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The characteristics, formation mechanism, and removal of large‐sized CaO–Al2O3–SiO2 inclusions in D2 high‐speed railway wheel steel are investigated. Large‐sized CaO–Al2O3–SiO2 inclusions are present in the continuous casting billet, with SiO2 and CaO content ranging from 5 to 15 and 30 to 65 wt%, respectively. The size mainly exceeds 10 μm. These inclusions originate from the calcium‐treatment stage of the refining process, during which liquid CaO–Al2O3–SiO2 inclusions are formed in the liquid steel. The contact angle between these inclusions and the liquid steel is below 40°, which results in excellent wettability. Consequently, the inclusions are difficult to remove from the liquid steel and are thus inherited into the billet. By reducing the SiO2 content and controlling CaO content between 8 and 30 wt%, small‐sized inclusions are formed. This requires reasonable control of the Al and Ca content in the liquid steel. When the Al and Ca content in liquid steel is maintained at 0.012 wt% and 8 ppm, respectively, the inclusions in the billet are mainly CaO·6Al2O3 (CA6) and CaO·2Al2O3 (CA2), both under 5 μm in size. These inclusions represent the suitable inclusions in steel. [ABSTRACT FROM AUTHOR]

In this study, the harmful effects of a banded structure on the bending fracture of electroplated fasteners are investigated. The origins and ways to control these structures are also explored. The fracture surfaces of the fasteners exhibit stepwise cracks and numerous intergranular facets associated with ductile tearing, which is a characteristic of hydrogen embrittlement. The origin of these stepwise cracks can be traced to the internal banded structures. Therefore, the internal banded structures promote fracturing during the bending test. The internal banded structures in the fasteners can be traced back to the internal cracks in the bloom, and inappropriate soft‐reduction processes lead to internal cracks. Notably, the carbon segregation at the internal cracks is severe. This is attributed to the proximity of the internal cracks to the center of the bloom, causing the enriched liquid in the interdendritic region to be sucked into these internal cracks. The optimal reduction position and amount are determined by employing heat‐transfer and soft‐reduction models. Fluctuations in superheat and casting speed significantly influence this optimal position and reduction amount. [ABSTRACT FROM AUTHOR]

Using an ultrasonic flaw detector, this study precisely located the dimensions and distribution characteristics of defects during the manufacturing process of 55NiCrMoV7 die steel forgings weighing hundreds of tons. The probing defects of the forging are predominantly concentrated in the central region, spanning approximately 4200 mm in length, 400 mm height and 1500 mm width. Detailed research on the characteristics of these probing defects was conducted through dissection sampling and scanning electron microscope analysis. The research results reveal that the maximum diameter of an individual defect can reach Φ10 mm, and localised cracking phenomena exist in specific areas within the forging. The flaws exhibit characteristics such as millimetre-sized, large, elongated MnS inclusions, with spacing between inclusions ranging from 200μm to 500μm. Furthermore, based on elemental analysis, the sulphur content from the surface to the core of the forging increased from 0.001% to 0.002% to a maximum of 0.0085%. Through analysis using thermodynamic and growth kinetics models, it was determined that in regions with higher sulphur content (S = 0.0085) at a certain distance from the steel ingot surface, the radius of MnS inclusions can reach up to 380μm under low cooling rates (1 K/min). After forging deformation, these can evolve into millimetre-sized MnS inclusions, leading to internal inspection defects. Therefore, we propose that reducing the sulphur content in the molten steel and increasing the degree of undercooling can effectively inhibit the precipitation and growth of MnS inclusions. Practical results indicate that this improvement strategy effectively enhances the inspection quality of the 55NiCrMoV7 die steel forgings weighing hundreds of tons. [ABSTRACT FROM AUTHOR]

The sigma‐phase distribution in a 2507 super duplex stainless steel continuous casting slab is analyzed using microstructural analyses and thermodynamic calculations. The sigma phase contains high Cr and Mo contents, and precipitates along the phase boundary between austenite and ferrite and grows toward the interior of the ferrite phase in the continuously cast slab. The sigma‐phase fraction at the center of the width direction is higher than that in the one fourth width. Furthermore, the sigma‐phase fraction at the center of the thickness direction is higher than those at the upper and lower surfaces, which is due to the lower cooling rate at the center. The Brinell hardness (HB) at the center increases to 313 HB with an increasing sigma‐phase fraction. The sigma‐phase fraction at the one fourth thickness near the lower surface is higher than that at the one fourth thickness near the upper surface at the same cooling rate. Additionally, Mo and Ni are severely segregated in this zone. A higher Mo content influences the precipitation and stabilization of the sigma phase. Furthermore, Scheil–Gulliver calculations indicate that compositional segregation results in an increase in the sigma‐phase precipitation temperature range. [ABSTRACT FROM AUTHOR]

The characteristics of large precipitates in H13 tool steel with 0.015wt% Ti were studied. The result shows that three types of phases larger than 1 μm exist in the as-cast ingot, that is, (Ti, V) (C, N) type phase, (V, Mo, Cr)C type phase and sulfide. (Ti, V) (C, N) type phase could be further classified as the homogeneous Ti-rich one and the Ti-V-rich one in which Ti/V ratio gradually changes. (V, Mo, Cr)C type phase contains the V-rich one and the Mo-Cr-rich one. The compositional characteristics in all of them have little relation with the cutting position or cooling rate. The precipitating process could be well described through calculation by Thermo-Calc software. During solidification, the primary phase (Ti, V)(C, N) first starts to precipitate in the form of Ti-rich carbonitride. With the development of solidification, the ratio of Ti decreases and that of V increases. Then the primary phase Ti-V-rich (Ti, V)(C, N) and V-rich (V, Mo, Cr)C appears successively. Mo-Cr-rich (V, Mo, Cr)C phase does not precipitate until the solidification process reaches to the end. Sulfide precipitates before (V, Mo, Cr)C type phase and it could act as the nucleus of (V, Mo, Cr)C.