A series of industrial trials of argon injection into ladle shroud (AIILS) with different argon-blowing rates were conducted in this study. Firstly, bubbles in actual liquid steel of the tundish were captured by the method of “cold steel plate dipping” and characterized by microscope examination. A detailed investigation on the three-dimension morphology of bubbles was carried out by using industrial computerized tomography (ICT). Then, the two-phase flow of liquid steel and argon gas in the tundish was numerically simulated to further investigate the motion behaviors of bubbles in liquid steel of the tundish. The simulated results showed that bubbles in the size range of this investigation had a large enough filtration rate to demonstrate a good performance on inclusion removal. Finally, the effect of AIILS on inclusion removal was analyzed by detecting the variations of inclusions as well as total oxygen content in steels taken from the ladle, tundish, and casting billet. The method of AIILS was more effective at improving the removal of inclusions in the range of 5 to 10 μm and obviously increased the removal rate of total oxygen content.

Using bubbles to remove inclusions in steel is rapidly becoming a popular method for refining. Fine bubbles are thought to be more effective on inclusion removal than big bubbles. The fine bubbles can be formed in molten steel using the argon injection into ladle shroud technology. There are two stages during the formation of fine bubbles in ladle shroud: bubble detachment from wall orifice and detached bubbles splitting into smaller ones in turbulent steel. Many reports have been published on the water model of the argon injection into ladle shroud technology, but industrial experimental research is in its early stage. In this study, high argon flow was injected into a ladle shroud and adopted in continuous casting production to produce fine argon bubbles in a tundish. The bubbles were captured by dipping a cold steel sheet into molten steel. The captured bubbles at the surface of a hot-dipped steel sheet, with a size of 1.0–3.0 mm, characterized the argon bubbles at steel/slag interface and slag phase in the upper part of a tundish rather than those inside molten steel in tundish. The bubbles inside molten steel in tundish were characterized by the captured bubbles in the interior of a hot-dipped steel sheet, and their morphology, size, and number were analyzed using scanning electron microscopy and confocal microscopy. The bubbles inside molten steel in tundish have a spherical shape and occasionally adhere to each other. These bubbles rang in size from 100 to 1000 μm, with an average of 500 μm. They are dispersed at the exit of a ladle shroud in its lower position, with a density of 15.2 cm–2. Moreover, it was observed that a bubble could adhere to inclusion, even multiple inclusions for part of bubbles. Bubbles adhered more strongly to Al2O3 inclusions than that to CaO(‒MgO)‒Al2O3‒SiO2 complex inclusions.

为了解决组合模型计算多流非对称中间包不同出口的停留时间分布曲线时出现的体积死区分率为负或结果偏差较大的问题,通过把"脉冲刺激-响应"试验所得的E曲线转换为F曲线,建立了基于F曲线的多流非对称中间包整体死区比例分析模型,并通过比较各流出口F曲线的差异量以及差异量极大值、最大值评价了多流非对称中间包钢液流动一致性.应用该模型对四流非对称中间包流场进行分析,水模型试验结果表明,无控流装置的中间包在3号、4号水口上方存在流动缓慢的区域,中间包钢水更新缓慢,流场不合理.增加挡墙后,中间包钢水流动一致性增强,2号、3号水口流动一致性相近;1号、4号水口流动一致性相近.使用挡墙+挡坝2的组合,各流一致性最好,死区比例最小.

中间包中生成微小气泡可显著促进夹杂物上浮去除.对中间包微气泡精炼技术进行了分析总结,并针对研究较为深入的长水口吹氩技术相关文献进行了详细分析.结果 表明,利用中间包中钢液湍动能破碎气泡可形成高效去除夹杂物技术,一些新技术正在研发过程中,长水口吹氩技术具有良好前景;"冷钢片沾钢"工业试验表明,中间包长水口吹氩可在中间包钢液中生成弥散细小氩气泡,生成的绝大部分氧气泡尺寸小于2 mm;长水口吹氩生成微小气泡的过程可分为气泡在长水口壁孔脱附和脱附气泡在湍急钢流中被剪碎成微小气泡两个阶段,其中钢液湍动能对氩气泡的剪切破碎作用十分明显;长水口吹氩技术水模型研究较多,数值模拟研究相对较少,工业试验研究才刚刚开始,有待更进一步的深入研究.

A novel ladle shroud was employed to produce small bubbles to remove inclusions smaller than 50 μm meanwhile inhibiting the formation of slag eye. Both water modeling and industrial tests were carried out to confirm the generation of small bubbles with the ladle shroud. The numerical model was developed to investigate slag-metal interfacial behaviors, considering the effects of bubble sizes and distributions. The gas flux through the surface was employed to evaluate the impact of bubble swarm on slag layer. The results show that reducing bubble sizes can effectively disperse bubbles passing through the slag layer, so as to inhibit the formation of slag eye, under the same gas flow rate. The diameter of slag eye matches well 1.7 times diameter of the region with gas flux higher than 0.0025 m/s. The heat loss of the entire tundish was estimated, considering the convection and radiation heat transfer of slag eye.

Dissolved gas flotation method has been developed to remove inclusions in molten steel. The principle is that bubbles formed on inclusions by vacuum treatment of nitrogen or hydrogen supersaturated molten steel can carry the inclusions to slag. A kinetic model was constructed to analyze the bubble growth and floating behavior during the degassing process of the method, and its accuracy was verified by related experiments. The results show that pretreatment pressure, bubble nucleation depth and gas type have significant effects on bubble growth and floating, while vacuum treatment pressure and inclusion radius have little effects on it. The growth rate and floating velocity of bubbles increase with the increase of pretreatment pressure or the decrease of bubble nucleation depth. The growth rate and floating velocity of hydrogen bubbles are much larger than those of nitrogen bubbles. Calculation results indicate that the diameters of the bubbles are mostly 0.2-10 mm during floating process via this method. Moreover, the distribution of the bubbles nucleating on the inclusions is dispersive. In addition to the bubbles carrying inclusions to slag directly, these dispersive fine bubbles have a high probability of inclusion adhesion resulting in an improvement of the inclusion removal.

Based on the theory of classical solidification nucleation, a thermodynamic model of bubbles nucleating on surfaces of convex spherical inclusions in molten steel was established. The expression of bubble critical-nucleation radius was derived. The differences between bubbles nucleating on convex spherical surface and flat substrate were discussed. The results show that the radii of spherical inclusions determine the degree of deviation on the two types of nucleation under same conditions and bubbles are more easily generated by flat substrate nucleation. In addition, the degree of deviation between the two types of nucleation decreases with the decrease of melt depth. The results also show that bubbles are easier to nucleate on the surface of inclusions with large radius or poor wettability. For a determined bubble critical-nucleation radius, the radii of spherical inclusions, the contact angles of spherical inclusions and the pressure of vacuum treatment have little influence on the depth range of bubble nucleation, while the pretreatment pressure shows significant impact on that. The probability of bubble nucleation reduced with the increase of melt depth, which was verified by related experiments.

Characteristics of molten metal heated with microwaves were the focus of this study. A series of experiments on the direct microwave heating of molten copper and molten iron were conducted in a MobileLab-W-R microwave workstation; both metals were effectively heated by direct microwaves. Effects of indirect versus direct heating were comparatively analyzed using different types of heating chambers. The direct heating method was then further investigated, taking microwave power, mass of molten metal, and temperature into consideration. The mechanism of direct microwave heating of molten metal was discussed. The results show that microwave can directly heat molten iron and molten copper at high rates that increase linearly with increasing microwave power. Heating rates of molten iron are similar to those of molten copper at constant mass and microwave power. However, the mass of molten iron has no clear linear relationship with heating rates due to the involvement of other factors, such as surface area of the molten iron and distribution of the microwaves. According to the theoretical analysis, when the states of copper and iron are transferred from solid to liquid, their resistivities increase, but their permeabilities drop significantly. As a result, the skin effect depths of microwave in molten copper and iron are clearly larger than those in the solid metals. Conductivity loss is the main mechanism of achieving direct microwave heating of molten metal. Microwave energy can be absorbed in four ways: collisions between electrons and nucleus, rapid liquid surface renewal, hindering of internal defects of electron movement, and atom movement and collision. Absorbed microwave energy can be transferred into the internal energy of the molten metal.

Combined model in which a tundish is divided into plug flow, well‐mixed flow, and dead region has been widely used for the analysis of residence time distribution curve in an actual tundish. However, the calculation of dead volume in conventional or Sahai's method does not follow its definition in the model assumption, which probably results in incorrect conclusions for flow field optimization of tundish. In this study, two methods for dead volume calculation are discussed, the well‐mixed volume is further divided into equivalent, active, and dead well‐mixed volume. Therefore, the method for dead volume calculation is revised by eliminating the contribution of dead well‐mixed flow. Water modeling experiments of single and multi‐strand tundish are conducted. The new calculation method is applied to analyze the residence time distribution curves and calculate the dead volume fraction in comparison with the other two methods. The results show after the dead well‐mixed flow is eliminated, the dead volume fraction calculated by the new method is smaller than those calculated by conventional and Sahai's method. According to the flow visualization experiments, the dead volume fraction calculated by the new method in this study is more rational than those calculated by the other two methods.

Abstract The influence of Ti addition (~0.10 wt%) on hot ductility of as-cast high-manganese austenitic steels has been examined over the temperature range 650–1,250 °C under a constant strain rate of 10−3 s−1 using Gleeble3500 thermal simulation testing machine. The fracture surfaces and particles precipitated at different tensile temperatures were characterized by means of scanning electron microscope and X-ray energy dispersive spectrometry (SEM–EDS). Hot ductility as a function of reduction curves shows that adding 0.10 wt% Ti made the ductility worse in the almost entire range of testing temperatures. The phases’ equilibrium diagrams of precipitates in Ti-bearing high-Mn austenitic steel were calculated by the Thermo-Calc software. The calculation result shows that 0.1 wt% Ti addition would cause Ti(C,N) precipitated at 1,499 °C, which is higher than the liquidus temperature of high-Mn austenitic steel. It indicated that Ti(C,N) particles start forming in the liquid high-Mn austenitic steel. The SEM–EDS results show that Ti(C,N) and TiC particles could be found along the austenite grain boundaries or at triple junction, and they would accelerate the extension of the cracks along the grain boundaries.

Effect of Mn (14.94,18.21, and 23.6wt%) and Al (0.002,0.75, and 1.47wt%) contents on hot ductility of five high alloy Fe-xMn-C-yAl austenitic Twinning induced plasticity (TWIP) steels were investigated by Gleeble-3500 thermo-mechanical simulator in the temperature range 700–1200℃ under a constant strain rate of 3 × 10−3s−1. The results indicated that the hot ductility of different Mn-containing TWIP steels are not appreciable with all the reduction of area (RAs) values lower than 30%, and RAs would be further decreased as the Mn content increased. The matrix of TWIP steel is inhomogeneous with severe Mn microsegregation in the interdendritic zone. Moreover, the C microsegregation ratio increases from 0.85 to 1.16, 0.76–1.22, to 0.74–1.32 when Mn concentration increases from 14.94wt%, 18.21wt%, to 23.6wt%, respectively. Additionally, the microstructure and the true stress-true strain curves suggested that dynamic recrystallization (DRX) took place in 14.94wt% Mn bearing TWIP steel, while the fraction of DRX grains decreased dramatically with increasing Mn content. Therefore, it is inferred that the high Mn addition inhibited DRX, together with the acceleration effect of C microsegregation by Mn addition should be the most predominant factor of the hot ductility loss with Mn content increases in TWIP steels. On the other hand, Al addition to TWIP steels resulted in a dramatic increase of AlN particles content. The AlN particle accounted for nearly 64% of the total precipitate content for the 1.59wt% Al containing TWIP steel. Compared with Al-free TWIP steel, the excessive number of fine AlN particles in the 1.59wt% Al containing steel effectively pinned the austenite grain boundaries, which inhibited the occurrence of DRX and simultaneously promote grain boundary sliding, resulting in the deterioration of hot ductility.

The solid-liquid interface structures and the Al segregation in the solidifying front were investigated by directional solidification and electron probe microanalysis. The results show that the interface structure would vary greatly with solidification velocity, the morphology of which changed from planar at 1 mu m/s, dendrite at 25 mu m/s, cellular-dendrite at 50 mu m/s, to hexagon cell at 100 and 150 mu m/s and finally high-velocity cell at 200 mu m/s. The stability of the planar interface was mainly influenced by Al segregation. The narrow solidification temperature range benefited the transformation of dendrite to cell-dendrite. Regarding Fe, Cr as solvents and Al as solute, the secondary dendrite arm spacing when solidification velocity is 25 mu m/s could be predicted by the formula which was widely adopted in binary alloy system, and the predicted values fit well with the experimental results. The Al partition coefficient was found to be 0.96, and it indicated that there existed slight Al segregation when solidification velocity is 1 mu m/s, which was the main stabilizing factor of the planar interface. Al segregation is not obvious at velocities above 25 mu m/s and the interface transformed to various structures when solidification velocities were higher than 25 mu m/s.

The phase transformations and precipitation behavior were investigated by using Thermo- Calc software in the Fe-(18-21)Cr-(3-5)Al-(0-0. 03)C-(0-0. 2)Si- (0-0. 2)Mn multicomponent system relevant to FeCrAl stainless steel during solidification. The vertical sections of this system were calculated by using the TCFE7 database. Based on these vertical sections, the influence of different elements was analyzed in the phase transformations during solidification and a diagram of the phase-transformation path of FeCrAl stainless steel was obtained during equilibrium solidification. The results indicate that the full-phase transformation path of FeCrAl stainless steel during the cooling process from 1600℃ to 300℃ is as follows: L→AlN+αδFe→AlN+αδFe+Cr7 C3→AlN+αδFe+Cr7 C3 +Cr23 C6→AlN+αδFe+Cr23 C6→AlN+αδFe+Cr23 C6 +σ→AlN+αδFe+Cr23 C6 +σ+α'→AlN+αδFe+Cr23 C6 +α'. The precipitation of Cr7 C3 and σ, during the solidification process mainly depends on the carbon and silicon contents in the system, respectively. Increasing the aluminum content can enlarge the stable region of αδFe +Cr7 C3 , lower the precipitation temperature of α', and restrain σ precipitation. Increasing the chromium content can reduce the stable region of αδFe+Cr7 C3 and enlarge the stable region of σ and α'.

The morphology, composition, and number of inclusions in Fe-Mn-C(-Al) twining-induced plasticity (TWIP) steels were investigated by scanning electron microscopy, energy-dispersive X-ray spectroscopy, and an automated program called "INCAFeature." The characteristics of the inclusions in four TWIP steels with different Al contents (0.002%-1.590%) as well as the influence of Al content on the precipitation of AlN inclusions were investigated.In addition, systematic thermodynamics calculations of AlN formed in TWIP steel were carried out using the appropriate thermodynamic data for high-Mn-Al TWIP steel.The results show that AlN would begin to precipitate and locally precipitate around the MnS(Se)-Al2O3 inclusions when the Al content in the steel reaches 0.75%.The thermodynamics calculations show that AlN could already form in the liquid TWIP steel at an Al content of 1.07%.Then, AlN would locally precipitate around the MnS(Se) inclusions, thus forming MnS(Se)-AlN aggregates.When the Al content increases to 1.59%, the precipitation temperature of AlN is 42℃ higher than the liquidus temperature of the TWIP steel.Furthermore, precipitated AlN inclusions in the liquid TWIP steel could act as heterogeneous nuclei for MnS(Se) inclusions, thus forming MnS(Se)-AlN inclusions.Moreover, according to the thermodynamics calculation, the lowest N content for AlN formation in the liquid Fe-18.21%Mn-0.64%C-1.59%Al steel is just 0.0043%.Therefore, the N content should be kept as low as possible to avoid the formation of excessive AlN inclusions during melting of Fe-Mn-C(-Al) TWIP steel.

In order to further investigate the solidification process of FeCrAl alloy for automobile exhaust gas purifying systems, the phase transformation and precipitation in FeCrAl alloy with different Al contents were studied using differential scanning calorimetry (DSC), Thermo-Calc calculation and high temperature in-situ X-ray diffraction (XRD). The Thermo-Calc calculation results manifests that the complete solidification process of FeCrAl alloy includes the transformations from L to alpha-(Fe, Cr) and from alpha-(Fe, Cr) to alpha (Fe-rich ferrite)+alpha' (Cr-rich ferrite) as well as the precipitation of (Fe, Cr)(7)C-3, (Fe, Cr)(23)C-6, and sigma. The calculation results with 6.49 wt% Al content are reasonable agreement with the DSC results. The solidification process investigation of FeCrAl alloy with different Al content indicates that the decrease of Al content in FeCrAl alloy increases the forming temperature of alpha' phase and the precipitation temperature of (Fe, Cr)(2)3C(6) and narrows the precipitation temperature range of (Fe, Cr)(7)C-3. In addition, the investigations using high temperature in-situ XRD further validate the separation of alpha (Fe-rich) and alpha' (Crrich) as well as the absence of ferrite to austenite transformation in FeCrAl alloy during solidification. (C) 2017 Elsevier B.V. All rights reserved.

The influence of Ti(mass fraction 0.10％)and the joint additions of Ti(mass fraction 0.11％)and V(mass fraction 0.20％)on the hot ductility of as-cast high manganese austenitic steels were studied using a Gleeble-3500 thermo-mechanical simulator over a temperature range of 700 to 1200℃.Fracture surfaces and particles precipitated at different testing temperatures were investigated via scanning electron microscopy(SEM)and X-ray energy dispersive spectrometry(EDS).The hot ductility curves as a function of temperature of high-Mn austenitic steels showed that Ti addition leads to loss of ductility in almost the entire testing temperature range.Moreover,the joint additions of Ti and V do not exhibit any improvement in the hot ductility,resulting in relatively poor hot ductility behavior.The phase diagrams of precipitates in Ti-and Ti-V-bearing high-Mn austenitic steels in the temperature range of 700 to 1600℃ were calculated via Thermo-Calc commercial software.The calculation results show that Ti(C,N)in Ti-bearing high-Mn steel precipitates at 1499℃,which is much higher than its liquidus temperature.This illustrates that Ti(C,N)particles form in the liquid steel.SEM-EDS results show that Ti(C,N)and(Ti,V)C particles form along the austenitic grain boundaries and the triple junction.These particles retard the occurrence of dynamic recrystallization and accelerate the extension of cracks near the grain boundaries.

The cleanliness evolution of SWRH82B steel produced by the BOF-LF-CC process was studied by systematic sampling and lab comprehensive analysis. The results show that the content of oxygen and the number of micro-inclusions significantly decrease after calcium treatment, while the content of nitrogen in molten steel slightly increases. Both the contents of oxygen and nitrogen sharp-ly increase during the casting process. 97% of the micro-inclusions in casting billets are smaller than 5μm. They mainly include trenchant angularity single Al2 O3 inclusions, spherical complex inclusions of oxides with CaS and MnS, and complex inclusions of sim-ple oxides or nitrides with MnS. Calcium treatment after LF refining removes oxygen from molten steel efficiently, but increase the amount of Al2 O3 inclusions in casting billets, which is harmful to the quality of SWRH82B steel. Due to serious lag entrapment, re-fractory erosion and reoxidation caused by unreasonable tundish structure, the cleanliness of liquid steel decreases.

A characterization of non-metallic inclusions in Fe-Mn-Si-Al twinning-induced plasticity (TWIP) steels during argon oxygen decarburization-electroslag remelting-forging (AOD-ESR-forging) process has been performed. The two main kinds of inclusions found in TWIP-AOD ingots are single Al(O) N and MnS(Se)-Al(O) N aggregates. After the ESR process, Al(O) N and MnS(Se) inclusions in all size ranges significantly decrease while other kinds of inclusions show only minor changes. In the present study, the precipitation, growth, and dissolution of AlN and MnS inclusions have been analyzed by thermodynamics and kinetics. It is found that AlN inclusions in TWIP steels can already precipitate in the liquid. Furthermore, AlN inclusions precipitated in the liquid TWIP steels can act as the heterogeneous nuclei of MnS inclusions, and thus forming Al(O) N and MnS(Se) . Al(O) N clusters. After ESR refining, the precipitation temperature of AlN and MnS inclusions will be significantly decreased, and AlN inclusions cannot be formed in the liquid TWIP steels in particular. The kinetic analysis show that the growth of AlN inclusions will be difficult during the ESR process, because their growth rate decreases significantly, while the dissolution rate increases.