1 Introduction

After discovery of superconductivity in the YBa2Cu3O7−δ (YBCO) system above liquid nitrogen temperature (77 K) [1], scientists have spent much effort to increase the superconducting transition temperature, T c, and critical current, J c. For this purpose, different preparation techniques, dopings/substitutions and additions to the system were performed.

Three different phases are presented in the YBCO system: Y-123, Y-124 and Y-247 with the T c values at 90, 80 and 80 K, respectively [2–5]. Y-123 system has a special importance due to high current carrying capacity at 77 K and high critical magnetic field, H c [6, 7]. J c of the system strongly depends on the Cu-chains which control the carrier density at the CuO planes [8] and so the T c value [9, 10]. It is expected that the increasing CuO planes in unit cell can cause an increase in the T c value of the system. However, the highest T c value was obtained only in Y-123.

The superconducting properties depend strongly on stoichiometry of the HT c YBCO system. Small changes in the stoichiometry, especially deficiency of Ba and Cu, can significantly influence the properties of YBCO. Yoshikate et al. investigated the Cu deficiency in the YBCO system. They found that the crystal structure changed significantly even with small change in the copper content of system and thus the superconducting properties was deteriorated [11]. Vasillev et al. [12] found that the change in the Ba and Cu contents of the system caused the change of the oxygen content and so the superconducting properties. Hattori et al. [13] studied the correlation between superconductivity and the Ba and Cu content. They found that the change of the Cu content from 3.26 to 2.82 with a small difference in the Ba content can cause the increase on room temperature resistivity and lead to the suppressed T c value. Schneider et al. investigated the effect of Ba/Y ratio on T c. They reported that the Ba/Y ratio between 0.5 and 2.4 did not have important effect on the T c value but T c decreased above 2.4. The highest J c was obtained when the ratio was 1.8 [14].

There are limited studies on the Se doped/substituted YBCO system [15–17]. In the studies, it was found that the Se substitution decreased drastically T c and changed the unit cell of YBa2Cu3O7 from orthorhombic to tetragonal. The c-axis decreased compared to undoped YBCO system.

In this paper, the samples with stoichiometry of Y2/3 Se1/3Ba2Cu3Ox were fabricated by conventional solid-state reaction technique. Structural/microstructural analyses of the samples were investigated by XRD, SEM-EDX. Electrical property of the samples was examined by R–T measurements. Magnetic properties, the critical current density and pinning properties of the samples were analyzed in the present paper.

2 Experimental details

The samples were prepared using solid-state reaction technique. High purity powders of Y2O3, SeO2, BaCO3 and CuO were weighted in the appropriate amounts to give nominal composition of Y2/3 Se1/3Ba2Cu3Ox. The powders were mixed using an agate mortar and then calcined at 900 °C for 24 h with intermediate grinding and mixing. After calcination, the powders were pressed into pellets at 4 tons. The pellets were sintered at temperatures between 900 and 980 °C for 24 h under oxygen atmosphere. It was seen that the samples heat treated at 980 °C were melted.

The structural characterization of the samples was performed by X-ray diffraction (XRD). Scan speed was selected as 2° min−1 in the range of 2θ = 20°–80°. Automated Rigaku RadB Dmax X-ray diffractometer with CuKα radiation was used for the XRD analysis.

The microstructural and compositional characterization of the materials were carried out with Leo EVO-40 XVP scanning electron microscope (SEM) and BRUKER X-flash detector 4010 energy dispersive x-ray spectroscope (EDX).

Temperature dependence of resistivity (ρ−T) of the samples fabricated was investigated by closed cycle Leybold LT-10 cryostat system combined with SRS ac bridge system using 17.7 Hz frequency. For resistivity measurements, the four probe electrical contacts were made by silver paste. The T c values of samples fabricated were determined by the resistivity differentiation dρ(T)/dT. The peak temperature in dρ(T)/dT plot was defined as the T c value. M–T and M–H measurements were performed using Quantum Design PPMS system at 10, 20, 30, 40 and 50 K up to 9 T. Critical current density, J c, was calculated with Bean model for all samples at 10–50 K up to 9 T. Pinning force, F p, depending on temperature and applied magnetic field was calculated using M–H data.

3 Results and discussion

3.1 XRD results

The XRD patterns of Y2/3 Se1/3Ba2Cu3Ox, depending on the heat treatment conditions, are given in Fig. 1. The samples heat treated at 900–960 °C for 24 h under O2 atmosphere consisted mainly of YBa2Cu3O7. However, impurity peaks such as BaSeO4 and CuSe2O5 were also detected, as seen in Fig. 1. The crystal symmetry of the samples was determined as orthorhombic. The unit cell parameters of the samples are listed in Table 1. The results indicated that the calculated unit cell parameters are close to the pure Y-123 system with a small difference [18–20]. We believe that the reason of the difference on the unit cell parameters can be due to occupation of the Se ions to interstitial sites. The interstitial occupation causes the change of the unit cell parameters.

Fig. 1
figure 1

XRD pattern of Y1Se0.33Ba1.67Cu2.67Ox compounds depending on heat treatment conditions

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Table 1 Crystal parameters of Y1Se0.33Ba1.67Cu2.67Ox samples depending on heat treatment conditions
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3.2 Micro-structural properties

Surface morphology of the samples fabricated at 900–960 °C for 24 h under O2 atmosphere is given in Fig. 2a–d. The main matrix of the sample fabricated at 900 °C for 24 h was identified to be Y-123 from the EDX analysis. Some dark regions were observed on the surface of the sample. EDX analysis showed that atomic composition of these dark regions corresponded to BaSeO4 and CuSe2O5, respectively. Different surface morphology was obtained with increasing the heat treatment temperature, Fig. 2b–d. The partial melted grains on the surface were observed on the samples. Especially, size and number of the partially melted grains increased with the heat treatment temperature. The samples were melted for the heat treatment temperatures above 960 °C.

Fig. 2
figure 2

Surface morphology of samples heat treated at (a) 900 °C, (b) 920 °C, (c) 940 °C and (d) 960 °C for 24 h under oxygen atmosphere

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3.3 Electrical properties

Temperature dependence of the electrical resistivity of the samples depending on the heat treatment conditions is presented in Fig. 3. All the samples showed metallic behavior above T c and the resistivity dropped to zero below T c. The T c and T 0 values of the samples heat treated at 900–960 °C are given in Table 2. The highest T c and T 0 value were obtained to be 94.9 and 90.3 K for the sample heat treated at 920 °C for 24 h. It was seen that normal state resistivity values in the samples heat treated above 920 °C increased compared to the samples fabricated at 900 and 920 °C. It is believed that this increase on the normal state resistivity value is attributed to deterioration of microstructure of the samples with the Se addition and/or with increasing the heat treatment temperature. Particularly, the occupation of the interstitial sites by the Se atoms can be caused the increase of the normal state resistivity of the samples, compared to the undoped YBCO material.

Fig. 3
figure 3

Temperature dependence of electrical resistivity of the samples fabricated at different heat treatment conditions

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Table 2 T c values of samples obtained from electrical and magnetic measurements depending on heat treatment cycles
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3.4 Magnetic properties

The temperature dependent magnetization was measured from 50 to 100 K under zero magnetic field. The temperature dependence of magnetization of the samples fabricated (M–T) is given in Fig. 4. The highest diamagnetic signal with sharp drop was obtained for the sample fabricated at 920 °C, suggesting strong grain connection in the sample. The onset temperature, \(T_{c}^{on}\), of diamagnetic signal of the samples is presented in Table 2.

Fig. 4
figure 4

Temperature dependence of DC-magnetization of the samples fabricated at different heat treatment conditions

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It is known that magnetic behavior of the superconducting materials is associated strongly to previous history of the applied magnetic field. When the applied magnetic field is increased from zero to a specific value and then returned through zero, the hysteresis behavior can be observed. The reason of this irreversible manner of magnetization can be due to flux trapping in superconducting samples, which has a crucial importance on critical current density.

Figure 5 shows the M–H curves of the samples at five different temperatures (10, 20, 30, 40, 50 K) up to 9 T. In general, the non-symmetric hysteresis was obtained for all the samples. The M–H curves decreased rapidly up to ~2 T and then varied smoothly with the magnetic field. The magnetization values for all the samples remained non-zero at 9 T for all the measurement temperatures. The highest zero-field magnetization value was obtained at 10 K for the sample fabricated at 920 °C for 24 h, as seen in Fig. 5b. But, the zero-field magnetization value at 10 K decreased for other samples. We believe that this is related to inter-grain and/or intra-grain magnetic behavior in the samples. Magnetization decreased with increasing the temperature and the applied magnetic field. At low temperatures, the pinning is strong and the flux mostly penetrates a material as the vortices. Magnetization caused by the pinning becomes smaller as temperature increases. The vortex motion in the material starts with increase in the temperature, which leads to deterioration in the magnetic configuration of fluxes [21, 22].

Fig. 5
figure 5

Magnetic hysteresis loops depending on temperature of the samples heat treated at (a) 900 °C, (b) 920 °C, (c) 940 °C and (d) 960 °C for 24 h under oxygen atmosphere

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To analyze M–H loops, we subtracted the linear paramagnetic contribution which is different for each samples, such that the M(H) dependencies have become almost independent of H at high fields. In contrast to [23, 24], the M(H) curves at different temperatures cannot be scaled for all the investigated sample. However, the scaling is observed for the M(H) curves of the different samples at the same temperature. The M(H) curves of the samples coincide then data are re-plotted at reduced units, M/M m versus H/H m, where H m is the field at which the virgin magnetization curve reaches its minimum M m. Figures 6 and 7 demonstrate the reduced magnetization M/M m versus the reduced magnetic field H/H m at 10 and 50 K for samples treated at 900, 920, 940, 960 °C.

Fig. 6
figure 6

Scaling of magnetization at T = 10 K. Reduced magnetization M / M m versus reduced magnetic field H/H m

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Fig. 7
figure 7

Scaling of magnetization at T = 50 K. Reduced magnetization M / M m versus reduced magnetic field H/H m

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It should be noted that the scaling cannot be reproduced to a parameter. Different scaling along M and H axes may be due to the inter-granular field compression [25] providing that the magnetic field near grains is larger than the external field.

The extended critical state model [26] combining the Valkov–Khrustalev approach [27] with the critical state model [28] was used to fit the M–H loops. The critical current density, J c, and the depth of the equilibrium surface layer l s, which are the main fitting parameters, determine the width and the asymmetry of magnetization loop. The fitting parameters at H = 0 are presented in Table 3, the depth l s is compared with the average grain size, D.

Table 3 Scaling coefficients H m, M m and fitting parameters J c, 2l s/D for M(H) curves at T = 10 and 50 K
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The critical current density obtained from M–H loops, \(J_{c}^{mag}\), was calculated using Bean’s formula [29] given by

$$J_{c}^{mag} = 20\frac{\Updelta M}{{a\left( {1 - a/3b} \right)}} ,$$
(1)

where ΔM = M + − M − is the width of magnetization hysteresis, a and b (a < b) are the dimensions of the samples. The applied magnetic field dependence of \(J_{c}^{mag}\), (\(J_{c}^{mag} - H\)), at 10, 20, 30, 40 and 50 K for the samples heat treated at 900, 920, 940 and 960 °C is presented in Fig. 8. The maximum \(J_{c}^{mag}\) value was calculated to be 9.76 × 105 A cm−2 at 10 K for magnetic field of 0 T for the sample fabricated at 920 °C for 24 h. The \(J_{c}^{mag}\) values dropped with increasing the applied magnetic field and also temperature as seen in Fig. 8a–d. \(J_{c}^{mag}\) for all the samples showed a plateau after the magnetic field of 1 T. We believe that intrinsic pinning of flux can be responsible for the plateau in the \(J_{c}^{mag} - H\) plots. The \(J_{c}^{mag}\) values of Y1Se0.33Ba1.67Cu2.67Ox samples are higher than that of undoped Y-123 system [30]. This result indicated that the Se adding and also Ba and Cu deficiencies in YBCO act as the pinning centers and increased the \(J_{c}^{mag}\) value of the YBCO system.

Fig. 8
figure 8

J c–H dependence’s at 10, 20, 30, 40 and 50 K for the samples fabricated (a) 900 °C, (b) 920 °C, (c) 940 °C and (d) 960 °C for 24 h under oxygen atmosphere

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When a current density, J c, is passed through a superconductor below T c, an electromagnetic interaction force emerges between fluxoids and the carriers which are responsible for the current. The Lorentz force, F L , [31] is

$$F_{L} = J \times B,$$
(2)

where J c is the current density and B the flux density in the sample. Fluxoids remain stationary and the superconducting behavior of the material maintains with zero resistance as long as the pinning force, F p, equals at least to F L . When F L acting the floxids exceeds F p, the floxids begin to move across the sample and the material becomes resistive. In that case, the current is no longer without loss. Figure 9 shows the magnetic field dependence of bulk pinning force, F p, of the samples at five different temperatures. It was found that F p decreased with increasing the temperature. The maximum value of F p was calculated to be 5.7 × 104 N/cm3 at 10 K for the sample heat treated at 900 °C for 24 h, as seen in Fig. 9a. F p with the applied magnetic field showed almost similar behavior for all the samples. Any peak in F p was not observed at temperatures of 20 and 50 K and F p increased almost linear up to 9 T. The maximum values of F p at 10 K are found to be 5.1 × 104 N/cm3, 7.1 × 103 N/cm3 and 1.75 × 103 N/cm3 for the sample fabricated at 920, 940 and 960 °C, respectively.

Fig. 9
figure 9

F p–H dependence’s at 10, 20, 30, 40 and 50 K for the samples fabricated (a) 900 °C, (b) 920 °C, (c) 940 °C and (d) 960 °C for 24 h under oxygen atmosphere

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To understand the pinning mechanism in the samples, the behavior of flux line lattice (FLL) and pinning centers in the sample should be considered. If the interaction force among flux lines is higher than F p, FLL is elastically stiff and the correlated flux grain size is large, which means that the bulk F p is small. On the other hand, if F p increases relative to the inter-flux line forces, the correlated grains become smaller and approach a limited grain size equal to average pinning center spacing, so that each flux lattice grain interacts with only one pinning center. In this state, the bulk F p has a maximum value, because each pinning center applies a maximal constraint on the FLL [32].

When the number of flux lines is equal to the number of pinning centers which penetrates to the sample at the applied magnetic field, the maximum F p value is reached. Therefore, the applied field value at the maximum F p depends strongly on the number of pinning centers penetrating to the samples and the peak value in F p related to the strength of the basic interactive forces (basic interactive forces are the forces between single, isolated flux lines and individual pinning centers). At the higher fields, there are more flux lines than pinning centers and therefore, a crossover from synchronization can be expected [ 33 ].

The grain size in the samples has also an importance on the pinning behavior depending on the applied magnetic fields. A decrease of grain size on the order of the flux lattice constant can cause the vortex channels to overlap, so the average direction of flux motion incorporates a component perpendicular to the grain boundaries. In this case, F p is much stronger than before and it is necessary to apply higher magnetic fields to break vortex in sample. This causes the peak of F p to shift toward higher applied magnetic fields [ 34 ].

4 Conclusion

Y2/3Se1/3Ba2Cu3Ox samples were successfully fabricated using solid state reaction technique. Crystal structure of samples was belonged to Y-123. Superconducting transition temperature was investigated both by R–T and M–T measurement technique. The value of superconducting transition temperature of samples were similar to Y-123 superconductor. The M–H loop and critical current density of samples were investigated and it is found that Se addition and Ba and Cu deficiency increased the J c value of the samples. The extended Valkov–Khrustalev model was applied to describe the magnetization loops at different temperatures. The behavior of estimated pinning force is argued for the samples.