A continuous casting protective slag special for Mn13 high wear-resistant steel containing nano composite components

Abstract

This invention relates to the field of special steel continuous casting metallurgical materials technology, and discloses a special continuous casting protective slag for Mn13 high wear-resistant steel containing nanocomposite components. The slag system comprises a matrix slag system containing at least CaO, SiO2, and MnO, and nanocomposite components dispersed therein. The nanocomposite components contain a first core-shell particle and / or a second core-shell particle. The first core-shell particle has hexagonal boron nitride as its core and oxide ceramic as its shell, providing gas barrier protection during the preheating stage and eliminating lubrication blind zones through super-wetting dispersion in the slag. The second core-shell particle has spinel as its core and graphitized carbon as its shell, converting the explosive combustion of carbon into controlled slow oxidation to stabilize the thickness of the slag pool. The matrix slag system pre-contains a high concentration of MnO to inhibit the replacement reaction of [Mn] with SiO2 in the steel from the source. This invention solves the problems of cracks and steel leakage caused by lubrication failure and heat transfer disorder in the continuous casting of high-manganese steel.

Description

Technical Field

[0001] This invention relates to the field of special steel continuous casting metallurgical materials technology, and in particular to a special continuous casting protective slag for Mn13 high wear-resistant steel containing nanocomposite components. Background Technology

[0002] High-manganese steel (such as the classic Mn13 grade) and high-manganese high-aluminum steel have long attracted attention from industry and academia due to their ability to undergo significant strain-induced martensitic transformation or twin-induced plasticity under strong impact loads, exhibiting excellent wear resistance and extremely low-temperature toughness. High-manganese steel products produced using the modern "continuous casting-hot rolling" process have finer and more dispersed internal grain structure and inclusions due to the directional cooling and solidification process, thus possessing higher yield strength and service life.

[0003] However, in the highly dynamic, high-temperature process of continuous casting involving multiphase flow and complex heat and mass transfer, the casting of ultra-low temperature high-manganese high-alumina steel is particularly challenging due to the extremely high concentration of active components [Mn] and [Al] in the molten steel, which possess extremely strong reducing thermodynamic driving potential. The core component of the traditional silicate-based CaO-SiO2-CaF2 system's protective slag network skeleton, SiO2, is extremely fragile in the face of these strong reducing agents, leading to violent displacement redox reactions at the slag-metal interface. These displacement reactions result in the significant consumption of SiO2, which originally acted as a glass network former, causing a sharp decrease in its mass fraction. Meanwhile, the reaction products MnO and Al2O3 continuously accumulate and enrich in the slag layer. From a microscopic perspective, Al2O3, in particular, as an amphoteric oxide, can act as a network former in alkaline slag systems. The large influx of Al2O3 into the slag macroscopically leads to an abnormally large increase in the viscosity of the protective slag, even causing crusting. Simultaneously, the destruction of the SiO2 network skeleton causes the slag composition to deviate significantly from the design range. In high-MnO protective slag systems, crystallization kinetics and phase evolution are reversed. Numerous tiny manganese olivine (2MnO·SiO2) crystals preempt nucleation sites and preferentially precipitate at lower temperatures, strongly inhibiting the traditional gunmetal crystallization process. This results in a significantly thinner solid slag film with increasingly finer crystal morphology, severely weakening the film's ability to scatter thermal radiation. Ultimately, this leads to an abnormal surge in localized heat flux density in the crystallizer, inducing dangerous heat transfer deterioration and causing numerous longitudinal and transverse cracks or leaks on the billet surface. Furthermore, the carbonaceous materials used to control the melting rate in traditional protective slags are prone to explosive macroscopic rapid consumption under high-temperature, high-oxygen conditions, causing drastic fluctuations in the thickness of the slag pool and further deteriorating lubrication and heat control effects.

[0004] Therefore, the most critical technical problems currently faced in high-manganese steel continuous casting are: First, at high temperatures, active [Mn] and [Al] undergo a strong slag-metal interface replacement reaction with SiO2 in the protective slag, leading to compositional evolution, increased polymerization degree of the protective slag melt, and abnormally high viscosity, thereby causing hydrodynamic lubrication failure; Second, the abnormal precipitation of manganese olivine due to the inversion of the precipitated phase system causes heat transfer deterioration and stress concentration on the solid slag film side of the crystallizer, making the initial billet shell prone to cracking or leakage due to thermal stress. Summary of the Invention

[0005] To address the aforementioned technical problems, this invention provides a continuous casting protective slag for Mn13 high wear-resistant steel containing nanocomposite components. The continuous casting protective slag comprises a matrix slag system and nanocomposite components dispersed within the matrix slag system. The matrix slag system includes at least CaO, SiO2, and MnO. Based on 100 parts by mass of the matrix slag system, the total addition amount of the nanocomposite components is 1 to 8 parts. The nanocomposite components include core-shell structured composite particles, selected from at least one of a first core-shell particle and a second core-shell particle. The core layer material of the first core-shell particle is hexagonal boron nitride, and the coating material is oxide ceramic. The core layer material of the second core-shell particle is spinel, and the coating material is a carbonaceous material.

[0006] The reason for controlling the total addition amount of nanocomposite components within the range of 1 to 8 parts per 100 parts of the matrix slag system is as follows: if the addition amount is less than 1 part, the volume fraction of core-shell particles in the slag film is too low, making it impossible to construct a continuous and uniform functional network in three-dimensional space, resulting in incomplete elimination of lubrication blind spots and insignificant carbon consumption control effect; if the addition amount is more than 8 parts, excessive solid nanoparticles will cause the high-temperature rheological properties of the matrix slag system to deviate from the design range, increase the apparent viscosity, and reduce the flow and spreading ability of liquid slag, which is not conducive to lubrication and inclusion adsorption.

[0007] The role of adding the above components and constructing the core-shell superstructure is as follows: First, by synergistically configuring CaO, SiO2, and a high concentration of initial MnO in the matrix slag system, it exerts a chemical potential suppression effect on the pathway of the active [Mn] element in high-manganese steel to deprive oxygen atoms from the underlying thermodynamics. Second, the introduction of the first core-shell particles, whose oxide ceramic shell provides a physical gas barrier for the inner hexagonal boron nitride during the entire process of protective slag product storage and transportation and slag melting from room temperature in the crystallizer, prevents atmospheric oxygen from contacting and oxidizing the hexagonal boron nitride during its most vulnerable medium-high temperature preheating stage (approximately 400°C to 900°C), and prevents it from suffering irreversible combustion degradation before the slag powder melts and seals the system. Meanwhile, the chemical homology between the outer oxide ceramic shell and the silicate slag matrix gives the core-shell particles excellent surface affinity when in contact with high-temperature liquid slag. This allows for instantaneous superwetting and nanoscale three-dimensional uniform dispersion within the liquid slag, fundamentally eliminating the lubrication blind zone caused by the agglomeration of exposed hexagonal boron nitride due to its extremely low surface polarity caused by silicate slag. Once the hexagonal boron nitride particles are uniformly dispersed into the liquid slag, the liquid slag layer itself constitutes a natural barrier to isolate atmospheric oxygen. The dissolved oxygen activity inside the slag is much lower than that of the external atmosphere. In this low-oxygen environment, hexagonal boron nitride can maintain structural stability for a long time and play a role in reducing friction through intergranular slip. In the even lower-temperature solid slag film region, the hexagonal boron nitride particles are permanently locked in the solidified glassy or crystalline matrix, which itself provides dense physical encapsulation protection. Furthermore, by imparting random orientation to the hexagonal boron nitride sheets during the initial dispersion stage of the outer oxide particles, the severe anisotropic heat conduction caused by the parallel orientation of two-dimensional sheet-like particles under shear force is cleverly broken. Numerous core-shell particles overlap each other in the matrix to form a highly uniform and isotropic heat-conducting network in three-dimensional space, eliminating local thermal stress concentration. Next, a second core-shell particle is introduced. This particle consists of a dense cluster of nanocrystalline spinels generated by high-energy ball milling, forming a rigid framework core. A dense carbon layer, graphitized by mechanical force, fills the intergranular gaps and coats the outer layer to form a carbon matrix shell. The ultra-high thermal stability and continuous heat sink effect provided by the spinel polycrystalline aggregate core, combined with the inherent high oxidation resistance and inertness of the graphitized carbon shell, and the indirect restriction effect of the rough micro-topography formed by the nanocrystals exposed on the surface of the polycrystalline aggregate after the surface carbon is gradually consumed, transform the carbon oxidation consumption kinetics from the explosive macroscopic rapid combustion of traditional free carbon black to a controlled surface-layer-by-layer slow oxidation mode. The physical melting rate of the slag powder is precisely controlled throughout the long casting cycle, ensuring a constant thickness of the liquid slag pool and eliminating dust pollution from flying free carbon black.

[0008] According to one aspect of the invention, preferably, the matrix slag system comprises, by mass percentage: CaO: 30% to 45%; SiO2: 30% to 45%; MnO: 15% to 20%; Al2O3: 6% to 10%; and the balance being flux.

[0009] The reason and mechanism for controlling the oxide components in the matrix slag system within the above-mentioned specific range is as follows: if the content of CaO and SiO2 is less than 30%, the protective slag cannot build a sufficient main aluminosilicate or fluorosilicate melt system skeleton; if the content of CaO and SiO2 is higher than 45%, it will cause the initial viscosity of the protective slag to be too high and the liquidus range to be too narrow during the cooling process, thereby losing the lubrication function required at high drawing speeds. Adding up to 15% to 20% initial MnO to the formulation beforehand can significantly increase the chemical potential of the manganese component in the liquid slag, based on the law of mass action in chemical reactions and Le Chatelier's principle. This fundamentally blocks the tendency of the positive interfacial chemical reaction of SiO2 in the active [Mn] reduction protective slag of steel. If the initial MnO content is less than 15%, it cannot generate sufficient reverse chemical potential suppression, making it difficult to block the spontaneous occurrence of the substitution reaction. If the initial MnO content is higher than 20%, it can easily lead to the excessive formation of high-melting-point mineral phases such as manganese olivine during slag cooling, causing a significant shift in crystallization temperature and unstable rheological properties. The purpose of controlling Al2O3 at 6% to 10% is to establish an initial concentration gradient buffer zone to resist the damage of trace [Al] elements to the protective slag network skeleton. If Al2O3 is below 6%, the buffer performance is insufficient and cannot resist the rapid increase in the degree of polymerization caused by Al2O3 enrichment in subsequent casting. If Al2O3 is above 10%, the apparent viscosity of the initial slag system is too high, which is not conducive to adsorbing and dissolving non-metallic deoxidation product inclusions floating from the depths of the molten steel.

[0010] According to one aspect of the invention, more preferably, the mass ratio of the binary basicity CaO / SiO2 of the matrix slag system is 0.8 to 1.3; the flux is selected from at least one of Li2O, Na2O, BaO, MgO and fluorides.

[0011] The mechanism of controlling the binary basicity of the matrix slag system within the low basicity range of 0.8 to 1.3 is as follows: maintaining this range helps the slag system accommodate more amphoteric oxides and effectively inhibits the premature precipitation of high-melting-point calcium-containing mineral phases on the high-temperature side, thereby maintaining a low initial viscosity and a wide liquidus range, ensuring the ability of the liquid slag film to flow downwards along the crystallizer wall into the gap between the billet shell and the copper plate under gravity and friction. If the binary basicity is below 0.8, the acidity of the slag is too strong, which will lead to excessively high overall polymerization of the silicon-oxygen tetrahedral network, excessively high initial viscosity, and reduced mass transfer efficiency for adsorbing inclusions in the molten steel; if the binary basicity is greater than 1.3, the crystallization temperature and crystallization growth rate increase monotonically thermodynamically, the long-range diffusion resistance of the rock-forming ion components decreases, and it is very easy to cross the abrupt change point in the upper region of the crystallizer and crystallize prematurely, which will cause a sharp increase in solid friction and tear the fragile primary billet shell. Within the range of 0.8 to 1.3, the binary basicity is further preferably narrowed to 0.9 to 1.2. This narrower range allows the basicity to form a better thermodynamic complementarity with the high concentration of MnO, further softening the silicon-oxygen framework, improving the diffusion mass transfer coefficient of adsorbed deoxygenated inclusions, and taking into account the appropriate crystallinity and heat transfer control capability of the solid slag film.

[0012] According to one aspect of the invention, more preferably, the flux contained in the matrix slag system, by mass percentage, satisfies the following limitations: when Na2O is contained, the content of Na2O is less than 3%; the content of Li2O is 3% to 5%; when MgO is contained, the content of MgO is not greater than 2%; when fluoride is contained, the content of fluoride, calculated as fluoride ions, is less than 5%.

[0013] This invention firmly pursues low-fluoride and low-sodium content, strictly controlling the fluoride ion content to less than 5% and the Na2O content to less than 3%. This is to reduce the volatilization of fluorides at high temperatures, prevent their corrosion of the crystallizer water jacket and secondary cooling zone components, and avoid pitting corrosion on steel surfaces with extremely high surface requirements caused by the strong alkaline NaOH hydration reaction. To compensate for the rheological losses caused by reducing fluoride and sodium content, this invention uses the more efficient and environmentally friendly flux Li2O. If the Li2O content is below 3%, the macroscopic melting point and viscosity of the slag cannot be effectively reduced, and the hemispherical melting point cannot be precisely controlled within the narrow range of 850℃ to 1000℃. If the Li2O content is above 5%, the excessive alkali metal ions will stimulate and locally recombine to form more complex short-range ordered silicate crystal precursor structures, greatly enhancing the sensitivity of the silicate network to crystalline evolution, leading to an abnormally high viscosity abrupt change temperature, causing the liquid slag to prematurely lose its fluid lubrication ability. The MgO content should be controlled to no more than 2%, primarily for fine-tuning the nucleation of crystalline mineral phases. Exceeding 2% will lead to uncontrolled crystallization temperature and volume fraction of the crystalline phase, impairing slag film performance. When BaO is used as a flux, its content should be controlled to no more than 10%. BaO acts as an alkaline network modifier, and its Ba... 2+Large ionic radius and low field strength can break the bridging oxygen connection of silicon-oxygen tetrahedra to reduce the degree of polymerization and viscosity of melt, and moderately widen the liquidus range. If the BaO content is higher than 10%, the excessive alkaline earth metal cations will cause the slag density and surface tension to increase significantly, and promote the precipitation of high melting point barium-containing mineral phase, which will damage the flow spreading and heat transfer control capabilities of the slag film.

[0014] According to one aspect of the invention, preferably, the protective slag further includes a silicate network depolymerizing agent, wherein the silicate network depolymerizing agent includes TiO2.

[0015] The role and microscopic physicochemical mechanism of introducing the depolymerizing agent TiO2 are as follows: When trace amounts of [Al] inevitably reduce SiO2 during the continuous casting of high-alumina steel, resulting in a large amount of Al2O3 entering the slag, Al ions will replace Si ions under the condition of providing free oxygen and form five-coordinate or four-coordinate bodies that seamlessly integrate into the original silicon-oxygen network. This leads to the consumption of charge balance, an increase in the degree of polymerization of the slag, and an abnormally large increase in viscosity. The introduction of TiO2 in this invention generates a large number of simple octahedral structural units in the silicate slag in the high-temperature molten state. Under the synergistic catalytic effect of the alkali metal oxide flux used in conjunction, these octahedral units with larger volume and unique charge density distribution will act like molecular wedges, forcibly inserting into and blocking the polymerization extension of the complex and huge aluminosilicate network, thereby greatly releasing the fluid units bound by the network and reshaping the melt at the microscopic topology. This powerful microscopic depolymerization effectively counteracts the increase in polymerization degree caused by Al2O3 enrichment, giving the protective slag immunity to viscosity deterioration caused by Al2O3 enrichment. This ensures that the protective slag can still maintain excellent macroscopic fluidity and a uniform and smooth liquid lubricating film when a moderate slag-metal reaction occurs.

[0016] According to one aspect of the present invention, preferably, the continuous casting protective slag further contains a crystallization regulator and a heterogeneous nucleating agent; the crystallization regulator includes Ce2O3; the heterogeneous nucleating agent includes TiN nanoparticles with a size not greater than 100 nm.

[0017] The mechanism of introducing the crystallization regulator Ce2O3 lies in: utilizing the rare earth cation Ce 3+The large ionic radius and unique 4f electron configuration fundamentally reverse the spatial coordination state and charge distribution of surrounding bridging and non-bridging oxygen ions, leading to long-range reorganization of the silicate network topology. This causes a systematic shift in the liquidus temperature, high-temperature apparent viscosity, and latent heat of solidification of the protective slag. As a highly effective endogenous crystallization regulator, it assists in the atomic-scale optimization and suppression of the extremely uneven high heat flux distribution in the crystallizer, forcibly breaking or mitigating the thermodynamic inhibition on traditional crystalline mineral phases caused by the preferential precipitation of fine and dense manganese olivine mediated by high MnO, thus reversing the deterioration of heat transfer. In fluorine-containing systems, Ce2O3 helps restore the competitive advantage of the lanceolate (3CaO·2SiO2·CaF2) crystallization pathway; in fluorine-free systems, Ce2O3 promotes the directional precipitation of calcium silicate mineral phases such as wollastonite at an appropriate rate and morphology, while also ensuring the ability of the solid slag film to regulate the scattering of thermal radiation. TiN nanoparticles with a size no greater than 100 nm are introduced as heterogeneous nucleating agents. During the service of the protective slag, they flow into the gap between the crystallizer wall and the primary billet shell along with the liquid slag and solidify to form a solid slag film. The TiN nanoparticles are embedded in the solid slag film that is closely attached to the outer surface of the billet shell. At the slag-steel interface in the meniscus region, the low interfacial mismatch caused by the very similar lattice constant between TiN and the austenite matrix plays a heterogeneous nucleation role in the solidification and crystallization of the outermost grains of the primary billet shell. This promotes the transformation of the surface region from coarse columnar crystals to finer equiaxed crystals, refines the initial solidification structure of the billet shell surface, and disperses the surface thermal stress among more fine grains, reducing the tendency for crack initiation and propagation.

[0018] According to one aspect of the invention, preferably, the coating material of the first core-shell particle is selected from at least one of amorphous silica and crystalline alumina.

[0019] The physicochemical mechanism for selecting the aforementioned specific coating materials lies in the fact that amorphous silica and crystalline alumina are intrinsic core network components of the matrix slag system. This chemical homology gives the coated hexagonal boron nitride particles excellent surface tension affinity when in contact with high-temperature liquid slag, transforming them from a hydrophobic state that is extremely repulsive to molten slag to a super-wetting state. This allows for nanoscale three-dimensional uniform dispersion in a very short time, eliminating solid lubrication blind zones. Simultaneously, during the storage and transportation of protective slag products and the preheating stage in the crystallizer from room temperature to complete slag melting, the dense silica or alumina shell provides a crucial physical gas barrier for the inner hexagonal boron nitride layer. This prevents contact between atmospheric oxygen and hexagonal boron nitride in the mid-to-high temperature range where oxidation and ablation are most severe, ensuring that the hexagonal boron nitride does not suffer irreversible structural degradation before the slag melts and seals the system. When the core-shell particles melt and enter the liquid slag pool along with the slag material, the hexagonal boron nitride particles are uniformly dispersed in the liquid slag body. The liquid slag layer itself becomes a natural protective environment for hexagonal boron nitride due to its extremely low dissolved oxygen activity. In the solid slag film region at a lower temperature, the hexagonal boron nitride particles are permanently locked in the solidified matrix, and the solidified glassy or crystalline slag body provides dense physical encapsulation protection, stably playing a two-dimensional interlayer slip friction reduction role throughout the entire working life.

[0020] According to one aspect of the present invention, preferably, the core layer material of the second core-shell particle is nanocrystalline magnesium aluminum spinel MgAl2O4, and the coating layer material is dense carbon with graphitization degree.

[0021] The mechanism for selecting this specific microstructure lies in the following: A large number of nanocrystalline spinel grains generated during the high-energy ball milling process form a rigid polycrystalline framework core in a densely aggregated state. A graphitized carbon layer fills the gaps between the spinel grains and coats the periphery of the aggregate, forming a carbon matrix shell. The ultra-high thermal stability of this polycrystalline aggregate core allows it to act as a continuous heat sink in a high-temperature slag surface environment, continuously absorbing and dissipating heat from the oxidation and combustion interface of the outer carbon shell, thus preventing thermal runaway and self-accelerated combustion of the carbon layer. Simultaneously, the graphitized carbon shell formed by high-energy ball milling, with its ordered graphite matrix facing outwards, exhibits a significantly lower density of active edge sites and a higher oxidation initiation temperature compared to amorphous carbon black, endowing the carbon shell with inherent high oxidation resistance and inertness. Crucially, as the carbon shell surface is gradually consumed, the densely packed spinel nanocrystals inside are exposed, forming a rough polycrystalline micro-topography on the carbon shell's retreating surface. The numerous grain protrusions and intergranular depressions create a tortuous surface channel, forcing oxygen molecules to travel a circuitous path to reach the deeper, unreacted carbon surface. This significantly limits the oxygen mass transfer rate at the microscale. The synergistic effect of the heat sink effect, graphitization inertness, and polycrystalline surface oxygen limitation transforms the carbon consumption pattern from explosive volumetric combustion to controlled, gradual surface retreat oxidation. This greatly stabilizes the melting rate fluctuation range of the protective slag during long-cycle casting, ensuring a constant thickness of the liquid slag pool.

[0022] According to one aspect of the present invention, preferably, the continuous casting protective slag is prepared by a method comprising the following steps: mixing and pre-melting matrix slag raw materials weighed according to a mass ratio to obtain a pre-melted matrix powder slag; and uniformly dispersing the nanocomposite components into the pre-melted matrix powder slag.

[0023] According to one aspect of the present invention, preferably, when the nanocomposite component includes the first core-shell particles, the first core-shell particles having an amorphous silica shell are prepared by the following steps: ultrasonically dispersing sheet-like micron-sized hexagonal boron nitride powder in a mixed solution of anhydrous ethanol and ammonia, fixing the pH value at 8.5, and adding tetraethyl orthosilicate dropwise in a constant temperature water bath at 45°C to carry out a hydrolysis and condensation reaction, so that silica heterogeneously nucleates on the surface of hexagonal boron nitride and self-assembles to form a coating layer.

[0024] In the chemical synthesis steps of the first core-shell particles described above, the introduction of anhydrous ethanol as a co-solvent is a necessary condition for the classic sol-gel process. Tetraethyl orthosilicate has extremely low solubility in pure aqueous phase, requiring an alcohol co-solvent to achieve homogeneous hydrolysis and subsequent heterogeneous nucleation and coating. The pH value is strictly fixed at 8.5 because pH is the decisive control valve for the hydrolysis and condensation rate of tetraethyl orthosilicate. If the pH value is below 8.5, the hydrolysis and polymerization rate of tetraethyl orthosilicate is too slow, making nucleation difficult in a short time; if the pH value is above 8.5, the condensation reaction is too vigorous, and silica tends to self-nucleate in solution to form independent monodisperse microspheres, failing to achieve perfect heterogeneous nucleation and self-assembly coating on the surface of the hexagonal boron nitride sheets. The temperature is controlled at a constant temperature water bath of 45°C because 45°C provides a suitable activation energy for the reaction. If the temperature is below 45℃, the hydrolysis reaction will be difficult to proceed completely, and the long-term aging time will be too long; if the temperature is above 45℃, the rapid evaporation of the solvent ethanol will disrupt the liquid phase mass transfer balance, resulting in uneven coating thickness or cracking and peeling. When the nanocomposite component contains the second core-shell particles, the second core-shell particles are prepared by the following steps: basic magnesium carbonate is pretreated to decompose it into magnesium oxide-containing active powder to obtain a magnesium source; the magnesium source and γ-alumina powder are used as precursor raw materials, carbon-containing substances and nano-nickel oxide as mineralizing agents are added, and ball milling is performed using a high-energy mechanical alloying method to generate nano-sized magnesium aluminum spinel through in-situ induced solid-phase reaction by mechanical force, while the carbon-containing substances undergo structural reorganization to form a dense carbon-coated shell layer on the surface of the spinel.

[0025] In the mechanical alloying synthesis of the second core-shell particles, both precursors used were oxidized raw materials—basic magnesium carbonate provided the magnesium source (Mg). 2+ γ-alumina provides the aluminum source (Al). 3+The two reactants directly to form MgAl2O4 spinel through a solid-state diffusion reaction driven by mechanical force. The reaction process involves no redox valence changes, fundamentally avoiding the uncontrollable exothermic reactions and safety hazards associated with using elemental metal precursors. γ-Al2O3 is chosen as the aluminum source instead of α-Al2O3 with its dense corundum structure because: γ-Al2O3 has a defective spinel structure, and its oxygen anion sublattice is cubically close-packed, just like the target product MgAl2O4 spinel. Furthermore, its lattice contains a large number of intrinsic cation vacancies. The transformation to the spinel phase only requires cations to rearrange and migrate within the nearly existing oxygen framework without reconstructing the oxygen framework. The solid-state reaction barrier is much lower than that of the thermodynamically stable and densely lattice-dense α-Al2O3, making it a highly active aluminum source for the mechanochemical synthesis of spinel at room temperature. If α-Al2O3 is used instead, the conversion rate to the spinel phase under the same ball milling conditions is significantly lower, making it difficult to obtain a perfectly crystallized spinel phase within a limited ball milling time. Adding nano-NiO powder directly as a mineralizer instead of aqueous hydroxide can completely avoid moisture release within the sealed ball mill jar and its potential side reactions. 2+ For d 8 The electronic configuration of Ni results in extremely high crystal field stabilization energy in an octahedral coordination field. Its octahedral position preference energy ranks among the highest of common divalent transition metal cations. 2+ When Al enters the forming spinel lattice, it preferentially occupies the octahedral coordination. 3+ Sites, rather than tetrahedral coordinated Mg 2+ Site. This substitution belongs to heterovalent substitution with a double mismatch of valence and size: on the one hand, Ni 2+ The effective octahedral ionic radius of Ni (approximately 0.069 nm) is significantly larger than that of Al³⁺ (approximately 0.054 nm), introducing strong local lattice distortion and strain fields around the substitution sites; on the other hand, divalent Ni… 2+ Occupying trivalent Al 3+ The lattice sites form acceptor-type point defects with negative effective charges, and the lattice must compensate for this charge by generating point defects with positive effective charges, such as oxygen vacancies. Magnesium-aluminum spinel itself belongs to a non-stoichiometric defect structure that readily generates Mg-Al antisite defects and cation vacancies. Under the synergistic effect of the high-density lattice distortion energy injected by repeated welding-fracture during high-energy ball milling and the formation of new interfaces, Ni... 2+ The charge-compensating oxygen vacancies and localized distorted strain fields associated with the occupied sites further exacerbate the degree of lattice distortion and non-stoichiometric deviation, thereby forming highly active diffusion channels rich in point defects such as oxygen vacancies and cation vacancies at the reaction interface. These defect sites greatly accelerate the diffusion of Mg. 2+ And Al 3+The solid-state diffusion rate across the oxide interface significantly reduces the nucleation activation energy of MgAl2O4 spinel. This allows the solid-state reaction, which would normally require high-temperature heat treatment, to be induced in situ in a macroscopic ambient temperature and low-pressure environment under the synergistic enhancement of the highly active γ-Al2O3 precursor and NiO mineralizer, as well as continuous mechanical energy injection, generating a nanocrystalline spinel phase. This results in a large number of nanocrystalline spinel grains. Due to the repeated welding-fracture kinetics of high-energy ball milling, these nanocrystals form polycrystalline clusters in a dense aggregated state, constituting a rigid framework core. Simultaneously, the enormous shear force forcibly reorganizes the lamellar framework of carbonaceous material, causing it to self-roll along the periphery of the nanocrystalline spinel aggregates, filling the intergranular gaps and forming a dense carbon shell with graphitization on the outer surface of the aggregates. If the NiO mineralizer is missing or the ball milling energy is insufficient, the solid-state diffusion rate is insufficient to overcome the high thermodynamic nucleation barrier of spinel, making it impossible to form a well-crystallized spinel phase and a dense, interpenetrating core-shell structure under room temperature ball milling conditions.

[0026] Compared with the prior art, the present invention has the following beneficial effects: This invention enables comprehensive dynamic control of the slag-metal interface, rheological behavior, crystallization kinetics, and heat transfer channels during the continuous casting of high-manganese and high-alumina steels by synergistically introducing specific matrix components into the matrix slag system and dispersing nanocomposite components with specific heterostructures.

[0027] By artificially adding an extremely high proportion of initial MnO to the initial formula, the chemical potential of the manganese component in the liquid slag is increased, which significantly changes the thermodynamic equilibrium constant at the high-temperature interface. This actively inhibits and blocks the positive redox substitution reaction in which the active [Mn] in high-manganese steel takes oxygen atoms from SiO2 in the protective slag. The chemical composition and microstructure of the protective slag can maintain extremely high stability during long-term continuous casting.

[0028] By introducing hexagonal boron nitride core-shell superstructure particles with a shell tightly coated by homologous oxide ceramics (silicon dioxide or alumina), phased sequential protection and full-process functional assurance are achieved. During storage, transportation, and preheating, the dense oxide shell acts as a physical gas barrier, effectively preventing high-temperature oxidation and burn-off of hexagonal boron nitride before the slag melts and seals. When the slag melts, the chemical homology of the oxide shell enables instantaneous superwetting of the core-shell particles, achieving nanoscale three-dimensional uniform dispersion in the liquid slag, completely eliminating the lubrication blind zone caused by the agglomeration of exposed hexagonal boron nitride due to its hydrophobic properties. After the hexagonal boron nitride particles enter the liquid slag, the slag layer itself acts as a natural anti-oxidation barrier due to its extremely low dissolved oxygen activity; in the solid slag film, it is permanently encapsulated by the solidified matrix. The hexagonal boron nitride dispersed in the slag film provides efficient boundary lubrication and friction reduction effects due to the unique interplanar slip characteristics of its layered crystal structure. In addition, the oxide shell imparts random orientation to the hexagonal boron nitride sheets during the initial dispersion stage, breaking the directional parallel stacking of sheet particles under the shear of the kinetic energy flow field. This transforms the unidirectional ultra-fast heat flow dispersion into a highly uniform and isotropic heat-conducting network in three-dimensional space, significantly eliminating local thermal stress concentration in the liquid / solid composite slag film and playing a key role in preventing surface cracking of the nascent shell.

[0029] By introducing second core-shell particles with a rigid framework of nanocrystalline spinel polycrystalline aggregates and a matrix shell of graphitized dense carbon, the carbon oxidation consumption kinetics are transformed from explosive volumetric combustion to a controlled, layer-by-layer, slow oxidation mode. This completely eliminates the explosive macroscopic rapid consumption of traditional free carbon black, resulting in an extremely stable melting rate of the protective slag framework during long-cycle casting. This achieves constant control of the slag pool thickness and eliminates on-site dust emission pollution. The components of this invention work synergistically, addressing the problems of longitudinal cracking and steel leakage caused by lubrication failure and heat transfer disturbances in the continuous casting of high-manganese, high-alumina special steel from multiple dimensions, including underlying thermodynamics, molecular network topology, and multiphase nanofluid rheology. This significantly improves the surface and matrix crystal structure quality of the cast billet. Attached Figure Description

[0030] To more clearly illustrate the technical solutions in the embodiments of the present invention or the prior art, the drawings used in the description of the embodiments will be briefly introduced below. Obviously, the drawings described below are only some embodiments of the present invention. For those skilled in the art, other drawings can be obtained based on these drawings without creative effort.

[0031] Figure 1 This is a transmission electron microscope (TEM) image of the first core-shell particle in an embodiment of the present invention; Figure 2 This is a transmission electron microscope (TEM) image of the second core-shell particle in an embodiment of the present invention; Figure 3 This is an X-ray diffraction (XRD) pattern of the second core-shell particle in an embodiment of the present invention. Detailed Implementation

[0032] The preparation process of the Mn13 high wear-resistant steel continuous casting protective slag containing nanocomposite components involved in this invention can be roughly summarized into three stages: pre-synthesis of nanocomposite components, batching and pre-melting of the matrix slag system, and uniform dispersion and composite of nanocomposite components into the pre-melted slag powder. The specific raw material specifications, operating conditions, and reaction mechanisms of each stage are described in detail below.

[0033] The first core-shell particle in the nanocomposite component, namely the composite particle with hexagonal boron nitride as the core and oxide ceramic as the coating layer, is prepared using a modified sol-gel heterogeneous nucleation method. A typical synthetic route using amorphous silica as the coating material is as follows: Flaky micron-sized hexagonal boron nitride powder (particle size D50 approximately 2μm to 5μm, purity not less than 99.0%, Dandong Xinlong Boron Chemicals Co., Ltd., grade BN-2M) is dispersed in an ultrasonic cleaner (300W power, 40kHz frequency) in a mixture of anhydrous ethanol and deionized water (ethanol to water volume ratio approximately 4:1) for 30 min to form a stable suspension. Subsequently, the pH of the system is adjusted to 8.5 with concentrated ammonia. The suspension is then transferred to a constant-temperature water bath magnetically stirred reactor, with the water bath temperature set at 45℃. Ammonia is added in small amounts to maintain the pH at 8.5. Under continuous stirring (350 r / min), tetraethyl orthosilicate (analytical grade, Sinopharm Chemical Reagent Co., Ltd.) was slowly added dropwise to the reaction system using a peristaltic pump at a mass ratio of hexagonal boron nitride to tetraethyl orthosilicate of 1:0.8 to 1:1.2, with the dropping rate controlled at 0.5 mL / min to 1.0 mL / min. Tetraethyl orthosilicate first dissolved in an alcohol-water co-solution system, then hydrolyzed under alkaline conditions. The resulting silanol intermediate underwent dehydration and condensation polymerization, achieving heterogeneous nucleation on the surface of the hexagonal boron nitride sheets, gradually self-assembling to form a continuous and uniform amorphous silica coating layer. The total reaction time was approximately 6 to 8 hours. After the reaction, the particles were centrifuged, washed three times with anhydrous ethanol, and vacuum dried at 60°C for 12 hours to obtain the first core-shell particle powder with a complete core-shell structure. Figure 1As shown, transmission electron microscopy characterization reveals that the first core-shell particle exhibits a clear core-shell structure with hexagonal boron nitride as the core and oxide ceramic as the shell. The coating thickness is approximately 15 nm to 30 nm, and it is in a continuous and dense state. This oxide shell provides physical gas barrier protection for hexagonal boron nitride during the storage and transportation of protective slag products and the preheating stage in the crystallizer, preventing it from being subjected to atmospheric oxidation and ablation in the medium-high temperature range (approximately 400℃ to 900℃). At the same time, the chemical homology between the shell material and the matrix silicate slag enables the core-shell particles to achieve instantaneous superwetting and nanoscale uniform dispersion upon contact with high-temperature liquid slag. When crystalline alumina is used as the coating layer, aluminum isopropoxide is used as the precursor to achieve hydrolytic deposition in a similar alcohol-water cosol gel system. The typical steps are as follows: Take the above-mentioned sheet-like micron-sized hexagonal boron nitride powder of the same specifications and ultrasonically disperse it in isopropanol for 30 min to form a stable suspension; weigh aluminum isopropoxide at a boron nitride to aluminum isopropoxide mass ratio of 1:0.8 to 1:1.2, dissolve it in an appropriate amount of isopropanol, and stir and reflux at an 80℃ water bath for 30 min to allow for complete depolymerization. Then, slowly add the aluminum isopropoxide solution dropwise to the boron nitride suspension; under constant temperature water bath conditions of 75℃ to 80℃ and continuous stirring (350 r / min), add a mixture of isopropanol and deionized water at a volume ratio of approximately 10:1 dropwise to the system to initiate controlled hydrolysis of aluminum isopropoxide, with the drop rate controlled between 0.5 mL / min and 1.0 mL / min. The system was soluble at a flow rate of mL / min and the pH of the system was adjusted to 4 to 5 with dilute nitric acid. The boehmite (AlOOH) sol generated by hydrolysis was heterogeneously nucleated and uniformly deposited on the surface of the hexagonal boron nitride sheets. The sol was aged for 2 to 4 hours. The resulting powder was centrifuged, washed three times with isopropanol, dried at 80°C, and then calcined at 550°C for 2 hours to dehydrate the surface coating layer and transform it into the γ-alumina crystalline phase. This yielded the first core-shell particles with crystalline alumina as the shell layer, with a shell thickness of approximately 15 nm to 30 nm.

[0034] The second core-shell particles were prepared using a high-energy mechanical alloying method. First, basic magnesium carbonate (4MgCO3·Mg(OH)2·5H2O, analytical grade, Tianjin Guangfu Fine Chemical Research Institute) was placed in a vacuum drying oven and pretreated at 300℃ for 4 hours to partially decompose it into a mixed active powder mainly composed of MgO and residual MgCO3. At the same time, water of crystallization and adsorbed water were removed to avoid free water interfering with the solid-phase reaction during ball milling. The pretreated magnesium source powder and γ-alumina powder (purity 99.5%, specific surface area approximately 150 m² / g to 250 m² / g, Shanghai Maclean Biochemical Technology Co., Ltd.) were mixed according to the stoichiometric ratio of magnesium aluminum spinel MgAl2O4 (MgO:Al2O3 molar ratio of 1:1). Graphite powder (carbon content not less than 99.0%, particle size approximately 10 μm, Qingdao Tianshengda Graphite Co., Ltd.) and nano-nickel oxide NiO powder (particle size approximately 20 nm to 50 nm, purity 99.0%, Aladdin Reagent Co., Ltd.) as mineralizing agents were also added. The amount of NiO added was 2% to 3% of the total mass of the precursors, and the amount of graphite powder added was 15% to 25% of the total mass of the precursors. γ-alumina was chosen as the aluminum source instead of α-alumina because γ-Al₂O₃ possesses a cubic close-packed defect spinel-type oxygen framework homologous to MgAl₂O₄ spinel and is intrinsically rich in cation vacancies. The transformation to the spinel phase requires only cation rearrangement without reconstructing the oxygen framework. Its room-temperature mechanochemical reactivity is far higher than that of dense corundum-structured α-Al₂O₃, allowing it to synergistically interact with the newly generated active MgO and NiO mineralizers from the pretreatment process, enabling the spinel phase to fully crystallize within a ball milling time of 20 to 30 hours. The use of nano-NiO powder instead of an aqueous hydroxide precursor completely avoids side reactions and pressure increases caused by in-situ water release within the sealed ball mill jar. The above raw materials were loaded into a stainless steel ball mill jar in a high-purity argon glove box, with a ball-to-material ratio of 20:1. The milling media were tungsten carbide cemented carbide balls. The ball mill jar was equipped with a safety pressure relief valve to release the small amount of CO₂ generated from the decomposition of residual carbonate ions in basic magnesium carbonate. The sealed ball mill jar was transferred to a planetary high-energy ball mill (Nanjing University Instrument Factory, QM-3SP4 type) and ball-milled at 400 r / min. An argon protective atmosphere was maintained throughout the process to prevent graphite oxidation. The total ball milling time was 20 to 30 hours, with a 10-minute pause every 30 minutes to control the macroscopic average temperature inside the jar and prevent large-area sintering and agglomeration of the powder. During ball milling, intense mechanical impact continuously injected high-density lattice distortion energy and micro-interfacial energy into the precursor powder. Residual magnesium carbonate was further decomposed into nano-sized, highly active MgO particles under mechanical force. The nano-NiO mineralizer was thoroughly mixed with the precursor and embedded in the reaction interface during ball milling. 2+ For d 8The electronic configuration of Ni results in extremely high crystal field stabilization energy in an octahedral coordination field. Its octahedral position preference energy ranks among the highest of common divalent transition metal cations. 2+ Al preferentially octahedral coordinated when entering the forming spinel lattice 3+ Sites, rather than tetrahedral coordinated Mg 2+ Site; Ni 2+ The effective octahedral ionic radius (approximately 0.069 nm) is significantly larger than that of Al. 3+ The octahedral effective ionic radius (approximately 0.054 nm) introduces strong local lattice distortion around the substitution sites, and the divalent Ni... 2+ Occupying trivalent Al 3+ The lattice sites form acceptor-type point defects with negative effective charges, and the lattice must compensate for this charge by generating point defects with positive effective charges, such as oxygen vacancies. Magnesium aluminum spinel itself belongs to a non-stoichiometric defect structure that readily generates Mg-Al antisite defects and cation vacancies. Under the synergistic effect of the high-density lattice distortion energy injected by high-energy ball milling and the newly formed interface, Ni... 2+ The charge-compensating oxygen vacancies and localized distorted strain fields associated with the occupancy sites further exacerbate the degree of lattice distortion and non-stoichiometric deviation, thereby forming highly active diffusion channels rich in point defects such as oxygen vacancies and cation vacancies at the reaction interface. These defect sites greatly accelerate the diffusion of Mg. 2+ And Al 3+ The solid-state interdiffusion rate across the MgO / Al2O3 interface significantly reduces the nucleation activation energy of MgAl2O4 spinel. This allows the solid-state reaction MgO + Al2O3 → MgAl2O4, which normally requires high-temperature heat treatment above 1200℃, to be induced in situ at room temperature in a macroscopic environment, generating a nanocrystalline spinel phase. This is achieved through the synergistic enhancement of the highly reactive oxygen framework of γ-Al2O3, the NiO mineralizer, and continuous mechanical energy injection. This results in a large number of nanocrystalline spinel grains with a particle size of approximately 20 nm to 50 nm. Due to the repeated welding-fracture kinetics of high-energy ball milling, these nanocrystals form densely aggregated polycrystalline clusters, constituting a rigid framework core. Simultaneously, the enormous shear force forcibly reorganizes the lamellar framework of graphite powder, causing it to bend and roll around the periphery of the nanocrystalline spinel aggregates, filling the intergranular gaps and forming a dense carbon shell with a certain degree of graphitization on the outer surface of the aggregates. Figure 2 High-resolution transmission electron microscopy characterization shows that the obtained second core-shell particles exhibit a composite morphology with polycrystalline spinel clusters as the core and a carbon matrix as the shell, with a total carbon shell thickness of approximately 5 nm to 15 nm; Figure 3 The X-ray diffraction pattern is shown. Figure 3The horizontal axis represents the diffraction angle 2θ in degrees (°), and the vertical axis represents the diffraction intensity in arbitrary units (au). The test used Cu Kα radiation (wavelength 0.15406 nm). The product exhibited a set of sharp diffraction peaks at approximately 19.0°, 31.3°, 36.9°, 44.8°, 55.7°, 59.4°, and 65.2° of 2θ, corresponding to the (111), (220), (311), (400), (422), (511), and (440) crystal planes of cubic MgAl₂O₄ spinel (space group Fd-3m, JCPDS standard card No. 21-1152), respectively. This confirms that the spinel core has crystallized perfectly and maintains a well-preserved cubic structure. Each spinel diffraction peak showed significant nanoscale broadening, with grain sizes estimated by the Scherrer formula to be 20 nm to 50 nm. Figure 2 The nanocrystal sizes observed by transmission electron microscopy are consistent. The broadened diffuse peaks located around 2θ (approximately 26°) are attributed to the (002) crystal plane diffraction of the disordered layered graphitized carbon, corresponding to the carbon shell layer encapsulating the outer surface of the spinel aggregates. No independent diffraction peaks for NiO were detected in the spectrum, indicating that the NiO mineralizer has been completely dissolved into the spinel lattice, rather than existing as a free second phase. If the NiO mineralizer is missing or the ball milling energy is insufficient, the solid-state diffusion rate is insufficient to overcome the high thermodynamic nucleation barrier of spinel, making it impossible to form a well-crystallized spinel phase and a dense core-shell encapsulation structure under room temperature ball milling conditions.

[0035] The matrix slag system was prepared using analytical grade oxide powders as the main raw material. Calcium oxide was obtained by calcining analytical grade calcium carbonate (Sinopharm Chemical Reagent Co., Ltd.) at 1100℃ for 4 h, and immediately sealed after calcination to prevent moisture absorption and carbonation. Silica was prepared using high-purity quartz powder (purity above 99.5%, Lianyungang Donghai Silica Powder Co., Ltd., particle size less than 75μm). Manganese oxide was prepared using chemically pure MnO powder (purity above 98.0%, Shanghai Maclean Biochemical Technology Co., Ltd.). Alumina was prepared using analytical grade α-Al2O3 powder (Sinopharm Chemical Reagent Co., Ltd.). Lithium oxide (Li2O) was prepared by decomposing lithium carbonate at 780℃ or directly using analytical grade lithium oxide powder. Fluorides were sourced from analytical grade calcium fluoride. The remaining fluxes, Na2O, MgO, and BaO, were all analytical grade reagents. After accurately weighing each component according to the designed mass ratio, mix and grind them thoroughly in an agate mortar, transfer them to a high-purity graphite crucible, and heat them to 1350°C to 1400°C in a high-temperature box-type resistance furnace heated by a silicon molybdenum rod at a heating rate of 5°C / min. Hold the temperature for 30 to 45 minutes to allow the material to be completely melted and homogenized. Then, quickly pour the melt into a stainless steel water-cooled mold and rapidly cool it into a glassy state. After coarse crushing by a jaw crusher, grind it in a vibrating ball mill until the particle size is less than 150μm to obtain a pre-melted matrix powder slag.

[0036] The pre-prepared nanocomposite components (first core-shell particles and / or second core-shell particles) and other functional additives (TiO2, Ce2O3, TiN nanoparticles, etc., all added externally, based on 100 parts by weight of the matrix slag system) are weighed according to the set ratio and loaded together with the pre-melted matrix slag powder into a V-type mixer. Dry mixing is carried out at 25 r / min for 2 to 3 hours to achieve macroscopic uniform dispersion of the nanocomponents in the matrix slag powder, thus obtaining the final continuous casting protective slag product. The TiO2 used is analytical grade anatase titanium dioxide (99.0% purity, Aladdin Reagent Co., Ltd.), the Ce2O3 is rare earth grade cerium oxide powder (99.9% purity, Baotou Ruixin Rare Earth Functional Materials Co., Ltd.), and the TiN nanoparticles are high-purity titanium nitride nanoparticles with a particle size of less than 100 nm (99.0% purity, Shanghai Chaowei Nanotechnology Co., Ltd.).

[0037] The performance testing methods and conditions involved in the following embodiments and comparative examples are described below. The initial viscosity and the viscosity after reaction were determined using a rotating cylinder method high-temperature viscometer (Bahr Thermoanalyse GmbH, Germany, VIS 403 model). The test temperature was 1300℃. Approximately 120 g of protective slag sample was placed in a molybdenum crucible. Under an argon protective atmosphere, the temperature was increased to 1400℃ at 5℃ / min and held for 10 min to completely melt the sample and eliminate bubbles. Then, the temperature was decreased to 1300℃ at 3℃ / min and stabilized for 5 min before the viscosity was read. The rotation speed was 200 r / min. The method for determining the viscosity change rate after the slag-gold reaction is as follows: 120 g of protective slag sample was pre-added to the molybdenum crucible of the viscometer and completely melted at 1300℃. Then, 20 g of pre-melted Mn13 high-manganese steel liquid (chemical composition by mass percentage: C 1.15%, Mn 12.8%, Si 0.45%, Cr 0.35%, balance Fe) was added above the slag surface. After reacting for 30 min under a weak argon protective atmosphere, the viscosity of the slag was measured again at 1300℃, and the change rate was calculated by comparing it with the initial viscosity. The SiO2 mass fraction loss rate was obtained by quantitatively analyzing the SiO2 content of the residue sample after the above slag-gold reaction experiment using an X-ray fluorescence spectrometer (Shimadzu XRF-1800, Japan) and comparing it with the initial design value. The hemispherical melting point was determined according to GB / T 219-2008 standard using an ash fusion property tester (Changsha Sande Instrument Equipment Co., Ltd., SDAF4000 model) under a weakly reducing atmosphere. Crystallization temperature and crystallinity were determined using differential scanning calorimetry (NETZSCH DSC 404 F3 model, Germany). Under argon protection, the temperature was cooled from 1400℃ to 600℃ at a rate of 20℃ / min. The starting temperature of the exothermic peak on the cooling curve was taken as the crystallization temperature, and the crystallinity was calculated by the ratio of the exothermic peak area to the peak area of ​​a known standard sample. The variation range of the liquid-slag pool thickness was determined on a laboratory simulated crystallizer: a copper water-cooled sleeve with an inner diameter of 60 mm was used. Under a holding temperature of 1300℃, protective slag powder was continuously replenished to the slag layer surface to simulate the actual casting process. The liquid-slag layer thickness was measured every 5 minutes using a heat-resistant steel probe method, continuously for 60 minutes. The difference between the maximum and minimum values ​​was taken as the variation range. The coefficient of friction of the slag film was determined by referring to the Bikerman interfacial friction test principle and was completed on a customized high-temperature reciprocating friction test device: a pre-melted slag film with a thickness of about 3 mm was sandwiched between a copper plate sample and a carbon steel sample, a normal load of 0.5 MPa was applied at 900℃, and the sample was reciprocated at a frequency of 0.5 Hz. The steady-state friction force was recorded and converted into the coefficient of friction.The surface crack index of the billet was determined on an industrial-grade continuous casting intermediate test platform. Mn13 high manganese steel was used as the casting steel, with a cross section of 150 mm × 150 mm and a casting speed of 0.8 m / min. The total number of longitudinal and transverse cracks visible on the surface of the billet per meter after exiting the crystallizer was counted. Example 1

[0038] The matrix slag system composition in this embodiment, by mass percentage, is as follows: CaO 36%, SiO2 33%, MnO 18%, Al2O 37%, Li2O 4%, Na2O 0.5%, MgO 1%, CaF2 0.5% (approximately 0.24% as F⁻). The binary basicity of the matrix slag system, CaO / SiO2, is 1.09. Based on 100 parts by mass of the matrix slag system, the following are added: 2 parts TiO2, 1 part Ce2O3, 0.3 parts TiN nanoparticles (approximately 60 nm in diameter), 1.5 parts of a first core-shell particle with amorphous silica as the shell, and 2.0 parts of a second core-shell particle (MgAl2O4@C) with graphitized carbon as the shell, for a total of 3.5 parts of nanocomposite components. The first core-shell particles were prepared using the aforementioned sol-gel method, with a hexagonal boron nitride to tetraethyl orthosilicate mass ratio of 1:1.0 and a reaction time of 7 h. The second core-shell particles were prepared using the aforementioned high-energy ball milling method, with a milling time of 24 h, a NiO addition of 2.5% of the total precursor mass, and a graphite powder addition of 20% of the total precursor mass. The matrix slag raw material was pre-melted at 1380℃ for 40 min, then water-quenched and ground to a particle size of less than 150 μm. It was then mixed with the aforementioned nanocomponents and functional additives in a V-type mixer for 2.5 h to obtain the finished protective slag. Example 2

[0039] This embodiment lowered the alkalinity and increased the MnO content to verify the synergistic effect between the low alkalinity end and the high MnO end. The composition of the matrix slag system by mass percentage is: CaO 32%, SiO2 38%, MnO 19%, Al2O3 6%, Li2O 3%, Na2O 0.5%, MgO 0.5%, CaF2 1% (approximately 0.49% as F⁻). The binary alkalinity of the matrix slag system, CaO / SiO2, is 0.84. Based on 100 parts by mass of the matrix slag system, the following components were added: 3 parts TiO2, 0.8 parts Ce2O3, 0.5 parts TiN nanoparticles (approximately 80 nm in diameter), 2.0 parts of the first core-shell particles with crystalline alumina as the shell, and 1.5 parts of the second core-shell particles (MgAl2O4@C) with graphitized carbon as the shell, for a total of 3.5 parts of nanocomposite components. The alumina shell of the first core-shell particles was prepared by hydrolytic deposition of aluminum isopropoxide followed by calcination at 550℃ for 2 h to obtain a γ-alumina crystalline phase, with a shell thickness of approximately 20 nm. The second core-shell particles were ball-milled for 22 h, with NiO added at 3% of the total precursor mass and graphite powder added at 18% of the total precursor mass. The pre-melting temperature was 1390℃, and the holding time was 35 min. This formulation had a relatively strong acidity and a relatively high initial viscosity; therefore, the TiO2 content was increased to 3 parts to enhance network depolymerization, and the amount of TiN added was further increased to strengthen the equiaxed crystalline nuclei on the surface. Example 3

[0040] This embodiment pushes the basicity to the higher upper limit while reducing MnO to near the lower limit to examine the comprehensive behavior of the high basicity formulation under low manganese reserve conditions. The matrix slag system composition by mass percentage is: CaO 42%, SiO2 33%, MnO 15%, Al2O3 6%, Li2O 3%, Na2O 0.5%, MgO 0.5%. No fluorides were added in this group, and fluxing was achieved entirely by Li2O. The binary basicity of the matrix slag system, CaO / SiO2, is 1.27. Based on 100 parts by mass of the matrix slag system, the following were added: 1.5 parts TiO2, 1.5 parts Ce2O3, 0.2 parts TiN nanoparticles (approximately 50 nm in diameter), 1.0 part of the first core-shell particles with amorphous silica as the shell, and 2.5 parts of the second core-shell particles (MgAl2O4@C) with graphitized carbon as the shell, for a total of 3.5 parts of nanocomposite components. The pre-melting temperature was 1400℃, and the holding time was 45 min. Due to the high basicity and low MnO content, the amount of Ce2O3 in this embodiment was increased to 1.5 parts to enhance the control over crystallization kinetics and prevent the solid slag film from becoming too thick due to the increased crystallization temperature in the high basicity system. Under fluorine-free conditions, Ce2O3 mainly promotes the directional precipitation of calcium silicate mineral phases such as wollastonite at a suitable rate and morphology, replacing the heat control function of gunmetal in the fluorine-containing system. Example 4

[0041] This embodiment selects proportions close to the optimal central values ​​for each component to achieve the best overall performance. The composition of the matrix slag system by mass percentage is: CaO 37%, SiO2 34%, MnO 17%, Al2O 38%, Li2O 3.5%, MgO 0.5%. No Na2O or fluorides are used, achieving a completely fluorine-free and sodium-free formulation. The binary basicity of the matrix slag system, CaO / SiO2, is 1.09. Based on 100 parts by mass of the matrix slag system, the following are added: 2.5 parts TiO2, 1.2 parts Ce2O3, 0.4 parts TiN nanoparticles (approximately 50 nm in diameter), 2.0 parts of a first core-shell particle with amorphous silica as the shell, and 2.0 parts of a second core-shell particle (MgAl2O4@C) with graphitized carbon as the shell, for a total of 4.0 parts of nanocomposite components. In the preparation of the first core-shell particles, the mass ratio of hexagonal boron nitride to tetraethyl orthosilicate was 1:1.1, and the reaction time was 8 h. The ball milling time for the second core-shell particles was extended to 28 h to obtain a denser carbon shell. The pre-melting temperature was 1380℃, and the holding time was 40 min. In this embodiment, the ratio of TiO2 to Al2O3 was also finely adjusted to 2.5:8 to achieve an optimal balance in the competitive intercalation between titanium oxide octahedra and aluminum oxide tetrahedra. Example 5

[0042] This embodiment uses only the first core-shell particles without adding the second core-shell particles to verify the contribution of a single type of nanocomposite component in lubrication and heat transfer. The melting control function of carbon is achieved solely by externally added ordinary carbon black (Tianjin Yiborui Chemical Co., Ltd., N330 type, particle size approximately 30 nm), but carbon black does not possess a core-shell constrained structure. The matrix slag system composition, by mass percentage, is: CaO 38%, SiO2 35%, MnO 16%, Al2O 37%, Li2O 3%, Na2O 0.5%, MgO 0.5%. The binary basicity of the matrix slag system, CaO / SiO2, is 1.09. Based on 100 parts by weight of the matrix slag system, the following components were added: 2 parts TiO2, 1 part Ce2O3, 0.3 parts TiN nanoparticles (approximately 70 nm in diameter), 3.0 parts of the first core-shell particles coated with a double layer of amorphous silica and crystalline alumina, and 2.5 parts carbon black (replacing the second core-shell particles). The total amount of nanocomposite components was 3.0 parts (considering only the core-shell particles). The pre-melting temperature was 1370℃, and the holding time was 35 min. Example 6

[0043] This embodiment pushes the formulation to the extreme regions of low basicity, high MnO, and high Li2O to investigate the tolerance behavior of the protective slag under multiple superimposed boundary conditions. The matrix slag system composition, by mass percentage, is: CaO 30%, SiO2 37%, MnO 20%, Al2O 36%, Li2O 5%, Na2O 1%, MgO 1%. The binary basicity of the matrix slag system, CaO / SiO2, is 0.81. Based on 100 parts by mass of the matrix slag system, the following components are added: 2 parts TiO2, 0.5 parts Ce2O3, 0.5 parts TiN nanoparticles (approximately 90 nm in diameter), 1.0 part of a first core-shell particle with an amorphous silica shell, and 3.0 parts of a second core-shell particle (MgAl2O4@C) with a graphitized carbon shell, for a total of 4.0 parts of nanocomposite components. The pre-melting temperature is 1360℃, and the holding time is 30 min. Since the Li2O content in this embodiment has reached the upper limit of 5%, the modification effect of alkali metal ions on the silicate network tends to be saturated. In terms of crystallization control, it mainly relies on the self-inhibition balance of high MnO concentration on manganese olivine precipitation, and the amount of Ce2O3 is reduced accordingly.

[0044] Comparative Example 1 The matrix slag system in this comparative example does not contain MnO and adopts the traditional CaO-SiO2-Al2O3 system formulation design approach. The matrix slag system composition, by mass percentage, is: CaO 40%, SiO2 40%, Al2O3 8%, Li2O 4%, Na2O 2%, MgO 1%, CaF 25% (approximately 2.44% as F⁻). The binary basicity of the matrix slag system, CaO / SiO2, is 1.00. Based on 100 parts by mass of the matrix slag system, 2 parts TiO2, 1 part Ce2O3, 0.3 parts TiN nanoparticles, 1.5 parts of the first core-shell particles, and 2.0 parts of the second core-shell particles are added. The preparation conditions of the nanocomposite components are exactly the same as in Example 1. This comparative example investigates the slag-metal reaction behavior and viscosity stability of the protective slag after contact with molten Mn13 steel under conditions lacking the initial chemical potential suppression of high-concentration MnO.

[0045] Comparative Example 2 The matrix slag formulation of this comparative example is exactly the same as that of Example 1 (CaO 36%, SiO2 33%, MnO 18%, Al2O 37%, Li2O 4%, Na2O 0.5%, MgO 1%, CaF2 0.5%), with the addition of 2 parts TiO2, 1 part Ce2O3, and 0.3 parts TiN in the same proportions as in Example 1, but without the use of any core-shell structured nanocomposite components. Instead, 1.5 parts of uncoated raw flake boron nitride powder and 2.0 parts of ordinary carbon black (same specifications as in Example 5) are added as conventional solid lubricants and carbonaceous fluxes. This comparative example is used to verify the core contribution of the core-shell structure design to lubrication performance, carbon consumption kinetics, and billet surface quality.

[0046] Comparative Example 3 The formulation of this comparative example is based on Example 1, but the TiO2 component has been removed. The composition of the matrix slag system is completely consistent with that of Example 1. Based on 100 parts by weight of the matrix slag system, 1 part Ce2O3, 0.3 parts TiN, 1.5 parts of the first core-shell particles, and 2.0 parts of the second core-shell particles are added, but TiO2 is not added. This comparative example aims to reveal the effect of the lack of a TiO2 network depolymerizer on the degree of polymerization and viscosity evolution behavior of the slag when the slag-gold reaction leads to the gradual enrichment of Al2O3 in the slag.

[0047] Comparative Example 4 In this comparative example, the MnO content was reduced to below the lower limit of 15%. The composition of the matrix slag system, by mass percentage, was: CaO 38%, SiO2 36%, MnO 10%, Al2O 38%, Li2O 4%, Na2O 1.5%, MgO 1%, CaF2 1.5% (approximately 0.73% as F⁻). The binary basicity of the matrix slag system, CaO / SiO2, was 1.06. The proportions of nanocomposite components and functional additives were exactly the same as in Example 1, based on 100 parts by mass of the matrix slag system. By setting the MnO content to only 10% (below the lower limit of 15%), the degree of deterioration of the slag-metal interface reaction under insufficient manganese chemical potential suppression was verified.

[0048] All the protective slag samples from the above embodiments and comparative examples were evaluated for system performance according to the aforementioned unified test method. The test results for each group are summarized in Tables 1 and 2.

[0049] Table 1. Composition of the matrix slag system and process parameters of each embodiment and comparative example.

[0050] Table 2 Key performance test results of each embodiment and comparative example

[0051] As can be seen from the test data in Table 2, the viscosity change rate after the slag-gold reaction in all six examples was controlled within 12%, the SiO2 mass fraction loss rate did not exceed 3%, the thickness fluctuation range of the liquid slag pool was between ±0.2 mm and ±0.8 mm, the slag film friction coefficient was in the low range of 0.034 to 0.049, and the surface crack index of the cast billet did not exceed 1.4 cracks / m, demonstrating excellent and stable comprehensive metallurgical performance. Among them, Example 4 had the most outstanding comprehensive performance because the components were in the optimal matching range, with a viscosity change rate of only 5.8%, a SiO2 loss rate of only 1.5%, a slag film friction coefficient as low as 0.034, and a crack index of only 0.4 cracks / m. Example 3 maintained an ideal crystallization temperature and a high crystallization rate under high alkalinity and fluorine-free conditions with a relatively high Ce2O3 content, verifying the adaptive adjustment capability of the system of the present invention within a wide alkalinity range. Although Example 6 was conducted under boundary conditions of lowest alkalinity and strongest acidity, resulting in a relatively high initial viscosity (0.25 Pa·s), the viscosity change rate and SiO2 loss rate after the reaction remained within acceptable limits, indicating that the initial MnO reserve of up to 20% had a significant reverse inhibitory effect on the slag-gold reaction. Example 5, lacking the second core-shell particles and replacing them with ordinary carbon black, exhibited significantly worse fluctuations in the liquid slag pool thickness compared to the other five examples containing both types of core-shell particles, deteriorating from ±0.2 to ±0.5 mm to ±0.8 mm. However, due to the increased amount of the first core-shell particles to 3.0 parts, its friction coefficient remained at a low level, indicating that the two types of core-shell particles each have their own functional focus and complement each other synergistically.

[0052] In contrast, all four comparative examples exhibited severe performance degradation in different dimensions. Comparative Example 1, which contained no MnO, showed a SiO2 mass fraction loss of up to 18.5% after 30 minutes of contact with molten Mn13 steel. A violent displacement redox reaction occurred at the slag-metal interface—the high concentration of [Mn] in the molten steel, driven by a large chemical potential, extracted bridging oxygen atoms from the silicon-oxygen network of the protective slag, leading to a large-scale collapse of the silicon-oxygen tetrahedral framework. Due to the significant consumption of SiO2, the core network formant, the slag composition deviated significantly from the originally designed low-viscosity eutectic range. Simultaneously, trace amounts of [Al] in the molten steel reduced SiO2 during the reaction, and the generated Al2O3, as an amphoteric oxide, acted as a network formant in the alkaline slag in the form of tetracoordinated aluminum, further enhancing the polymerization degree of the residual slag. The rapid loss of SiO2 and the massive enrichment of Al2O3 resulted in a viscosity surge of 86.3%, nearly causing the slag to lose its flowability. Although the MnO generated in the reaction itself, as an alkaline network modifier, has a certain depolymerization and viscosity-reducing effect, this localized effect is far from sufficient to offset the catastrophic viscosity increase caused by the overall collapse of the network skeleton and the re-polymerization of Al2O3. The influx of a large amount of exogenous MnO also drastically altered the crystallization pathway, causing the crystallization temperature to plummet to 982℃ and the crystallization rate to only 26%. This means that the solid slag film is extremely thin and discontinuous, significantly weakening the thermal radiation shielding effect, and causing the surface crack index of the cast billet to soar to 5.3 cracks / m. This comparison fully demonstrates the irreplaceable role of pre-positioning a high concentration of initial MnO in the matrix slag system for passivating the slag-gold interface reaction from a thermodynamic source.

[0053] The matrix slag system and functional additive formulation of Comparative Example 2 were completely identical to those of Example 1, with the only difference being that uncoated exposed hexagonal boron nitride and ordinary carbon black replaced the first and second core-shell particles, respectively. The differences in SiO2 loss rate and viscosity change rate were minimal, indicating that the inhibition of the slag-metal reaction was mainly borne by MnO in the matrix slag system, rather than the core-shell components. However, Comparative Example 2 showed a significant deterioration in two indicators: lubrication friction and carbon consumption control. The slag film friction coefficient nearly doubled, increasing from 0.042 to 0.083. This was because the exposed hexagonal boron nitride, due to its strong hydrophobic properties, severely agglomerated in the silicate slag, failing to achieve three-dimensional uniform dispersion, resulting in large-area lubrication blind zones within the slag film. Simultaneously, the uncoated hexagonal boron nitride suffered severe atmospheric oxidation and erosion during the preheating stage of the protective slag from room temperature to slag melting, significantly shortening its effective lubrication life. The thickness fluctuation of the slag pool worsened sharply from ±0.3 mm to ±2.6 mm. This was because ordinary carbon black, lacking core-shell constraint in the high-temperature, oxygen-rich environment of the slag surface, underwent explosive and rapid combustion and consumption. The instantaneous consumption of a large amount of carbon caused a sudden thinning of the slag layer in some areas, which was subsequently restored by the replenishment of powder. This resulted in drastic periodic fluctuations in the melting rate and the thickness of the slag pool. The crack index increased from 0.8 cracks / m to 3.9 cracks / m, fully demonstrating the core guarantee role of the core-shell structure design in ensuring lubrication uniformity, carbon consumption kinetic stability, and billet surface quality.

[0054] Comparative Example 3 removed the TiO2 depolymerizing agent. Before the slag-gold reaction, its initial viscosity, hemispherical melting point, crystallization temperature, and crystallinity were almost identical to those of Example 1. This is because the role of TiO2 is not manifested in the initial state, but rather in the dynamic process of Al2O3 gradually enriching during the slag-gold reaction. After 30 min of reaction, the viscosity change rate of Comparative Example 3 was as high as 43.5%, far exceeding the 8.2% of Example 1. This result clearly reveals its underlying mechanism: when trace amounts of [Al] inevitably reduce some of the SiO2 in the protective slag, and the generated Al2O3 enters the slag layer, aluminum ions seamlessly integrate into the silicon-oxygen network in a four- or five-coordinated manner, playing the role of network forming organisms, leading to a sharp increase in the degree of polymerization of the molten slag and an abnormal increase in viscosity. In Example 1, the large number of simple titanium-oxygen octahedral structural units provided by TiO2, under the synergistic catalysis of alkali metal oxide Li2O, forcefully insert into and cut off the polymerization extension of the aluminum-silicon-oxygen network in the form of molecular wedges, effectively releasing the bound fluid units and controlling the viscosity change to an acceptable level. In contrast, Example 3 lacked this microscopic depolymerization mechanism, and the degree of polymerization of the aluminosilicate network increased unchecked. Although the total loss of SiO2 was comparable to that of Example 1 (2.3% vs. 2.1%), the rheological behavior of the slag had deteriorated significantly, ultimately manifested as a higher coefficient of friction and an increase in the crack index to 2.6 cracks / m.

[0055] Comparative Example 4 reduced the MnO content to 10%. Although it did not completely remove MnO as in Comparative Example 1, this concentration was insufficient to thermodynamically suppress the active [Mn] in the high-manganese steel by a sufficient reverse chemical potential. After the slag-gold reaction, the SiO2 loss rate climbed to 9.8%, the viscosity change rate reached 36.2%, the crystallization temperature dropped to 1018℃, and the crystallization rate fell to 31%. Compared with the groups in the examples where the MnO content was in the range of 15% to 20%, the degree of performance degradation showed an approximately positive correlation with the magnitude of MnO deficiency—the SiO2 loss rate was only 2.1% when MnO was 18% in Example 1, and only 1.9% when MnO was 15% in Example 3 (high basicity helps to further inhibit the reaction), while the SiO2 loss was close to double digits when MnO was only 10% in Comparative Example 4. This set of gradient data indicates that there is a critical threshold for MnO concentration. When it is below this threshold, the chemical potential of the manganese component in the slag is no longer able to resist the erosion of the silicon-oxygen network of the protective slag by [Mn] in the molten steel. The slag-metal reaction gets out of control, leading to a chain reaction of subsequent viscosity changes, crystallization path shifts, and deterioration of billet quality.

[0056] Based on the data comparison and trend analysis of the above embodiments and comparative examples, the technical effects of the present invention can be explained in depth from the perspectives of multi-level microstructure and reaction thermodynamics. At the slag-metal interface reaction level, the present invention utilizes Le Chatelier's principle to pre-establish a high-concentration MnO chemical potential barrier in the matrix slag system, resulting in a significant thermodynamic reduction of the equilibrium constant of the substitution reaction at the interface. MnO plays the role of a network modifier in the CaO-SiO2-MnO ternary slag system—its Mn... 2+ Ions embed themselves into non-bridging oxygen sites in a six-coordinated octahedral form, locally disrupting the bridging oxygen connections of silicon-oxygen tetrahedra, moderately reducing network connectivity and viscosity. Simultaneously, the increased activity of MnO in the slag causes the Gibbs free energy change of the forward reaction [Mn] + (SiO2) → (MnO) + [Si] to approach zero or become positive, suppressing SiO2 consumption at its source. When the MnO content is below 15% (e.g., 10% in Comparative Example 4 and 0% in Comparative Example 1), this chemical potential barrier is insufficient to counteract the extremely strong reducing driving force generated by the high [Mn] content (up to 12.8%) in molten steel above 1500℃. The interfacial reaction proceeds spontaneously and violently, with a large amount of SiO2 being reduced and consumed. Meanwhile, the Al2O3 that floods into the slag (due to the simultaneous reduction of SiO2 by [Al]) acts as a network forming body, causing an irreversible increase in the degree of polymerization of the melt, ultimately resulting in a catastrophic surge in viscosity.

[0057] In terms of solid lubrication and heat transfer uniformity, the oxide coating layer of the first core-shell particles plays a phased sequential function. During the medium-high temperature stage (approximately 400°C to 900°C) of protective slag product storage and transportation and preheating in the crystallizer, the dense oxide shell provides a crucial physical barrier for the inner hexagonal boron nitride—this temperature range is precisely the stage where hexagonal boron nitride oxidation and ablation are most severe. At this time, the slag material is still in a solid powder state and cannot form a liquid protective layer. Without shell protection (as in Comparative Example 2 using exposed hexagonal boron nitride), the hexagonal boron nitride will suffer irreversible oxidation and degradation before the slag material melts. In the uncoated state, the basal polarity of hexagonal boron nitride is extremely low, and its surface energy is much lower than that of silicate slag, thus it is repelled and agglomerated by the slag (Comparative Example 2 is such a case). The chemical homology of the oxide shell enables the core-shell particles to achieve instantaneous superwetting when in contact with the high-temperature liquid slag, allowing hexagonal boron nitride to achieve nanoscale three-dimensional uniform dispersion in the liquid slag. When hexagonal boron nitride particles are uniformly dispersed into the liquid slag, the liquid slag layer itself forms a natural barrier to isolate atmospheric oxygen—the dissolved oxygen activity inside the liquid silicate slag is much lower than the partial pressure of atmospheric oxygen above the slag surface, allowing hexagonal boron nitride to maintain structural stability in this low-oxygen environment. In the even lower-temperature solid slag film region, the hexagonal boron nitride particles are permanently locked in the solidified glassy or crystalline matrix, physically encapsulated and protected by the dense solidified slag. During service, the hexagonal boron nitride cores uniformly dispersed in the slag film rely on the extremely weak van der Waals forces between the basal planes unique to its layered crystal structure to provide efficient boundary lubrication at the copper plate / slag film / billet shell interface, keeping the coefficient of friction stable between 0.034 and 0.049. More importantly, the oxide shell imparts a random spatial orientation to the hexagonal boron nitride sheets during the initial dispersion stage, preventing the sheet particles from tending to stack in a parallel orientation under shear flow. This causes heat conduction to no longer form an extremely fast channel along a certain direction, but to be uniformly distributed in three-dimensional space, eliminating the sharp peak of local heat flux density and fundamentally suppressing the risk of primary shell cracks induced by thermal stress concentration.

[0058] At the carbon consumption kinetics level, the second core-shell particles form a rigid framework core with polycrystalline clusters densely aggregated with a large number of nanocrystalline MgAl2O4 spinels. Graphitized dense carbon fills the intergranular gaps and coats the aggregates to form a carbon matrix shell. The ultra-high thermal stability of this polycrystalline aggregate core allows it to act as a continuous heat sink in high-temperature slag surface environments, continuously absorbing and dissipating heat from the carbon shell oxidation interface, effectively preventing thermal runaway and self-accelerated combustion of the carbon layer. The graphitized carbon shell, with its ordered graphite basal surface facing outwards, has significantly fewer active edge sites and a higher oxidation initiation temperature compared to amorphous carbon black, endowing the carbon shell with inherent high oxidation resistance and inertness. Crucially, as the carbon shell surface is gradually consumed, the densely packed spinel nanocrystals inside are exposed, forming a rough polycrystalline micro-topography on the carbon shell's retreating surface. The numerous grain protrusions and intergranular depressions create a tortuous surface channel, forcing oxygen molecules to travel a circuitous path to reach the deeper, unreacted carbon surface, significantly limiting the oxygen mass transfer rate at the microscale. The synergistic effect of the heat sink effect, graphitization inertness, and polycrystalline surface oxygen limitation significantly smooths out the carbon consumption rate, maintaining a highly consistent physical melting rate of the protective slag throughout the casting cycle, and compressing the liquid slag pool thickness fluctuation to an extremely narrow range of ±0.2 to ±0.5 mm. In contrast, in Comparative Example 2, the directly exposed ordinary carbon black, lacking the aforementioned triple protective mechanism on the high-temperature slag surface above 1300℃, undergoes explosive and rapid oxidation. The rapid consumption of carbon leads to a sudden thinning of the local slag layer and drastic fluctuations in the liquid slag pool thickness of ±2.6 mm, subsequently causing lubricating film rupture and uneven heat transfer. Although ordinary carbon black was also used in Example 5 to replace the second core-shell particles, its fluctuation range (±0.8 mm) was between that of the full core-shell configuration example and Comparative Example 2. This indicates that although the presence of the first core-shell particles cannot make up for the fundamental defect of uncontrolled carbon consumption, the uniformly dispersed skeleton formed by them inside the slag film alleviates the unevenness of slag film thickness caused by the fluctuation of melting rate to a certain extent.

[0059] The effectiveness of TiO2 as a silicate network depolymerizer was verified in Comparative Example 3. In the example containing TiO2, even though the slag-metal reaction resulted in the enrichment of a certain amount of Al2O3 into the slag, TiO2 occupied the weak nodes in the network at high temperature in the form of titanium-oxygen octahedra. This molecular-scale wedging effect blocked the infinite extension of the aluminosilicate polymer chains, thus controlling the viscosity change rate after the reaction to a single-digit percentage level. However, Comparative Example 3 lacked this microscopic depolymerization mechanism. The enriched Al2O3 smoothly embedded itself into the silica-oxygen network in the form of tetracoordinated aluminum, enhancing its degree of polymerization. The viscosity change rate jumped to 43.5%, and the slag macroscopically tended to be in a stagnant state, with a significant decline in lubrication and flowability. The rare earth regulation effect of Ce2O3 is reflected in the fine tuning of crystallization temperature and crystallization rate. In Example 3, under high alkalinity and fluorine-free conditions, the Ce2O3 content was increased to 1.5 parts, successfully stabilizing the crystallization temperature at 1132℃ and increasing the crystallization rate to 57%, avoiding the premature crystallization and excessively thick solid slag film problems usually associated with high alkalinity systems. Rare earth ions Ce 3+ The large ionic radius and specific interaction of Ce₂O₃ with the 4f subshell electrons on the coordinated oxygen atoms cause long-range coordination distortion and local charge redistribution in the silicate network, inhibiting the preferential and abundant precipitation of manganese olivine. In fluorine-free slag systems, this network topology regulation of Ce₂O₃ promotes the directional precipitation of calcium silicate mineral phases such as wollastonite at appropriate rates and crystal morphologies, ensuring that the solid slag film has sufficient thickness and appropriate crystallinity to scatter thermal radiation and regulate heat transfer; in fluorine-containing slag systems, it helps to restore the thermodynamic competitive advantage of the lanceolate crystallization path. TiN nanoparticles, as heterogeneous nucleating agents, embed themselves in the solid slag film layer closely adhering to the outer surface of the billet shell during the service of the protective slag. At the slag-steel interface in the meniscus region, the low interfacial mismatch caused by the extremely similar lattice constant between TiN and austenite leads to heterogeneous nucleation of the solidification and crystallization of the outermost grains of the primary billet shell, promoting the transformation of the initial solidification structure of the surface layer from coarse columnar crystals to finer equiaxed crystals. The refinement of surface grains allows the thermal stress borne by the shell during solidification to be more evenly distributed among the fine grains, significantly reducing the probability of surface crack initiation and propagation, which is indirectly reflected in the overall decrease in the crack index.

[0060] The above description is merely a preferred embodiment of the present invention and does not limit the patent scope of the present invention. Any equivalent structural or procedural transformations made based on the content of the present invention specification, or direct or indirect applications in other related technical fields, are similarly included within the patent protection scope of the present invention.

Claims

1. A special continuous casting protective slag for Mn13 high wear-resistant steel containing nanocomposite components, characterized in that, The continuous casting protective slag includes a matrix slag system and nanocomposite components dispersed in the matrix slag system; The matrix slag system includes at least CaO, SiO2 and MnO; Based on 100 parts by weight of the matrix slag system, the total amount of the nanocomposite components added is 1 to 8 parts; The nanocomposite component includes core-shell structured composite particles, wherein the core-shell structured composite particles are selected from at least one of a first core-shell particle and a second core-shell particle. The core layer material of the first core-shell particle is hexagonal boron nitride, and the coating material is oxide ceramic; The core layer material of the second core-shell particle is spinel, and the coating material is carbonaceous material.

2. The continuous casting protective slag for Mn13 high wear-resistant steel containing nanocomposite components according to claim 1, characterized in that, The matrix slag system comprises the following components by mass percentage: CaO: 30% to 45%; SiO2: 30% to 45%; MnO: 15% to 20%; Al2O3: 6% to 10%; And the remaining amount of flux.

3. The continuous casting protective slag for Mn13 high wear-resistant steel containing nanocomposite components according to claim 2, characterized in that, The mass ratio of the binary basicity CaO / SiO2 in the matrix slag system is 0.8 to 1.

3. The flux is selected from at least one of Li2O, Na2O, BaO, MgO and fluorides.

4. The continuous casting protective slag for Mn13 high wear-resistant steel containing nanocomposite components according to claim 3, characterized in that, The flux contained in the matrix slag system, by mass percentage, meets the following limitations: When Na2O is present, the Na2O content is less than 3%; The Li2O content is 3% to 5%; When MgO is present, the MgO content shall not exceed 2%; When BaO is present, the BaO content shall not exceed 10%; When fluoride is present, the fluoride content, expressed as fluoride ions, is less than 5%.

5. The continuous casting protective slag for Mn13 high wear-resistant steel containing nanocomposite components according to claim 1, characterized in that, The protective slag also includes a silicate network depolymerizing agent, which includes TiO2.

6. The continuous casting protective slag for Mn13 high wear-resistant steel containing nanocomposite components according to claim 1, characterized in that, The continuous casting protective slag also contains crystallization regulators and heterogeneous nucleating agents; The crystallization regulator includes Ce2O3; The heterogeneous nucleating agent comprises TiN nanoparticles with a size not exceeding 100 nm.

7. The continuous casting protective slag for Mn13 high wear-resistant steel containing nanocomposite components according to claim 1, characterized in that, The coating material of the first core-shell particle is selected from at least one of amorphous silica and crystalline alumina.

8. The continuous casting protective slag for Mn13 high wear-resistant steel containing nanocomposite components according to claim 1, characterized in that, The core layer material of the second core-shell particle is nanocrystalline magnesium aluminum spinel MgAl2O4, and the coating layer material is dense carbon with graphitization.

9. The continuous casting protective slag for Mn13 high wear-resistant steel containing nanocomposite components according to claim 1, characterized in that, The continuous casting protective slag is prepared by a method including the following steps: The matrix slag raw materials weighed according to the mass ratio are mixed and pre-melted to obtain pre-melted matrix powder slag; The nanocomposite components are uniformly dispersed into the pre-melted matrix powder.

10. The continuous casting protective slag for Mn13 high wear-resistant steel containing nanocomposite components according to claim 9, characterized in that, When the nanocomposite component contains the first core-shell particles, the first core-shell particles with an amorphous silica shell are prepared by the following steps: ultrasonically dispersing sheet-like micron-sized hexagonal boron nitride powder in a mixed solution of anhydrous ethanol and ammonia, fixing the pH value at 8.5, and adding tetraethyl orthosilicate dropwise in a constant temperature water bath at 45°C to carry out a hydrolysis and condensation reaction, so that silica heterogeneously nucleates on the surface of hexagonal boron nitride and self-assembles to form a coating layer; When the nanocomposite component contains the second core-shell particles, the second core-shell particles are prepared by the following steps: Basic magnesium carbonate is pretreated to decompose it into magnesium oxide-containing active powder to obtain a magnesium source; the magnesium source and γ-alumina powder are used as precursor raw materials, carbon-containing substances and nano-nickel oxide as a mineralizing agent are added, and ball milling is performed using a high-energy mechanical alloying method. Nano-sized magnesium aluminum spinel is generated by in-situ induced solid-phase reaction by mechanical force, while the carbon-containing substances undergo structural reorganization to form a dense carbon-coated shell layer on the surface of the spinel.

Images