rY (square dotted line) and rSm (triangular dotted line) at different DTs are calculated, see Fig. 35.

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131

Fig. 35. DTs dependent supersaturation (r) which sustains growth of the b phase, calculated from Dc = cL  cE, combining the experiment results of DTsc values [73].

Obviously, at a given DTs, rY < rSm, since r is in direct proportion to the t.c.s. value. As the DTs increases, both rY and rSm which sustain the RE211 phase growth increase. According to the relation 211 R R123 melt / Rgrow / r, as r rises, the RE211 phase grows faster and RE123 melt more quickly. Hence, as s DT rises, the corresponding increase in rSm is larger than that in rY, owing to the t.c.s. effect. Then the growth rate of Sm211 phase exceeds, which favors the complete decomposition of the Sm123 films at stage II, resulting in a lower DTsc of Sm123 films compared with that of Y123 films. Furthermore, as the substrate materials change from MgO to LAO, the bonding at film/substrate interface strengthens, which is equivalent to a higher DTs that RE123 films can endure. That is to say, melting of RE123 phase delays to a higher DTs due to that substrate variation. For Sm123 with a large t.c.s. value, a slight increase of DTs relates to a substantial increase of r, leading to a distinctive enhancement of SmGrow 211 s as well as SmMelt 123 . Thus, substrate-change-caused increased-DT gives rise to fast melting of Sm123, leading to a small increase in DTsc value. On the other hand, for Y123 with a small t.c.s. value, an increase of DTs has an insignificant effect on the supersaturation. Thus, the substrate effect dominates the superheating capacity of Y123 phase, resulting in a large increase in DTsc with the substrate change. The experimental results, marked in Fig. 35, are in good agreement with the above discussions, i.e., the enhancement of the film thermal stability of YBCO (with a small t.c.s. value) is much more significant to the substrate variation than that of SmBCO (with a large t.c.s. value). In brief, combining experimental data and t.c.s. decided r-DTs curve, the correlation among the bonding of RE123/sub., t.c.s. value and the capability of superheating is clarified. 4. Exploits in fundamental study and engineering application of thin films with high thermal stability In the previous sections, the mechanism of film thermal stability was reviewed. Significant progresses have been obtained in figuring out the influences in the thermal stability of lowdimensional materials. These universal understandings provide effective approach to enhance the film thermal stability, which is of great importance for applications in both extending science research and technological industries. 4.1. New field for basic research: unconventional phase transition in high-superheating state As discussed in Section 3.1.2, it was indicated that YBCO/LAO thin film have a high superheating capability, remaining solid even if it is heated up to 100 K over the temperature of its peritectic

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melting point. This finding opens a new field for fundamental study in highly-superheating state. In this respect, for instance, unconventional phase transition, as a crucial subject for both fundamental study and practical applications, can be studied. More than 100 years ago, Ostwald mentioned an anomalous phase transition that the metastable phase with a smaller free-energy barrier was able to nucleate prior to the stable phase during the crystallization process [74,75]. After that, intensive research focused on the unconventional phase transition and phase competition in various materials, including metals and oxides [76,77,46,78–84]. However, most of these studies concentrated on deeply undercooled melts. From the aspect of achieving a highly superheating state, an unconventional phase transition was reported by employing HTOM [85]. Then the research interest on the unconventional phase transition and stabilization of materials was extended to more scientific and technological fields. Fig. 36 shows the typical morphologies of the YBCO film on the LAO substrate during the entire melting process. Fig. 36(d and e) shows the fact that instead of needle-like Y211 grains, an unknown phase with a square shape nucleated directly from the rest part of isolated YBCO films. The sides of these square grains are parallel to the h1 1 0i orientation of the LAO substrate. The surface morphology of the unknown square grains was observed by optical microscopy after the melting process, as shown in Fig. 37. It is obvious that most of those grains show identical epitaxial relationships with the LAO substrate. The complete decomposition of YBCO film occurred at a temperature up to 1108 °C, indicating it has a significantly high superheating capacity. By further heating up, the Y211 grains decomposed at 1180 °C as shown in Fig. 36(f), the resultant grains are similar to those ones that directly grew from the YBCO film. In addition, these grains grew very fast in the Y211 phase front. The possible reason should be that decomposition of needle-like Y211 phase provided continuous nutrient for the growth of those grains. According to the phase diagram, the grains resulted from Y211 phase decomposition should be the well-known cubic Y2O3 (Y200) phase. It seems that most grains originated from Y123 and Y211 decomposition are of similar shapes, with the same epitaxial orientation on the LAO substrate. Energy-dispersive spectrometer (EDS) was employed to identify the unknown substance which directly resulted from the YBCO film. Fig. 38(a) shows the scanning electron microscopy (SEM) image of the melted YBCO film. The typical square grains are substance directly nucleated from YBCO matrix,

Fig. 36. Morphologies of the YBCO/LAO thin film captured during the melting process at: (a) 965 °C, (b) 1000 °C, (c) 1065 °C, (d) 1080 °C, (e) 1108 °C, (f) 1180 °C [85].

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Fig. 37. The optical micrographs of the unknown square grains resulted from YBCO/LAO thin film, captured after the melting process [85].

Fig. 38. SEM micrograph and EDS spectrum of the YBCO/LAO film after melting. (a) The surface morphology of the obtained grains: the typical square grains are directly nucleated from YBCO matrix, while the one at the lower right corner without a regular square shape is the Y200 phase nucleated from Y211 phase; (b) the EDS mapping result of the whole region indicates those grains with gray color are of the same composition; (c) the EDS result of the region spectrum 1 shows the square grain consists of Y and O elements [85].

while the one at the lower right corner without a regular square shape is the Y200 phase nucleated from Y211 phase. The EDS result of the region spectrum 1 (see Fig. 38(c)) shows that the square grain consists of only Y and O elements. Since the EDS analysis was applied under low vacuum condition, i.e., the atmosphere contains oxygen element, it does not show a proportion of Y:O = 2:3. Moreover, the EDS element mapping result of the whole region is shown in Fig. 38(b). It is obvious that, all

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the square grains originated from the YBCO films are of the same composition, amazingly as well as the Y200. As a consequence, it was confirmed that those square grains nucleated directly from the YBCO film are definitely the same phase as the Y200 grains. It is known that Y211 is the only thermodynamically stable phase during the peritectic melting of YBCO phase at this stage. In other words, an unexpected peritectic melting: Y123 ? Y200 + liq. certainly occurred. For comparison, the phase formation behaviors of the YBCO films on STO and MgO were studied. However, those two films just perform a thermodynamically peritectic melting: Y123 ? Y211 + liq. The main difference between the LAO related film and the others is the superheating capability, which implies that the unconventional peritectic melting may correlate with the superheating capacity. Aiming at clarifying the distinct phase transitions in the superheating state during the peritectic melting, first of all, the phase diagram is taken into consideration. As shown in Fig. 39, the supersaturation for Y211 phase growth associated with the superheating status of Y123 phase can be described as

r1 ¼ DC 1 =C Y211 E On the other hand, as shown in Fig. 39, when Y123 phase is superheated above the temperature TcY200, at which the two extended liquidus curves intersect, the superheating involved supersaturation also supplies the driving force for the Y200 phase growth. The supersaturation for Y200 phase growth can be expressed as

r2 ¼ DC 2 =C Y200 E We term such a superheating state when Y123 phase is superheated above TcY200, a high superheating. According to the fitting result, the critical temperature TcY200 for the system in this work is about 1360 K as shown in Fig. 39. Based on the experimental results, the Y123 film on LAO can be superheated up to 100 K, which is apparently above the critical temperature TcY200. Therefore, it is the highly superheated Y123 film involved supersaturation that provides driving force for the Y200 phase growth. In other words, the Y200 phase could nucleate in this situation. In contrast, although the Y123 films on other substrates possess superheating capacities, they cannot survive above TcY200. Thus

Fig. 39. An illustration of the driving force for the growth of Y211 and Y200. Similar to undercooling, the superheating status can be represented by the extended parts of the liquidus curve, which equilibrates with the Y123 phase, into Y211+L region [85].

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there is no driving force for the Y200 phase nucleation. As a result, it can be concluded that the high superheating was a decisive precondition for the unconventional phase transition in the superheating state. Furthermore, it should be noticed that DC1 is larger than DC2, while CY211 is smaller than CY200 , as E E shown in Fig. 39, leading to r2 < r1. That is to say, the supersaturation r caused driving force for Y211 phase nucleation is obviously larger than the one for the Y200 phase nucleation. It means that the Y211 phase nucleation should be preferential to the Y200 phase, which is opposite to the present result. Consequently, there must be other factors affecting the phase transition in the highly superheated Y123 film on LAO. To search the other factors that affect the phase transition in high superheating, the nucleation competition concept [76,77] is adopted, which are commonly applied in the case of phase selection in deep undercooling states. In view of this concept, forming an interface with low free energy is equivalent to reaching a more stable state. In this way, the metastable phase might nucleate preferentially over the stable phase, which is the so-called epitaxial stabilization phenomenon [46,78–81]. Consequently, as a key factor relating to the interface energy, lattice mismatch between phases and substrates should be taken into account. The calculation of Y211/substrate is based on the relationship reported in previous work [31,32]. The calculation of Y200/substrate is based on the dominant epitaxial relationship Y200h1 0 0i//substrateh1 1 0i observed in the experiment. The calculated results for various interfaces are given in Table 9. Compared with the cubic Y200 phase, despite the fact that the Y211 phase in the a-axis orientation has a smaller lattice mismatch with respect to LAO, the lattice mismatch of Y211 with LAO is larger in the c-axis orientation. Therefore, the Y200/LAO interface may have a lower free energy, which is equivalent to a small formation work for nuclei. In terms of the nucleation competition concept, if the formation work for Y211 nuclei is larger than that for Y200 nuclei, the metastable phase Y200 may nucleate prior to the stable one Y211. As a result, on the LAO substrate, the metastable Y200 phase formed as a primary phase in the highly superheating state. In short, the lattice fit at the Y200/substrate interface, i.e., the formation work for nuclei is a crucial factor that affects the phase selection in the high superheating state. In summary, an unexpected phase transition during the peritectic melting in the high superheating state was observed. From the highly superheated Y123 film, the metastable Y200 phase grew epitaxially on LAO with the orientation of Y200h1 0 0i//LAOh1 1 0i, while the conventional stable Y211 phase appeared from the superheated Y123 films on both STO and MgO. It is clear that this phase transition happens only if the Y123 film can be superheated above a critical temperature (TcY200), and the lattice mismatch of Y200/substrate is smaller than that of Y211/substrate. This approach suggests a possible method, i.e., for materials with a substrate, studying their transition evolutions and possible metastable phases in the highly superheating state. More interesting subjects can be extended in this recently-opened realm for basic research. 4.2. As seed materials: effectively inducing the growth of high performance REBCO crystals 4.2.1. Top-seeded melt-growth and progresses of seed materials searching It is well known that thin films can act as a seed material in the preparation process, such as TopSeeded Solution-Growth (TSSG), Top-Seeded Melt-Growth (TSMG), and Liquid Phase Epitaxy (LPE). Among these techniques, TSMG process is a mature one for batch preparation of well-oriented REBCO crystals (c-axis oriented bulks) with large sizes. Up to now, REBCO (RE = rare earth elements except Ce and Tb) crystals have shown an extraordinary ability in carrying high critical current density and trapping large magnetic fields [86-88]. As a

Table 9 Comparison of the lattice mismatch of Y211/substrate and Y200/substrate. Phase/sub. a c

Y211h0 0 1i//LAOh1 1 0i 0.0465

3Y211h1 0 0i//4LAOh1 1 0i

Y200h1 0 0i//2LAOh1 1 0i

0.0019

0.0194

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Fig. 40. Schematic illustration of the equipment of top-seeded melt growth (TSMG).

result, RE-Ba–Cu–O crystals have great potential for large-scale engineering applications, such as magnetic bearings and hysteresis motors. Due to the large anisotropy and quantum flux penetration mechanism of copper-based superconductor, it is necessary to control the growth orientation of REBCO bulks. In order to obtain large REBCO c-axis oriented bulks, the TSMG method is commonly used [89,90], as illustrated in Fig. 40. The seed placed at the top surface of the compacted pellets plays a crucial role in the process to achieve single-nucleation and uniform orientation. Normally, a qualified seed material should mainly follow three standards: (i) a higher melting point than the processing temperature, (ii) a similar structure and a good lattice match with the REBCO crystal, (iii) small reactivity with respect to the REBCO precursor. Many efforts have been made in searching suitable seeds. Firstly, non-superconducting seed materials, like MgO and SrLaGaO4, have remarkably high Tm, but the superconducting performance of the resulting REBCO bulk could be weakened by contamination and the large lattice misfit between seed and REBCO material [91]. Secondly, SmBCO and NdBCO single crystals, as traditional seed materials, are most commonly utilized to grow c-axis oriented REBCO bulks. The advantage of small lattice misfit and the growth restriction of lower Tm system coexist in this method. Alternatively, a special ‘‘hot seeding’’ process (see Fig. 41) offers the self-seeding pathway to grow crystals (NdBCO seeded NdBCO crystals, etc.) [92]. Since the seed is placed on preheating pellet at a temperature below the peritectic temperature of the RE123 phase, the melting point of seed does not have to be higher than the target grain. However, the requirement of a special furnace construction and the complicated seeding operation limit an extensive industrial application of ‘‘hot seeding’’. In most cases, the convenient cold seeding method is much more popular. From Fig. 42 one can see that the seed is placed on the pellet at room temperature before the heating process. During the slow-cooling process, the Y211 phase reacts with liquid to produce Y123 according to the peritectic solidification. Obviously, this cold seeding method is more convenient and beneficial for batch growth. However, it requires a sufficiently high thermal stability for the seed material to endure a long holding time at a maximum processing temperature (Tmax). Generally, Tmax is higher than the melting point of the precursor by at least 30 K, for complete melting of the precursor. Based on Table 4.2.1 [70,93–95], the peritectic temperature of REBCO materials increase with increasing RE ionic radii. Since SmBCO and NdBCO have higher melting points than YBCO, and similar

Fig. 41. A schematic diagram of hot-seeding TSMG of bulks.

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Fig. 42. A schematic diagram of cold-seeding TSMG of bulks.

crystal structures as YBCO, their bulks and single crystals are commonly used as seed materials for YBCO bulk growth [96–98]. On the other hand, compared with YBCO, LREBCO (LRE123, LRE = light rare earth element, e.g., Nd, Sm, Gd) bulk possess better superconductive performance. However, suitable bulks or single crystals are not available for cold-seeding melt-growth of the LREBCO materials due to their high melting points. Hence, large-sized LREBCO materials cannot be fabricated reliably, which severely hindered their developments. Recently, Babu and Cardwell et al. [99–103] developed a novel kind of generic seed which is valid for seeding most REBCO superconductors. As shown in Fig. 43 [63], Tp of NdBCO bulk is increased by about 20 K with Mg-doping, becoming the highest one among those listed in Table 10. Using the Mg-doped NdBCO as a seed for cold-seeding, a SmBCO bulk superconductor was successfully fabricated in air. Besides crystals, thin films also acted as suitable seed materials for bulk growth. Cai et al. first used highly c-oriented NdBCO/MgO thin films in cold-seeding melt-growth of low-Tp materials, i.e., YBCO/ Ag bulks [104,105]. Single domain growth with a relationship of YBCOh1 0 0ibulk//NdBCOh1 0 0ifilm was obtained, indicating the possibility of applying film-seed in melt growth. One important feature in this work is that Tmax in melt growth is lower than Tp of the film material: NdBCO. In short, for cold-seeding melt-growth (TSMG) of REBCO, there have been various difficulties caused by inferior properties of the traditional seed materials, particularly, insufficiently high thermal stability. In recent 10 years, following the discovery of its superheating phenomenon, the REBCO films have become superior seed materials for growing REBCO crystals in TSMG, leading to solving a series problems, from which researchers suffered for long by using traditional seeds. Firstly, homo-seeds are available in the cold-seeding method. Secondly, fabrication of difficult systems can be expected, such as growing high Tp LREBCO materials, recycling failed samples. Thirdly, a high Tmax (60 K + Tp) can be realized, which is beneficial to suppress heterogeneous nucleation, leading to growing large-sized REBCO crystals. Fourthly, long tolerance of film-seed at Tmax can be also expected, which favors the batch growth of REBCO crystals. Finally, more controllable TSMG processes can be attained with applications of film-seeds, such as tuning the contact angle of grains in multi-seeded

Table 10 Peritectic temperatures measured by DTA and solution method, the RE ionic radius, electrical and magnetic properties of various RE123 materials. RE

La

Tp (K) [70,93,94]

1363

Pr

Nd

1363 1244 ± 1 1359 ± 5 1341 1341 I.R. (Å) [70] 1.160 1.109 Tc (K) [94,95] 98 95 Jc (lT  KA/cm2); 77 K [94] 40–50 Birr (T); 77 K [94] 8–10

Sm

Eu

Gd

Dy

1333 1333 ± 5 1327 1.079 93.5 40–50 6–8

1323

1303

1283

Y

1273 1278 ± 5 1319 1303 1283 1278 1.066 1.053 1.027 1.019 93 92.5 92 92 50–60 20–30 6–7 5–6

Ho

Er

Yb

1263

1253

1173

1278 1268 1238 1.015 1.004 0.985 92 92 92

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Fig. 43. Differential thermal analysis of NdBCO samples. It shows that the Tp values of Mg-doped NdBCO crystals shift to higher temperatures by about 20 K [63].

Fig. 44. Optical top view of an YBCO crystal seeded by an Y123 thin film [106].

melt-growth, gaining the large c-growth sector in crystals, fabricating single crystal in low supersaturated melt and so on. 4.2.2. Demonstrations of superheating capacity: thin-film as homo-seed in REBCO cold-seeding process Using an Y123/MgO thin film as a homo-seed, a success in cold-seeding melt-growth of a c-axis oriented YBCO bulk was reported for the first time [106]. Fig. 44 shows the optical top view of a grown crystal seeded by a c-axis Y123/MgO thin film, presenting a good orientation correlation of YBCOh1 0 0ibulk//YBCOh1 0 0ifilm//MgOh1 0 0isubstrate. In this process, the YBCO film withstood a maximum processing temperature (Tmax = 1045 °C) over its Tp for a long period of 1.5 h and acted as a seed for the YBCO bulk growth. It is the first experiment in cold-seeding melt-growth that demonstrated the unique superheating property of the REBCO film-seed since Tmax in that work is distinctively higher than Tp of the film-seed material. This magic success provides even stronger evidence than ever before with respect to the superheating phenomenon of the YBCO thin film. It should be noticed that different from the one in the HTOM and YSNG experiments, the film endure a much longer time at high temperatures. How the seed films survived after such a long time

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attracts interests. Taking dissolvability into consideration, one reason that may explain the long time surviving of the seed film is the composition of the liquid. In the YSNG method, before the film being put into the liquid, there was no yttrium in the liquid at all. The yttrium element can easily dissolve into the Ba–Cu–O liquid. The superheating property could not prevent the YBCO film from dissolving, which may be the main factor resulting in the quick disappearing of Y123 grains. In the MG experiment, however, the composition of the precursor is Y:Ba:Cu = 1.7:2.35:3.35, which means high level of yttrium exists in the liquid. In such a Y-saturated liquid, the Y diffusion and the Y123 decomposition from the film is effectively suppressed. Another possible reason that may strengthen the stability of Y123 thin film against melting is the high ratio of Ba:Cu in the melt. A tendency was reported that the dissolvability of rare earth element, such as Nd, depresses with the increasing ratio of Ba:Cu in the melt [107]. This tendency may also exist in the Y–Ba–Cu–O system even if its influence is not distinct. The precursor composition of Y1.7Ba2.35Cu3.35Ox features a relatively high ratio of Ba:Cu, and corresponds to a lower Y dissolvability in the process of melting growth. Thus the occurrence of the Y diffusion and the Y123 decomposition from the film become difficult. The third possible reason is the coarsening of YBCO grains in the seed film during the long time heating. In the YSNG method, the seed film was heated within 5 s. The coarsening of crystal could not manifest itself in such a short period of time. But in the MG experiment, because the Y123 film was heated for several hours, the selfcoarsening may become remarkable. This effect could reduce the grain-boundary density of the seed film, and even the film itself may become thicker by absorbing nutrient from the precursor below. According to Griffith et al. [108], a small grain-boundary density leads to low nucleation rate, which may thwart the melting of YBCO/MgO thin film and enhance its superheating capacity. So the coarsening effect seemingly enhance the thermal stability of the seed film. In a word, YBCO thin films maintain solid status during the whole MG process and can act as seeds effectively. Following above YBCO film-seeded growth of YBCO bulk, Oda et al. employed SmBCO films as homo-seeds, and firstly reported a fully grown SmBCO/Ag crystal with 30 mm in diameter, under low oxygen partial pressure [109]. The Tmax used in the heating procedure is 1045 °C, about 15 °C higher than Tp of the film material: SmBCO (1030 °C in reduced oxygen atmosphere, about 35 °C lower than its Tp in air), and the holding time is one hour. Fig. 45 shows the top view picture of the as-grown superconductor. The clear fourfold growth facet lines indicate that a single grain was successfully gained in an optimized melt-growth process, confirming the superheating capacity of the SmBCO film-seed. It should be noticed that the trapped field of the sample is 1.52T at 77 K (see Fig. 46). This is the first report of the exciting success that a large-sized SmBCO/Ag crystal with such a high performance was fully grown, while a high Tmax (>Tp of the film material) was applied by taking advantage of superior superheating nature of the SmBCO film-seed.

Fig. 45. The top view of the as-grown SmBCO crystal by using SmBCO/MgO thin film as a cold seed, with 30 mm in diameter [109].

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Fig. 46. The trapped field distribution at 77 K of the as-grown SmBCO crystal [109].

It should be pointed out that, when applying thin film as seeds, the as-grown REBCO crystals are inevitably contaminated by Mg from the MgO substrate, which causes a decrease of Tc. To solve this problem, a mini-pellet should be used between the film-seed and the main pellet [110]. In brief, developments of REBCO thin films with superheating capacities as homogeneous seed materials certainly contribute to the reliable fabrication of large-sized bulks with high superconductive properties. 4.2.3. YBCO-buffered NdBCO film with enhanced thermal stability for growing higher processingtemperature required REBCO crystals So far, by employing REBCO/MgO thin films as seed materials, most kinds of REBCO (RE = Y, Gd, Sm) bulks have been successfully induced by the TSMG method. Particularly, LREBCO bulk superconductors have gained widespread attentions for superconducting applications, since they exhibit higher Tc, Jc and trapped field values even in an external field in comparison with YBCO [111]. However, even though NdBCO has the highest Tp among the commonly applied REBCO materials, as it is in the structure of a thin film, it has a superheating upper limit at about 1098 °C in melt growth as mentioned above, only 14 °C higher than its Tp. This demerit preclude NdBCO/MgO film from being used to seed the growth of high-Tp LREBCO crystals, e.g., RE = Sm or Nd. Recently, by growing an YBCO buffer between the NdBCO film and the MgO substrate, namely, gaining NdBCO/YBCO/MgO film, the manufacturer recognized its effect on broadening the growth window of high-quality NdBCO films [112]. With this special NdBCO-film/YBCO-buffer-layer/ MgO-substrate structure, better crystallinity and higher thermal stability of the NdBCO film can be expected. Fig. 47 shows the pole figure result of the NdBCO films with and without an YBCO buffer layer. Small unexpected peaks (satellite peaks) in the NdBCO/MgO film do not appear in the YBCO

Fig. 47. X-ray pole figure of: (a) NdBCO/MgO thin film; (b) NdBCO/YBCO/MgO thin film with a 25-nm YBCO buffer layer [114].

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buffered NdBCO film, revealing a good fourfold symmetry. Convincingly, it is the YBCO intermediate layer that plays a role in improving the in-plane alignment of the NdBCO film. Similar to the sandwich structure of REBCO/buffer/substrate in coated conductors, the YBCO buffer layer in NdBCO/YBCO/MgO thin film performs functions in: (1) compensating the lattice misfit between NdBCO and MgO; (2) alleviating the problems related to the interface from the substrate [113]. Due to the excellent in-plane alignment, the grain boundaries in the NdBCO/YBCO/MgO thin film are much fewer than those in the NdBCO/MgO thin film. Therefore, an enhanced superheating capability of NdBCO/YBCO/MgO thin film can be expected. For the first time, Xu HH et al. applied an NdBCO/YBCO/MgO thin film as a cold seed to induce homo-epitaxy, and the full growth of a single-grain NdBCO crystal without the Ag addition was gained (16 mm in diameter) [114]. Remarkably, this NdBCO/YBCO/MgO film seed has endured the Tmax of 1117 °C for 1.5 h, which is about 20 °C higher than the Tmax (1098 °C) that NdBCO/MgO thin film could survive [115]. It is evident that the NdBCO/YBCO/MgO thin film possesses a higher superheating upper limit. More importantly, using a higher Tmax in the MG process is of great assistance in growing large-size and high-performance REBCO bulks [105]. Fig. 48 shows the top view of the as-grown NdBCO crystal with a size of 16 mm in diameter, presenting fourfold growth facet lines clearly. Furthermore, Fig. 49 shows a top view of the as-grown SmBCO sample, which has a world-record size of 32 mm in diameter in the air-processed cold-seeding method [116]. It can be seen that the bulk has been grown from the NdBCO/YBCO/MgO thin film seed to the mini-pellet, and then to the big pellet. The NdBCO/YBCO/MgO thin film seed underwent a high Tmax of 1106 °C for growing that large sized SmBCO crystal. A homogeneous mini-pellet was placed between the seed and the sample pellet to impede the diffusion of Mg ions from the seed. Fig. 50 displays the superconducting transition of the SmBCO crystal, showing high Tc over 94 K, with a transition width less than 1 K (the difference in temperature between 10% and 90% of the

Fig. 48. The as-grown NdBCO crystal seeded by an NdBCO/YBCO/MgO thin film, using a Tmax of 1117 °C [114].

Fig. 49. Top view of the SmBCO crystal with a size of 32 mm in diameter [116].

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Fig. 50. The temperature dependence of the magnetization of Sm123 crystal grown in air [116].

Fig. 51. The critical current density as a function of the applied field for the Sm123 crystal at 77 K [116].

maximum diamagnetic momentum). The resulting high Tc should be ascribed to the addition of Sm2Ba4Cu2O9 (Sm242). Critical current density, Jc, as a function of the applied field at 77 K is shown in Fig. 51, exhibiting a maximum value of 68 kA/cm2 at a zero applied field. Meanwhile, the peak effect can be clearly observed at an external field, 38 kA/cm2 at 1.7 T (see Fig. 51). The enhanced Jc values indicate that the addition of Sm2Ba4CuBiOz (Sm2411) plays a significant role in improving Jc. On the other hand, in the recycling process of failed bulks, high Tmax is strongly required for a complete melting of solid bulk materials. Using YBCO-buffered NdBCO films for seeding and applying a Tmax ranging from 1105 °C to 1115 °C, various failed YBCO bulks were successfully recycled by a cold-seeding MG process [117]. Fig. 52 shows photographs of a failed YBCO bulk and the corresponding recycled bulk processed at Tmax of 1105 °C with the NdBCO/YBCO/MgO film seed. The

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Fig. 52. (a) A failed YBCO bulk. (b) The corresponding recycled bulk processed at Tmax of 1105 °C with a NdBCO/YBCO/MgO film seed [118].

Table 11 Various kinds of REBCO crystals seeded by the NdBCO/YBCO/MgO thin film and the Tmax of each process we succeeded in. RE123 systems

Tmax (°C)

Doped RE123 systems

Tmax (°C)

SmBCO GdBCO NdBCO

1115 1120 1117

Ag-GdBCO Ag-NdBCO

1115 1115

corresponding recycled bulk displays a good seed-induced single domain characterized by fourfold growth facet lines, indicating that the recycling process was successful. In short, the NdBCO/YBCO/MgO film presented a higher superheating level, about 20 °C higher (Tmax, up to 1120 °C) than that of non-buffered NdBCO film. Considering the similar crystal structure in the whole REBCO family, it is reasonable to assume that almost all kinds of REBCO materials could be seeded by NdBCO/YBCO/MgO thin films. Table 11 lists the various kinds of REBCO crystals effectively seeded by the NdBCO/YBCO/MgO thin film and the Tmax of each process we succeeded in. In addition, Fig. 53 shows the types of film seed that have been practically used and their reported Tmax values in the MG processes. Remarkably, Tmax as high as 1120 °C was tolerated, which is a Tmaxtolerating record in the homo-seeding melt growth of REBCO crystal to our knowledge. It can be seen that intrinsic Tp of REBCO significantly affects Tmax a film-seed can tolerate. What is more, the film/sub. construction affects Tmax significantly, as Tmax for NdBCO/YBCO/MgO film is about 20 °C higher than that for NdBCO/MgO film. In brief, YBCO-buffered NdBCO film with enhanced thermal stability favors the growth of higher processing-temperature required REBCO crystals. Seeding by the NdBCO/YBCO/MgO thin film, higher Tmax can be applied in the melt growth, which is beneficial to suppress the heterogeneous nucleation, broaden the growth window, and grow large-sized and high-Tp REBCO crystals. Besides, using NdBCO/ YBCO/MgO film seeds, a novel, convenient and effective process for recycling the failed REBCO bulks can be realized.

4.2.4. Long tolerability of film-seed with melt and two-layer batch growth of YBCO Nowadays, bulk superconducting magnets [118,119] are expected to enter the commercial market as basic elements of industrial applications in the near future. Batch processing has drawn much attention recently as a cost-effective method of REBCO materials with a uniform quality. Conventionally, samples placed in one layer were grown in one batch [120–126]. It can be imagined that processing samples placed in multi-layers in one batch is an even more commercially favorable way to fabricate REBCO bulk superconductors since the furnace capacity can be more effectively used. However, in the multi-layer batch growth, the temperature distribution is generally not so uniform

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Fig. 53. The types of film seeds that have been practically used and the Tmax they have been reported in the MG processes [118].

especially among the layers. Therefore, the seed ought to tolerate a high temperature for a long holding time. By employing the NdBCO/YBCO/MgO film seed with a high thermal stability, a two-layer batch growth of YBCO crystals was carried out [127]. Since the existence of a temperature gradient, the actual slow-cooling starting temperature of the top layer samples is about 15 °C lower than that of the bottom layer samples according to the pre-set heating procedure. Therefore, a seed/mini pellet/ main pellet construction was adopted for the bottom layer samples, because this construction has been confirmed to tolerate for a higher Tmax [128]. Three such samples were put on the bottom layer of the self-made SiC shelf. On the other hand, three pellets with the seed directly put on the top surface were placed on the top layer. Fig. 54 is a schematic drawing showing how the samples were arranged. 1100 °C was chosen as Tmax to widen the growth window and to make sure dissolution of the film did not occur at the bottom layer. Fig. 55 presents photographs of the as-grown YBCO bulks with sizes of 16 mm in diameter, bottom layer samples in Fig. 55(a), and top layer samples in Fig. 55(b), respectively. Obviously, a difference in the top view exists between the samples from the two layers. Furthermore, the trapped field distributions of two typical samples at both layers were characterized at 77 K as shown in Fig. 54. The maximum trapped fields were approximately 0.4 T for the top layer sample and 0.45 T for the bottom layer one, respectively. The distinct properties and macrostructures suggest that he growth mode of YBCO has some distinctions in the two layers. On the other hand, both field profiles exhibit a single peak at the center position of the sample, indicating that no weak-link exists in those bulk samples. Furthermore, both the clear fourfold growth

Fig. 54. A schematic drawing illustrating the self-made SiC shelf with two panels and the arrangement of the samples. Filmseeds with and without mini-pellets are put on the precursors in the bottom layer and the top layer, respectively [128].

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Fig. 55. Top views of the batch-processed YBCO crystals seeded by NdBCO/YBCO/MgO films. (a and b) Represent bottom layer and top layer samples, respectively [128].

Fig. 56. The trapped field distribution of the batch-processed YBCO crystals seeded by an NdBCO/YBCO/MgO film at 77 K. (a and b) Show the trapped field distribution of top and bottom layer samples, respectively [128].

Fig. 57. X-ray diffraction pattern of the top surface of a crystal from the top layer [128].

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Table 12 Growth conditions and properties of the samples prepared in top and bottom layers.

a

Sample

Tmax(°C)

DTa

Growth time (h)

Maximum trapped filed (T)

Top sample Bottom sample

1093 1108

Small Large

40 25

0.4 0.45

DT represents the temperature difference between the up-side and down-side of each sample.

facet lines of the bulks at the bottom layer (see Fig. 56) and the XRD results of the top layer ones shown in Fig. 57 indicate that the samples are c-axis oriented. In other words, it means that the asgrown crystals were single-domain, i.e., the film seed efficiently induced the growth of YBCO. From Table 12 we know that the film seed at bottom layers underwent a Tmax value of 1108 °C and a long holding time over 15 h on the surface of the molten pellet. It is amazing that the seed films were capable to endure both a high processing temperature and a long holding time. As mentioned before, thermal stability of the seed film for a long holding time on the surface of the molten pellet is of great significance in the REBCO bulk growth, especially in the batch process. To be specific, firstly, the capability of enduring a long period can make up for the lack of perfect temperature uniformity inside the furnace, since film-seeds with high thermal stability suit for both early and late inducement of the REBCO growth. The additional advantage of this property is that Tp does not has to be determined precisely anymore for the growth of a REBCO crystal. One can safely use a slow-cooling starting temperature (the one set after Tmax in the heating profile, see Fig. 42) higher than the conventional one in the temperature profile, which makes the optimization of the process easy, and leads to the reduction of the failure rate. Due to these benefits, the mechanism of this property attracts interests. In the following, the reason why thin films tolerate a long holding time on the surface of the molten pellet was investigated. A schematic illustration (Fig. 58) helps to explain the mechanism. There are two factors that enhance the thermal stability of film-seeds, as shown in the upper left corner of Fig. 58. Firstly, the so-called coarsening effect, as discussed in Section 4.2.2. Secondly, during the heating process, owing to the diffusion of Nd elements from films into the precursor below, an interlayer of Nd–YBCO solid solution may form between the film and the pellet. This interlayer material crystallizes at a higher temperature than YBCO, and protects the film from further dissolving into the liquid. On the other hand, there are also two factors that impair the thermal stability of film-seeds, as shown in the lower right corner of Fig. 58. Firstly, the film gradually dissolves into the Ba–Cu–O melt underneath, resulting in a decrease of its thermal stability. Secondly, the concentration of Nd

Fig. 58. Schematic illustration of the factors affecting the thermal stability of film-seeds. The y-axis represents the effective strength, the x-axis represents the time. Positive and negative effects on the thermal stability of film-seeds with time are represented in the upper left corner and the lower right corner, respectively [128].

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element in the YBCO melt near the film-seed would become lower due to its diffusion, accelerating the dissolution of the film. Thus the competition between the factors plays a crucial role in influencing the thermal stability of film-seeds in the bulk growth. As illustrated in Fig. 58, the effects of grain coarsening and Y–NdBCO crystallization are predominant in a short holding time, while the effects of dissolution and diffusion are overwhelming in a long holding time. In brief, a cost-effective process was applied to grow two layers of YBCO crystals in one batch. The mechanism of the long tolerability of the film seed on the surface of the molten pellet was explained by the competition between the dissolution/diffusion of the film materials into the liquid and NdBCO grain coarsening/Nd–Y crystallization.

4.2.5. Thermal stability of film-seed in low supersaturated melt for growing YBCO crystals Besides bulk materials, large size and high quality YBCO single crystals are extremely necessary for both practical applications and fundamental studies. In the latter case, particularly, for the investigation on magnetic order and related phenomena by neutron scattering experiments, sizable single crystals are essential [129]. As the simplest method, the self-flux technique [67] is widely used in the growth of single crystals. However, for the YBCO system, it cannot produce sufficiently large crystals due to its low growth rate and small growth window. Although by Top-seeded solution growth (TSSG), high quality YBCO single crystals with large sizes can be obtained, however, the growth rate is limited by a low solubility of Y element in the Ba–Cu–O solvent [130–132]. Recently, Babu et al. [133] and Wang et al. [134] reported a modified TSMG, in which a single phase of Y123 as a precursor material was used, instead of the conventional Y123 + Y211. As a result, a well c-axis-oriented YBCO crystal was gained with a relative high growth efficiency, which includes about 2 vol% Y211, much less than commonly grown ones (about 30 vol% Y211 or more). Furthermore, the modified TSMG is characterized by the growing crystal from a low supersaturated melt, with which the film-seed certainly suffers from the problem of reduced thermal stability. Firstly, in the normal TSMG process [89], as illustrated in Table 13 (a⁄), the starting material of Y123 + 30 mol%Y211 is employed. Undergoing a heating route up to 1100 °C and a holding period of over 1.5 h, the Y123 decomposition and the Y211 coarsening take place, as discussed in the above section. And then the precursor becomes a mixture of liquid and Y211 particles. The latter is composed of small-sized Y211 caused by the Y123 peritectic decomposition, and large sized Y211 derived from the Y211 primary phase coarsening. Both together supply the yttrium solute in the crystal growth stage, leading to a high supersaturation status in the melt. In such a growth process, REBCO (RE = Y and rare earth element) film-seeds are applicable to effectively induce the growth of large-sized and high-performance REBCO bulks due to both intrinsically high superheating capacity and extrinsically high supersaturated melt [114,116,134,135]. However, in the modified TSMG without the addition of Y211 into a precursor, as shown in Table 13 (b⁄), the crystal grows in an Y211-less melt, which possesses a much lower supersaturation level because the yttrium solute is provided only by the decomposed Y211 particles, resulting from Y123 peritectic melting. Therefore the film-seed, effective in inducing crystal growth in conventional TSMG, may have diffusion and dissolution tendency in this new approach, leading to a low thermal stability problem. Fig. 59(a) shows the top-views of a conventional YBCO bulk induced by an NdBCO/YBCO/MgO filmseed after enduring Tmax of 1100 °C. However, when the same Tmax was exploited on the Y211-free precursor pellets, self-nucleation occurred and we obtained polycrystals, indicating that the NdBCO film

Table 13 Characteristics of three TSMG processes.

a⁄ b⁄ c⁄

Crystal

Starting composition

Supersaturation

Tmax (°C)

Film thermal stability

Single Poly Single

Y123 + 30 mol% Y211 Y123 Y123 with mini-pellet a⁄

High Low High (mini-pellet)

1100 1040 1080

High Low High

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Fig. 59. Top-views of the YBCO crystals (a) with the primary Y211 addition, (b) without the primary Y211 addition, (c) using a mini-pellet (with Y211)/main pellet (without Y211) [134].

dissolved into the melt and could not play a role in inducing the crystal growth. That is to say, the NdBCO film is not stable in the supersaturation-reduced melt. To search suitable Tmax of film-seeds for inducing the growth of single crystals, different series of heating procedure with gradually decreased Tmax were performed. To our surprise, even Tmax was decreased to 1040 °C (the lower limit to grow MG samples in our experiment), we still can see an imperfect top-view of the as-grown crystal (shown in Fig. 59(b)), indicating that the crystal is not highly c-axis oriented. Thus we conclude that the thermal stability of film-seeds is significantly reduced in the case of using the low-supersaturated melt. As we know, during the heating process, there are several factors that affect the thermal stability of the film seed. On the one hand, as negative factors, the film dissolves into the Ba–Cu–O melt due to the Nd solubility, while Nd elements diffuse from the film-seed owing to the Nd concentration gradient [128]. On the other hand, as positive factors, small grains in the film-seed become large arising from the coarsening phenomenon [108,128,135], while the Nd–YBCO interlayer (with higher crystallization temperature) forms between the film and the pellet, protecting the film from dissolving. In this modified MG process, because of low supersaturation in the melt, the dissolution/diffusion effect is predominant over the grain coarsening and Y–NdBCO crystallization one. As a result, the film-seed cannot tolerate high Tmax in this low supersaturated melt. To solve the problem of reduced thermal stability for film-seeds, a mini-pellet was inserted between the main bulk and the film-seed, with a composition of Y123 and additional Y211 for

Fig. 60. Y123 single crystals grown by modified TSMG method [134].

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enhancing the level of supersaturation. Thus the thermal stability of the film-seed for preparing YBCO single crystals was remarkably improved in the newly-used structure of film-seed/mini-pellet (with Y211)/main bulk (without Y211). As a consequence, a highly c-axis oriented single crystal (undergoing Tmax of 1080 °C) was successfully grown as shown in Fig. 59(c). More importantly, using this modified process in a simple one-side-heating furnace, batch processed single crystals, cost-effectively and high-efficiently, were fabricated. Fig. 60 presents those crystals with a total weight of approximately 100 g, supplied for neutron scattering experiments. More remarkably, for various fundamental interests, the achievements also promise us on synthesizing large chemically-doped YBCO single crystals, which are confronting serious technical difficulties by other crystal growth methods.

4.2.6. Highly oriented a/b-direction and a–b plane of film-seed in multi-seed process for controlling grain boundary Multi-seeded melt growth (MSMG) [136–138] is an effective alternative to conventional TSMG, involves placing several seeds of the same orientation simultaneously on the top surface of the precursor pellet prior to melt processing. As a result, the entire growth time can be reduced greatly due to simultaneous epitaxial growth from each seed independently. Unfortunately, the MSMG method may result in the trapping of unreacted liquid phase material at grain boundaries, which can degrade the field trapping ability of the bulk sample [139–143]. The presence of unreacted liquid has been observed normally at the (1 0 0)/(1 0 0) grain boundary, but rarely at the (1 1 0)/(1 1 0) grain boundary [140–142]. However, with conventional SmBCO and NdBCO crystal-seeds, the contact angle of grains cannot be controlled accurately in a multi-seeding process. REBCO thin film seeds, with highly controlled orientation (i.e. a well-defined a–b plane and precisely known a-direction), can be easily cleaved with a precise orientation, which makes them potentially suitable to overcome the major limitations of the multi-seeding process. Furthermore, a variety of grain boundaries can be obtained by varying the angle between the seeds. By detailed study, Li et al. [135] has established that the (1 1 0)/(1 1 0) grain interface is linked more strongly than other grain boundaries, and exhibits improved electrical and microstructural homogeneity. The results show that the residual melt at (1 0 0)/(1 0 0) grain boundaries increases with increasing inter-seed distance, leading to the formation of a weak link (weakly-superconducting regions) at the grain boundary. Moreover, the trapped magnetic field has been reported to degrade as the number of seeds increases due specifically to the poor connection of each domain [144–147]. These results suggest that an optimized process is needed to improve the connectivity between grains for a multi-seeding growth process using more than two seeds. After, Cheng et al. [148] develop a more reliable thin film seed arrangement for the growth of multi-seeded GdBCO bulk samples. 1. The effect of seed separation on a four-seed MSMG process

Fig. 61. (a) Photograph of the top view of a GdBCO bulk sample fabricated by MSMG using an asymmetric four-seed arrangement with the angle of the adjacent seeds set precisely to 90°. The white frame depicts the combined domain produced by the adjacent seeds, whereas the arrows indicate the direction of the growth front. (b) Schematic diagram of an asymmetrical (1 1 0)/(1 1 0) seed arrangement produced using four thin film seeds [147].

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The optical micrograph of the as-grown GdBCO bulk sample of diameter 42 mm is shown in Fig. 61(a). This multi-seeded sample was grown using four NdBCO thin film seeds of square planar shape with their edges aligned parallel to each other resulting in the formation of (1 1 0)/(1 1 0) grain boundaries. Each of the four individual GdBCO domains extends to the edge of the sample, indicating that a large multi-seeded bulk sample has been fabricated successfully. A schematic illustration of the arrangement of the seeds prior to melt processing is shown in Fig. 61(b). Initially, single grain growth occurs at the position of the two seeds separated by the shorter distance. The residual melt is rejected ahead of the growth front, as indicated by the arrows, resulting in the formation of a precipitate-free grain boundary. These two domains subsequently impinge, as depicted by the light gray square in Fig. 61(b), at which point further growth occurs perpendicularly, or nearly perpendicularly, to the grain boundary. The excess melt liquid is pushed out from the growth front as growth proceeds, leading to the formation of a clean grain boundary. On the other hand, the conventional arrangement of seeds in a multi-seed, melt process is illustrated in Fig. 62(a and b). Here the multi-grain samples are fabricated from four seeds with symmetrical arrangement of (1 0 0)/(1 0 0) and (1 1 0)/(1 1 0), respectively, and results in the presence of much Barium-rich phase at the impinging point (or dirty grain boundary). Obviously, it is not possible to avoid the formation of (1 0 0)/(1 0 0) grain boundaries for a symmetrical arrangement of multi-seeds, regardless of whether a pre-determined seed arrangement of (1 0 0)/(1 0 0) or (1 1 0)/(1 1 0) is employed. Therefore, the multi-seed effect represents a significant barrier to the optimization of the field trapping ability of the bulk sample for a symmetrical arrangement of seeds. Furthermore, compared to the conventional, symmetrical multi-seed arrangement in forming the (1 1 0)/(1 1 0) grain boundaries [146], the asymmetrical arrangement leads to some interesting configurations of the pre-oriented seeds, such as of three and five seeds shown schematically in Fig. 62(c and d). The (1 1 0)/(1 1 0) grain boundaries can always form during growth by specific placement of the seeds. Residual melt can then be pushed out along the intersection of adjacent domains, forming strongly coupled grain boundaries. Here the effect of seed separation is eliminated via an asymmetrical arrangement of multi-seeds, which, again, has the potential to improve the trapped magnetic field of large MSMG bulk samples. 2. The effect of seed separation for an asymmetrical four-seed MSMG process Fig. 63 shows a photograph of the as-grown GdBCO bulk sample using four seeds arranged asymmetrically. The locations of the seeds on the left of the sample tilt slightly from the (1 1 0)/ (1 1 0) orientation. It can be seen from Fig. 63(a) that the domain nucleating from the pair of seeds

Fig. 62. Schematic illustration of the growth process using multi-seeds: (a) the formation of the ‘‘dirty’’ (1 0 0)/(1 0 0) grain boundaries; (b) the formation of the ‘‘dirty’’ impinging point among grains and the (1 1 0)/(1 1 0) grain boundaries. Schematic illustration of the growth process of an asymmetrical arrangement of multi-seeds of the (1 1 0)/(1 1 0) grain boundaries: (c) three seeds and (d) five seeds. The arrows represent the direction of grain growth [147].

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Fig. 63. (a) Photograph of the top view of a GdBCO bulk sample grown by MSMG using an asymmetric arrangement of fourseeds. The arrows indicate the growth facet lines of the grains on the left of the photograph. The inset gives an enlarged image of the pair seeds region, with the white diamond frame depicting the unexpected crystal plane of the grain on the right side. (b) Schematic illustration of the seed locations and growth process corresponding to the two growth modes [147].

on the right of the photograph is much larger than that nucleating from the pair on the left. In general, two kinds of crystal plane can be obtained from domains grown from seeds of different orientation, as illustrated in Fig. 63(b). Two types of domain growth may be studied by considering the growth of two domains from seeds with different orientations. The growth induced by the two seeds to the left of the sample is shown schematically in Fig. 64(a). The geometry of the domain is maintained as it grows, since the (1 0 0) and (0 1 0) faces grow at an equal rate. As a result, the two domains grow separately, as depicted by the white arrows in Fig. 63(a). At the same time, the residual melt is gradually rejected at the angle of intersection of the growing domains to form a clean grain boundary. Eventually, a combined c-axis oriented domain was grown successfully, with clear fourfold growth front facet lines evident relative to the crystallographic orientation of the seed crystal.

Fig. 64. Schematic illustration of growth evolution corresponding to the two growth modes: (a) the growth induced by the two seeds to the left of the sample; (b) the growth induced by the two seeds to the right of the sample [147].

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Fig. 64(b) illustrates the various stages of domain growth from the seeds on the right of the sample, for which the seeds are aligned to precisely 90°, leading to a formation of a single (1 1 0) crystal plane. More specifically, in the early stage of the solidification process, the two domains were induced separately from seeds, then impinged to form the (1 1 0) face, as indicated in Fig. 64(b) by the hexagon with blue lines. In essence, high index faces grow faster than low ones, so the growth rate of (1 1 0) face grows more rapidly than that of the (1 0 0) face. As a result, the appearance of this diamond shaped domain leads to a much larger combined product domain. This fact is potentially important for multi-seeded, melt processed GdBCO for which a large growth rate is required to fabricate large, heterogeneously nucleated bulk samples. In short, the use of multi-seeds arranged asymmetrically with a (1 1 0)/(1 1 0) orientation is particularly promising for further enlargement of the domain size and for potentially enhancing the superconducting performance of multi-seeded melt processed REBCO bulk samples. In this respect, the use of film-seed has a great advantage, since it can be readily sliced with a precise a-direction, which especially benefits the control of the contact angle between grains in the multi-seed process. 4.2.7. Large-sized film-seed inducing the growth of REBCO crystal with large c-growth sector Usually, REBCO superconductor crystals prepared by Top-Seeded Melt-Growth (TSMG) characteristically possess two types of growth sectors: four a-growth sectors (a-GS) and one c-growth sector (c-GS), associated with grain sector boundaries (a/c-GSB and a/a-GSB), as illustrated in Fig. 65. It is well known that the crystal with a larger single domain of c-GS should have a stronger capacity to trap magnetic fields. Thus, it is necessary to optimize the preparation of the RE-Ba–Cu–O crystal for enhancing its superconducting performance by exploring an effective method to obtain crystals with a larger c-GS. It is clear that different sizes of seeds lead to different volume fractions of a-GS and c-GS as shown in Fig. 65, which certainly affect the growth and the performance of REBCO bulk superconductors. In this point of view, the use of large-sized seed (LSS) to induce the growth of large sized crystal with the high volume fractions of c-GS is commonly considered due to its several advantages [148,149]. Firstly, the use of LSS with large-area epitaxial-growth results in fully-growth along the a-axis direction shortly, and prevents the crystal from self-nucleation on the a–b plane surface, which readily occurs during the long period of growing time. Secondly, Y211 coarsening in the melt and Y211 segregation at the growth front commonly occur, which become increasingly severe with time and cause the problem in the growth of Y123, such as slowing-down and terminating ahead of time. The employ of LSS with a time-saving process, results in a continuous growth by suppressing these two effects. Finally

Fig. 65. Schematic illustration showing virtual growth modes of Y-123 c-axis oriented crystals induced by a small and a large seed, respectively. The a-GS and the c-GS are represented as blue and green entities, respectively [151].

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Fig. 66. Top and side views of YBCO grown c-axis oriented crystals induced by NdBCO film seeds with different sizes of 2 ⁄ 2 mm2, 4 ⁄ 4 mm2, 7 ⁄ 7 mm2 and 9 ⁄ 9 mm2, respectively [151].

and most importantly, following a rapid completion of the growth of four small a-GS by applying LSS, the YBCO crystal with large c-GS and high performance is possible to be achieved. In this regards, by employing a large Sm123 seed with a size of 14 ⁄ 14 mm2 on the a–b plane, Xu KX et al. [150] succeeded in fabricating a large YBCO bulk (53 mm in dia.), which demonstrated a higher levitation force than that induced by the small seed of 2 ⁄ 2 mm2. However, the preparation of LSS for Sm123 material requires a complex and time-consuming route: being grown in an optimized procedure with the long period of time, and then being carefully cut down along the a–b plane and polished. The main disadvantage is that the seed materials of REBCO bulks cannot endure a high Tmax of over 1100 °C, which is absolutely necessary for seeding the growth of high Tp REBCO materials, such as SmBCO and NdBCO. Instead of conventional LSS of REBCO bulks, NdBCO films (NdBCO film/YBCO buffer layer/MgO substrate) were used as seed materials [118]. The main advantages of the use of NdBCO films are their high superheating capacity (enduring Tmax up to 1120 °C), and conveniently commercial availability of large-sized film-seeds with reliable high quality. As a result, a series of seeds (2 ⁄ 2 mm2, 4 ⁄ 4 mm2, 7 ⁄ 7 mm2 and 9 ⁄ 9 mm2) were applied to study the effects of the seed size on the growth of YBCO crystals [151]. As shown in Fig. 66, top views of all samples without appearance of self-nucleated/grown grains, present clear ‘‘X-type’’ a/a growth sector boundaries (a/a GSB), indicating that filmseed induced crystals have four a-GS with well c-axis orientations. Most importantly, as shown in side views, the triangle region (representing c-GS) becomes increasingly large with the seed size, indicating that the grown YBCO crystal with the large-size seed possesses a larger volume fraction of c-GS, which tends to be beneficial to achieve stronger capacity to trap magnetic field for engineering applications. This success demonstrates a possible approach to grow large-sized REBCO crystals with large c-growth sections. 5. Conclusion and prospects As is well known, materials with good quality and unique properties are highly desired to promote social progress. Recent reports have indicated that superheating of solid, an unconventional phenomenon, can be achieved by suppressing the heterogeneous nucleation of melt at surfaces. However, it is practically difficult to prepare such confined structures. More recently, significant progresses have been made in the superheating mechanism and applications of REBCO thin films with free surfaces.

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By using various methods including liquid phase epitaxy, high temperature in-situ observation and melt growth, the intrinsic superheating property of these functional oxides have been verified. Due to the developed and easy preparation process of these thin film materials, it is a suitable object to investigate the superheating mechanism and a potential material for wide applications in technology. With respect to the superheating mechanism, the key factors which significantly influence the thermal stability of thin films are reviewed. Firstly, the origin of the superheating phenomenon is that the free surface of REBCO films has a low Gibbs energy, which is known as a non-melting surface. The development and understanding of these low-dimensional films are attractive for more functional oxides which possess strong anisotropy. Secondly, the film thermal stability relates to interface structure such as out-of-plane and in-plane alignment as well as lattice mismatch between film and substrate. Finally, the temperature coefficient of solubility (t.c.s., i.e., RE elements in Ba–Cu–O liquid) correlated supersaturation contributes to affect film thermal stability. The research shed novel lights on the criterions of superheating, which is crucial to extend the understanding of melting and superheating to the field of functional oxides. More importantly, the systematic understanding of superheating mechanism of functional oxides with the perovskite structure is advantageous to explore new film/ substrate constructions with higher thermal stability. Furthermore, the REBCO thin films with superheating property have wide applications, such as in extending the field of phase stabilization and seeding the growth of REBCO materials. On the one hand, under the state of high superheating, phase transformation behaviors can be investigated. It was found that if the film can be highly superheated above a critical temperature, an unexpected metastable phase epitaxially nucleated prior to the stable phase. This phenomenon is popularly reported in a deep undercooling state. Therefore, the research interests of phase competition is extended to the high superheating state, which favoring the study on the phase relations and exploring possible metastable phases. Similar to the metastable phase transition in the deep undercooling state, this extension carries great theoretical and practical importance in both fundamental study and industrial applications. On the other hand, functional oxides of thin films with high thermal stability have wide applications in seed-required processes. For instance, applying thin films as seed materials in cold-seeding melt growth, and making use of their superheating property, a higher maximum processing temperature (Tmax) and a wider window for single grain growth can be achieved, which is beneficial to the growth of large-sized and high-performance REBCO superconductor bulks. Remarkably, the quality and thermal stability of NdBCO films can be further enhanced by buffering an YBCO nano-layer between film and substrate, enduring a maximum processing temperature (Tmax) up to 1120 °C in cold-seeding melt-growth. This development provides an excellent seed material for realizing the growth of even higher processing-temperature required REBCO crystals: fabricating high Tp REBCO (RE = Nd, Sm) bulks, and recycling the failed YBCO bulks in a simple process. It should be pointed out that it used to be extremely difficult to find a suitable seed for NdBCO due to its highest Tp among practicallygrown REBCO superconductors. Up to date, around ten research laboratories worldwide have been benefiting from the superheating nature of REBCO films, in seeding growth of superconductor crystals. For instance, using NdBCO film-seed, Noudem and Izumi et al. have succeeded in growing c-oriented GdBCO bulks by the seeded infiltration and growth (SIG) method [152]. With high-quality NdBCO/ MgO film-seeds, Muralidhar et al. have fabricated multiple GdBCO bulks in the batch process [115]. Sawh and Weinstein et al. [153] have fabricated 2 cm diameter YBCO bulks with high tapped field. Yang CM and Chen et al. studied the nucleation density of YBCO as a function of cooling rate, using the NdBCO film as heterogeneous cold-seed in the TSMG process [154]. In brief, thanks to the development of film-seed with superheating nature, fabrication of all LREBCO materials by cold-seeding method becomes achievable and reliable. Apart from its superheating property, the REBCO thin films have three other advantages for applications: (1) they are easy to get sliced. With highly controlled orientation, i.e., a well-defined a–b plane and precisely known a-direction, the films especially benefits the control of grain boundaries in the multi-seed process, while the cutting step for precisely-oriented REBCO crystal-seed is complex and time-consuming. (2) They are commercially available to conveniently gain large-sized film-seeds due to its developed preparing process, which also helps to get a full-growth of large sized REBCO bulk with high performance. (3) They are reliable and stable in quality, favoring the batch growth process for industrial production. In brief, the extensive employment of REBCO-films with excellent superheating capacity as seed materials

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makes essential progresses for producing and recycling industry of currently-grown REBCO superconductors, which are important for resource-conservation and economic-development. On the other hand, superior (but process-sensitive) materials which used to lack suitable seed can be expected as well, such as SmBCO, NdBCO and LaBCO superconductors. Apparently, theoretical study as well as experimental research on the thermal stability of REBCO thin film has guiding significance to both science and techniques, such as to set up the superheating criterion, to explore novel structures with high thermal stability, to seek universal seeds for preparing more materials and so on. As a fascinating theme in material science, further efforts on superheating are required to: (1) Search low-dimensional materials with larger anisotropy compared with REBCO. Their free surfaces with the lowest Gibbs energy are expected to possess higher thermal stability. (2) Explore new film/substrate structures with high thermal stability based on the superheating criterions. (3) Develop suitable buffer layers with high melting point and small lattice misfit with both films and substrates, to enhance the superheating capability of REBCO films. (4) Investigate unique phase transformations and novel metastable phases in highly-superheating state. (5) Realize commercialized production of more materials using these film-seeds with superheating capacity. Finally and most importantly, scientific understandings of the magic superheating phenomenon will stimulate more basic interests and open a new prospect, for instance, study on other functional oxides that have the same nature as the construction of REBCO film/substrate. On the other hand, great successes in engineering application of film-seeds with unconventionally-high thermal stability may bring about a breakthrough in the seeding growth technology. Acknowledgments The authors are grateful for financial support from the Ministry of Science and Technology of China (Grant No. 2012CB821404), and National Nature Science Foundation of China (Grant No. 51172143). The authors are very thankful for support from the key laboratory of artificial structures & quantum control, state key laboratory for metal matrix composites and instrumental analysis center of Shanghai Jiao Tong University and THEVA Company. The authors acknowledge the collaborators for their great contributions to this work including Y. Shiohara, T. Izumi, D.X. Huang, F.H. Li, W. Wan, H. Ikuta, M. Oda, Y. Yoshida, B.W. Tao, J. Xiong and Y. Wang. The valuable work of students involved in this work are appreciated, particularly J. Hu, Y.Q. Cai, C.Y. Tang, L.J. Sun, S.B. Yan, L. Cheng, F. Qin, X. Wang, T.Y. Li, H.H. Xu, H. Li, L.S. Guo, B.N. Peng, D.J. Yu, W. Wang, H.C. Li and W.S. Fan, Y.F. Zhuang and Y.M. Lei. We also thank Dr. Hari Babu for his helpful suggestions. References [1] Cahn RW. Materials science: melting and the surface. Nature 1986;323:668. [2] Däeges J, Gleiter H, Perepezko JH. Superheating of metal crystals. Phys Lett A 1986;119:79. [3] Sheng HW, Ren G, Peng LM, Hu ZQ, Lu K. Superheating and melting-point depression of Pb nanoparticles embedded in Al matrices. Philos Mag Lett 1996;6:417. [4] Saka H, Nishikawa Y, Imura T. Melting temperature of in particles embedded in an Al matrix. Philos Mag A 1988;57:895. [5] Goswami R, Chattopadhyay K. The superheating of Pb embedded in a Zn matrix the role of interface melting. Philos Mag Lett 1993;68:215. [6] Zhang DL, Cantor B. Melting behaviour of In and Pb particles embedded in an Al matrix. Acta Metall Mater 1991;39:1595. [7] Zhong J, Zhang LH, Jin ZH, Lu K. Superheating of Ag nanoparticles embedded in Ni matrix. Acta Mater 2001;49:2897. [8] Rossouw J, Donnelly SE. Superheating of small solid-argon bubbles in aluminum. Phys Rev Lett 1985;55:2960. [9] Rösner H, Wilde G. The impact of altered interface structures on the melting behaviour of embedded nanoparticles. Scripta Mater 2006;55:119. [10] Jin ZH, Sheng HW, Lu K. Melting of Pb clusters without free surfaces. Phys Rev B 1999;60:141. [11] Konrad H, Weissmüller J, Birringer R, Karmonik C, Gleiter H. Kinetics of gallium films confined at grain boundaries. Phys Rev B 1998;58:2142.

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