The measurement of mixing time in a water model of soft-stirring steelmaking ladles is practically facing a problem of bad repeatability. This uncertainty severely affects both the understandings of transport phenomenon in ladles and the measurement accuracy. Scaled down by a ratio of 1:4, a water model based on an industrial 260-ton ladle is used. This paper studies the transport process paths and mixing time of salt solution tracers in the water model of eccentric gas-stirred ladles with a low gas flow rate. After a large number of repeated experiments, the different transport paths of the tracer and the error of the mixing time in each transport path are discussed and compared with the numerical simulation results. The results of a large number of repeated experiments on the water model show that there are five transport paths for the tracer in the ladle. The tracer of the first path is mainly transported by the left-side main circulation flow, which is identical to the numerical simulation results. The tracer of the second and third paths are also mainly transported by the left-side circulation flow, but bifurcations occur when the tracer in the middle area is transported downward. In the third path, the portion and intensity of the tracer transferring to the right side from the central region is higher than in the second path. The fourth path is that the tracer is transported downward from the left, middle, and right sides with a similar intensity at the same time. While the tracer in the fifth path is mainly transported on the right side, and the tracer forms a clockwise circulation flow on the right side. The mixing times from the first transport path to the fifth transport path are 158.3 s, 149.7 s, 171.7 s, 134 s and 95.7 s, respectively, among which the third transport path and the fifth transport path are the maximum and minimum values among all transport paths. The error between the mixing time and the averaged mixing time at each monitoring point in the five transport paths of the tracer is between −34.7% and 40.9%. Furthermore, the error of the averaged mixing time of each path and the path-based average value is between 5.5% and 32.6%.

This study investigates the distribution characteristics of residual ferrite in 304L austenitic stainless steel continuous casting slab and the impact of heat treatment processes on its content. Through optical microscopy (OM), thermodynamic calculation software (Thermo–Calc) and heat treatment experiments, it is found that the residual ferrite content along the thickness direction at the width center of the slab exhibits an “M”-shaped distribution—lowest at the edges (approximately 3%) and highest near the center (approximately 13%). Within the triangular zone of the slab, the residual ferrite content varies between 1.8% and 12.2%, with its average along the thickness direction also showing an “M”-shaped distribution; along the width direction, the average residual ferrite content is lower at the edge positions, while within the internal triangular zone, it ranges between 8% and 10%. The ferrite morphology changes significantly across solidification zones: elongated in the surface fine-grain zone, lath-like and skeletal in the columnar grain zone and network-like in the central equiaxed grain zone. Thermodynamic calculations indicate that the solidification mode of the 304L continuous casting slab follows the FA mode. Heat treatment experiments conducted across the entire slab thickness demonstrate effective reduction in residual ferrite content; the optimal reduction is achieved at 1250 °C with a 48 min hold followed by air cooling while preserving the original “M”-shaped distribution characteristic after treatment. Increasing the heat treatment temperature, prolonging the holding time and reducing the cooling rate all contribute to reducing residual ferrite content.

Rare earth can modify inclusions in non-oriented silicon steel which is harmful to magnetic properties. This study focused on the 3.1% Si non-oriented silicon steel under industrial production conditions. Samples were taken during the stages before and after addition of rare earth ferrosilicon alloy in Ruhrstahl-Heraeus (RH) unit, different pouring time in tundish, and continuous casting slab. This study systematically examined the morphology, composition, and size distribution of inclusions throughout the smelting process of non-oriented silicon steel by scanning electron microscopy with energy-dispersive spectroscopy (SEM/EDS), and thermodynamic analysis at liquid steel temperature and thermodynamic analysis of equilibrium solidification. The research results demonstrated that the rare earth treatment ultimately modifies the original Al₂O₃ inclusions in the non-oriented silicon steel into REAlO₃ and RE₂O₂S inclusions, while also aggregating AlN inclusions to form composite inclusions. After rare earth modification, the average size of the inclusions decreases. In the RH treatment process, the inclusions before the addition of rare earth ferrosilicon alloy are mainly AlN and Al₂O₃. After the addition of rare earth ferrosilicon alloy, the inclusions are mainly RES and REAlO₃. In the tundish and continuous casting, the rare earth content decreased, and the rare earth inclusions transform into RE₂O₂S and REAlO₃. For the size of inclusions, after adding rare earth ferrosilicon alloy, the average size of inclusions rapidly decreased from 16.15 μm to 2.65 μm and reach its minimum size 2.16 μm at the end of RH treatment. When the molten steel entered the tundish, the average size of inclusions increased slightly and gradually decreased with the progress of pouring. The average size of inclusions in the slab is 5.79 μm. Phase stability diagram calculation indicates the most stable rare earth inclusion is Ce₂O₂S in molten steel. Thermodynamic calculations indicated that Al₂O₃, Ce₂O₂S, Ce₂S₃, AlN, and MnS precipitate sequentially during the equilibrium solidification process of molten steel.

The service performance of austenitic stainless steel is substantially affected by the presence of residual ferrite. In this paper, as-cast 12.5% Ni 316L austenitic stainless steel billets are taken as the research object, and samples are selected from the edge, quarter-thickness, and center positions of the billet, as well as two directionally solidified specimens prepared at different withdrawal speeds. By means of metallographic analysis, Thermo-Calc thermodynamic calculations, and EBSD phase analysis, the characteristics of residual ferrite and precipitated phases in the two types of as-cast 12.5% Ni 316L stainless steel were systematically investigated. The results show that the ferrite morphologies at the edge, quarter-thickness, and center positions of the billet are granular and short-rod, skeletal, and clustered net-like and lath-like, respectively. The ferrite morphologies of the two directionally solidified specimens are similar, both being predominantly skeletal structures; the main difference is that in the high-withdrawal-speed directionally solidified specimen (No. 2), the ferrite is finer and more densely distributed. The residual ferrite contents measured at the edge, quarter-thickness, and center positions of the billet are 4.88%, 5.90%, and 8.99%, respectively; those of directionally solidified specimens No. 1 and No. 2 are 6.4% and 7.7%, respectively. For the billet, the ferrite content increases progressively from the edge to the center. Regarding precipitated phases, the edge of the billet exhibits a mixed microstructure of secondary precipitates, namely Sigma phase and Chi phase; at the quarter-thickness position, the coupled precipitation of these two phases is more pronounced; at the center, part of the ferrite has completely decomposed, with the Chi phase disappearing and only the Sigma phase remaining. In the two directionally solidified specimens, only a small amount of the Sigma phase is precipitated as secondary phases, and the ferrite remains relatively intact. Based on the morphology analysis of the ferrite structure, the solidification mode of the billet is determined to be the FA mode, which is consistent with both the Scheil calculation results and the chromium-nickel equivalent calculation results; however, it differs from the thermodynamic equilibrium solidification results obtained using Thermo-Calc.