Exploration of new superconductors and functional materials, and fabrication of superconducting tapes and wires of iron pnictides

Although approximately 100 IBSCs have been reported, their parent materials may be classified into seven types in terms of crystal structures (figure 1). These materials contain a common structural unit of the FePn (or FeCh) layer formed by the square net of Fe²⁺ (as the formal charge), which is tetrahedrally coordinated by four pnictogen (Pn) and/or chalcogen (Ch) atoms (see figure 1(a)). Unlike cuprate superconductors, where the parent materials are Mott insulators, this layer shows metallic conductivity without doping. An insulating blocking layer is composed of M, MO or MF etc, where M indicates a metallic element such as an alkali, alkaline earth, or rare earth metal that lies between FePn (or FeCh) layers. Similar to cuprates, this layered structure provides quasi-two-dimensional carrier transport properties, although the magnitude of anisotropy rather differs depending on the blocking layer. The local structure of the FePn layer is affected directly by the atomic (or ionic) size of M because M elements in the blocking layer bond to Fe elements. The crystal structures and the brief introductions of the seven different parent materials for IBSCs are described below.

• (i) 1111-type materials (LnFePnO, Ln: lanthanide, AEFeAsF, AE: alkaline earth, Pn: P, As).

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1111-type compounds have the same structure as LaFeAsO, and are the prototype version of IBSCs. Due to their atom composition ratios, they are called '1111-type'. Figure 1(b) shows their crystal structure, which is a ZrCuSiAs-type structure [128, 129] with a tetragonal P4/nmm space group. Although LaFePO and LaFeAsO, along with their crystal structures, were identified by Zimmer et al 20 years ago [130], they were discovered to be superconductors in 2006 [3] and 2008 [4], respectively. Moreover, their two-dimensionality is relatively high among the seven types, and only this group has T_c values above 50 K as a bulk form.

The 1111-type compound is composed of an alternating stack of positively charged LnO layers and negatively charged FeAs (or FeP) layers along the c-axis. As mentioned above, the local structure of the FeAs layer is the same in all types of IBSCs. The distance between the FeAs layers corresponds to the length of the c-axis (∼0.8–0.9 nm). The formal valence state of each atom is Ln³⁺, Fe²⁺, As³⁻, and O²⁻. Fe²⁺ contains six electrons in its 3d orbital, and these electrons play an essential role in driving the superconductivity and magnetism. The electronic and magnetic properties of LaT_MPnO (T_M: 3d transition metal (Cr–Zn)) are summarized in table 3 [120]. The 1111-type compounds composed of Fe and Ni reveal superconductivity.

Table 3.  Summary of electromagnetic properties of LaT_MPnO. (T_M: 3d transition metal, and Pn=P or As.)

The lanthanide elements from La to Gd can occupy the Ln site for the 1111-type of material with Pn = P [130]. In the case of Pn = As, La to Ho and Y can also occupy the Ln site [136]. Additionally, Ca(Fe_1−xCo_x)AsF (T_c = 22 K) is a fluoride-containing superconductor of this type [138].

In this project, the effect of a hydride ion as a new electron donor to this type was studied earnestly and the results are described in section 3.1.2.

• (ii) 122-type materials (MFe₂Pn₂, M: alkali, alkaline earth or Eu)).

122-type materials have a 'ThCr₂Si₂' type crystal structure with a tetragonal I4/mmm space group [137]. This group contains the largest number of compounds among the five parent families.

In the case of AEFe₂P₂, not only alkaline earth elements but also lanthanides (La–Pr, Eu) can occupy the AE site. In AEFe₂As₂, the AE site can be occupied by alkaline earth, alkali metal, or Eu²⁺. Figure 1(c) shows the crystal structure of the 122-type. The layer composed of AE ions, which is thinner than the Ln-–O layer of 1111-type, is sandwiched between the FeAs conducting layers. The distance between the FeAs layers of 122-type (0.5–0.6 nm) is shorter than that of 1111-type (0.8–0.9 nm). Because the nearest FeAs layers face each other with a mirror plane, the lattice parameter c is twice the FeAs–FeAs distance. The lattice parameter a (∼0.4 nm) is almost the same as that of 1111-type. Consequently, both 1111 and 122-type materials have similar Fe–Fe distance in the FeAs layer. Since single crystals of several millimeters can be obtained using Sn or FeAs as a flux, the physical properties of 122-type are well evaluated compared to other types of IBSCs. Johrendt's group of Germany was the first to report superconductivity for 122-type materials [119].

In this project, lanthanide element doped 122 superconductors were prepared and evaluated in their bulk and thin-film forms, and these results are described in section 3.1.3.

• (iii) 111-type materials (AFePn, A: alkali metal).

While an AE ion (alkaline earth ion with formal charge of 2+) is alternately sandwiched between FePn layers in 122-type, 111-type compounds contain two A ions (A: Li⁺, Na⁺) between FePn layers in a unit cell. The crystal structure of this type is known as 'CeFeSi' type, with a tetragonal P4/nmm space group (figure 1(d)). This type is compatible with the structure of 1111-type where all the oxygen atoms are removed, and the Ln site is occupied by Li⁺ or Na⁺. Wang et al [139] (T_c = 18 K: LiFeAs) and Parker et al [140] (T_c = 10 K: NaFeAs) first reported superconductivity for 111-type materials.

• (iv) Materials with thick blocking layer (32522-type (AE₃M₂O₅Fe₂Pn₂, M: Al, Sc)), (42622-type (AE₄M₂O₆Fe₂Pn₂, M: Sc, V, Cr)), (homologous type (Ca_n+1Sc_nO_y Fe₂As₂: n = 3, 4, 5)).

The distance between the FePn/Ch layers is in the order of the 1111, 122, 111 and 11-types. In contrast, these three types of iron oxy-pnictide have a thick blocking layer composed of a quasi-perovskite structure assembled by MO₅ pyramids and AE (see figure 1(e) for the 42622-type (Sr₄Sc₂O₆Fe₂As₂)). The FeAs–FeAs distance is 1.55 nm and 2.45 nm for Sr₄Sc₂O₆Fe₂As₂ and Ca₆(Sc_0.4Ti_0.6)₅O_yFe₂As₂, respectively. The groups of Kishio and Shimoyama have studied these types of materials systematically [141–143]. The highest T_c reported so far is 43 K [143]. Considering the thick blocking layer, this type should have the highest two-dimensionality, but the concrete value of anisotropic properties has not yet been reported because of difficulty of single crystal growth. The 32522-type has been proposed as a promised parent material [144, 145], and the emergence of superconductivity in the 32522-type was reported by Shirage et al in 2011 for (Ca₃Al₂O_5−y)(Fe₂Pn₂) (Pn = As (T_c = 30.2 K) and P (T_c = 16.8 K)) [146].

• (v) Materials containing additional arsenic (Ca_1−xLa_xFeAs₂), (Ca₁₀(M₄As₈)(Fe₂As₂)₅, (M: Pr, Ir)).

These new types of iron pnictide superconductor were found by Nohara's group of this project. The details for (Ca_1−xLa_xFeAs₂) (T_c = 43 K) and (Ca₁₀(M₄As₈)(Fe₂As₂)₅, (M: Pr, Ir)) (T_c = 38 K) are described in sections 3.1.4 and 3.1.5, respectively.

• (vi) 11-type materials (Fe_1+xCh, Ch: Se, Te).

The 11-type crystal has the simplest structure among the parent compounds and is essentially the alkali metal-free 111-type. This crystal structure is known as 'α-PbO' type with a tetragonal P4/nmm space group (figure 1(f)). A typical 11-type superconductor is β-FeSe (T_c = 8 K) [147]. Medvedev et al reported the 11-type may exhibit a high T_c (=37 K) under 8.9 GPa [148].

Furthermore, FeSe attracts attention as one of the candidates showing higher T_c than boiling temperature of liquid N₂. Several groups in China reported that the monolayer of FeSe deposited on a SrTiO₃ substrate showed high T_c (65 K) in 2012 and they raised T_c to 100 K [149–153]. Though this superconductivity emerges so far only for monolayers of FeSe deposited on a SrTiO₃ substrate, a new route to high T_c materials is expected to be found.

• (vii) 245-type materials (A_1−xFe_2−ySe₂: A = K, Cs, Rb, Tl).

In 2010, Guo et al reported a potassium-intercalated iron selenide superconductor with relatively high T_c value (30 K) [154]. The crystal structure changed from 11 to quasi-122-type upon intercalation, of which the space group is assigned to I4/m due to vacancy ordering as shown in figure 1(g). Though Guo et al noted its chemical notation as K_xFe₂Se₂, the detailed chemical and structural analyses for its optimal material showed the composition to be A_0.8Fe_1.6Se₂ (=A₂Fe₄Se₅) and ordering of Fe vacancies with a $\sqrt{5}\times \sqrt{5}\times 1$ supercell in the 122-type crystal structure [155]. This type of material shows a wide range of non-stoichiometry, and with low Fe concentration is an antiferromagnetic insulator. The superconductivity in A_xFe_ySe₂ emerges in the proximity of an antiferromagnetic (AFM) Mott insulating state, similar to the cuprate high temperature superconductors [156]. Many unique properties let us classify this as an independent type apart from the 122-type. Ivanovskii reviewed this material [157].

In this project, we intercalated Na to FeSe employing the ammonothermal method, which cannot be prepared using a conventional thermal treatment at high temperatures. The result is described in section 3.1.6. Using this type of material with low Fe content as a Mott insulator, we examined the effect of the electric field on its electrical transport properties. The results are described in section 4.2.5.

Figure 2 shows the photoemission spectra of LaFeAsO and LaFeAsO_0.94F_0.06, and calculated partial density of states (PDOS) for Fe 3d and As 4p [158]. The Fermi level (E_F) controlling the transport property is primarily formed by a complex tangle of five Fe 3d orbitals, due to the small contribution of As, which is unlike cuprate superconductors where only Cud_x2−y2 contributes to the E_F. With five bands comprising E_F, multi-pockets, i.e., a disconnected Fermi surface (FS), appear on the FS. The energy levels of d_x2−y2, d_xy, and d_yz are sensitive to both changes in the symmetry of the FeAs₄ tetrahedron and the carrier density. Such an electronic structure dominates high T_c with the unique pairing mechanism.

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At the early stage of the theoretical approach for the pairing mechanism, several physicists [159–166] suggested the possibility of spin fluctuation mediated pairing, where the spin fluctuation arises around the nesting vector (π, 0) (see figure 3 [160]). The spin fluctuation mediates s_±-wave pairing, where the gap function has s-wave symmetry, but its sign is reversed between the electron and hole Fermi surfaces.

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In contrast, recent experimental results showed that high T_c is revealed when the nesting is degraded, or even in the absence of the nesting by heavy doping of impurities [11, 149, 151, 154, 167]. To explain the robust superconducting state against impurities, Kontani and Onari [168] proposed a mechanism of the S₊₊-wave superconducting state induced by orbital fluctuations, due to the phonon-mediated electron–electron interaction. On the other hand, Suzuki et al [169] succeeded in reproducing the general trend of composition dependence of T_c in LnFeAsO_1−xH_x (Ln: La, Ce, Sm and Gd) by the diagonal (next nearest neighbor) electron hopping model, where the next nearest neighbor (diagonal) hoppings between iron sites dominate over the nearest neighbor ones, plays an important role in the enhancement of the spin fluctuation and thus superconductivity. The theoretical and experimental evaluation for the superconducting mechanism will continue from now on.

Unlike cuprate superconductors whose parent materials are Mott insulators, the parent materials of IBSCs are antiferromagnetic metals with sufficient conduction carriers. Hence, it is considered that carrier doping into IBSCs mainly alter the FS, which in turn leads to suppression of antiferromagnetism.

Here we mainly describe the electronic phase diagram for the 1111-type (LnFeAsO, Ln: rare earth element). The parent materials for the 1111-type have tetragonal crystal structure at room temperature, but transform into orthorhombic structure at lower temperatures. In LaFeAsO, Pauli paramagnetism (PM) is shown around room temperature, and changes into AFM at a slightly lower temperatures (T_N ∼ 140 K) than that of the structural transitions (T_s ∼ 160 K) [170, 171].

Generally, superconductivity occurs in the tetragonal phase and not in the orthorhombic phase due to antiferromagnetic ordering in the orthorhombic phase. With doping (e.g., substituting O for F in LnFeAsO), the tetragonal–orthorhombic transition temperature decreases and is accompanied by suppression of the AFM state and superconductivity emerges in succession. Electrons are doped into the bulk, when an element with more valence electrons is substituted. In contrast, holes are doped by substituting an element with fewer valence electrons. In many cases of both 1111 and 122-type, it is possible to substitute Fe or As in the conducting layer and Ln, AE, O or F in the blocking layer for other elements. The former and the latter are called 'direct doping' and 'indirect doping', respectively.

The critical temperature (T_c) increases, reaches a maximum, and then decreases as the dopant level increases. Since the decrease in T_c in the over doping level is due to the precipitation of the secondary phase as SmOF in SmFeAsO_1−xF_x, the proposed phase diagram for 1111-type doped with F does not show the correct T_c behavior in the over doping region [172–176]. In contrast, Hanna et al [9] prepared SmFeAsO_1−xH_x and showed its optimal T_c (=55 K) at x = 0.20 and decrease in T_c by additional doping (over doping) without precipitation of the secondary phase, indicating a wide superconducting dome in 1111-type. Figure 4(a) shows the schematic phase diagram for the 1111-type.

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The first IBSC reported was formed by electron doping as LaFeAsO_1−xF_x where F substituted the O site as (${{{\rm F}}_{{\rm O}}}^{\bullet }~+~e^{\prime} )$ [4]. In addition to the substitution of oxygen sites by F (T_c = 55 K for SmFeAsO_0.9F_0.1) [115], various routes for electron doping have been reported, including the formation of an oxygen vacancy (${{{\rm V}}_{{\rm O}}}^{\bullet }~+~e^{\prime} :$ T_c = 55 K for SmFeAsO_0.85) using a high-pressure synthesis [117, 177, 178] substitution of H⁻ for an O²⁻ site (${{{\rm H}}_{{\rm O}}}^{\bullet }~+~e^{\prime} :$ T_c = 55 K for SmFeAsO_0.8H_0.2) [9], substitution of Th for Ln (${\rm T}{{{\rm h}}_{Ln}}^{\bullet }~+~e^{\prime} :$ T_c = 56 K for Gd_0.8Th_0.2FeAsO) [118, 179], and substitution of Co, Ni or Ir for Fe (${\rm C}{{{\rm o}}_{{\rm Fe}}}^{\bullet }+e^{\prime} :$ T_c = 14 K for LaFe_1−xCo_xAsO [180, 181], T_c = 22 K for I_1−xCo_xAsF [137], ${\rm N}{{{\rm i}}_{{\rm Fe}}}^{\bullet \bullet }~+~2e^{\prime} :$ T_c = 6 K for LaFe_1−xNi_xAsO [182], ${\rm I}{{{\rm r}}_{{\rm Fe}}}^{\bullet \bullet }~+~e^{\prime} :$ T_c = 18 K for SmFe_1−x Ir_xAsO [183]). The optimal amount of doping is 0.1–0.2/Fe atoms for each route, and indirect doping appears to be more effective than direct doping in achieving a high T_c, which should be due to less structural perturbation to the conducting layer.

For the 122-type, the shape of the electronic phase diagram is similar to the 1111-type as a general trend. The remarkable difference between the 1111- and 122-types is whether the AFM and superconducting phases are distinctly overlapped or not as shown in figure 4(b). In the 1111-type, the regions showing AFM and superconductivity are separated or barely overlap, whereas the 122-type materials have AFM regions with a high T_c. The optimal T_c is apparently located around the temperature corresponding to the extrapolation of the SDW curve to zero temperature, i.e., a superconducting dome appears around the quantum critical temperature of SDW [184]. The emergence of SC by doping of isoelectronic dopant, such as P for As, is also a unique property of the 122-type.

The comparison in doping between the 122- and the 1111-type is shown in table 4.

Compared to MgB₂ and cuprates, IBSCs have several distinct characteristics. It is included in the unique characteristic of IBSCs that the Fe 3d multi-orbital form Fermi surface described in (b) and the parent material is the antiferromagnetic metal described in (c).

Generally, T_c decreases upon doping with magnetic impurities such as Fe, Ni, and Co. In the case of cuprate superconductors, T_c of YBa₂Cu₃O_7−y decreases from 90 K to 50 K by substituting Ni (17%) for Cu, and that of La_1.85Sr_0.15CuO_4−y also decreases from 40 K to 4.2 K by substituting Ni (5%) for Cu [191]. The substitution of such elements for Fe on FeSCs with an optimal state shows a similar trend. The superconductivity of NdFeAsO_0.89F_0.11 (T_c = 48 K) disappears by substituting Co (>11%) or Mn (>4%) [192]. In contrast, the emergence of superconductivity by substitution of Co²⁺ (3d⁷), Ni²⁺ (3d⁸) or other transition metals for Fe²⁺ (3d⁶) in the non-superconducting parent material described in (c) is also a unique nature for IBSC.

The high T_c, large upper critical field (H_c2) and small anisotropy are important merits in applying IBSCs practically. Table 5 summarizes these values of IBSCs along with those of MgB₂ and cuprates. Only cuprates achieve higher T_c than the boiling point of liquid N₂ (77 K). It has been reported that the anisotropic ratio of the resistivity (γ_ρ) of the 122-type IBSCs is compatible with that of MgB₂ and smaller than that of cuprates. The H_c2(0) of IBSCs is higher than that of MgB₂, but is smaller than that of a typical cuprate. The anisotropic ratio of the H_c2, γ_H, of IBSCs is smaller than those of MgB₂ and cuprates. The H_c2(0) is defined as the upper critical field at 0 K. The γ_ρ means the ratio of the resistivity along the crystal axes directions, a (ρ(a)) and c (ρ(c)) measured just above T_c. The application of IBSCs to superconducting wires and devices will be described in section 4.

In alkali metal (A) intercalated compounds, A atoms occupy the octahedral or trigonal prism sites between chlorine layers of β-MNCl [58, 311, 312]. The largest metal concentration attained is expected to be x = 0.5 for A_xMNCl (M = Zr, Hf). We are interested in the electron doping to a much higher concentration using multivalent metals. Alkaline earth (AE = Ca, Sr, Ba) [59] and rare earth metals (RE = Eu, Yb) [60] have been successfully introduced into the interlayer space of parent materials by using liquid ammonia solutions, which make it possible to study the effect of high doping concentration and additional magnetic spin, respectively. The dependence of the metal doping concentration x of AE_x(Solv)_yHfNCl (AE = Ca, Sr, Ba; Solv = NH₃, THF) onto the T_c is shown in figure 36 [59]. Note that the T_c is hardly influenced by the doping concentration x up to x ∼ 0.4, corresponding to x ∼ 0.8 for monovalent alkali metal intercalated compounds A_x(Solv)_yHfNCl. The superconductors are not yet over-doped. The T_c decreases in the following order of the increasing basal spacing (d) with intercalation of Sr, Ba (d = ~10 Å) > Ca_x(NH₃)_y, Ba_x(NH₃)_y, Sr_x(NH₃)_y (d = 12 Å) > Ca_x(THF)_y (d = 15 Å), although a minimum level of doping is certainly necessary for the superconductivity. The as-prepared compounds are cointercalated with ammonia used as a solvent, which can be replaced with THF. With varying electron-doping concentrations and interlayer spacings, the highest T_c of 26.0 K was obtained for the Ca and THF cointercalated compound Ca_0.11(THF)_yHfNCl, a new record of high T_c in the electron-doped metal nitride chloride system [59].

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Multivalent rare-earth metals (Eu and Yb) can be solved in liquid ammonia, and intercalated into β-MNCl (M = Zr, Hf) from the solutions [60]. The T_c values of the ammonia cointercalated compounds RE_x(NH₃)_yHfNCl (RE = Eu, Yb) are comparable with those of AE_x(NH₃)_yHfNCl (figure 36); 24.1 and 23.0 K for Eu_0.13(NH₃)_y and Yb_0.11(NH₃)_y, respectively. The temperature dependence of the magnetic susceptibility of RE_x(NH₃)_yMNCl measured in a temperature range of 100–300 K suggests that the RE metals exist as the paramagnetic ions Eu⁺² and Yb⁺³. The anisotropic magnetoresistance measurement has evidenced that the paramagnetism of Eu²⁺ and Yb³⁺ can coexist with the superconductivity even under high magnetic fields up to 14 T, as shown in figure 37 for Eu_0.08(NH₃)_yHfNCl. The anisotropic upper critical fields ($H_{c2}^{\parallel ab}$ and $H_{c2}^{\parallel c})$ were determined with the magnetic field parallel to the ab plane and c axis, respectively, on a uniaxially oriented pellet sample. The anisotropy parameter γ = (d$H_{c2}^{\parallel c}$/dT)/(d/dT) = 4.1, comparable to those of other electron-doped β-MNCl superconductors, 3.7 for Li_0.48(THF)_yHfNCl [313] and 4.5 for ZrNCl_0.7 [314].

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All the T_c values of the electron doped β-HfNCl so far determined can fit on a single curved line as a function of the basal spacing (d) as shown in figure 38 [59], suggesting that the alkali, alkaline-earth and rare-earth metals act as similar electron dopants, and the T_c is hardly dependent on the doping concentration. A very similar trend was also observed in the electron doped β-ZrNCl [315]. It should be noted that the T_c increases with increasing d upon cointercalation of NH₃ and THF, and then decreases gradually with the further increase of d upon cointercalation of PC.

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In a previous study we prepared the alkali metal intercalation compounds of the α-form structured HfNBr, A_x(THF)_yHfNBr (A = Li, Na) [316]. The intercalation was successful, and the color of the layered crystals changed from pale yellow to black. However, the resulting compounds were found to not be superconductors, but insulators with resistivity >10⁷ Ωcm. This is quite a contrast to the fact that the electron doped β-HfNBr shows high-T_c superconductivity, like β-HfNCl. Later we have used α-form layered crystals TiNCl, which are prepared by the ammonolysis of TiCl₄ at elevated temperatures, followed by purification via chemical transport [317]. TiNCl can be intercalated with alkali metals as well as neutral organic molecules such as pyridine and diamines. The intercalation compounds become superconductors. The reason for the quite different behaviors of the two kinds of α-form crystals TiNCl and HfNBr is not clear.

• (i) Alkali metal intercalation in TiNCl.

In the first attempt of the intercalation, TiNCl was subjected to reaction with various kinds of metal azides AN₃ (A = Li, Na, K, Rb) at elevated temperatures under vacuum [317]. The azides are thermally decomposed to metal and nitrogen, and the resulting metal is intercalated into the interlayer space between chloride layers. Part of the alkali metal is also used to extract or deintercalate chloride ions from the interlayer space, forming a metal chloride

$$\begin{eqnarray*}&&\begin{array}{lllllllllllllll} {\rm TiNCl}+\left( x+y \right){\rm M}{{{\rm N}}_{3}}\to {{{\rm M}}_{x}}{\rm TiNC}{{{\rm l}}_{1-y}} \\ \quad +y{\rm MCl}+3/2\left( x+y \right){{{\rm N}}_{2}}. \\ \end{array}\end{eqnarray*}$$

The metal intercalated compounds show superconductivity with T_c ∼ 16 K, although the superconducting volume fractions determined from the diamagnetic expulsion were found to be as low as 5–30%. The Rietveld analysis of the x-ray powder diffraction data revealed that the TiNCl crystalline layers are mutually shifted to accommodate the metal atoms between the chloride layers as shown in figure 39. The space groups of the resulting new polytypes are Bmmb for Na, and Immm for K or Rb intercalated compounds. The space group of Li_xTiNCl is unchanged from Pmmn of the pristine crystal.

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A mild intercalation reaction using metal naphthalene solutions in THF can produce superconductors with a much higher volume fraction >60% [56]. The as-prepared compound is co-intercalated with THF. As shown in figure 40, the basal spacing (d) of the compounds varies depending on the cointercalation conditions. Na_0.16(THF)_yTiNCl has a basal spacing of 13.10 Å (T_c = 10.2 K), which decreases to 8.44 Å (T_c = 18.0 K) upon removal of THF by prolonged evacuation at 90 °C. The THF molecules can be replaced with larger size solvent molecules PC (propylene carbonate); the basal spacing increases to 20.53 Å with decreasing T_c to 7.4 K. Figure 41 shows the T_c values as a function of 1/d for alkali metals and solvent cointercalated compounds. The T_c decreases with the increase of the basal spacing (d). The data fit on a linear line passing through the origin, suggesting the importance of the Coulomb interlayer coupling in the pairing mechanism in this system. It should also be noted that the T_c of non-cointercalated compounds A_xTiNCl (A = Na, K, Rb) also fit on this line except Li_0.13TiNCl. The as-prepared sample Li_0.13(THF)_yTiNCl has a basal spacing of 13.1 Å similar to that of Na_0.16(THF)_yTiNCl, and a T_c = 10.2 K. The non-cointercalated compound Li_0.13TiNCl was obtained by evacuation at 150 °C, which has the smallest basal spacing of 7.8 Å, the same as that of the pristine TiNCl. Unexpectedly, T_c was found to be ∼6.0 K, much lower than the value expected for the small basal spacing of figure 41. Li ions are small enough in size to penetrate into chlorine layers, forming double LiCl layers between [TiN]₂ layers, [TiN]₂(Li_0.13Cl)(ClLi_0.13)[TiN]₂, in which Li ions are located close to TiN superconducting layers in parallel with Cl atoms. On the other hand, in the THF cointercalated compound, Li ions can be coordinated with THF molecules between chlorine layers [ClTi₂N₂Cl]Li_0.26(THF)_y[ClTi₂N₂Cl]. The low T_c of Li_0.13TiNCl against the small d suggests that the location of positive centers may also influence the Coulomb interlayer coupling for superconductivity. The linear relation shown in figure 41 appears to be applied to the structure where the positive centers are located between the chlorine layers as shown in figure 40 for the Na_0.16(THF)_y cointercalated compound. Another systematic study on the relation between the basal spacing and T_c has been performed on TiNBr, and again a similar linear relation was found [57].

• (ii) Anisotropic superconducting properties.

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The anisotropic magnetic susceptibility of the layered superconductor Na_0.16(THF)_yTiNCl was measured on the highly oriented pellet sample with the magnetic field parallel and perpendicular to the c and ab plane [56]. Figure 42 shows the anisotropic H_c2 thus determined as a function of temperature. The anisotropy parameter was calculated to be γ = 1.5 and 1.2 for Na_x(THF)_yTiNCl and Na_xTiNCl, respectively. Note that the TiNCl superconductors exhibit rather isotropic or 3D character. The expansion of the basal spacing from 8.44 to 13.10 Å by cointercalation of THF has little effect on the anisotropy. The superconductor derived from TiNBr also shows a similar γ value [57]. The characteristic superconducting parameters of the α- and the β-structured layered nitride superconductors are compared in table 12. The anisotropy parameter γ = ξ_ab/ξ_c (ratio of the coherence lengths in the ab plane and along the c-axis) for K_0.21TiNBr is calculated to be ∼1.3, close to that found in electron-doped TiNCl; Na_0.16TiNCl (γ = 1.2, T_c = 18.1 K) and Na_0.16(THF)_yTiNCl (γ = 1.5, T_c = 10.2 K) [56]. In contrast, the β-structured nitrides show the anisotropy parameter γ as large as 3.7–4.5 [59, 60, 313, 314]. The small anisotropy parameter γ appears to be one of the characteristic features of the α-structured superconductors. The coherence length along the c-axis (ξ_c) of β-Li_0.48(THF)_yHfNCl is about 16 Å, comparable with the basal spacing 17.8 Å, i.e., the separation of the superconducting layers. This suggests that the superconducting β-form layers may be weakly Josephson coupled. On the other hand, in the α-form TiNBr and TiNCl, the ξ_c are more than three times larger than the basal spacing, implying that the nitride layers are more strongly coupled.

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High-T_c layered cuprate superconductors have large anisotropy parameters γ on H_c2, varying in the range of 3–30 [318]. Iron pnictide superconductors LnFeAs(O, F) (Ln = La, Sm, Nd) recently discovered also show a large anisotropy of γ = 4–9 [319–323]. The large anisotropy parameters of the layered compounds have been considered to be important to realize high-T_c superconductivity. The electron-doped β-ZrNCl and β-HfNCl are also classified into this category. However, the Ba-122 superconductors such as (Ba, K)Fe₂As₂ with T_c = 38 K have been developed, which have a small isotropic parameter γ = 1.5–1.9 [231, 319, 324–327]. The layer coupling through intervening (Ba, K) atoms seems to be stronger than those of the 1111 pnictides coupled through metal oxide layers. It is interesting to note that β-ZrNCl_0.7 (table 12), which is electron-doped by a partial deintercalation of chlorine atoms from the interlayer space, has an anisotropy parameter as large as 4.5. In β-ZrNCl_0.7 with d = 9.8 Å, the nitride layers should be directly coupled without intervening alkali atoms. Nevertheless, the anisotropy parameter is comparable to, or even larger than, that of the cointercalated compound β-Li_0.48(THF)_yHfNCl. It is evident that the interlayer separation is not a decisive parameter for the anisotropy on H_c2. The small anisotropy on H_c2 and the Coulomb coupling between the superconducting layers should be the relevant nature of the superconductivity of the α-form layered nitrides [57]. For more discussion on the superconducting mechanisms and the anisotropy of the two different kinds of layered nitrides, a theoretical study including the electric band structure is required [328].

In the formation of the intercalation compound of β-form layered compounds, organic solvent molecules are cointercalated with metal atoms; organic molecules alone cannot be intercalated. In contrast, α-form TiNCl can intercalate neutral organic molecules without metal atoms. TiNCl can form an intercalation compound with pyridine from liquid and gas phases, Py_0.25TiNCl. The basal spacing increases to 13.5 Å with the molecular plane oriented perpendicular to the layers as schematically shown in figure 43 [317]. The compound becomes a superconductor with T_c = 8.6 K. The T_c is different from those of the alkali metal intercalated compounds. It is interesting to develop different kinds of organic compounds which can be intercalated to obtain high-T_c superconductors. Although the doping mechanism is not yet clear in the intercalation compound with pyridine, it would be reasonable to estimate that the lone pair electrons of nitrogen atoms in pyridine may act as electron donors to the TiNCl layers. It was reported that FeOCl isotypic with TiNCl forms an intercalation compound with pyridine, and the electrical conductivity increases by about seven orders due to the charge transfer from the organic Lewis base to the FeOCl layers [329]. In that study, various kinds of aliphatic amines have been intercalated into TiNCl to develop new superconductivity.

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n-alkyl monoamines (C_nH_2n+1NH₂, 3 ≤ n ≤ 12) can form intercalation compounds with TiNCl, expanding the basal spacing to a value in the range of 12.0 to 37 Å, with the alkyl chains oriented in various ways. All of the compounds with n-alkyl monoamines are not superconductors down to 2 K [55]. It is interesting that ethylene diamine (NH₂CH₂CH₂NH₂) can form a similar intercalation compound with TiNCl with a basal spacing of 11.12 Å, which shows superconductivity with T_c = 10.5 K [55]. Systematic studies have been done using alkylene diamines with different numbers of carbon atoms, NH₂C_nH_2nNH₂ (2 ≤ n ≤ 12). The results are shown in figure 44 [55]. Most of the diamine intercalation compounds are superconductors. The basal spacings are near 12.5 Å irrespective of the chain length of diamines, suggesting that the alkylene chains are aligned with the molecular axis parallel to the layers, and oriented along the b-axis as shown in figure 45. Diamines with an even number of carbon atoms appear to have larger superconducting volume fractions, and the diamines with longer alkylene chains are suitable for higher T_c. The compound with n = 10 (decamethylene diamine, DMDA) shows a large volume fraction > 50%, and T_c = 17.1 K, which is comparable with T_c = 18.1 K of Na_0.16TiNCl. Mechanisms for the superconductivity are not clear, and remain open problems for physicists as well as chemists.

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Ca–Al–Si ternary metal mixtures were pre-melted using an arc furnace, followed by remelting at 1000–1200 °C using radio frequency induction heating under an Ar atmosphere in an h-BN crucible. The cooled ternary mixtures were supplied for the HPHT treatment up to 5–13 GPa and 600–1000 °C by using a Kawai-type multianvil apparatus. A new ternary compound Ca₂Al₃Si₄ was obtained above 650 °C under a pressure of 5 GPa. It crystallizes with the space group Cmc2₁ and the lattice parameters a = 5.8846(8), b = 14.973(1), and c = 7.7966(8) Å [45]. The structure is composed of an aluminum silicide framework [Al₃Si₄] and layer structured [Ca₂] network interpenetrating with each other as shown in figure 46. The [Ca₂] subnetwork has an isomorphous structure with black phosphorus. The compound shows superconductivity with T_c of 6.4 K.

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Under a higher pressure of 13 GPa at 1000 °C, solid solutions Ca(Al_1−xSi_x)₂ (0.35 ≤ x ≤ 0.75) isomorphous with the cubic Laves phase were obtained [71]. As shown in figure 47 the structure can be regarded as a kind of clathrate compound composed of face-sharing truncated tetrahedral cages with Ca atoms at the center, Ca@(Al,Si)₁₂. The compound with a stoichiometric composition CaAlSi shows superconductivity with T_c of 2.6 K [71]. This is the first superconducting Laves phase compound composed solely of commonly found elements.

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Although the ternary system Al–Mg–Si is an important subject in the development of commercial Al alloys [335], Mg₂Si with antifluorite structure is the only compound known as bulk phase in the ternary system under ambient pressure. However, various kinds of fine coherent precipitates are formed in the Al matrix during low temperature aging, which are considered to play an important role for the hardening of a commercial Al alloy. It is reasonable to assume that the ternary and binary precipitates found in Al-based alloys are formed under high pressure generated by Al matrix of the alloy. New binary and ternary compounds in the ternary system Al–Mg–Si have been prepared using HPHT conditions, and the structures are determined using single crystals. Some of them become superconductors.

Ternary compounds Mg(Mg_1–xAl_x)Si (0.3 < x < 0.8) have been prepared under HPHT conditions of 5 GPa at 800–1100 °C. The single crystal study revealed that the compound (x = 0.45) is isomorphous with the anticotunnite, or the TiNiSi structure, and crystallizes with space group Pnma, with lattice parameters a = 6.9242(2), b = 4.1380(1), c = 7.9618(2) Å, and Z = 4. The compound with x > 0.5 shows superconductivity with a T_c ∼6 K [70]. The compound is a peritectic solid solution associated with other phases such as Mg₉Si₅, Al, and Si, depending on cooling protocols in the preparation. The band structure calculation on the composition of MgAlSi suggests that the Al and Mg orbitals mainly contribute to the density of states near the Fermi level, and the substitution of Mg with Al favors superconductivity.

Two kinds of magnesium-based compounds Mg₉Si₅ and Mg₄AlSi₃ have been prepared under a similar HPHT condition. Single crystal study revealed that Mg₉Si₅ crystallizes in space group P6₃ (no. 173) with the lattice parameters a = 12.411(1) Å, c = 12.345(1) Å, and Z = 6 [46]. The structure can be derived from the high pressure form Mg₂Si with anticotunnite structure; excess Si atoms of Mg₉Si₅ form Si–Si pairs in the prismatic cotunnite columns running along the c-axis. Mg₄AlSi₃ is obtained by a rapid cooling of a ternary mixture Mg:Al:Si = 1:1:1 from ∼800 °C to room temperature under a pressure of 5 GPa. The compound crystallizes in space group P4/ncc (no. 130) with the lattice parameters a = 6.7225(5) Å, c = 13.5150(9) Å, and Z = 4 [46]. As shown in figure 48, the structure consists of an alternate stacking of [AlSi₂] layers having a Cairo pattern and [Mg₄Si] antitetragonal prismatic layers. It can be viewed as composed of hexa-Si-capped tetragonal prismatic cages Mg₈Si₆ with an Al atom at the center of each cage, Al@Mg₈Si₆. The compound shows superconductivity with a transition temperature T_c = 5.2 K. The formation regions of the two kinds of new magnesium-based compounds have been proposed.

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Silicon and carbon chemistries are often discussed comparatively from the viewpoint of belonging to the same group in the periodic table [336, 337]. It is interesting to prepare metal doped carbon clathrate compounds, in which carbon forms a sp³ clathrate network-like open diamond framework. A hole-doped diamond was found to become a superconductor [338]. If the synthesis of an electron doped carbon clathrate is realized, it is expected that the carbon sp³ framework with a high Debye temperature should exhibit high-T_c superconductivity [339].

We have already obtained 3D carbon frameworks with cages by polymerization of C₆₀ crystals under HPHT conditions [340, 341]. An attempt has been made to prepare a Ba doped clathrate-like structure from Ba-doped fulleride, Ba₃C₆₀, using HPHT conditions.

The powder Ba₃C₆₀ sample was compressed using a Kawai-type multianvil press at 5–15 GPa and 500–1150 °C [97]. The x-ray powder diffraction (XRD) pattern showed that the Ba₃C₆₀ compressed under 15 GPa at 900 °C is changed into an amorphous solid, which was found to be chemically stable in air and even in water. The high Vickers micro-hardness (1700 kg mm⁻²) and the Raman spectra of the solids suggest that the C₆₀ molecules are collapsed to form an amorphous 3D polymer encapsulating the Ba atoms. As shown in figure 49, the solid obtained by compression under 15 GPa at 900 °C shows a semi-metallic conductivity. It is interesting to note that the Hall coefficient of this sample is positive in the whole temperature range <300 °C, indicating that the dominant carriers are holes. As shown in figure 50, the covalent diameter of a Ba atom is too large to substitute one C atom in the carbon matrix. The diameter is rather comparable with the covalent diameter of a six-membered carbon ring. If a six-membered carbon ring with 12 electrons in the sp³ hybridized orbitals is substituted by a Ba atom with a similar size having only two valence electrons, the carbon matrix should be efficiently hole-doped. The electrical characteristics such as conductivity, Hall coefficient, carrier density, and mobility of the carbon matrix encapsulating Ba atoms were found to be comparable with those of B-doped diamond [97]. The Ba-encapsulated carbon matrix obtained in this study is amorphous. It is interesting to prepare crystalline carbon analogs for silicon clathrate compounds.

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The SrPt₂Sb₂ had been known to have the CaBe₂Ge₂-type structure with tetragonal cell of a = 4.603 and c = 10.565 Å [351] but any physical properties had not been reported. In the FIRST Project, this material was revisited to elucidate its physical properties [41]. Samples were synthesized starting from Sr, Pt and Sb in two-step procedures: arc melting and re-melting of the arc-melted specimen. Though electron probe microanalysis confirmed the composition of SrPt₂Sb₂ (122) for the major part of the ingot obtained, the powder x-ray pattern was not consistent with the CaBe₂Ge₂-type tetragonal lattice. Thus, SrPt₂Sb₂ has a different structure which may be some derivative of the CaBe₂Ge₂-type. As shown in figure 52, electrical resistivity, magnetization and specific heat measurements confirmed a bulk superconducting transition at T_c = 2.1 K for SrPt₂Sb₂. It is a type-II superconductor with a lower critical field (H_c1) of 6 Oe and upper critical field (H_c2) of 1 kOe at 1.8 K. Debye temperature (Θ_D) and electronic specific heat coefficient (γ) were derived from specific heat data to be Θ_D = 183 K and γ = 9.2 mJ (mol K²)⁻¹. The normalized specific heat jump at T_c was calculated as ΔC(T_c)/γT_c = 1.29, consistent with the BCS weak coupling limit of 1.43. Normal state electrical resistivity of SrPt₂Sb₂ exhibited anomalies around 250 K with thermal hysteresis which corresponded to a certain structural transition, though its detail has not yet been elucidated. SrPt₂Sb₂ is a 122-type superconducting antimonide discovered for the first time, and detailed studies for the structure and the phase transition are greatly desired.

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The Ba-derivative of the aforementioned compound, BaPt₂Sb₂ had not been reported, and was prepared first in the FIRST Project by the arc melting method for a Ba, Pt and Sb mixtures [42]. The x-ray pattern of the BaPt₂Sb₂ sample was consistent with a monoclinic lattice having parameters of a = 6.702 Å, b = 6.752 Å, c = 10.47 Å and β = 91.23°. The structure of BaPt₂Sb₂ is shown in figure 53, which can be interpreted as a monoclinic variant of the CaBe₂Ge₂-type structure (note that the a-axis and b-axis of BaPt₂Sb₂ correspond to diagonal of the a₀-axis and b₀-axis of the original CaBe₂Ge₂-type lattice with the relationship of a(b) ≈ √2a₀). Figure 54 gives the temperature dependency of the resistivity of BaPt₂Sb₂, which confirms superconducting transition at 1.8 K. Specific heat data gave Θ_D = 146 K and γ = 8.6 mJ (mol K²)⁻¹, deriving ΔC(T_c)/γT_c = 1.37, which is comparable with the BCS weak coupling limit. Magnetization measurements revealed type-II superconductivity with μ₀H_c2(0) = 0.27 T and Ginzburg–Landau (GL) coherent length ξ_GL(0) = 350 Å. Band structure calculations revealed that Fermi surfaces of BaPt₂Sb₂ resemble those of SrPt₂As₂ with two 2D-like Fermi surfaces and two 3D-like ones. Thus, it has more 3D-like character compared with the AFe₂As₂ system, which may account for the relatively lower T_c of BaPt₂Sb₂.

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Pd-based CaBe₂Ge₂-type antimonides such as LnPd₂Sb₂ (Ln = La, Ce, Pr, Nd, Eu) [353] were known to exist but superconductivity was not reported for this class of compounds. In the FIRST Project, LaPd₂Sb₂ was selected among them and reinvestigated in detail to find the superconducting transition [39]. The LaPd₂Sb₂ sample was prepared at 900 °C starting from La, Sb and Pd in an evacuated silica tube. The resulting specimen was a single phase of the CaBe₂Ge₂-type compound though its tetragonal lattice parameters (a = 4.568 Å, c = 10.266 Å) were slightly different from the previous report [353]. Figure 55 shows the temperature dependence of resistivity for LaPd₂Sb₂ under zero and various magnetic fields. The superconducting transition occurs at onset temperature of 1.4 K and zero resistivity temperature of 1.2 K at H = 0. Magnetization measurements revealed type-II nature superconductivity with μ₀H_c2(0) = 0.86 T and ξ_GL(0) = 233 Å. Bulk superconductivity was confirmed by specific heat measurements which gave Θ_D = 210 K, γ = 6.89 mJ (mol K²)⁻¹ and ΔC(T_c)/γT_c = 1.325, consistent with the BCS weak coupling limit. Figure 56 gives the total density of states for La, Pd1, Pd2, Sb1 and Sb2 (see figure 51 on the atom sites with the suffixes of 1 and 2) calculated by DFT. It is seen that Pd contributes the most to the total DOS at the Fermi level, consistent with a general tendency that the DOS at the Fermi level are dominated by the d-band of the M atom in CaBe₂Ge₂-type AM₂X₂. Figure 56 reveals hybridization between Pd 4d of Pd1 (Pd2) and Sb 5p of Sb2 (Sb1) as demonstrated by synchronized modulation in partial DOS of Pd1 (Pd2) and Sb2 (Sb1). The total DOS at the Fermi level is nearly 40% of that of SrPt₂As₂, which may account for relatively lower T_c of LaPd₂Sb₂ compared with SrPt₂As₂.

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The As-derivative of the aforementioned compound, LaPd₂As₂, which was known to have a ThCr₂Si₂-type structure [353], was also studied in the FIRST Project [40]. In the ThCr₂Si₂-type AEM₂X₂, interlayer X–X bonding is sometimes formed between the adjacent M₂X₂ layers to give the collapsed tetragonal (cT) structure with a 3D network. Superconductivity in the collapsed Fe₂As₂ planes is often attained under high pressure taking well known examples of AEFe₂As₂ (AE = Ca, Sr, Ba, Eu) which undergo both superconducting and collapsed transitions under high pressure [354–357]. Superconductivity under ambient pressure in the cT structure is rather rare and has been reported for AEPd₂As₂ (AE = Ca, T_c = 1.27 K; AE = Sr, T_c = 0.92 K) [358]. In the FIRST Project, it was elucidated that LaPd₂As₂ has the cT structure under ambient pressure with the interlayer As–As distance of 2.318 Å, slightly shorter than the covalent single bond of 2.38 Å for As. Moreover, it was confirmed for the first time that this cT phase shows type-II superconductivity below 1 K. The superconducting and physical parameters obtained for LaPd₂As₂ are μ₀H_c2(0) = 0.402 T (by the Werthamer–Helfand–Hohenberg (WHH) theory using the ρ_90% point as T_c), ξ_GL(0) = 137 Å, Θ_D = 261 K, γ = 5.56 mJ (mol K²)⁻¹ and ΔC(T_c)/γ T_c = 1.17. The density of states at the Fermi level calculated from the specific heat data is as small as ∼0.84 states per eV per formula unit (fu) which may explain the relatively low T_c of this cT phase.

Only few reports are available for Co-based superconducting compounds [359, 360]. In the FIRST Project, LaCo₂B₂, which was reported first in 1973 [361], was re-visited and was found to be superconducting after isovalent or aliovalent substitution for the constituent cations [34]. Polycrystalline samples of (La_1−xY_x)Co₂B₂, La(Co_1−xFe_x)₂B₂ and LaCo₂(B_1−xSi_x)₂ as well as the mother compound were prepared by the arc-melting method for La, Co, B, Y, Fe mixtures. Powder x-ray diffraction of the arc-melted LaCo₂B₂ sample was consistent with the tetragonal ThCr₂Si₂-type structure as reported previously [361], having lattice parameters of a = 3.610 Å and c = 10.20 Å. As shown in figure 57, the 10% Y-doped sample showed superconductivity below 4.4 K. Magnetic susceptibility data revealed Pauli paramagnetism for the normal state of this compound and the M–H curve (figure 57(b) inset) at 2 K reveals type-II nature of superconductivity. Superconductivity with T_c ∼ 4 K was also seen in the Fe-doped system of La(Co_1−x Fe_x)₂B₂ for x ≥ 0.1 while the Si-doped samples did not show superconductivity.

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Electronic structure of LaCo₂B₂ was investigated theoretically by the DFT method; calculated DOS is shown in figure 58. From the DFT calculation, it was confirmed that La ions take +3 state and metallic conduction occurs in the CoB layer composed of highly covalent Co and B. This strong covalency suppresses the spin moment of the Co ion resulting in the Pauli paramagnetic state. Such a situation is caused by the relatively shallow B 2p level compared with the As 4p level in LaFeAsO where antiferromagnetic state is realized rather than the paramagnetic one.

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BaNi₂As₂ has a tetragonal ThCr₂Si₂-type structure at room temperature [348, 362, 363] but it undergoes a structural transition at ∼130 K to a triclinic form where alternate Ni–Ni bonds are formed in the Ni plane (see figure 60(a)) [362]. Below 0.7 K, the triclinic phase shows superconductivity, which is believed to be of conventional BCS-type [348, 363–365]. In the FIRST Project, BaNi₂As₂ was reinvestigated because of this unique structural transition, i.e., this material was expected to offer a stage for studying chemical tuning of soft phonons by elemental substitution and its effects on superconductivity [35].

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Single crystals of BaNi₂(As_1−xP_x)₂ were grown by a self-flux method starting from a mixture of Ba, NiAs, Ni and P [35]. The solubility limit was determined to be x = 0.13 and powder x-ray diffraction confirmed the single-phase nature for all specimens with x ≤ 0.13. Figure 59 shows the temperature dependence of the electrical resistivity parallel to the ab plane of the BaNi₂(As_1−xP_x)₂ crystal. The tetragonal-to-triclinic transition is clearly seen in this figure as the resistivity anomaly with thermal hysteresis. The transition temperature decreases with increasing x and finally the triclinic phase disappears for x ≧ 0.07. Enhancement of superconductivity by phosphorus doping is striking; T_c, which is below 0.7 K for the triclinic phase with x < 0.07, is suddenly increased to 3.33 K in the tetragonal phase with x = 0.077. Figure 60 shows the phase diagram of the BaNi₂(As_1−xP_x)₂ system. Triclinic phase formation is suppressed by phosphorus doping and instead superconductivity is enhanced drastically following the disappearance of the triclinic phase. The Debye frequency ω_D and logarithmic averaged phonon frequency ω_ln calculated from the specific heat of the tetragonal phase exhibit significant softening near the tetragonal-to-triclinic phase boundary. The low-lying soft phonons seem to play, being strongly coupled with acoustic modes, an important role in the enhancement of superconductivity in the tetragonal phase. Indeed, the normalized specific heat jump ΔC(T_c)/γT_c is increased from 1.3 in the triclinic phase to 1.9 in the tetragonal phase with x = 0.077, i.e., the system is transformed from the weak coupling regime to the strong coupling one by phosphorus doping. Such a mechanism of T_c enhancement is common to other systems of CaC₆ [366–368] and Te [369] under high pressure.

IrTe₂ has a trigonal CdI₂-type structure (see figure 61) with the space group of P$\bar{3}$m1. The Ir atom is octahedrally coordinated by the Te atoms and the edge-sharing of the IrTe₆ octahedron forms the Te–Ir–Te composite layers which are stacked along the c-axis of the trigonal lattice. Both the Ir and Te atoms form 2D regular triangular lattices within the ab-plane with three equivalent Ir–Ir (Te–Te) bonds. IrTe₂ undergoes a first-order transition at ∼250 K transforming to a low temperature phase for which a monoclinic structure was first proposed [370]. Recently, a CDW-like superlattice modulation with wave vector of q = (1/5,0,−1/5) was observed in electron diffraction patterns for the low temperature phase [371]. Considering the partly filled d-orbitals of the Ir atoms, the orbital degree of freedom is believed to play an important role in this transition [372, 373]. In this project, the low temperature structure was analyzed in detail using x-ray data for single crystals [374]. The structure at 20 K was found to be triclinic (space group P$\bar{1})$ as shown in figure 62. In this triclinic structure, one out of five Ir–Ir bonds along the trigonal a-axis shrink considerably forming Ir–Ir dimers as illustrated by the yellow hatch in figure 62. The plane of the dimers propagates with the vector of q = (1/5, 0, −1/5) consistent with the aforementioned electron diffraction data. This dimerization seems to affect the physical properties of the system seriously as the first-principles band calculations indicated that tilted two-dimensional Fermi surfaces emerge in the triclinic phase with a possible switching of the conduction plane from the vassal (ab) plane in the trigonal phase to tilted plane normal to q in the triclinic phase [374].

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It was found in 2011 that partial substitution of Pt for Ir in IrTe₂ suppresses the formation of the low temperature phase resulting in the appearance of superconductivity [375, 376]. A phase diagram of the Ir_1−xPt_xTe₂ system was first determined for the polycrystalline samples [63] and then by single crystal data in the FIRST Project [377]. Figure 63 is the phase diagram based on single crystal data; superconductivity appears for ∼0.04 < x < ∼0.14 in Ir_1−xPt_xTe₂ with the highest T_c of ∼3.2 K for x = 0.04 at the phase boundary of the trigonal and triclinic phases (the monoclinic phase in this figure corresponds to the triclinic phase in figure 62, representing the simplified symmetry). For the polycrystalline sample with x = 0.04, type-II superconductivity with μ₀H_c2(0) = 0.17 T and ΔC(T_c)/γT_c = 1.5 has been reported [63]. The Pt substitution works as the electron doping shifting the Fermi level upward and affecting the DOS near the Fermi level (it causes an increase of DOS near the Fermi level in the triclinic phase, and a decrease of DOS in the trigonal phase). Another effect of the substitution, which seems to be more essential, is the suppression of the triclinic phase by breaking the Ir–Ir dimers.

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Similar results on suppression of the phase transition and appearance of superconductivity have been reported for several systems such as Ir_1−xPd_xTe₂ [371], Pd_xIrTe₂ [371] and Cu_xIrTe₂ [378], where substitution or intercalation of Pd or Cu brings about essentially the same effect as the Pt substitution. In the FIRST Project, two more advances were made on IrTe₂ and related systems. The first is the isovalent Rh doping for IrTe₂ [64]. The Ir_1−xRh_xTe₂ system shows similar features to Ir_1−xPt_xTe₂, i.e., triclinic phase formation is suppressed by Rh doping, resulting in the appearance of superconductivity with T_c ∼ 2.6 K, despite that band filling is unchanged by the isovalent Rh doping as long as a rigid-band picture is concerned. A distinct difference was, however, seen between Ir_1−xRh_xTe₂ and Ir_1−xPt_xTe₂; in the former phase, a doping of x ∼ 0.1 is needed for the complete suppression of the triclinic phase, which is three times larger than x = 0.03 in the latter. This difference seems to be caused by the lower volume expansion and the aforementioned unchanged band filling in the case of Rh doping

Another advance was attained from the study of a related system of Au_1−xPt_xTe₂ [62]. AuTe₂ has a monoclinically distorted CdI₂-type average structure with a space group of C2/m where Te–Te zigzag chains run along the a-axis [379]. In actuality, the structure is associated with incommensurate modulation with a wave vector q = −0.4076a* + 0.4479c* and due to this modulation, Te–Te dimers with a short distance of 2.88 Å exist in the real structure instead of the zigzag chains [380, 381]. It was found that Pt doping brings about structural change from the distorted CdI₂-type for x = 0 to distortion-free CdI₂-type without Te–Te dimers for x = 0.35, via the two-phase mixed region for x = 0.1 and 0.15 [62]. Type-II superconductivity was observed for the distortion-free x = 0.35 sample with T_c = 4.0 K, H_c2(0) = 12.9 kOe and ξ_GL(0) = 160 Å. The ΔC(T_c)/γT_c is 1.57, exceeding the BCS weak coupling limit. It should be noted that superconductivity appears by breaking of the Te–Te dimers, which shows striking similarity with the Ir_1−xPt_xTe₂ system where superconductivity appears by breaking of the Ir–Ir dimers. It was also found that superconductivity is induced in AuTe₂ by application of pressure instead of Pt doping [61]. As shown in the phase diagram of figure 64, application of mechanical pressure causes a structural change from the distorted CdI₂-type to the distortion-free CdI₂-type via the two-phase mixed region. Superconductivity appears above 2.12 GPa with the highest T_c of 2.3 K at 2.34 GPa.

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Iridium dichalcogenides IrX₂ (X = Se and Te) sometimes take pyrite-type (FeS₂-type) structures after introduction of vacancies for the Ir sites [382]. It is also known that application of high pressure is effective to stabilize the pyrite-type form against the Cd₂I₂-type one [383]. In the Ir-deficient phase Ir_xX₂, Ir vacancies are distributed randomly forming a cubic pyrite-type structure with a space group of Pa$\bar{3}$ (figure 65). This structure can be interpreted as the NaCl-type constructed by the face centered cubic sublattice of Ir and the X–X dimers located at the center of each edge as well as at the body center of the cubic lattice. In the Ir_xX₂ phase with the particular value of x = 0.25, Ir vacancies tend to be distributed in an ordered way in which one of four Ir atoms are regularly removed [382]. The vacancy ordering results in a stoichiometric phase of Ir₃X₈ having a rhombohedral structure with the space group of R$\bar{3}$ (in Ir₃X₈, the rhombohedral distortion is far less pronounced and the vacancy ordering may not be perfect compared with the corresponding phase of Rh₃X₈ [382]).

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In the FIRST Project, the pyrite-type Ir_xX₂ samples were synthesized under high pressure and their superconductivity was studied for the first time [65]. Figure 66 gives the electronic phase diagram obtained for the Ir_xSe₂ system. By increasing x, the system changes from insulating to metallic nature and superconductivity starts to emerge at the Ir content of x ∼ 0.75 with the highest T_c of ∼6.4 K at x = 0.91 (0.91 is the highest Ir content attained experimentally). The Se–Se distance of the Se dimer increases linearly with increasing x as shown in figure 65. A similar phase diagram was obtained for the Ir_xTe₂ system, with the highest T_c of ∼4.7 K for x = 0.93. DFT calculations were carried out for the vacancy-ordered structural model of Ir₃Se₈, where there exist one long Se₁–Se₁ dimer (the edge site dimer with r_Se1–Se1 = 2.61 Å) and three short Se₂–Se₂ dimers (the body center site dimer with r_Se2–Se2 = 2.50 Å). The DFT calculations revealed that the band crossing the Fermi level consists mainly of the σ* (anti bonding) orbital of the Se₁–Se₁dimer and d_z2 orbitals of the nearest Ir atoms with far less contribution from the σ* orbitals of the Se₂–Se₂ dimers. Such a nature of the electronic structure results in the half-filled narrow conduction band which easily becomes insulating by the electron–electron correlation, electron–lattice interactions and/or the disordered Ir vacancies. DFT calculations were also performed for the vacancy-free pyrite-type structure which is composed of equivalent Se–Se dimers. In this situation, the σ* orbitals of the Se–Se dimers contribute equally to form a wider conduction band. It should be noted that the Ir vacancy introduction causes a linear increase of the Se–Se distance of the Se dimer in correlation with the monotonous increase of T_c. In the CdI₂-type IrTe₂, breaking of the Te–Te dimer is essential for the appearance of superconductivity, while in pyrite-type Ir_xX₂, control of the bonding state of the X–X dimers by elongation and equalization is indispensable for inducing superconductivity. Superconducting parameters obtained for Ir_0.91Se₂ are μ₀H_c2(0) = 14.3 T (type-II superconductor), ξ_GL(0) = ∼48Å and ΔC(T_c)/γT_c = 3.1 (strong coupling superconductor).

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The strong correlation between the Se–Se distance and T_c was also confirmed for the Ir_0.94−xRh_xSe₂ system in the FIRST Project [66]. Figure 67 indicates the phase diagram of the Rh-doped system in question. With increasing Rh content, the system undergoes changes from the non-metal state to the normal-metal state with T-square resistivity via the strange-metal state with T-linear resistivity. Accompanied by this alteration, the Se–Se distance first increases then decreases, taking the maximum at x ∼ 0.4. The striking correlation of the Se–Se distance and T_c, Θ_D and ΔC(T_c)/γT_c is worth noting; both T_c and ΔC(T_c)/γT_c have maximum values, while Θ_D has the minimum value when the Se–Se distance is longest. This suggests strengthening of the electron–phonon coupling and softening of phonon due to the structural instability at the edge of weak dimer states when the Rh content is x ∼ 0.4.

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A similar study has been carried out for the corresponding telluride system of Ir_0.95−xRh_xTe₂ [67]. In this system, the Te–Te distance of the Te-dimer shows a parabolic change taking the maximum between x = 0.2 and 0.3. Nevertheless, aforementioned correlation was not observed between the Te–Te distance and T_c nor ΔC(T_c)/γT_c; both parameters decrease almost linearly with x. This result may be related to the reduced electron–electron repulsion in the telluride system compared with the selenide system, which is caused by the wider width of the conduction band in the telluride system consisting of the more spatially spread Te 5p orbital.

A wide variety of silicide superconductors have been known and most of them crystallize in ceontrosymmetric structures showing conventional s-wave superconductivity. Some exceptions are the Ce-containing heavy-fermion superconductors described above. In the FIRST Project, we found two new noncentrosymmetric silicide superconductors, SrAuSi₃ [49] and Li₂IrSi₃ [50]. Here, their structural and physical properties are overviewed.

SrAuSi₃ is the first noncentrosymmetric superconductor containing Au, which is a heavy element and may cause strong spin–orbit coupling. The SrAuSi₃ is stable only under high pressure and its polycrystalline sample was prepared under high temperature–high pressure conditions [49]. In figure 68, crystal structure of SrAuSi₃ is shown; it is a BaNiSn₃-type tetragonal structure in the space group of I4mm. The structure can be interpreted as sequence of the atom planes along the c-axis as Sr–(Au–Si₂–Si)–Sr–(Au–Si₂–Si)–Sr... which has a close relationship with the ThCr₂Si₂-type and CaBe₂Ge₂-type structures (see figure 51). Figure 69 shows the temperature dependence of electrical resistivity measured by varying the magnetic field (H). T_c of SrAuSi₃ is 1.6 K at H = 0 and decreases with increasing magnetic field, giving H_c2(0) ∼ 2.2 kOe which is much lower than the Pauli limit (H_p(0) ∼ 30 kOe). This result suggests that H_c2 is governed by the orbital pair breaking mechanism. Indeed, the orbital limit estimated from the WHH theory is ∼1.5 kOe, comparable with the experimental value. However, H_c2 of SrAuSi₃ increases almost linearly with temperature deviated substantially from the WHH convex upward curve. This deviation may suggest a nonspherical Fermi surface or gap anisotropy in SrAuSi₃. The GL coherent length and penetration depth were estimated to be ξ_GL(0) = 390 Å and λ(0) = 4400 Å giving the GL parameter κ_GL = 11 consistent with the type-II nature of superconductivity.

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Specific heat data gave γ = 6.0 mJ (mol K²)⁻¹ and Θ_D = 410 K. The normalized specific heat jump is calculated to be ΔC(T_c)/γT_c = 1.92 and the electron–phonon coupling to be λ_ep ≈ 0.97 using density of states at the Fermi level, 1.3 states (eV f.u.)⁻¹ from the band calculation. These parameters indicate that SrAuSi₃ is a moderately strong coupling superconductor. T_c was estimated by the McMillan formula [391], T_c = (Θ_D/1.45) × exp{–1.04(1 + λ_ep)/[λ_ep – μ*(1 + 0.62 λ_ep)]}, using the standard value of 0.13 for the Coulomb repulsion parameter μ*, resulting in ∼19 K. The large difference between this value and experimental T_c of 1.6 K is striking; it may be caused by the parity mixing or the gap anisotropy in the noncentrosymmetric superconductivity. The DFT band calculation for SrAuSi₃ revealed that two-types of carrier exist on multiple Fermi surfaces. The major carriers conduct in the Si layers while other types of carriers conduct through the 3D network in the structure. Some of the latter carriers looked to have asymmetric spin–orbit coupling, suggesting that they may cause the unusual behaviors in SrAuSi₃.

IrSi₃ crystallizes in a hexagonal structure with the space group P6₃mc [392] (figure 70) where planar layers of four-fold Si (a distorted kagome network) are stacked along the c-axis sandwiching the Ir atoms. The Ir atoms are placed with unequal distances from the upper and lower neighboring Si planes leading to a polar structure with no central symmetry. In the FIRST Project, Li atoms were intercalated to this structure for the first time to obtain Li₂IrSi₃ [50]. Li₂IrSi₃ has a trigonal structure with the space group of P31c as given in figure 70 where the essential nature of the Si plane stacking is preserved. In Li₂IrSi₃, the positions of Ir and Li are not symmetric concerning the distances from the upper and lower neighboring Si planes, leading to the nonpolar asymmetry. However, displacement from the symmetrical equivalent position is Δz/c ∼ 0.007, which is one order of magnitude smaller than that of the mother compound of IrSi₃.

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Figure 71 shows resistivity in magnetic field as a function of temperature. Li₂IrSi₃ exhibits type-II superconductivity with T_c of 3.8 K. Temperature dependence of μ₀H_c2 gave μ₀H_c2(0) = 0.3 T and ξ_GL(0) = 330 Å (see the inset in figure 71) while the specific heat data gave γ = 5.3 mJ (mol K²)⁻¹ and Θ_D = 484 K, which are worth comparing with γ = 0.73 mJ (mol K²)⁻¹ and Θ_D = 516 K for IrSi₃. The normalized specific heat jump at T_c is ΔC(T_c)/γT_c = 1.41, consistent with the BCS weak coupling limit. All these experimental data suggest the conventional nature of superconductivity in Li₂IrSi₃; in particular, μ₀H_c2(0) (0.3 T) is much lower than the Pauli limit of μ₀H_P(0) = 6.9 T. The weakened inversion symmetry breaking seems to account for the less noticeable unique nature of superconductivity. The appearance of superconductivity in Li₂IrSi₃ may be concerned with the enhancement of γ, i.e., a seven-fold increase of the electronic DOS at the Fermi level after the intercalation of Li for IrSi₃.

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The ternary equiatomic pnictides with the general formula, M'MPn (M' is a large sized electropositive transition metal, M is a smaller transition metal, Pn is a pnictogen) form a large family of compounds. In the FIRST Project, the equiatomic pnictides with 4d and 5d transition metals for the M site were studied because previous studies have been rather confined to the systems with 3d transition metals for M. Consequently, type-II superconducting transitions were observed for the first time in LaIrP, LaIrAs and LaRhP, which were prepared using the high-pressure synthesis technique [51]. Lattice parameters for LaRhAs and LaIrP were in good agreement with previous reports [393, 394]. In figure 72, the crystal structure is shown for LaRhP, which has a tetragonal lattice with the space group of I4₁md and is an ordered ternary derivative of the α-ThSi₂-type structure. In the structure, the Rh and P atoms are linked forming a 3D network with a trigonal planar coordination, and the La atoms are placed in the cavities of the network. There are two sets of Rh–P zigzag chains running toward the perpendicular directions.

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In table 13, superconducting parameters and the lattice parameters are shown for LaIrP, LaIrAs and LaRhP. Among them, LaIrP has the highest T_c of 5.3 K and the highest H_c2(0) of 13.8 kOe (WHH value) or 16.4 kOe (GL value) corresponding to the highest value of electron–phonon coupling, λ_ep = 0.67. The ΔC(T_c)/γT_c values are less than unity for all three compounds indicating the weak electron–phonon coupling regime. For every compound, the experimental H_c2(0) is much lower than the Pauli limit of field, suggesting that the asymmetric spin–orbital coupling induced by the lack of central symmetry is not significant in the LaMPn system.

As described in the previous section, the 122 compounds seem the most suitable for application to superconducting wires or tapes among IBSCs. The 122 compounds exhibit superconductivity by substituting Co for the Fe site, K for the Ba(Sr) site, or P for the As site. K-doped Ba(Sr)122 has the highest T_c and upper critical field (H_c2) as summarized in table 14. Although synthesis of K-doped Ba-122 (Ba-122:K) epitaxial films by a molecular beam epitaxy (MBE) method was reported [402], the existence of volatile K makes it difficult to fabricate its films by a PLD method. It was also reported that the films were not stable in ambient atmosphere [402]. P-doped Ba-122 (Ba-122:P) has a higher T_c than Ba-122:Co of approximately 30 K [403] and is expected to be stable in ambient atmosphere. Adachi et al of ISTEC group chose this compound as a material candidate for production of superconducting tapes by a PLD method using a second-harmonic Nd:YAG laser, and examined the film preparation conditions on MgO single crystal substrates [404].

High-purity targets with a nominal composition of BaFe₂(As_0.6P_0.4)₂ were carefully synthesized by a conventional solid-state reaction method. Epitaxial films were obtained at a substrate temperature of approximately 800 °C. The energy density on the target was relatively high, approximately 10 J cm⁻², leading to a deposition rate of 5 nm s⁻¹ at the repetition rate of 10 Hz and the substrate–target distance of 7 cm. The average FWHM value of the peaks in the ϕ-scan was about 1.5°. The obtained film exhibited a ${{T}_{{\rm c}}}^{{\rm onset}}$ and ${{T}_{{\rm c}}}^{{\rm zero}}$ of 26.5 and 24.0 K, as seen in the resistive transition curve of figure 76. An even higher T_c of 27.0 K was also observed for a film without patterning. The first synthesis of Ba-122:P epitaxial films was previously achieved by an MBE method [405]. The observed T_c values of the PLD films are comparable to those reported for the MBE films [405] and the single crystals [403]. Figure 77 shows the dependence of J_c on the applied field along the c-axis at different temperatures for the Ba-122:P epitaxial film with the ${{T}_{{\rm c}}}^{{\rm onset}}$ of 26.5 K and a self-field J_c value at 4.2 K of 3.5 MA cm⁻². The film exhibited rather high in-field J_c values, for example, approximately 1 MA cm⁻² at 4.2 K, 3 T and 10 K, 1 T, which are higher than those for Ba-122:Co films [406].

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The dependence of superconducting properties of BaFe₂(As_1−xP_x)₂ (Ba-122:P) epitaxial films on the P content was systematically investigated by Miura et al by using Ba-122:P targets with nominal P content x of 0.25, 0.33, 0.40 and 0.50 [407]. The x values in the films analyzed using an electron probe micro analyzer (EPMA) were found to be slightly smaller by approximately 0.05 than the nominal x values in bulk targets. Figure 78 shows the temperature dependence of H_c2 for Ba-122:P films with various analyzed x values in magnetic fields up to 12 T. The x = 0.28 film exhibits a maximum ${{T}_{{\rm c}}}^{{\rm zero}}$ of 26.5 K, and a further increase in x leads to a reduction in ${{T}_{{\rm c}}}^{{\rm zero}},$ while the x = 0.19 film showed a broad transition and ${{T}_{{\rm c}}}^{{\rm zero}}$ of about 12 K. This x dependence of T_c is similar to that observed in single crystals [403]. The x = 0.28 film also exhibits the highest H_c2 in both field directions. The inset shows the anisotropy of H_c2, γ_H = ${{H}_{{\rm c}2}}^{{\rm ab}}$/${{H}_{{\rm c}2}}^{{\rm c}}$ for the films. The x = 0.28 film with the optimal T_c shows the smallest γ_H value of 1.54. This is different from the case of cuprate superconductors [408], and preferable from the viewpoint of wire or tape application. Figure 79 shows the angular dependence of J_c for the x = 0.28 film at 10 K under different magnetic fields. This film also exhibits high in-field J_c values. It is also found that J_c shows minimum values in the c-axis direction at all fields from 0.5 to 7 T, indicating that no c-axis-correlated pinning centers are included in this film.

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The Ba-122:P epitaxial films on MgO substrates fabricated by a PLD method showed relatively high in-field J_c values. However, these values do not seem to be high enough for high-field application. Since the Ba-122 superconductors have very high H_c2(0) values over 60 T, a remarkable improvement in in-field J_c would be expected, if effective vortex pinning centers could be introduced into the materials. Actually, amorphous tracks induced in Ba-122:K single crystals by heavy-ion irradiation significantly enhanced their in-field J_c and resulted in a matching field as high as 21 T without T_c degradation [409], although heavy-ion irradiation is not practical for fabrication of long wires or tapes. For the case of Ba-122 epitaxial films, naturally formed defects in Ba-122:Co films enhanced the J_c in the fields parallel to the c-axis to some extent [400], as described in the previous section. It was also reported that naturally formed nanopillars along the c-axis in the Ba-122:Co thin films prepared by PLD using a SrTiO₃ buffer layer and oxygen-rich targets significantly enhanced the in-field J_c around the c-axis direction [410, 411]. However, the J_c around the ab plane direction was not much improved.

For application to superconducting tapes, it is desirable to find a controllable and practical way of enhancing vortex pinning in Ba-122 films in an isotropic way. In cuprate superconductor films such as REBa₂Cu₃O_y (REBCO; RE = Y or rare-earth elements) films, the addition of nanoparticles consisting of second oxide phases has been found to be effective in enhancing their in-field J_c in an isotropic way [412–414]. However, in order to apply a similar technique to Ba-122 epitaxial films, the second phases need to be chemically stable and crystallographically compatible with the Ba-122 phase, and its size should be small enough to avoid blocking of the current path.

Miura et al tried introducing BaZrO₃ (BZO) nanoparticles, which are known to be effective pinning centers in REBCO films, into Ba-122:P epitaxial films on MgO by a PLD method [415]. Approximately 80 nm thick BZO-doped Ba-122:P epitaxial films were grown from 1 mol.% and 3 mol.% BZO-doped Ba-122:P targets with the nominal P content of 0.33. The deposition conditions were similar to those employed by Adachi et al for the synthesis of undoped Ba-122:P films [404]. The cross-sectional elemental maps and x-ray diffraction pattern shown in figure 80(a) of the film grown from the 3 mol.% BZO-doped target indicated the presence of homogenously dispersed BZO nanoparticles. The average nanoparticle size and the average spacing were found to be 8 nm and 24 nm, respectively, leading to a density n of approximately 6.8 × 10²² m⁻³. The cross-sectional TEM image of a typical BZO nanoparticle is shown in figure 80(b). The size of the nanoparticle is about 5 and 10 nm parallel and perpendicular to the c-axis, respectively. The periodicity of the Ba-122 planes around the nanoparticles is only perturbed by the creation of stacking faults. However, nano-beam diffraction patterns for the nanoparticle and the Ba-122:P matrix indicated that the BZO nanoparticles are not epitaxially oriented along the Ba-122:P matrix. The ${{T}_{{\rm c}}}^{{\rm zero}}$ values for the undoped and doped films are 26.3 and 25.0 K, respectively, indicating that introduction of the BZO nanoparticles does not induce significant T_c degradation.

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Figure 81(a) shows the field dependence of J_c (H//c) at 5 K for the undoped Ba-122:P and BZO added Ba-122:P films. The self-field J_c (${{J}_{{\rm c}}}^{{\rm s}.{\rm f}.})$ increases monotonically with the amount of BZO additive and the field decay of J_c is greatly reduced within the measured field range. The Ba-122:P film shows a characteristic crossover field H* (90% of J_c(H)/${{J}_{{\rm c}}}^{{\rm s}.{\rm f}.}),$ which is indicated by the arrows, followed by a power-law regime (J_c ∝ H^−α) with α ∼ 0.40 at intermediate fields. A more rapid decay of J_c is observed as H approaches the irreversibility field, H_irr. For the Ba-122:P + BZO films, H* also increases with the BZO content. At intermediate fields, we find that both films with BZO show a slower decay of J_c(H), indicating the importance of BZO nanoparticles to enhance J_c(H) in magnetic fields. The non-power-law dependence observed for the Ba-122:P + BZO films is similar to that observed in RE123 films with strong pinning coming from uniformly dispersed nanoparticles [413, 414, 416].

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Figure 81(b) shows the angular dependence of J_c, J_c(θ), curves measured for the three films at 1 T, 15 K. By adding BZO the J_c increases for all orientations with respect to the J_c of the Ba-122:P film. In particular, the Ba-122:P + 3 mol% BZO film exhibits an almost isotropic J_c with the value for H//c 2.6 times higher than that of the Ba-122:P film. The minimum value of J_c(θ), J_c,min, of 1.5 MA cm⁻² at 1 T, 15 K is over 28 times and 7 times higher than that of Ba-122:Co films with c-axis columnar defects [410] and Ba-122:Co films with super-lattice structures [417], respectively, in very similar field and temperature conditions, indicating strong isotropic pinning by the BZO nanoparticles.

In figure 82, the pinning force, F_p = J_c(H) × μ₀H is compared with that of several superconductor materials. At 15 K, F_p for the Ba-122:P + 3 mol% BZO film is over 3 times higher than that for the Ba-122:P film and higher than NbTi [418] at 4.2 K at all magnetic fields. Comparing with MgB₂ data at 15 K [419], F_p is clearly higher for μ₀H > 0.5 T. At T = 5 K, the F_p of the Ba-122:P + 3 mol.% BZO film, which is the minimum value in all field directions, reaches ∼59 GN m⁻³ for μ₀H > 3 T up to the highest field we measured (9 T), a 50% increase over Nb₃Sn at 4.2 K [420].

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Miura et al also found that the characteristic magnetic field, where the maximum relative J_c enhancement by BZO addition was observed, increased monotonically by increasing the nanoparticle density. From the comparison with the previous results of REBCO films with BZO nanoparticles [421], it was deduced that the effective way to optimize the J_c performance of IBSCs as well as cuprates is dispersing nanoparticles with the average size smaller than ∼3*(2ξ(T)) and with densities such that the nanoparticle spacing matches the intervortex distance. These results suggested the possibility of further enhancing in-field J_c properties, in particular at higher fields, by optimizing the landscape of nanoparticles or nanoscale defects in Ba-122:P epitaxial films.

Sato et al carefully investigated the properties of Ba-122:P epitaxial films on MgO (001) single-crystal substrates prepared by PLD using the second harmonic of an Nd:YAG laser and found that films with very high and less anisotropic in-field J_c could be obtained at certain deposition conditions [422]. They employed a semiconductor infrared diode (λ = 975 nm, and maximum power = 300 W) for substrate heating and achieved high substrate temperature (T_s) up to 1400 °C. High-purity Ba-122:P targets with the nominal P content x of 0.30 were used. It was found that 150–200 nm thick epitaxial films with high crystallinity (Δϕ and Δω well below 0.8°) were grown at rather high T_s of 1000–1100 °C (optimum at 1050 °C) and low growth rate of 0.2–0.4 nm s⁻¹ for laser fluence of 3.0–3.5 J cm⁻². Their lattice parameters slightly different from the single crystal data [403] indicated existence of tensile strain in the films. The optimum film exhibited a T_c of 26.5 K, a narrow transition width ΔT_c of 1.5 K, and a very high self-field J_c up to 7 MA cm⁻². This J_c value is comparable to the recently reported high value for a film grown by an MBE method [423].

Figure 83 shows the angular dependence of J_c at 12 K, 3 T for Ba-122:P epitaxial films grown at the optimum T_s. The J_c(θ_Η) curves exhibit a broad peak around the c-axis direction (θ_Η = 0°) in addition to the intrinsic J_c peak at θ_Η = 90°, indicating existence of pinning centers along the c-axis. With decreasing growth rate from 0.39 nm s⁻¹ to 0.22 nm s⁻¹, the θ_Η = 0° peak becomes more prominent and the J_c values in all directions become remarkably higher, resulting in less anisotropic angular dependence. These results indicate that the vortex pinning properties and J_c anisotropy of Ba-122:P epitaxial films can be controlled by tuning the growth rate.

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In figure 84, the magnetic field dependence of ${{J}_{{\rm c}}}^{H//ab}$ (closed symbols) and ${{J}_{{\rm c}}}^{H//c}$ (open symbols) of the optimum Ba-122:P epitaxial films is compared with those reported for other Ba-122 epitaxial films with high J_c [415, 417, 424–426] as well as SmFeAsO_1−xF_x and Fe(Se,Te) films [427, 428]. The optimum Ba-122:P film exhibits an even higher ${{J}_{{\rm c}}}^{H//{\rm c}}$ than the Ba-122:P film with BZO nanoparticles at H > 4 T and a ${{J}_{{\rm c}}}^{H//c}$ as high as 0.8 MA cm⁻² at 4 K, 9 T, resulting in the highest pinning force of 72 GN m⁻³. It would also be the noteworthy that the ${{J}_{{\rm c}}}^{H//ab}$ value of 1.1 MA cm⁻² at 9 T giving a pinning force of 99 GN m⁻³ is the highest obtained for IBSC films.

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Figure 85 shows the cross-sectional bright-field scanning TEM (STEM) images for Ba-122:P epitaxial films grown at the growth rate of 0.22 nm s⁻¹ and 0.39 nm s⁻¹. As indicated by the vertical white arrows in figure 85, there are many vertical defects with a substantially higher density than that observed in the Ba-122:Co epitaxial films by Katase et al. It is also found that most of the defects in the film grown at 0.22 nm s⁻¹ start appearing at mid-thickness and are oriented parallel to the c-axis, while the defects in the latter film originate just at the substrate surface and are tiled with respect to the c-axis. Other planar or line defects in the ab plane, such as stacking faults, were not observed. Energy dispersive x-ray spectroscopy (EDX) combined with STEM revealed that the chemical composition of the defects is the same as that of the matrix region and that the impurity oxygen concentration in the films is less than the detection limit. These results indicated that the defects are not an impurity phase such as BaFeO₂ [429] but may be edge or threading dislocations and/or domain boundaries. The stronger vortex pinning along the c-axis for the film grown at the lower rate can be explained by a high density of straight defects with a lateral size (4 nm) close to the double of the ξ_ab of Ba-122:P at 4 K [407]. The very high in-field J_c observed in these epitaxial films indicates a high potential of Ba-122:P epitaxial films for application to superconducting tapes or wires, though the very high T_s would not be favorable to film fabrication on metal substrates.