Compositional Dependence of Curie Temperature and Magnetic Entropy Change in the Amorphous Tb–Co Ribbons

The Curie temperature (T_c) and magnetic entropy change (−ΔS_m), and their relationship to the alloy composition of Tb–Co metallic glasses, were studied systematically in this paper. It was found that, in contrast to the situation in amorphous Gd–Co ribbons, the dependence of Tc on Tb content and the maximum −ΔS_m vs. T_c -2/3 plots in Tb–Co binary amorphous alloys do not follow a linear relationship, both of which are supposed to be closely related to the non-linear compositional dependence of Tb–Co interaction due to the existence of orbital momentum in Tb.     

Superionic Solid Electrolyte Li7La3Zr2O12 Synthesis and Thermodynamics for Application in All-Solid-State Lithium-Ion Batteries

Solid-state reaction was used for Li₇La₃Zr₂O₁₂ material synthesis from Li₂CO₃, La₂O₃ and ZrO₂ powders. Phase investigation of Li₇La₃Zr₂O₁₂ was carried out by x-ray diffraction (XRD), scanning electron microscopy (SEM) and energy-dispersive x-ray spectroscopy (EDS) methods. The thermodynamic characteristics were investigated by calorimetry measurements. The molar heat capacity (C_p,m), the standard enthalpy of formation from binary compounds (Δ_oxH_LLZO) and from elements (Δ_fH_LLZO), entropy (S⁰₂₉₈), the Gibbs free energy of the Li₇La₃Zr₂O₁₂ formation (∆_f G⁰₂₉₈) and the Gibbs free energy of the LLZO reaction with metallic Li (∆_rG_LLZO/Li) were determined. The corresponding values are C_p,m = 518.135 + 0.599 × T − 8.339 × T⁻², (temperature range is 298–800 K), Δ_oxH_LLZO = −186.4 kJ·mol⁻¹, Δ_fH_LLZO = −9327.65 ± 7.9 kJ·mol⁻¹, S⁰₂₉₈ = 362.3 J·mol⁻¹·K⁻¹, ∆_f G⁰₂₉₈ = −9435.6 kJ·mol⁻¹, and ∆_rG_LLZO/Li = 8.2 kJ·mol⁻¹, respectively. Thermodynamic performance shows the possibility of Li₇La₃Zr₂O₁₂ usage in lithium-ion batteries.     

Some Thermomagnetic and Mechanical Properties of Amorphous Fe75Zr4Ti3Cu1B17 Ribbons

The microstructure, revealed by X-ray diffraction and transmission Mössbauer spectroscopy, magnetization versus temperature, external magnetizing field induction and mechanical hardness of the as-quenched Fe₇₅Zr₄Ti₃Cu₁B₁₇ amorphous alloy with two refractory metals (Zr, Ti) have been measured. The X-ray diffraction is consistent with the Mössbauer spectra and is characteristic of a single-phase amorphous ferromagnet. The Curie point of the alloy is about 455 K, and the peak value of the isothermal magnetic entropy change, derived from the magnetization versus external magnetizing field induction curves, equals 1.7 J·kg⁻¹·K⁻¹. The refrigerant capacity of this alloy exhibits the linear dependence on the maximum magnetizing induction (B_m) and reaches a value of 110 J·kg⁻¹ at B_m = 2 T. The average value of the instrumental hardness (HV_IT) is about 14.5 GPa and is superior to other crystalline Fe-based metallic materials measured under the same conditions. HV_IT does not change drastically, and the only statistically acceptable changes are visibly proving the single-phase character of the material.     

Influence of Cu Content on Structure,Thermal Stability and Magnetic Properties in Fe72−xNi8Nb4CuxSi2B14 Alloys

The effect of substitution of Fe by Cu on the crystal structure and magnetic properties of Fe_72−xNi₈Nb₄Cu_xSi₂B₁₄ alloys (x = 0.6, 1.1, 1.6 at.%) in the form of ribbons was investigated. The chemical composition of the materials was established on the basis of the calculated minima of thermodynamic parameters: Gibbs free energy of amorphous phase formation ΔG^amorph (minimum at 0.6 at.% of Cu) and Gibbs free energy of mixing ΔG^mix (minimum at 1.6 at.% of Cu). The characteristic crystallization temperatures T_x1onset and T_x1 of the alpha-iron phase together with the activation energy E_a for the as-spun samples were determined by differential scanning calorimetry (DSC) with a heating rate of 10–100 °C/min. In order to determine the optimal soft magnetic properties, the wound cores were subjected to a controlled isothermal annealing process in the temperature range of 340–640 °C for 20 min. Coercivity Hc, saturation induction Bs and core power losses at B = 1 T and frequency f = 50 Hz P_10/50 were determined for all samples. Moreover, for the samples with the lowest Hc and P_10/50, the magnetic losses were determined in a wider frequency range 50 Hz–400 kHz. The real and imaginary parts of the magnetic permeability µ′, µ″ along with the cut-off frequency were determined for the samples annealed at 360, 460, and 560 °C. The best soft magnetic properties (i.e., the lowest value of Hc and P_10/50) were observed for samples annealed at 460 °C, with Hc = 4.88–5.69 A/m, Bs = 1.18–1.24 T, P_10/50 = 0.072–0.084 W/kg, µ′ = 8350–10,630 and cutoff frequency at 8–9.3 × 10⁴ Hz. The structural study of as-spun and annealed ribbons was carried out using X-ray diffraction (XRD) and a transmission electron microscope (TEM).     

Structure,Magnetocaloric Effect and Critical Behavior of the Fe60Co12Gd4Mo3B21 Amorphous Ribbons

The aim of the paper was to study the structure, magnetic properties and critical behavior of the Fe₆₀Co₁₂Gd₄Mo₃B₂₁ alloy. The X-ray diffractometry and the Mössbauer spectroscopy studies confirmed amorphous structure. The analysis of temperature evolution of the exponent n (ΔS_M = C·(B_max)ⁿ) and the Arrott plots showed the second order phase transition in investigated material. The analysis of critical behavior was carried out in order to reveal the critical exponents and precise T_C value. The ascertained critical exponents were used to determine the theoretical value of the exponent n, which corresponded well with experimental results.     

Large Electrocaloric Responsivity and Energy Storage Response in the Lead-Free Ba(GexTi1−x)O3 Ceramics

Ferroelectric property that induces electrocaloric effect was investigated in Ba(Ge_xTi_1−x)O₃ ceramics, known as BTGx. X-ray diffraction analysis shows pure perovskite phases in tetragonal symmetry compatible with the P4mm (No. 99) space group. Dielectric permittivity exhibits first-order ferroelectric-paraelectric phase transition, confirmed by specific heat measurements, similar to that observed in BaTiO₃ (BTO) crystal. Curie temperature varies weakly as a function of Ge-content. Using the direct and indirect method, we confirmed that the adiabatic temperature change ΔT reached its higher value of 0.9 K under 8 kV/cm for the composition BTG6, corresponding to an electrocaloric responsivity ΔT/ΔE of 1.13 × 10⁻⁶ K.m/V. Such electrocaloric responsivity significantly exceeds those obtained so far in other barium titanate-based lead-free electrocaloric ceramic materials. Energy storage investigations show promising results: stored energy density of ~17 mJ/cm³ and an energy efficiency of ~88% in the composition BTG5. These results classify the studied materials as candidates for cooling devices and energy storage applications.     

Study on the Magnetocaloric Effect of Room Temperature Magnetic Refrigerant Material La0.5Pr0.5(Fe1−xCox)11.4Si1.6 and the Effect Arising from Co Doping on Its Curie Temperature

The arc-melting method was adopted to prepare the compound La_0.5Pr_0.5(Fe_1−xCo_x)_11.4Si_1.6 (x = 0, 0.02, 0.04, 0.06, 0.08), and the magnetocaloric effect of the compound was investigated. As indicated by the powder X-ray diffraction (XRD) results, after receiving 7-day high temperature annealing at 1373 K, all the compounds formed a single-phase cubic NaZn₁₃ crystal structure. As indicated by the magnetic measurement, the most significant magnetic entropy change |∆S_M(T)| of the sample decreased from 28.92 J/kg·K to 4.22 J/kg·K with the increase of the Co content under the 0–1.5 T magnetic field, while the Curie temperature T_C increased from 185 K to the room temperature 296 K, which indicated that this series of alloys are the room temperature magnetic refrigerant material with practical value. By using the ferromagnetic Curie temperature theory and analyzing the effect of Co doping on the exchange integral of these alloys, the mechanism that the Curie temperature of La_0.5Pr_0.5(Fe_1−xCo_x)_11.4Si_1.6 and La_0.8Ce_0.2(Fe_1−xCo_x)_11.4Si_1.6 increased with the increase in the Co content was reasonably explained. Accordingly, this paper can provide a theoretical reference for subsequent studies.     

Effects of the Addition of Co or Ni Atoms on Structure and Magnetism of FeB Amorphous Alloy: Ab Initio Molecular Dynamics Simulation

The effects of the substitution of Fe by Co or Ni on both the structure and the magnetic properties of FeB amorphous alloy were investigated using first-principle molecular dynamics. The pair distribution function, Voronoi polyhedra, and density of states of Fe_80−xTM_xB₂₀ (x = 0, 10, 20, 30, and 40 at.%, TM(Transition Metal): Co, Ni) amorphous alloys were calculated. The results show that with the increase in Co content, the saturation magnetization of Fe_80−xCo_xB₂₀ (x = 0, 10, 20, 30, and 40 at.%) amorphous alloys initially increases and then decreases upon reaching the maximum at x = 10 at.%, while for Fe_80−xNi_xB₂₀ (x = 0, 10, 20, 30, and 40 at.%), the saturation magnetization decreases monotonously with the increase in Ni content. Accordingly, for the two kinds of amorphous alloys, the obtained simulation results on the variation trends of the saturation magnetization with the change in alloy composition are in good agreement with the experimental observation. Furthermore, the relative maximum magnetic moment was recorded for Fe₇₀Co₁₀B₂₀ amorphous alloy, due to the induced increased magnetic moments of the Fe atoms surrounding the Co atom in the case of low Co dopant, as well as the increase in the exchange splitting energy caused by the enhancement of local atomic symmetry.     

Tailoring the Stability of Ti-Doped Sr2Fe1.4TixMo0.6−xO6−δ Electrode Materials for Solid Oxide Fuel Cells

In this work, the stability of Sr₂(FeMo)O_6−δ-type perovskites was tailored by the substitution of Mo with Ti. Redox stable Sr₂Fe_1.4Ti_xMo_0.6−xO_6−δ (x = 0.1, 0.2 and 0.3) perovskites were successfully obtained and evaluated as potential electrode materials for SOFCs. The crystal structure as a function of temperature, microstructure, redox stability, and thermal expansion properties in reducing and oxidizing atmospheres, oxygen content change, and transport properties in air and reducing conditions, as well as chemical stability and compatibility towards typical electrolytes have been systematically studied. All Sr₂Fe_1.4Ti_xMo_0.6−xO_6−δ compounds exhibit a regular crystal structure with Pm-3m space group, showing excellent stability in oxidizing and reducing conditions. The increase of Ti-doping content in materials increases the thermal expansion coefficient (TEC), oxygen content change, and electrical conductivity in air, while it decreases the conductivity in reducing condition. All three materials are stable and compatible with studied electrolytes. Interestingly, redox stable Sr₂Fe_1.4Ti_0.1Mo_0.5O_6−δ, possessing 1 μm grain size, low TEC (15.3 × 10⁻⁶ K⁻¹), large oxygen content change of 0.72 mol·mol⁻¹ between 30 and 900 °C, satisfactory conductivity of 4.1–7.3 S·cm⁻¹ in 5% H₂ at 600–800 °C, and good transport coefficients D and k, could be considered as a potential anode material for SOFCs, and are thus of great interest for further studies.     

Synthesis Method and Thermodynamic Characteristics of Anode Material Li3FeN2 for Application in Lithium-Ion Batteries

Li₃FeN₂ material was synthesized by the two-step solid-state method from Li₃N (adiabatic camera) and FeN₂ (tube furnace) powders. Phase investigation of Li₃N, FeN₂, and Li₃FeN₂ was carried out. The discharge capacity of Li₃FeN₂ is 343 mAh g⁻¹, which is about 44.7% of the theoretic capacity. The ternary nitride Li₃FeN₂ molar heat capacity is calculated using the formula C_p,m = 77.831 + 0.130 × T − 6289 × T⁻², (T is absolute temperature, temperature range is 298–900 K, pressure is constant). The thermodynamic characteristics of Li₃FeN₂ have the following values: entropy S⁰₂₉₈ = 116.2 J mol⁻¹ K⁻¹, molar enthalpy of dissolution Δ_dH_LFN = −206.537 ± 2.8 kJ mol⁻¹, the standard enthalpy of formation Δ_fH⁰ = −291.331 ± 5.7 kJ mol⁻¹, entropy S⁰₂₉₈ = 113.2 J mol⁻¹ K⁻¹ (Neumann–Kopp rule) and 116.2 J mol⁻¹ K⁻¹ (W. Herz rule), the standard Gibbs free energy of formation Δ_fG⁰₂₉₈ = −276.7 kJ mol⁻¹.     

Interplay between superconductivity and magnetism in Fe1?xPdxTe

The attractive/repulsive relationship between superconductivity and magnetic ordering has fascinated the condensed matter physics community for a century. In the early days, magnetic impurities doped into a superconductor were found to quickly suppress superconductivity. Later, a variety of systems, such as cuprates, heavy fermions, and Fe pnictides, showed superconductivity in a narrow region near the border to antiferromagnetism (AFM) as a function of pressure or doping. However, the coexistence of superconductivity and ferromagnetic (FM) or AFM ordering is found in a few compounds [RRh₄B₄ (R = Nd, Sm, Tm, Er), R′Mo₆X₈ (R′ = Tb, Dy, Er, Ho, and X = S, Se), UMGe (M = Ge, Rh, Co), CeCoIn₅, EuFe₂(As_1−xP_x)₂, etc.], providing evidence for their compatibility. Here, we present a third situation, where superconductivity coexists with FM and near the border of AFM in Fe_1−xPd_xTe. The doping of Pd for Fe gradually suppresses the first-order AFM ordering at temperature T_N/S, and turns into short-range AFM correlation with a characteristic peak in magnetic susceptibility at T′_N. Superconductivity sets in when T′_N reaches zero. However, there is a gigantic ferromagnetic dome imposed in the superconducting-AFM (short-range) cross-over regime. Such a system is ideal for studying the interplay between superconductivity and two types of magnetic (FM and AFM) interactions.Since the first discovery of superconductivity (SC) a century ago, the effects of magnetic impurities and the possibility of magnetic ordering in superconductors has been a central topic of condensed matter physics. Due to strong spin scattering (1, 2), it has generally been believed that the conduction electrons cannot be both magnetically ordered and superconducting. Even though it is thought that Cooper pairs in cuprates, heavy fermions, and Fe-based compounds are mediated by spin fluctuations (3–5), SC generally occurs after suppressing the magnetic ordering either through chemical doping or the application of hydrostatic pressure (6–10). However, there is growing evidence for the coexistence of superconductivity with either ferromagnetic (FM) (11–20) or antiferromagnetic (AFM) ordering (21–24). With the decrease of temperature (T), some of these systems show magnetic ordering before the superconducting transition (T_c) (14–17, 20), some are ordered in a reversed sequence (11–13, 18, 19, 22), some have the two orderings occur concomitantly (22, 25), and some show reentrant superconductivity (partially) overlapping with a magnetically ordered phase (11–13, 26). Despite extensive investigations of interaction between SC and magnetic moments, there is so far no unified theory for the coexistence of SC and magnetism. With the lack of theoretic guidance, the existing experimental findings lead to two schools of thought: one is that both orders result from the same conduction electrons as evidenced by their synchronized magnetic and superconducting orders (22), and the other is that there are two separate sets of electrons responsible for magnetic ordering and superconductivity, respectively (19, 21, 25). What remains incomprehensible is the case where superconductivity and magnetic ordering coexist but are in competition with each other, as seen in Fe-based systems (23, 24, 27).The discovery of superconductivity in Fe-based compounds has sparked enormous interest in the scientific community. Although Fe is the most well known ferromagnet, all parent compounds of Fe-based superconductors exhibit AFM ordering. Though superconductivity is induced after suppressing the AFM ordering, it can coexist with either remaining AFM ordering (23, 24, 27) or new FM ordering (18, 19), and this provides an ideal platform for studying the interplay between superconductivity and magnetism. Among the Fe-based superconductors, the chalcogenide FeTe_1−xSe_x is unique in several aspects: (i) it is the only compound composed of slabs of Fe(Te/Se)₄ stacked together without an interlayer spacer; (ii) it becomes superconducting via isovalent doping of Se for Te, with the highest T_c occurring at the 50% doping level; and (iii) the parent compound Fe_1+yTe shows nonmetallic electrical conduction and forms (π, π)-type AFM ordering with a large magnetic moment (28, 29), in contrast to the (π, 0)-type ordering with a small magnetic moment seen in Fe pnictides. The unusual AFM order in Fe_1+yTe cannot be explained by a simple Fermi-surface nesting picture, thus leading to arguments for a correlated local-moment scenario (27). Because of these differences, the application of hydrostatic pressure or partial doping on the Fe site, such as with Co or Ni, does not generate the generic phase diagram as seen for other Fe-based superconductors (30–33).To gain insight into the relationship between magnetism and superconductivity, we choose to substitute Pd for Fe in FeTe. This decision is motivated by the following: (i) Pd is known to exhibit FM instability (34, 35) even though it is paramagnetic in bulk, and (ii) PdTe is a superconductor with T_c ∼4.5 K (36–42). The partial replacement of Fe by Pd will allow us to understand the roles of Fe, Pd, and Te in both magnetism and superconductivity. Based on electrical transport, and magnetic and thermodynamic property measurements, we show that the ground state of Fe_1−xPd_xTe varies from AFM to FM ordering, to a superconducting state; this is one of rare cases where superconductivity flirts with both AFM and FM.Single crystals of Fe_1−xPd_xTe were grown via high-temperature melting, with the procedure described in SI Text. The crystal structure and the phase purity were measured by both powder and single-crystal X-ray diffraction. Fig. 1 shows powder X-ray diffraction patterns (Fig. 1A) and lattice parameters obtained from single-crystal X-ray refinement (Fig. 1B) for different doping levels; it confirms that the undoped FeTe forms a tetragonal structure, belonging to the P4/nmm space group. At room temperature, the lattice parameters are a = 3.8202 Å and c = 6.2686 Å. Similar to previous observations (43), our single-crystal refinement result indicates that there are ∼9% extra Fe atoms (T2 site of Fig. 1B) incorporated at interstitial sites of the Te layers. As soon as Pd is introduced into the system, the interstitial sites of Fe_1−xPd_xTe are occupied by Pd atoms (i.e., T2 = Pd when x > 0); though it remains in the same tetragonal space group, its lattice parameters a and c both increase with increasing Pd doping concentration for x ≤ x_s ∼0.6 (Fig. 1B). This finding indicates that Pd doping creates negative pressure, which is surprising because Pd²⁺ (∼0.80 Å) has a nearly identical ionic radius as that of Fe²⁺ (∼0.77 Å). The crystal structure of Fe_1−xPd_xTe remains tetragonal up to x_s (∼0.6). Above x_s, Fe_1−xPd_xTe crystallizes in a hexagonal structure (space group P6₃/mmc), and its lattice parameters also increase with increasing x (Fig. 1B); this results in an increase of the unit cell volume with increasing x in the entire doping range (Fig. 1B). In the hexagonal structure, the system no longer incorporates any interstitial sites (Tables S1–S4). Though the hexagonal structure at x > x_s can be regarded as a deformed FeTe structure, the local environment of Fe/Pd is transformed from a tetrahedron (x < x_s) to an octahedron (x > x_s) (44).Open in a separate windowFig. 1.Fe_1−xPd_xTe: (A) Powder X-ray diffraction patterns and (B) its unit cell parameters at room temperature obtained from single-crystal X-ray refinement.To correlate the structural information with physical properties, we show, in Fig. 2, the doping dependence of structural, electrical, and magnetic properties of Fe_1−xPd_xTe, with specifics presented later. The undoped (x = 0) Fe_1.09Te undergoes both an AFM ordering and a structural transition from paramagnetic (PM) tetragonal (high temperatures) to an AFM-ordered monoclinic (low temperatures) phase at T_N/S ∼70 K, with the first-order characteristic. Upon Pd doping, T_N/S is suppressed (solid circles), reaching zero at x = x_N/S ∼0.15 (dash line). In this doping region, the electrical resistivity (ρ) changes from nonmetallic (NM) character (dρ/dT < 0) above T_N/S to metallic (M) behavior (dρ/dT > 0) below T_N/S. When the characteristic feature for the first-order transition is absent at x_N/S, the magnetic susceptibility of Fe_1−xPd_xTe exhibits a peak at T^′_N. The value of T^′_N decreases with increasing x (Fig. 2, open circles), which eventually reaches zero at x_c ∼0.88. The magnetic susceptibility peak is a signature for short-range (SR) AFM correlation developed below T^′_N, with no obvious fingerprint in the electrical resistivity. Above x_c, superconductivity appears (Fig. 2, solid diamonds). The transition temperature T_c increases with increasing x. Remarkably, the system shows ferromagnetism (Fig. 2, solid squares) in the doping range of 0.78 ≤ x ≤ 0.98 with the maximum T_FM (∼190 K) at x_c, where AFM and SC vanish. In addition to the variable magnetic phases, there is doping-induced structural phase transition, being tetragonal at x < x_s ∼0.6 (Fig. 2, dark blue) and hexagonal at x > x_s (light blue). Such a structural transition is responsible for a change of electronic structure, manifested by NM behavior at x < x_s and a metallic character at x > x_s.Open in a separate windowFig. 2.Phase diagram of Fe_1−xPd_xTe. AFM, antiferromagnetic; AFM-M, antiferromagnetically ordered metallic state; FM, ferromagnetically ordered state; M, metallic; NM, nonmetal; PM, paramagnetic; PM-M, paramagnetic metal; SC, superconductivity; SR, short-range; SR-AFM, short-range AFM correlation.We now present the detailed experimental results in support of the phase diagram. Fig. 3A shows the temperature dependence of ρ of Fe_1−xPd_xTe for 0 ≤ x ≤ 0.5 (the experimental technique is described in SI Text). Similar to previous observations (45), the electrical resistivity of undoped (x = 0) Fe_1.09Te increases with decreasing temperature at high temperatures, but sharply decreases below T_N/S ∼70 K; this is caused by the first-order coupled structural and AFM phase transition (45, 46). Below T_N/S, the structure of Fe_1.09Te becomes monoclinic. Upon Pd doping, T_N/S is suppressed, and the first-order transition is no longer observed at x = 0.2. The smooth and monotonic temperature dependence of ρ indicates the absence of both structural and AFM transitions at 0.1 < x < 0.2. Though it sustains a nonmetallic temperature dependence (dρ/dT < 0) down to 2 K, the magnitude of the electrical resistivity deceases with increasing x in the entire doping range above T_N/S (Fig. 3 A and C). For 0.5 < x ≤ 1.0, the electrical resistivity of Fe_1−xPd_xTe continues to decrease with increasing x, as shown in Fig. 3B. More interesting is that ρ shows metallic behavior (dρ/dT > 0) in the entire temperature range measured for 0.5 < x ≤ 1.0. We believe that the cross-over from nonmetallic to metallic character is associated with the structural change at x_s. Furthermore, there is a sharp drop of resistivity at low temperatures for 0.88 < x ≤ 1.0, as shown in Fig. 3D. This drop is due to the emergence of superconductivity because the magnetic susceptibility shows diamagnetism as well in the temperature range (Fig. 3H). Note that the superconducting transition temperature T_c increases with increasing x.Open in a separate windowFig. 3.(A and B) Temperature dependence of the electrical resistivity (ρ) of Pd_1−xFe_xTe at the indicated compositions. (C and D) Temperature dependence of ρ near the first-order AFM/structural transition in the low-doping region, and the superconducting transition in the high-doping region plotted as ρ/ρ(5 K) vs. T, respectively. (E–G) Temperature dependence of the magnetic susceptibility (χ) of Fe_1−xPd_xTe measured at a magnetic field of 1 T. G Inset shows χ vs. x = 0.99 and 1.0; H is the χ(T) measured at 20 oersted (Oe) for 0.97 ≤ x ≤ 1.0.With Pd doping, the magnetic properties of Fe_1−xPd_xTe also change. Fig. 3 E–H shows the temperature dependence of the magnetic susceptibility, χ, at indicated values x. For 0 ≤ x < 0.2, χ initially increases with decreasing temperature, and then drops steeply at T_S/N (Fig. 3E); this confirms the first-order nature of the structural/AFM transition at 0 ≤ x < 0.2. Such a feature is absent for x ≥ x_N/S ∼0.15. Instead, there is a characteristic peak in magnetic susceptibility at T^′_N (Fig. 3F). Because there is no anomaly in electrical resistivity (Fig. 3 A–C) or specific heat, the susceptibility peak cannot be due to a true phase transition, but rather marks the onset of SR AFM correlation as seen in other materials (47). With increasing x, T^′_N moves to lower temperatures, and the magnitude of χ decreases as well (Fig. 3F). Though the characteristic peak is expected to completely vanish around x = x_c ∼0.88, the magnetic susceptibility reveals a dramatic increase below another characteristic temperature T_FM at x ≥ 0.78, as shown in Fig. 3G, which indicates the onset of ferromagnetic ordering at T_FM. With increasing x, T_FM initially increases then decreases after reaching the maximum (∼190 K) at x = x_c ∼0.88. Interestingly, the low-temperature susceptibility peak is still present for samples with x = 0.78, 0.80, 0.83, and 0.88, which suggests the coexistence of FM ordering and AFM interactions in this doping region. Above x_c, χ(T) continues to show FM behavior at high temperatures, but becomes negative at low temperatures (below T_c) under low magnetic fields, as depicted in Fig. 3H. The negative χ below T_c indicates the emergence of superconductivity, because their electrical resistivities also drop sharply (Fig. 3D). Whereas T_FM decreases, T_c increases with increasing x (above x_c). As shown Fig. 3G Inset, χ(T) exhibits paramagnetic behavior above T_c for x = 0.99 and 1.0, indicating the complete suppression of FM. However, T_c continues to increase with x (Fig. 3H).Clearly, there is a coexistence of AFM and FM ordering at 0.78 ≤ x ≤ 0.88, and of FM ordering and SC at 0.88 ≤ x ≤ 0.98 at low temperatures for Fe_1−xPd_xTe, which is further supported by the following results. First, the high-temperature magnetic susceptibility data allows us to extract both Curie–Weiss temperature (θ_CW) and effective magnetic moment (μ_eff) by fitting our experimental data to the Curie–Weiss law , where N_A is the Avogadro’s constant and k_B is Boltzmann’s constant. The x dependences of θ_CW and μ_eff are shown in Fig. 4 A and B, respectively. Note that θ_CW < 0 for x ≤ 0.8, indicating that the dominant magnetic interaction is AFM in this doping region. The fact that | θ_CW | >> T_N/S implies 2D AFM ordering for x < 0.2, whereas a small θ_CW value at 0.2 ≤ x ≤ 0.8 is consistent with our interpretation that there is short-range AFM correlation. And θ_CW > 0 for 0.8 ≤ x < 0.98, which confirms the ferromagnetic interaction in this doping range. Although θ_CW is comparable to the corresponding transition temperature T_FM, the effective magnetic moment is small, as shown in Fig. 4B, which implies that the FM ordering results mainly from Fe, even though Pd has to be the vehicle of coupling between Fe atoms (see discussion below). According to the single-crystal X-ray refinement results (Tables S1–S3), the actual Fe concentration is very close to the nominal value for 0.8 ≤ x ≤ 0.98. We thus estimate the effective magnetic moment of Fe using the nominal Fe concentrations; this gives μ_eff = 1.60μ_B/Fe (x = 0.80), 2.42μ_B/Fe (x = 0.83), 3.79μ_B/Fe (x = 0.88), 4.21μ_B/Fe (x = 0.92), 3.53μ_B/Fe (x = 0.95), and 2.52μ_B/Fe (x = 0.98). Note that the highest magnetic moment for x = 0.92 is close to the Fe moment in Fe_1.09Te (∼ 4.65μ_B/Fe), as plotted in Fig. 4B. In the latter case, the actual ordered magnetic moment is ∼3.31μ_B/Fe (28), smaller than that obtained from the Curie–Weiss constant; this is explained as due to the itinerant nature of the electrons in FeTe (28). For Fe_1−xPd_xTe, we may estimate the ordered magnetic moment from the magnetization. Fig. 4C shows the field dependence (H) of magnetization (M) at T = 3 K for x between 0.80 and 0.98. Note that M(H) reveals a well-defined hysteresis loop, confirming the nature of FM ordering for 0.8 ≤ x ≤ 0.98. Though M(H) is not saturated up to 6 T, we may estimate the lower bound of the ordered magnetic moment using M(H = 6 T) data shown in Fig. 4B. At T = 3 K, μ ∼0.33μ_B/Fe (x = 0.80), 0.48μ_B/Fe (x = 0.83), 1.49μ_B/Fe (x = 0.88), 1.85μ_B/Fe (x = 0.92), 1.67μ_B/Fe (x = 0.95), and 0.22μ_B/Fe (x = 0.98). In addition to the fact that M (H = 6 T) < M_sat, the small ordered magnetic moment could result from (i) the itinerant nature of the electrons and (ii) the reduced concentration of Fe in the highly Pd doped region. Nevertheless, the specific heat reveals anomalies in the compounds with high T_FM (Fig. 4D), indicating a true phase transition at T_FM.Open in a separate windowFig. 4.Fe1-xPdxTe (A) Doping dependence of Curie–Weiss temperature. (B) Doping dependence of the effective magnetic moment and magnetization at 6 T. (C) Magnetization vs. magnetic field at 3 K for x = 0.8, 0.88, 0.92, 0.95, and 0.98. (D) Specific heat vs. temperature for x = 0.92 and 0.88.From the results presented here, it is most likely that the FM dome is intrinsic, due to the ordering of the Fe magnetic moments. Within the dome, our single-crystal X-ray refinement indicates that the structure remains the same as that of pure PdTe with actual Fe concentration close to the nominal value, and there are no interstitial sites of either Pd or Fe, which strongly suggests that the FM ordering is not due to the formation of Fe clusters. In early studies, FM was detected in alloys of Fe in Pd (i.e., Pd_1−xFe_x) at compositions down to 0.08% Fe (48), because magnetic moments on the Fe sites polarize the surrounding Pd matrix (48–51). It seems that Fe doping in PdTe compound gives rise to a similar effect, i.e., Pd matrix mediates the magnetic interactions between Fe atoms.The central question is, how can both FM ordering and superconductivity coexist in the region of 0.88 ≤ x ≤ 0.98? Having excluded the possibility of structural phase separation, one may consider the scenario of electronic phase separation with either macroscopic coexistence of singlet SC and FM in different regions or microscopic coexistence of triplet SC and FM (11, 21, 52). If Fe_1−xPd_xTe were a triplet superconductor, a small amount of impurity or disorder would completely suppress superconductivity, as seen in Sr₂RuO₄ (53). Though we observe anomalous specific heat below T_c of PdTe (42), the insensitivity to Fe doping does not support triplet SC. For the former case, further experimental investigations are necessary, such as combined scanning electron microscopy and tunneling microscopy to directly probe both SC and FM regions. Theoretically, almost all Pd d orbitals in PdTe are occupied due to the covalency of the Pd–Te bonding, according to recent first-principle calculations (44), and this leads to a strong suppression of the local magnetic moment (44). As shown in Fig. 3G (Inset), the magnetic susceptibility of PdTe (x = 1) indeed exhibits paramagnetic behavior above T_c. Upon substitution of Pd by Fe, T_c decreases nearly linearly with increasing Fe concentration (1 − x; Fig. 2). As Fe concentration reaches ∼3% (1 − x = 0.03), both SC and FM ordering coexist. The fact that FM ordering does not immediately kill superconductivity is likely due to the unique electronic structure of PdTe. First-principles calculations indicate that the states in the proximity of the Fermi level consist mainly of Te p electrons, which are weakly hybridized with Pd e_g orbitals (44); this is in dramatic contrast to the electronic structure of FeTe, in which the states near the Fermi level derived from Fe with direct Fe–Fe interactions, and the Te p states lie well below the Fermi level and hybridized weakly with the Fe d states (54). For PdTe, the partial doping of Fe may lead to an even weaker hybridization between Te p and Pd e_g orbitals, due to the shift of Fermi level. Our experimental observation is in support of this scenario: with increasing Fe concentration, superconductivity is gradually suppressed to zero, at which the FM ordering reaches the maximum (Figs. 2 and ​and4).4). After T_c reaches zero at x_c, the disappearance of the superconducting energy gap allows the rearrangement of electronic structure, which is in favor of antiferromagnetic interaction; the latter apparently competes with the existing FM interaction, thus leading to the complete suppression of FM at x < 0.78.In summary, the substitution of Fe by Pd results in an extremely rich phase diagram for Fe_1−xPd_xTe. Powder and single-crystal X-ray diffraction measurements both indicate that there is doping-induced structural transition from a tetragonal phase at x < x_s ∼0.6 to a hexagonal phase at x > x_s. Correspondingly, the electrical resistivity changes from nonmetallic character at x < x_s to metallic behavior at x > x_s. Magnetically, the system undergoes a first-order AFM transition at T_N/S for x ≤ x_N/S ∼0.15, with the structure transition occurring at the same temperature from tetragonal to monoclinic. When T_N/S approaches zero at x_N/S, there is no longer a temperature-induced structural transition, whereas short-range AFM correlation exists for x_N/S < x ≤ 0.88. Superconductivity sets in at x ≥ x_c ∼0.88. In addition, a gigantic FM dome with 0.78 ≤ x ≤ 0.98 is centered at the critical concentration x_c, where both AFM and SC vanish. The coexistence of FM ordering with either SR AFM or SC results most likely from the unique electronic structure of Fe_1−xPd_xTe. With the low Pd concentration, both electrical conduction and magnetism of Fe_1−xPd_xTe are determined by Fe/Pd d electrons, because they are near the Fermi level. Due to the ferromagnetic instability and the extended orbitals of 4d electrons of Pd, both AFM interactions and electrical resistivity in Fe_1−xPd_xTe decrease with increasing x. Before the complete suppression of SR AFM, FM ordering occurs due to the Fe magnetic moment polarization in the Pd matrix. However, the hexagonal structure at x > x_s with Pd/Fe in an octahedral environment leads to more localized d electrons from Pd/Fe but itinerant p electrons from Te. It seems that the weak hybridization between p–d electrons allows for the magnetic ordering of Fe/Pd moments and superconducting ordering of p conduction electrons. Nevertheless, these two orderings compete with each other. Hence, the physics of Fe_1−xPd_xTe in the hexagonal phase is in contrast to known Fe-based superconductors, in which physical properties are mainly determined by Fe d states. Fe_1−xPd_xTe provides another and rare example for studying the interplay between SC, FM, and AFM ordering, where two separate sets of electrons are responsible for FM ordering and superconductivity, respectively.     

Theoretical Prediction and Experimental Validation of the Glass-Forming Ability and Magnetic Properties of Fe-Si-B Metallic Glasses from Atomic Structures

Developing new soft magnetic amorphous alloys with a low cost and high saturation magnetization (B_s) in a simple alloy system has attracted substantial attention for industrialization and commercialization. Herein, the glass-forming ability (GFA), thermodynamic properties, soft magnetic properties, and atomic structures of Fe_80+xSi_5−xB₁₅ (x = 0–4) amorphous soft magnetic alloys were investigated by ab initio molecular dynamics (AIMD) simulations and experiments. The pair distribution function (PDF), Voronoi polyhedron (VP), coordination number (CN), and chemical short- range order (CSRO) were analyzed based on the AIMD simulations for elucidating the correlations between the atomic structures with the glass-forming ability and magnetic properties. For the studied compositions, the Fe₈₂Si₃B₁₅ amorphous alloy was found to exhibit the strongest solute–solute avoidance effect, the longest Fe-Fe bond, a relatively high partial CN for the Fe-Fe pair, and the most pronounced tendency to form more stable clusters. The simulation results indicated that Fe₈₂Si₃B₁₅ was the optimum composition balancing the saturation magnetization and the GFA. This prediction was confirmed by experimental observations. The presented work provides a reference for synthesizing new Fe-Si-B magnetic amorphous alloys.     

Reinforcing Increase of ΔTc in MgB2 Smart Meta-Superconductors by Adjusting the Concentration of Inhomogeneous Phases

Incorporating with inhomogeneous phases with high electroluminescence (EL) intensity to prepare smart meta-superconductors (SMSCs) is an effective method for increasing the superconducting transition temperature (T_c) and has been confirmed in both MgB₂ and Bi(Pb)SrCaCuO systems. However, the increase of ΔT_c (ΔT_c = T_c ‒ T_cpure) has been quite small because of the low optimal concentrations of inhomogeneous phases. In this work, three kinds of MgB₂ raw materials, namely, ^aMgB₂, ^bMgB₂, and ^cMgB₂, were prepared with particle sizes decreasing in order. Inhomogeneous phases, Y₂O₃:Eu³⁺ and Y₂O₃:Eu³⁺/Ag, were also prepared and doped into MgB₂ to study the influence of doping concentration on the ΔT_c of MgB₂ with different particle sizes. Results show that reducing the MgB₂ particle size increases the optimal doping concentration of inhomogeneous phases, thereby increasing ΔT_c. The optimal doping concentrations for ^aMgB₂, ^bMgB₂, and ^cMgB₂ are 0.5%, 0.8%, and 1.2%, respectively. The corresponding ΔT_c values are 0.4, 0.9, and 1.2 K, respectively. This work open a new approach to reinforcing increase of ΔT_c in MgB₂ SMSCs.     

Microstructure and Piezoelectric Properties of Lead Zirconate Titanate Nanocomposites Reinforced with In-Situ Formed ZrO2 Nanoparticles

Lead zirconate titanate (PZT)-based ceramics are used in numerous advanced applications, including sensors, displays, actuators, resonators, chips; however, the poor mechanical characteristics of these materials severely limits their utility in composite materials. To address this issue, we herein fabricate transgranular type PZT ceramic nanocomposites by a novel method. Thermodynamically metastable single perovskite-type Pb_0.99(Zr_0.52+xTi_0.48)_0.98Nb_0.02O_3+1.96x powders are prepared from a citrate precursor before both monoclinic and tetragonal ZrO₂ nanoparticles ranging from 20 to 80 nm are precipitated in situ at a sintering temperature of 1260 °C. The effects of ZrO₂ content on the microstructure, dielectric, and piezoelectric properties are investigated and the mechanism, by which ZrO₂ toughened PZT is analyzed in detail. The ZrO₂ nanoparticles underwent a tetragonal to monoclinic phase transition upon cooling. The fracture mode changed from intergranular to transgranular with increasing ZrO₂ content. The incorporation of ZrO₂ nanoparticles improved the mechanical and piezoelectric properties. The optimized piezoelectric properties (ε^T₃₃/ε₀ = 1398, tan δ = 0.024 d₃₃ = 354 pC N⁻¹, k_p = 0.66 Q_m = 78) are obtained when x = 0.02. T_c initially increased and subsequently decreased with increasing ZrO₂ content. The highest T_c = (387 °C) and lowest ε^T₃₃/ε₀ was obtained at x = 0.01.     

On the attribution and additivity of binding energies

It can be useful to describe the Gibbs free energy changes for the binding to a protein of a molecule, A—B, and of its component parts, A and B, in terms of the “intrinsic binding energies” of A and B, ΔG_Aⁱ and ΔG_Bⁱ, and a “connection Gibbs energy,” ΔG^s that is derived largely from changes in translational and rotational entropy. This empirical approach avoids the difficult or insoluble problem of interpreting observed ΔH and TΔS values for aqueous solutions. The ΔGⁱ and ΔG^s terms can be large for binding to enzymes and other proteins.     

Linking the Electrical Conductivity and Non-Stoichiometry of Thin Film Ce1−xZrxO2−δ by a Resonant Nanobalance Approach

Bulk ceria-zirconia solid solutions (Ce_1−xZr_xO_2−δ, CZO) are highly suited for application as oxygen storage materials in automotive three-way catalytic converters (TWC) due to the high levels of achievable oxygen non-stoichiometry δ. In thin film CZO, the oxygen storage properties are expected to be further enhanced. The present study addresses this aspect. CZO thin films with 0 ≤ x ≤ 1 were investigated. A unique nano-thermogravimetric method for thin films that is based on the resonant nanobalance approach for high-temperature characterization of oxygen non-stoichiometry in CZO was implemented. The high-temperature electrical conductivity and the non-stoichiometry δ of CZO were measured under oxygen partial pressures pO₂ in the range of 10⁻²⁴–0.2 bar. Markedly enhanced reducibility and electronic conductivity of CeO₂-ZrO₂ as compared to CeO_2−δ and ZrO₂ were observed. A comparison of temperature- and pO₂-dependences of the non-stoichiometry of thin films with literature data for bulk Ce_1−xZr_xO_2−δ shows enhanced reducibility in the former. The maximum conductivity was found for Ce_0.8Zr_0.2O_2−δ, whereas Ce_0.5Zr_0.5O_2-δ showed the highest non-stoichiometry, yielding δ = 0.16 at 900 °C and pO₂ of 10⁻¹⁴ bar. The defect interactions in Ce_1−xZr_xO_2−δ are analyzed in the framework of defect models for ceria and zirconia.     

Influence of Dissolving Fe–Nb–B–Dy Alloys in Zirconium on Phase Structure,Microstructure and Magnetic Properties

This paper refers to the structural and magnetic properties of [(Fe₈₀Nb₆B₁₄)_0.88Dy_0.12]_1−xZr_x (x = 0; 0.01; 0.02; 0.05; 0.1; 0.2; 0.3; 0.5) alloys obtained by the vacuum mold suction casting method. The analysis of the phase contribution indicated a change in the compositions of the alloys. For x < 0.05, occurrence of the dominant Dy₂Fe₁₄B phase was observed, while a further increase in the Zr content led to the increasing contribution of the Fe–Zr compounds and, simultaneously, separation of crystalline Dy. The dilution of (Fe₈₀Nb₆B₁₄)_0.88Dy_0.12 in Zr strongly influenced the magnetization processes of the examined alloys. Generally, with the increasing x parameter, we observed a decrease in coercivity; however, the unexpected increase in magnetic saturation and remanence for x = 0.2 and x = 0.3 was shown and discussed.     

Structural,Magnetic, Dielectric and Piezoelectric Properties of Multiferroic PbTi1−xFexO3−δ Ceramics

PbTi_1−xFe_xO_3−δ (x = 0, 0.3, 0.5, and 0.7) ceramics were prepared using the classical solid-state reaction method. The investigated system presented properties that were derived from composition, microstructure, and oxygen deficiency. The phase investigations indicated that all of the samples were well crystallized, and the formation of a cubic structure with small traces of impurities was promoted, in addition to a tetragonal structure, as Fe³⁺ concentration increased. The scanning electron microscopy (SEM) images for PbTi_1−xFe_xO_3−δ ceramics revealed microstructures that were inhomogeneous with an intergranular porosity. The dielectric permittivity increased systematically with Fe³⁺ concentration, increasing up to x = 0.7. A complex impedance analysis revealed the presence of multiple semicircles in the spectra, demonstrating a local electrical inhomogeneity due the different microstructures and amounts of oxygen vacancies distributed within the sample. The increase of the substitution with Fe³⁺ ions onto Ti⁴⁺ sites led to the improvement of the magnetic properties due to the gradual increase in the interactions between Fe³⁺ ions, which were mediated by the presence of oxygen vacancies. The PbTi_1−xFe_xO_3−δ became a multifunctional system with reasonable dielectric, piezoelectric, and magnetic characteristics, making it suitable for application in magnetoelectric devices.     

Study of Bulk Amorphous and Nanocrystalline Alloys Fabricated by High-Sphericity Fe84Si7B5C2Cr2 Amorphous Powders at Different Spark-Plasma-Sintering Temperatures

The new generation of high-frequency and high-efficiency motors has high demands on the soft magnetic properties, mechanical properties and corrosion resistance of its core materials. Bulk amorphous and nanocrystalline alloys not only meet its performance requirements but also conform to the current technical concept of integrated forming. At present, spark plasma sintering (SPS) is expected to break through the cooling-capacity limitation of traditional casting technology with high possibility to fabricate bulk metallic glasses (BMGs). In this study, Fe₈₄Si₇B₅C₂Cr₂ soft magnetic amorphous powders with high sphericity were prepared by a new atomization technology, and its characteristic temperature was measured by DSC to determine the SPS temperature. The SEM, XRD, VSM and universal testing machine were used to analyze the compacts at different sintering temperatures. The results show that the powders cannot be consolidated by cold pressing (50 and 500 MPa) or SPS temperature below 753 K (glass transition temperature T_g = 767 K), and the tap density is only 4.46 g·cm⁻³. When SPS temperature reached above 773 K, however, the compact could be prepared smoothly, and the density, saturation magnetization, coercivity and compressive strength of the compacts increased with the elevated sintering temperature. In addition, due to superheating, crystallization occurred even when the sintering temperature was lower than 829 K (with the first crystallization onset temperature being T_x1 = 829 K). The compact was almost completely crystallized at 813 K, resulting in a sharp increase in the coercivity of the compact from 55.55 A·m⁻¹ at 793 K to 443.17 A·m⁻¹. It is noted that the nanocrystals kept growing in size as the temperature increased to 833 K, which increased the coercivity remarkably but showed an enhanced saturation magnetization.     

Large Low-Field Reversible Magnetocaloric Effect in Itinerant-Electron Hf1−xTaxFe2 Alloys

First-order isostructural magnetoelastic transition with large magnetization difference and controllable thermal hysteresis are highly desirable in the development of high-performance magnetocaloric materials used for energy-efficient and environmental-friendly magnetic refrigeration. Here, we demonstrate large magnetocaloric effect covering the temperature range from 325 K to 245 K in Laves phase Hf_1−xTa_xFe₂ (x = 0.13, 0.14, 0.15, 0.16) alloys undergoing the magnetoelastic transition from antiferromagnetic (AFM) state to ferromagnetic (FM) state on decreasing the temperature. It is shown that with the increase of Ta content, the nature of AFM to FM transition is gradually changed from second-order to first-order. Based on the direct measurements, large reversible adiabatic temperature change (ΔT_ad) values of 2.7 K and 3.4 K have been achieved under a low magnetic field change of 1.5 T in the Hf_0.85Ta_0.15Fe₂ and Hf_0.84Ta_0.16Fe₂ alloys with the first-order magnetoelastic transition, respectively. Such remarkable magnetocaloric response is attributed to the rather low thermal hysteresis upon the transition as these two alloys are close to intermediate composition point of second-order transition converting to first-order transition.