Although lath martensite has generally been considered resistant to recrystallization during tempering, the present study demonstrates that bulging recrystallization can occur even in low-carbon (LC) lath martensite after prolonged tempering. The recrystallization behavior of ultra-low-carbon (ULC) and LC steels was systematically investigated to clarify the microstructural factors governing this phenomenon. Optical microscopy and electron backscatter diffraction (EBSD) analyses revealed that bulge recrystallization initiated at prior austenite grain boundaries in both steels, although its onset was significantly delayed in LC steel. KAM profiles across the prior austenite grain boundaries showed no pronounced difference between the unrecrystallized matrices on the two sides, suggesting that a pre-existing dislocation-density difference across the boundary is not indispensable for the initiation of bulging. The recrystallized grains inherited the crystallographic orientation of the adjacent martensitic matrix and preferentially grew along the longitudinal direction of martensite blocks, demonstrating that recrystallization proceeds through protrusive migration of pre-existing grain boundaries guided by the anisotropic boundary network of lath martensite. Increasing carbon content refined the martensitic substructure and increased the dislocation density, while simultaneously enhancing carbide precipitation and delaying recrystallization. Phase-field simulations demonstrated that intersections between prior austenite grain boundaries and martensitic substructure boundaries can generate an initial protrusion even without a pre-existing dislocation-density difference, whereas the stored-energy difference associated with dislocations promotes its subsequent growth. Carbide particles suppress recrystallization through boundary pinning. In the present low-alloy LC steel, rapid coarsening of cementite during prolonged tempering progressively reduced the pinning force, allowing recrystallization to occur despite the higher carbide density. These results demonstrate that bulging recrystallization in low-carbon lath martensite is governed by the interplay among martensitic substructure, stored energy, and carbide pinning, providing a unified mechanistic framework for recrystallization during tempering.
Charging direct reduced iron (DRI) into the blast furnace (BF) can couple hydrogen metallurgy with conventional ironmaking, but the BF behavior of industrial high-gangue DRI is unclear. Two industrial hydrogen-based DRI samples were characterized, and the representative high-gangue DRI was examined by thermogravimetric-mass spectrometry, high-temperature deformation, and mixed burden softening–melting–dripping tests. Although both DRI had similar metallization degrees of 90.45 pct, the high-gangue DRI had a lower total iron content of 84.82 pct and a higher gangue loading of 9.1 pct, including 6.71 pct SiO2. Their porous structures favored gas–solid reactions, but their compressive strengths were far below that of conventional fired pellets, indicating a high breakage risk during BF use. During inert heating, mass loss was governed by coupled reactions among Fe3C, FeO, and carbon-bearing species, producing CO and CO2. Pronounced deformation occurred at 1305 °C to 1320 °C, mainly due to localized softening/melting of carbon-bearing metallic regions. In mixed burden tests containing 10 pct high-gangue DRI, ΔPmax decreased from 18.96 to 15.72 kPa and S from 1599 to 1028 kPa·°C, indicating improved high-temperature permeability but a wider softening–melting–dripping interval. These results support BF-oriented evaluation and efficient utilization of industrial high-gangue DRI.
This study used a solidification thermal simulation method to investigate the distribution and morphological evolution of titanium-containing precipitates during the non-equilibrium solidification of high-titanium (H-Ti) steel. The main titanium-containing precipitates were TiN, TiC, and oxide–TiN composite precipitates, which exhibited distinct precipitation stages, spatial distributions, and morphologies. TiN met the conditions for precipitation over a wide solid-fraction range and appeared mainly as cubic particles, aggregated particles, or nested structures; the cubic/nested morphology was associated with faceted growth, heterogeneous nucleation, and subsequent particle contact rather than simply with early precipitation. In contrast, TiC precipitated mainly in the interdendritic residual liquid during the final stage of solidification (fs > 0.8), where Ti and C enrichment, confined liquid channels, and local eutectic-like reactions promoted lamellar, Chinese-script, and dendritic morphologies. Along the solidification direction of the thermal simulation (TS) sample, the mean TiN size increased from 3.8 μm near the cold end to 5.1 μm in the columnar-to-equiaxed transition (CET) region and then decreased to 4.5 μm, whereas the TiC area fraction varied from 0.84 to 4.6 pct with a pronounced local increase in the CET/equiaxed regions. Thermodynamic and interfacial-energy analyses indicate that TiN, with its anisotropic interfacial energies and high interfacial energy with molten steel, tends to undergo faceted growth along low-energy planes, forming a cubic morphology. TiC shows good wettability with the Fe melt. When TiC precipitates, the dendrite structure is already well developed, and solute diffusion in the interdendritic residual liquid facilitates continuous growth along the solid–liquid interface, yielding more complex morphologies. Solute segregation affects precipitation more strongly than cooling rate, while higher cooling rates suppress growth. The distinct morphologies of TiN and TiC arise from the coupled effects of precipitation stage, local solute enrichment, wettability, interfacial-energy anisotropy, and cooling conditions. These results link the morphology and spatial distribution of TiN and TiC to the local solidification structure, solute segregation, and cooling conditions, thereby providing guidance for controlling titanium-containing precipitates during the continuous casting of H-Ti steel.
