CN117897507A

CN117897507A - Alloy steel

Info

Publication number: CN117897507A
Application number: CN202180101652.2A
Authority: CN
Inventors: 亚伦·约翰·米德尔顿
Original assignee: Wantai Alloy Co ltd
Current assignee: Wantai Alloy Co ltd
Priority date: 2021-10-06
Filing date: 2021-10-06
Publication date: 2024-04-16
Also published as: KR20240075821A; WO2023057062A1; CA3228109A1

Abstract

A vanadium alloy steel comprising carbon and vanadium is provided, wherein the microstructure of the steel comprises coherent interphase precipitates comprising vanadium. Furthermore, a method for producing such vanadium alloy steel is provided, comprising: providing a starting steel, heating the starting steel to at least partially austenitize the starting steel; the temperature is maintained at 650 ℃ ± 200 ℃ for about 25 minutes or less.

Description

Alloy steel

Technical Field

The present disclosure relates to alloy steels and methods for producing the same. And more particularly to vanadium alloy steels with enhanced elastic modulus.

Background

In view of the decarburization tendency of using a structural steel sheet having a stronger and thinner thickness, there is a need to develop a steel providing improved properties such as fire resistance, buckling resistance and bending resistance.

The development and use of refractory structural steels in civil engineering is a particular problem in the construction industry. Furthermore, there is a strong trend in the steel industry to reduce the use of implied carbon.

The development of lean-alloyed high strength steels with excellent fire resistance would have great economic benefits, enabling the use of more slim column and beam designs and speeding up the progress towards global decarburization targets. However, fire resistance requirements and decarbonization targets tend to be incompatible.

Furthermore, it would be beneficial to provide steel with enhanced buckling and bending resistance, for example, in both the automotive industry and in the construction of wind towers.

It is an object of the present disclosure to solve or at least mitigate the problems associated with the prior art.

Disclosure of Invention

In a first aspect, a vanadium alloy steel is provided comprising, preferably consisting essentially of, more preferably consisting of carbon and vanadium; wherein the microstructure of the steel comprises coherent interphase precipitates containing vanadium (coherent interphase precipitate).

In some embodiments, a vanadium alloy steel is provided that contains, preferably consists essentially of, and more preferably consists of carbon, vanadium, and balance iron; wherein the microstructure of the steel comprises coherent interphase precipitates containing vanadium.

In some embodiments, the vanadium alloy steel comprises, preferably consists essentially of, more preferably consists of (in weight percent):

carbon in the range of about 0.06 to about 1.1; and

vanadium in the range of about 0.1 to about 1.5;

wherein the steel microstructure comprises coherent interphase precipitates containing vanadium.

carbon in the range of about 0.06 to about 1.1;

vanadium in the range of about 0.1 to about 1.5; and

the balance of iron is iron,

wherein the microstructure of the steel comprises coherent interphase precipitates containing vanadium.

In some embodiments, the vanadium alloy steel comprises, preferably consists essentially of, more preferably consists of: vanadium and carbon; optionally one or more of nitrogen, molybdenum, copper, silicon, chromium and/or manganese; and the balance of iron; wherein the microstructure of the steel comprises coherent interphase precipitates containing vanadium.

carbon in the range of about 0.06 to about 1.1, and

vanadium in the range of about 0.1 to about 1.5;

optionally one or more of the following:

about 0.015 or less of nitrogen,

about 1.6 or less of molybdenum,

about 1 or less of the copper is present,

about 1.2 a or less of silicon,

about 0.3 or less of chromium, and/or

About 1.6 or less manganese; and

the balance of iron;

As will be described below, in the steels disclosed herein, the average elastic modulus (i.e., young's modulus) of the steel is enhanced. The coherent vanadium precipitates are believed to create a large strain field in the steel, resulting in a change in lattice parameters that can alter the bulk modulus of the steel.

The enhancement of the modulus of elasticity is considered to be applicable at high temperatures due to the greater stability of coherent precipitates to tempering and coarsening, thus enhancing the buckling resistance of the steel. In this way, improved fire resistance is achieved, for example, with lower alloy embodiments.

Providing a higher modulus steel as disclosed herein may also be beneficial to provide enhanced structural benefits, such as greater buckling resistance, e.g., embodiments using higher alloys. This may enable the construction of higher wind turbines.

Providing a higher modulus steel as disclosed herein may also be beneficial in providing enhanced resistance to flexing, such as embodiments using higher alloys. This may be beneficial in making lighter, more rigid automotive structures.

As shown in fig. 1, the lattice parameter variation effect is believed to be a direct result of matrix expansion and contraction caused by the elastic adaptation of the precipitates in the surrounding steel matrix. The vanadium-containing precipitates 2 have a lattice parameter that is different from the lattice parameter of the surrounding steel (e.g. ferritic) matrix 4. In other words, there is a lattice parameter mismatch between the vanadium-containing precipitates and the surrounding steel, which causes a large strain field (shown by the hatched area denoted by reference numeral 6 in fig. 1) to be generated when coherent precipitates containing vanadium are formed. Thus, the average lattice parameter of the overall steel is changed, thereby changing the bulk modulus of elasticity of the steel. In contrast, where non-coherent precipitates were formed, no such strain fields were observed.

This change in lattice parameter is considered anisotropic. However, as would be expected in polycrystalline materials such as steel, random grain orientation can neutralize this anisotropy, which contributes to modulus enhancement on a macroscopic scale.

In combination with this physical effect is a change in the electron distribution in the vicinity of the coherent interface, which further contributes to an enhancement of the bulk modulus.

Fig. 2 shows that vanadium is unique in having a relatively low mismatch between the lattice parameters of, for example, vanadium-containing precipitates compared to ferrite and austenite. Data points relating to mismatch in ferrite are shown by filled circles and triangles, and data points relating to mismatch in austenite are shown by open circles and triangles. Furthermore, fig. 2 shows the high limit dimensions of coherent precipitates comprising vanadium (in this case in ferrite) compared to precipitates comprising other elements commonly used in steel. This is believed to promote the generation of coherent precipitate formation, as a larger coherent limit has the effect of easier maintenance of the coherent property by established steel production methods.

Turning to fig. 3, the relationship between strengthening and the size of the precipitates is shown. As can be seen, the shear strength increases with increasing particle size, up to the critical particle size. As described above, this is due to coherent strain hardening. When the size of the particles exceeds the critical particle size, non-coherent precipitates are formed and the modulus strengthening effect is no longer seen.

The vanadium-containing precipitates are believed to be very effective in enhancing the elastic modulus of the steel by virtue of their high co-occurrence limit and high shear modulus compared to other precipitates.

When the precipitates are small and coherent with the surrounding matrix, then the dislocation mechanism transitions from a well-known ashby-orown bypass mechanism (ashby-orowan looping mechanism), which is applicable to non-coherent precipitates, to a grain shear mechanism. The strengthening mechanism due to dislocation shearing of the ordered coherent particles produces operation of the ordered and modulus strengthening mechanism. The resulting modulus enhancement increment can be described by Knowles-Kelly equation (1).

Where b.apprxeq.0.248 nm is the Berger vector of the dislocation, r is the average radius of the vanadium containing precipitates, e.g. VC particles (.apprxeq.3.1.+ -. 1.0 nm), and G.apprxeq.81.6 GPa is the shear modulus of the ferritic matrix. Δg is the difference in shear modulus between the matrix and the precipitate (175.7 GPa-81.6 gpa=94.1 GPa). f= (4/3) pi nr ³ Is the volume fraction of VC nano-precipitates, where n is the number density of the precipitates (. Apprxeq.1.18.+ -. 0.03X10) ^-4 nm ^-3 )。

This shows a direct relationship between the number density of the vanadium-containing precipitates (e.g. carbides), their average size, their shear modulus (which depends on the V: C ratio of the vanadium precipitates) and the subsequent modulus strengthening.

For example, fig. 5a shows a simulation of a ferritic steel as disclosed herein and shows the elastic strain field associated with coherent interphase precipitates comprising vanadium, which in this embodiment are formed as nano-sized platelets. The platelets are arranged in columns and the spacing between the columns is indicated by reference numeral 22. For clarity, only two columns of interphase precipitates are shown.

In precipitates containing vanadium, such as vanadium carbides and/or vanadium carbonitrides, there is a large relative difference in the atomic mismatch parallel and perpendicular components relative to the orientation of the precipitates. There is a very low mismatch parallel to the surface of the platelet precipitate, which translates into a large coherent critical size limit and swelling effect. A relatively large mismatch and a low coherent size limit are observed perpendicular to the precipitate, which results in a shrinking effect on the lattice spacing perpendicular to the broad platelet surface of the precipitate. This shrinkage reduces the average lattice spacing, the bergs vector, which in turn increases the elastic modulus of the unit cell, which results in a significant increase in young's modulus on a macroscopic level when applied to multiple grains on a macroscopic scale.

A variation of the lattice parameter along the line A-A is shown. This shows that the average lattice spacing shrinks due to anisotropic misfit strain, which changes the bulk modulus of elasticity.

In some embodiments, the steel is a refractory steel. The enhancement of the modulus of elasticity as described above is applicable at high temperatures due to the greater stability of coherent precipitates to tempering and coarsening, thereby enhancing the buckling resistance of the refractory steel.

The stability of buildings and other structures in a fire depends on the extent to which the steel structure softens when heated to the temperature created by the fire. A steel is generally considered refractory if its short-term strength when heated to such a temperature is about 0.6 to 0.7 of its strength at room temperature.

In an exemplary embodiment, the average elastic modulus is 120GPa or more at a temperature in the range of about 600℃to about 700 ℃.

In some embodiments, the steel comprises vanadium in the range of about 0.1 to about 0.3 (wt%). For example, the steel is a ferritic steel (e.g., refractory alloy HSLA (high strength low alloy, high strength low alloy)).

In some embodiments, the steel comprises vanadium in the range of about 0.3 to about 1.5 weight percent. For example, the steel is ferritic steel (e.g., AHSS (advanced high strength steel, advanced high strength steel)).

In some embodiments, the steel comprises vanadium in the range of about 0.1 to about 1.5 (wt%). For example, the steel is a ferritic steel (e.g., hydrogen-resistant steel).

In some embodiments, the steel comprises vanadium in the range of about 0.1 to about 1 (wt%). For example, the steel is pearlitic steel.

The use of vanadium to form interphase precipitates is considered advantageous compared to other alloying elements, such as niobium and titanium, because vanadium has a relatively low impact on the transformation temperature (i.e. the temperature at which the steel transforms from austenite to ferrite). Lowering the transition temperature has the adverse effect of increasing the likelihood of random precipitates rather than interphase precipitates. Thus, the use of vanadium is beneficial in maintaining the transition temperature within the desired range.

The presence of carbon contributes to the hardness, strength and hardenability of the steel. Carbon also acts as an austenite stabilizer.

In some embodiments, the steel comprises carbon in the range of about 0.06 to about 0.2 weight percent. For example, the steel is a ferritic steel (e.g., refractory alloy HSLA (high strength low alloy)).

In some embodiments, the steel comprises carbon in the range of about 0.06 to about 0.3 (wt%). For example, the steel is a ferritic steel (e.g., AHSS (advanced high strength steel)).

In some embodiments, the steel comprises carbon in the range of about 0.06 to about 0.3 (wt%). For example, the steel is a ferritic steel (e.g., hydrogen-resistant steel).

In some embodiments, the steel comprises carbon in the range of about 0.6 to about 1.1 weight percent. For example, the steel is pearlitic steel.

In some embodiments, the steel is a modulus-enhanced steel having an average elastic modulus (i.e., measured bulk elastic modulus) of greater than about 210Pa, such as about 220GPa or greater, such as in the range of about 210GPa to about 300GPa, such as in the range of about 220GPa to about 300 GPa.

The elastic modulus can be measured using a dynamic resonance method. A sample of steel having a uniform cross-section (e.g., circular, square, or rectangular) is prepared and the characteristic vibration frequency of the steel is determined and related to the elastic modulus using a known formula. ASTM standard C1259 specifies standard test methods for dynamic young's modulus, shear modulus and poisson's ratio of advanced ceramics by pulse-excited vibration for bars of rectangular and circular cross section.

Optionally, the steel is modulus-enhanced AHSS steel. Optionally, the steel is a refractory steel (e.g., HSLA steel) having an average elastic modulus of about 120GPa or more at a temperature in the range of about 600 ℃ to about 700 DEG C

Optionally, the particle size of the coherent interphase precipitate is about 9nm or less, optionally in the range of about 5nm to about 9nm, optionally greater than about 5nm. This may be measured using high resolution transmission electron microscopy, X-ray diffraction or any other suitable method.

The particle size of the coherent interphase precipitate was taken as the average of the maximum particle sizes of representative particle samples.

For coherent interphase precipitates in ferrite, a Focused Ion Beam (FIB) microscope is typically used to prepare the lamellar sample. Analysis of the lamina layer is also typically performed using a high resolution scanning transmission electron microscope (high-resolution scanning transmission electron microscope, HR STEM) calibrated with a probe having a cold field emission gun in situ. By this method, a high-angle annular dark field (HAADF) image can be obtained at a suitable collection angle. Given such HAADF and STEM micrographs, the size, morphology and number density of coherent precipitates can be obtained, whereby the average number density is derived from the average taken from the different grains. Concentration profiles can also be obtained from HAADF data that can verify the properties of the nanoparticles we combine. Atom probe tomography can also be used to verify and provide improved 3D spatial resolution of the atomic concentration profile (from HAADF micrographs) of core-shell nanoparticles.

As shown in fig. 2 and described above, this is relatively large compared to precipitates formed by other commonly used elements and is believed to contribute to the increase in elastic modulus of the overall steel (as shown in fig. 3).

Optionally, the interphase precipitate comprises vanadium carbide, optionally wherein the interphase precipitate comprises a compound of formula V _x C _y With increased vanadium levels, where x > y (e.g., coherent interphase precipitates include V ₄ C ₃ 、V ₆ C ₅ And/or V ₅ C ₃ ). During the transition from coherent to non-coherent particles, the precipitate-matrix interface energy increases, so that a larger fraction of vacancies is required for the precipitate stabilization, resulting in a slightly gradual decrease in the carbon to vanadium ratio and thus an increase in its stability.

It is also believed that a small increase in vanadium content in the ferritic steel may reduce the carbon to vanadium ratio in the precipitate, which helps to increase the shear modulus of the precipitate, resulting in modulus strengthening.

In some embodiments, the steel is a hydrogen-resistant steel, such as embodiments using higher alloys. Vanadium carbides within the steel are believed to act as hydrogen traps that can irreversibly trap hydrogen.

By increasing the ratio of V: C in the interphase precipitates, a large number of carbon vacancies are created at the coherent interface of the steel (e.g., ferrite) lattice matrix. These vacancies act as hydrogen traps, resulting in enhanced hydrogen embrittlement resistance.

This is shown in fig. 4a and 4 b. FIG. 4b shows V ₄ C ₃ 3D representation of hydrogen atoms 8 occupying carbon vacancies in the co-phase interphase precipitate 2. FIG. 4a shows the reaction with TiC precipitates 14, nbC precipitates 12 and V ₄ C ₃ The potential energy well associated with the precipitate 10. As can be seen, with V ₄ C ₃ The potential energy well associated with the precipitate 10 is deeper than the potential energy wells of the other precipitates and thus acts as a more efficient hydrogen trap.

Furthermore, vanadium carbide precipitates containing enhanced vanadium are believed to contain carbon vacancies that are inherently more thermally stable upon exposure to fire.

Optionally, the steel comprises a ferrite phase and coherent interphase precipitates are formed in the ferrite phase, for example to form spheroidal or nodular ferrite (sintered ferrite).

In some embodiments, coherent interphase precipitates are formed when the austenite phase is transformed into the ferrite phase, in other words, under conditions corresponding to the transition phase boundaries between austenite and ferrite.

Upon cooling, to accommodate the cumulative elastic strain caused by the formation of coherent interphase precipitates, it is believed that spherical ferrite (NF) may form which changes its orientation during growth. Spherical ferrite due to lattice parameter mismatch due to iron carbide (i.e., cementite) has only been previously found in eutectic cast steel. It has been found that spherical ferrite may also be produced due to the mismatch created by the coherent alloy carbides. Spherical ferrite is characterized by a disordered grain orientation and a disordered grain boundary orientation. Each grain is not a isomorphous body and is similar to a nodular appearance. Ferrite transformation with interphase precipitation is a eutectoid transformation similar to degenerate pearlite in high carbon steel, in which alloy carbides are formed instead of iron carbides, cementite. Therefore, it can be considered that such spherical or nodular ferrite is similar to pearlite observed in high carbon steel.

This branching of NF, and hence the change in ferrite/austenite interface plane of the interphase precipitates, is believed to optimize the microstructure texture towards better hydrogen embrittlement resistance. It is believed that the low angle grain boundaries and the substitutional lattice (Coincidence Site Lattice, CSL) which exhibit lower relative grain boundary energies can significantly improve the hydrogen embrittlement resistance of ferritic steels. One possible mechanism for this is an increase in the number of atoms arranged at the CSL grain boundaries (e.g., between Σ5 and Σ13), which can reduce the vacancy density and further mitigate segregation of hydrogen atoms at the grain boundaries.

In some embodiments, the steel composition comprises Mo, cr, and/or Cu. In this way, it is believed that the steel may benefit from core-shell nanoparticles that exhibit higher coarsening and lattice parameter shrinkage resistance to capillary actuation. Referring to fig. 7b, in some exemplary embodiments, the core-shell nanoparticle 16 may include a VN core 18. For example, the core-shell nanoparticle may include a shell 20 comprising Mo, cr, and/or Cu.

Capillary driven coarsening of particles is hindered due to interfacial segregation of Mo, cr and/or Cu, which reduces the driving force for further coarsening of particles due to the reduction of interfacial energy. Coarsening resistance aids in grain refinement during austenitization and prior austenite grain size control. The growing core-shell nanoparticles are inherently resistant to coarsening and a zener pinning (zener pinning) mechanism is applied on the austenite grain boundaries to limit the grain size.

Thus, improved toughness is achieved.

Copper is believed to have a low mismatch with the ferrite lattice, which enables it to contribute to elastic mismatch strain by forming high density coherent interphase particles. The low magnetic moment of Cu is believed to play a positive role in increasing the number density of the composite nanoparticle. Without being bound by any particular theory, it is believed that the application of a magnetic field increases the Gibbs free energy of the system. The increase in gibbs free energy promotes precipitation of alloying elements including Cu, and since Cu has a low magnetic moment, the precipitates are not hindered by a magnetic field.

In some embodiments, the steel comprises copper in the range of greater than 0 to about 0.5 weight percent. For example, the steel is a ferritic steel (e.g., refractory alloy HSLA (high strength low alloy)).

In some embodiments, the steel comprises copper in the range of greater than 0 to about 1 (wt%). For example, the steel is a ferritic steel (e.g., AHSS (advanced high strength steel)).

In some embodiments, the steel comprises copper in the range of greater than 0 to about 1 (wt%). For example, the steel is a ferritic steel (e.g., hydrogen-resistant steel).

In some embodiments, the steel comprises copper in the range of greater than 0 to about 0.5 weight percent. For example, the steel is pearlitic steel.

The presence of manganese is believed to assist in the refinement of ferrite grains, for example when higher coiling temperatures are used.

In some embodiments, the steel comprises manganese in the range of about 0.4 to about 1.6 (wt%). For example, the steel is a ferritic steel (e.g., refractory alloy HSLA (high strength low alloy)).

In some embodiments, the steel comprises manganese in the range of greater than 0 to about 0.8 wt.%. For example, the steel is a ferritic steel (e.g., AHSS (advanced high strength steel)).

In some embodiments, the steel comprises manganese in the range of greater than 0 to about 1.6 wt.%. For example, the steel is a ferritic steel (e.g., hydrogen-resistant steel).

In some embodiments, the steel comprises manganese in the range of about 0.6 to about 1.6 (wt%). For example, the steel is pearlitic steel.

The presence of silicon is believed to increase the number density of coherent interphase precipitates, for example when higher coiling temperatures are used. Silicon retards pearlite formation and can allow a slower cooling rate, which promotes the formation of coherent interphase precipitates.

In some embodiments, the steel comprises silicon in the range of greater than 0 to about 0.5 weight percent. For example, the steel is a ferritic steel (e.g., refractory alloy HSLA (high strength low alloy)).

In some embodiments, the steel comprises silicon in the range of greater than 0 to about 0.5 weight percent. For example, the steel is a ferritic steel (e.g., AHSS (advanced high strength steel)).

In some embodiments, the steel comprises silicon in the range of greater than 0 to about 1.5 weight percent. For example, the steel is a ferritic steel (e.g., hydrogen-resistant steel).

In some embodiments, the steel comprises silicon in the range of greater than 0 to about 1.2 weight percent. For example, the steel is pearlitic steel.

The presence of chromium is believed to increase the solubility of other microalloying elements and reduce the size of the precipitates. In this way, the average precipitate size can be kept below the limit size of the commonality, resulting in a larger proportion of the coherent precipitate.

In some embodiments, the steel comprises chromium in the range of greater than 0 to about 0.3 weight percent. For example, the steel is a ferritic steel (e.g., refractory alloy HSLA (high strength low alloy)).

In some embodiments, the steel comprises chromium in the range of greater than 0 to about 0.3 weight percent. For example, the steel is a ferritic steel (e.g., AHSS (advanced high strength steel)).

In some embodiments, the steel comprises chromium in the range of greater than 0 to about 0.3 weight percent. For example, the steel is a ferritic steel (e.g., hydrogen-resistant steel).

In some embodiments, the steel comprises chromium in the range of greater than 0 to about 0.1 weight percent. For example, the steel is pearlitic steel.

The presence of molybdenum is believed to help increase the number density of the co-interphase particles by forming a complex with the precipitate that can improve the diffusion kinetics of the precipitate.

Optionally, the interphase precipitate comprises Mo, optionally wherein the interphase precipitate comprises a compound of formula (Mo, V) _x C _y Wherein x > y, e.g. (Mo, V) ₄ C ₃ /(Mo,V)C。

This is believed to have the effect of reducing the lattice mismatch between the interphase precipitates and the ferrite phase, resulting in improved commonality.

Furthermore, the presence of Mo is believed to reduce the migration rate of austenite to ferrite phase during transformation, which enables the formation of a greater number of phase-to-phase precipitates.

In some embodiments, the steel comprises molybdenum in the range of greater than 0 to about 0.2 wt.%. For example, the steel is a ferritic steel (e.g., refractory alloy HSLA (high strength low alloy)).

In some embodiments, the steel comprises molybdenum in the range of greater than 0 to about 0.5 weight percent. For example, the steel is a ferritic steel (e.g., AHSS (advanced high strength steel)).

In some embodiments, the steel comprises molybdenum in the range of greater than 0 to about 0.5 weight percent. For example, the steel is a ferritic steel (e.g., hydrogen-resistant steel).

In some embodiments, the steel comprises molybdenum in the range of greater than 0 to about 0.2 wt.%. For example, the steel is pearlitic steel.

Optionally, the ferrite phase comprises grains having an average size of less than about 20 μm. For example, the average grain size is in the range of about 5 μm to about 20 μm. The grain size may be measured by any suitable method, for example using a scanning electron microscope and electron back-scattering diffraction. The average grain size can then be calculated using a cross-sectional method. In this method, a straight line is drawn on a micrograph, and the number of grain boundaries intersecting the line is counted. The average grain size is obtained by dividing the number of crossing points by the actual line length.

Optionally, the steel microstructure comprises VN precipitates, optionally wherein the microstructure comprises intragranular VN nucleated acicular ferrite (e.g. V (C, N)). Optionally, VN precipitates form in austenite as intragranular nucleating agents for acicular ferrite formation. This is believed to be due to the lower solubility of VN in austenite relative to vanadium carbides.

It is believed that the intra-crystalline VN nucleated acicular ferrite structure improves the commonality of inter-phase particles and refines the ferrite grains, resulting in corresponding resultant benefits to the elastic modulus and toughness of the steel.

In this way, needle-like microstructures are formed simultaneously with the system of nanoscale interphase precipitates. The needle-like microstructure is schematically shown in fig. 7 a.

In some embodiments, the steel comprises nitrogen in the range of greater than 0 to about 0.015 (wt%). For example, the steel is a ferritic steel (e.g., refractory alloy HSLA (high strength low alloy)).

In some embodiments, the steel comprises nitrogen in the range of greater than 0 to about 0.015 (wt%). For example, the steel is a ferritic steel (e.g., AHSS (advanced high strength steel)).

In some embodiments, the steel comprises nitrogen in the range of greater than 0 to about 0.015 (wt%). For example, the steel is a ferritic steel (e.g., hydrogen-resistant steel).

In some embodiments, the steel comprises nitrogen in the range of greater than 0 to about 0.01 (wt%). For example, the steel is pearlitic steel.

In some embodiments, where not all of the nitrogen is used to form VN precipitates as an intra-crystalline nucleating agent for acicular ferrite formation, the inter-phase vanadium carbonitrides may form as coherent inter-phase precipitates. Nitrogen will further increase the driving force for precipitation, which is beneficial for obtaining high number density interphase precipitates.

Optionally, the steel comprises single phase ferritic steel, such as HSLA, AHSS, and/or hydrogen resistant steel.

Optionally, the steel comprises pearlitic steel. Thus, the steel includes ferrite phase and cementite phase. The modulus of the ferrite phase is enhanced by the presence of coherent phase particles and the above mechanism.

Optionally, the pearlite steel comprises vanadium enhanced cementite phase, optionally the vanadium enhanced cementite phase comprises vanadium dissolved in cementite, e.g. to form Fe ₂ VC and/or FeV ₂ C。

The average lattice spacing in the ferrite phase shrinks due to misfit strain from both eutectic interphase precipitates containing vanadium and cementite. Furthermore, given a suitable isothermal holding state, vanadium can be readily dissolved in cementite, which can further contribute to mismatch-induced lattice strain due to the variation of cementite lattice parameters.

Cementite is well known for its anisotropy, and this also applies to its anisotropy of the elastic modulus component as a function of crystal orientation. Dissolution of vanadium in cementite reduces this anisotropy. It is believed that this will evolve as a function of the isothermal holding duration, as the composition of cementite changes with aging.

The lattice strain associated with elastic mismatch of alloy carbide or V-enhanced cementite can be verified by using high resolution scanning transmission electron microscopy (HR STEM) and X-ray diffraction (XRD), respectively. Atomic probe tomography can also be used as an additional method of quantifying atomic concentration in three dimensions.

For the formation of lattice strains associated with iron carbides and strains induced by V-enhanced cementite in higher carbon pearlite steels, it is suggested to use electron back scattering diffraction (XRD). The elastic strain was quantified by XRD measured full width at half maximum (full width at half maximum, FWHM) of the ferrite diffraction peak. The lattice strain in the sheet ferrite quantified by XRD measurement will also be related to the experimental stress. Modulus enhancement will also contribute to test stress, which can be conventionally obtained by tensile testing equipment.

Referring to fig. 5b, an exemplary embodiment is shown wherein the steel comprises cementite agglomerate structures. Additional strain applied to the ferrite lattice due to cementite is shown. This results in a decrease in the average lattice spacing of the ferrite matrix in the region adjacent to the cementite aggregation structure.

A variation of the lattice parameter along line A-A is shown. This shows that the average lattice spacing shrinks due to the mismatch strain of both the coherent vanadium-alloy carbide and cementite, which changes the bulk modulus of elasticity.

Fig. 5b also shows a schematic representation of the microstructure of pearlite containing V-enhanced cementite.

Optionally, the tensile strength of the steel is in the range of about 690MPa to about 2000MPa, for example in the range of about 690MPa to about 1800 MPa.

Tensile strength can be measured by clamping the ends of a properly prepared standard specimen in a tensile tester and then applying a continuously increasing uniaxial load until failure occurs. ASTM E8/E8M-13 may be used: "Standard test method for tensile test of metallic Material (Standard Test Methods for Tension Testing of Metallic Materials)".

The end result of the greater number density and optimized precipitate composition may indeed impact the suitability of the steels disclosed herein for use in refractory or hydrogen transportation/storage applications.

The structure of the steel alloys described herein may be determined by conventional microstructural characterization techniques, such as, for example, optical microscopy, TEM, SEM, AP-FIM, and X-ray diffraction (including combinations of two or more of these techniques).

In another aspect, a method for preparing a vanadium alloy steel as disclosed herein is provided, the method comprising:

a. providing a starting steel;

b. heating the starting steel to at least partially austenitize the steel;

c. maintaining the temperature of the steel at 650 ℃ ± 200 ℃ for about 25 minutes or less;

wherein the vanadium alloy steel comprises coherent interphase precipitates containing vanadium.

Optionally, the starting steel comprises, preferably consists essentially of, more preferably consists of: vanadium and carbon.

Optionally, the starting steel comprises, preferably consists essentially of, more preferably consists of: vanadium, carbon and balance iron.

Optionally, the starting steel comprises, preferably consists essentially of, more preferably consists of (in weight-%):

carbon in the range of about 0.06 to about 1.1, and

vanadium in the range of about 0.1 to about 1.5.

Carbon in the range of about 0.06 to about 1.1, and

vanadium in the range of about 0.1 to about 1.5; and

the balance of iron.

Optionally, the starting steel comprises, preferably consists essentially of, more preferably consists of: vanadium and carbon; one or more of nitrogen, molybdenum, copper, silicon, chromium, and/or manganese; and the balance iron.

carbon in the range of about 0.06 to about 1.1, and

vanadium in the range of about 0.1 to about 1.5;

one or more of the following:

about 0.015 or less of nitrogen,

about 1.6 or less of molybdenum,

about 1 or less of the copper is present,

about 1.2 a or less of silicon,

about 0.3 or less of chromium, and/or

About 1.6 or less manganese; and

the balance of iron.

Optionally, heating the starting steel to at least partially austenitize the steel includes increasing the temperature to at least about 900 ℃, optionally to a temperature in the range of about 920 ℃ to about 1300 ℃, for example to a temperature in the range of about 1100 ℃ to about 1300 ℃.

In some embodiments, the starting steel may be provided in the form of a billet, slab, bloom, or any other suitable form.

In some embodiments, the heating may be performed by induction heating and/or a roller hearth furnace.

Maintaining the temperature of the steel at 650 ℃ ± 200 ℃ for about 25 minutes or less (i.e., isothermal holding) increases the number density of coherent interphase precipitates, which results in an enhanced elastic modulus (according to formula (1) above).

It has been found that the formation of precipitates at the ferrite/austenite interface is hindered when the interface progresses rapidly. Thus, isothermal maintenance in this way increases interphase precipitation (e.g. precipitation of vanadium carbides and/or vanadium carbonitrides), which consequently improves the properties of the steel.

The plate spacing between columns of inter-phase precipitates decreases with increasing transition time. Therefore, isothermal holding has the effect of reducing the spacing between columns of inter-phase precipitates, thereby increasing the number density of inter-phase precipitates.

Furthermore, isothermal maintenance in this way is believed to have the effect of increasing the V: C ratio in the interphase precipitates, such that a large number of carbon vacancies are created at the coherent interface of the steel (e.g., ferritic) lattice matrix. These vacancies act as hydrogen traps, resulting in enhanced hydrogen embrittlement resistance (see, e.g., fig. 4a and 4 b).

In some embodiments, the temperature is maintained at 650 ℃ ± 200 ℃ for about 20 minutes or less.

In some embodiments, the temperature is maintained by placing the steel in an insulated tank.

In some embodiments, batch annealing is performed, for example using an insulated box to control the temperature applied to the steel. For example, a hood type furnace may be used to apply heat treatment to the entire coil.

In some embodiments, the temperature is maintained by using an induction heating system, such as an induction furnace.

Optionally, the temperature is maintained at 650 ℃ ± 200 ℃ for about 15 minutes or less. In some embodiments, the temperature is maintained at 650 ℃ ± 200 ℃ for about 5 minutes to about 20 minutes, for example, in the range of about 10 minutes to about 20 minutes.

Optionally, after isothermal holding, the steel is air cooled.

Optionally, the temperature is maintained at 650 ℃ ± 100 ℃.

Optionally, the steel is cooled at a cooling rate in the range of about 2 ℃/sec to about 80 ℃/sec, optionally at a cooling rate in the range of about 2 ℃/sec to about 50 ℃/sec, prior to maintaining the temperature at 650 ℃ ± 200 ℃.

It has been found that the formation of precipitates at the ferrite/austenite interface is hindered when the interface progresses rapidly. Thus, controlling the cooling rate in this way increases interphase precipitation (e.g. vanadium carbide and/or vanadium carbonitride precipitation), which consequently improves the properties of the steel.

Furthermore, rapid cooling of the steel in this way prevents or counteracts the formation of undesired phases and/or promotes the formation of acicular ferrite.

In some embodiments, cooling is performed by air cooling, forced air cooling, laminar flow cooling (lamella cooling), and/or any other suitable cooling method.

In the presence of Mn in the steel composition, this can promote higher number densities at faster cooling rates.

In some embodiments, after heating the starting steel, the steel is hot rolled at a temperature above the recrystallization stop temperature (recrystallisation stop temperature, RST), optionally by recrystallization controlled rolling and/or V (C, N) precipitation controlled rolling.

By performing hot rolling at a temperature higher than RST, grain deformation is restricted, thereby promoting the generation of interphase precipitates.

Optionally, the steel is hot rolled before cooling the steel to 650 ℃ ± 200 ℃.

In some embodiments, hot rolling is performed at a temperature in the range of about 820 ℃ to about 1200 ℃.

In some embodiments, hot rolling includes a rough rolling step, an intermediate step, and/or a finish rolling step.

Optionally, V (C, N) precipitation control rolling may be used for compositions containing relatively large amounts of nitrogen. This achieves a low level of grain deformation.

In some embodiments, hot rolling the steel includes hot rolling at a substantially V (C, N) precipitation temperature time inflection point (ase).

For example, in the case of using V (C, N) precipitation control rolling, this may be performed at a substantially V (C, N) precipitation temperature time inflection point.

At the inflection point of V (C, N) precipitation temperature and time, the system has the optimal free energy of VN precipitation. Therefore, performing hot rolling at this temperature promotes the formation of acicular ferrite.

The V (C, N) precipitation temperature time inflection point depends on the specific composition of the steel, but is typically about 850 ℃.

In some embodiments, a magnetic field is applied to the steel composition during or after maintaining the temperature.

Optionally, the magnetic field is a static magnetic field in the range of about 0.2T to about 16T. Optionally, the magnetic field is a varying magnetic field, such as that applied by an induction heating system. In some embodiments, the varying magnetic field is applied by an induction heating system that also controls the temperature of the steel (e.g., for maintaining the temperature at 650 ℃ ±200 ℃).

The application of the magnetic field is advantageous for increasing the number density of coherent precipitates and also for optimizing the composition of the constituent nanosized precipitates.

In practice, long isothermal holding durations may be detrimental to productivity. The proposed environmentally friendly method for accelerating the isothermal holding duration is to use a magnetic field that can act to slow down the austenite-to-ferrite transformation kinetics (thereby refining the spacing between the precipitates and maximizing the number density of interphase precipitates). This may also increase the extent of carbide at higher V: C ratios due to paramagnetic effects.

Prior to precipitation, the vanadium in the solution acts to raise the curie temperature Tc (but exerts a small influence on the magnetic moment and small interaction energy with an external magnetic field) such that the temperature range between Tc and a given transition temperature below the curie point increases and thus the magnetization at the transition point is greater. The external magnetization energy serves to significantly increase the gibbs free energy of the system, increasing the free energy of total nucleation of the nanoparticles, which in turn increases the nucleation rate and significantly increases the quantitative number density of the co-mingled vanadium precipitates.

In steel compositions containing higher amounts of carbon (i.e., pearlite steels), the applied magnetic field also acts to refine the inter-pearlite layer spacing (due to the increase in nucleation rate of cementite). As previously described, this may also contribute to enhancement of lattice strain and modulus due to strain applied to the ferrite lattice in the region adjacent to the cementite agglomerate structure.

Optionally, after the temperature has been maintained at 650 ℃ ± 200 ℃ for about 25 minutes or less, the steel is reheated to at least partially austenitize the steel.

Optionally, the reheating step includes increasing the temperature to a range of about 920 ℃ to about 1150 ℃.

Optionally, after the steel has been reheated, the temperature is again maintained at 650 ℃ ± 200 ℃ for about 25 minutes or less.

In some embodiments, the temperature is maintained at 650 ℃ ± 200 ℃ for about 20 minutes or less. Optionally, the temperature is maintained at 650 ℃ ± 200 ℃ for about 15 minutes or less. In some embodiments, the temperature is maintained at 650 ℃ ± 200 ℃ for a range of about 5 minutes to about 20 minutes, such as a range of about 10 minutes to about 20 minutes. Optionally, after isothermal holding, the steel is air cooled.

Optionally, the temperature is maintained at 650 ℃ ± 100 ℃.

As previously mentioned, isothermal holding in this manner improves interphase precipitation (e.g., vanadium carbide and/or vanadium carbonitride precipitation).

Optionally, after the steel has been reheated, it is cooled, optionally to a temperature of 650 ℃ ± 200 ℃ at a cooling rate in the range of about 2 ℃/sec to about 80 ℃/sec.

Optionally, the cooling rate is in the range of about 2 ℃/sec to 50 ℃/sec. As previously described, controlled cooling enhances interphase precipitation (e.g., vanadium carbide and/or vanadium carbonitride precipitation) and inhibits the formation of undesirable phases.

At least a part of the precipitates previously produced in the steel will remain upon reheating, and thus reheating, subsequent cooling and/or subsequent isothermal holding steps will lead to an increase in the number density of the precipitates in the final steel.

Drawings

FIG. 1 shows a schematic diagram showing interactions of coherent and non-coherent precipitates with a surrounding ferritic lattice matrix;

FIG. 2 shows a comparison of lattice parameter mismatch and limiting coherent precipitate size for precipitates containing Ti, V, zr or Nb;

FIG. 3 shows the relationship between strength and precipitate particle size;

FIGS. 4a and 4b are schematic diagrams of the ability of coherent interphase precipitates to act as hydrogen traps;

FIG. 5a shows the effect of the strain field in the ferritic matrix on the average lattice parameter in ferritic steel;

FIG. 5b shows the effect of the strain field in the ferritic matrix on the average lattice parameter in the pearlitic steel;

fig. 6 shows three typical beam compositions (according to european specification 3);

fig. 7a shows a schematic view of a VN nucleated acicular ferrite structure;

FIG. 7b shows a schematic of a core-shell nanoparticle;

FIG. 8 illustrates a process diagram representative of a method for manufacturing steel in accordance with the present disclosure;

FIG. 9 shows a flow chart depicting further steps that may be applied to the method of FIG. 8;

FIG. 10 illustrates a process diagram representative of a method for manufacturing modulus enhanced hot rolled high carbon wire steel in accordance with the present disclosure; and

Fig. 11 shows a flow chart describing further steps that may be applied to the method of fig. 10.

Examples

The vanadium alloy steels and production methods disclosed herein will now be described with reference to the following non-limiting examples.

Example 1

Four exemplary steel compositions according to the present disclosure are described in table 1. Each composition also contains the balance iron.

Table 1 (all wt.%)

It has been determined that the ferrite compositions listed in table 1 will have the characteristics and properties listed in table 2.

Characteristics and properties of steel	Maximum value	Minimum value
			Average modulus of elasticity	300Gpa	210Gpa
Ultimate tensile Strength	1800Mpa	690Mpa

TABLE 2

Example 2

Typical compositions of reinforced modulus AHSS steels are provided in table 3. Such a composition also contains the balance iron. In this case, the balance iron was 97.6955 wt%.

Table 3 (all wt.%)

Example 3

Conventional steel grade S690MC (according to EN 10051) was compared with nanostructured reinforced modulus steel as disclosed herein. The nominal elastic modulus of conventional S690MC steel is considered to be 210GPa (according to European Standard EN 1993-1-1: european Specification 3: steel structural design (European Standard EN 1993-1-1:Eurocode 3:Design of steel structures), european Standard EN 1993-1-12: reviewed-high strength steel (European Standard EN 1993-1-12: general-High strength steels)). The average elastic modulus of the modulus-reinforced steel is considered to be 230GPa, in other words, 20GPa is increased.

Fig. 6 shows three typical beam compositions (according to european specification 3). The results with respect to the decrease in deflection at the increase in elastic modulus are listed in table 4. These results were obtained by computer simulation.

TABLE 4 Table 4

Example 4

With reference to fig. 8, one exemplary method for producing the reinforced modulus steel disclosed herein will now be described.

The cast steel slab (or alternatively the billet) is heated to a reheat temperature (Treh) in the range of about 1100 ℃ to about 1300 ℃. This is indicated by reference numeral 102 on fig. 8. This temperature is maintained until the billet is fully heated over its thickness (i.e., at Teq as shown in fig. 8). By heating the steel blank in this way, the steel is heated to form austenite (γ) having a grain microstructure indicated by reference numeral 104 as shown in the schematic diagram.

The heated steel is then hot rolled at a temperature in the range of about 820 ℃ to about 1200 ℃, first in a roughing mill as indicated by reference numeral 106, followed by hot rolling in a finishing mill as indicated by reference numeral 108. During hot rolling, the steel is cooled at a temperature of standard air cooling, for example at about 0.7 ℃/sec.

The hot rolling 108 in the finishing mill may be performed using conventional recrystallization controlled rolling. Alternatively, particularly in the case of steel having a composition containing a relatively large amount of nitrogen, V (C, N) precipitation controlled rolling may be performed. This is done at the V (C, N) precipitation time temperature inflection point 110 where there is the best free energy for VN precipitation. The V (C, N) precipitation time temperature inflection point varies depending on the specific steel composition, but is typically about 850 ℃.

When hot rolling steel, especially in finishing mills, VN precipitates in austenite as an intragranular nucleating agent for acicular ferrite formation. At this stage, the steel has a microstructure as schematically shown at 112.

Before the Recrystallization Stop Temperature (RST), and after hot rolling has been completed, the steel is rapidly cooled 114 at a cooling rate in the range of about 2 ℃/sec to about 80 ℃/sec until a temperature of 650 ℃ ± 100 ℃ is reached. This may be done by forced air cooling, laminar cooling or any other suitable cooling means. Rapid cooling of the steel in this manner prevents or impedes the formation of undesirable phases and/or promotes the formation of acicular ferrite schematically shown at 116.

When a temperature of 650 ℃ ± 100 ℃ is reached, the steel is formed into a coil. The temperature of the steel is then maintained 118 at 650 c + 100 c for a period of about 10 minutes to about 20 minutes, after which the steel is air cooled. As described above, isothermal holding in this way promotes the generation of interphase precipitates.

In this way, a modulus-enhanced nanostructured steel is produced.

Optionally, the method may further comprise the steps listed in fig. 9 to increase the number density of interphase precipitates.

After the isothermal holding step 118, the coiled steel is cooled and then the blank is stamped or cut 120 from the uncoiled strip. Optionally, this step may be omitted.

The steel may then be austenitized 122 a second time by heating to a temperature in the range of about 920 ℃ to about 1150 ℃, for example using a roller hearth furnace and/or induction heating apparatus. The reheated steel is then cooled 124 at a rate in the range of about 2 deg.c/sec to about 50 deg.c/sec until a temperature of 650 deg.c + 40 deg.c is reached. Cooling may be performed using forced air cooling, laminar flow cooling, or any other suitable method.

A static magnetic field in the range of about 0.2T to about 16T may then be applied 26 to the steel. The temperature is maintained at 650 c 40 c for about 15 minutes or less while the magnetic field is applied. This may be done using an insulated cabinet, for example.

Alternatively, after stamping or cutting 120 the blank from the uncoiled strip, a varying magnetic field is applied 128 to the steel. This may be applied using an induction heating system, for example. The temperature is maintained at 650 ℃ ± 100 ℃ for 15 minutes or less while the varying magnetic field is applied.

In an alternative embodiment, during the isothermal holding step 118, a magnetic field is applied to the steel by using a static magnetic field or a varying magnetic field.

The resulting modulus-enhanced nanostructured steel is then air cooled 130 and ready for cold stamping and/or forming as desired.

Alternatively, instead of uncoiling the steel and stamping the blank, the entire coil may be reheated (i.e., austenitized) 122, cooled 124, and then held at 650 ℃ ± 100 ℃ for the required time. This can be done using an insulated box (e.g., a hood-type furnace suitable for applying heat treatment to all coils) as a batch annealing process.

The method is applicable to both ferritic and pearlitic steels.

Example 5

Fig. 10 illustrates a method of manufacturing a modulus-enhanced hot rolled high carbon wire steel. As described in relation to example 4, the cast steel billet is heated 202 to a reheating temperature.

The heated steel is then hot rolled 206 using a roughing mill, an intermediate mill, a finishing mill, and/or a torsion-free V-wire mill at a temperature in the range of about 840 ℃ to about 1200 ℃. Hot rolling is performed above the recrystallization stop temperature so that grain deformation is kept to a minimum.

After hot rolling is completed, the steel is rapidly cooled 214 at a cooling rate in the range of about 2 ℃ per second to about 80 ℃ per second until a temperature of 650 ℃ ± 100 ℃ is reached.

The temperature is then maintained 218 at 650 ℃ ± 100 ℃, for example using an on-line induction heating coil. Referring to fig. 11, the steel is then air cooled 230.

Optionally, the steel may then be cold drawn 132. In this process, the steel wire is induction heated and austenitized, air cooled to 600 ℃ ± 50 ℃, and then maintained at 650 ℃ ± 200 ℃ for less than 5 minutes. These steps are applied in a continuous, slow-going manner. Optionally, in order to refine the interlayer spacing of the pearlite, a static magnetic field may also be applied at this stage, which may further contribute to lattice strain and modulus enhancement.

Unless otherwise indicated, each of the integers described herein may be used in combination with any other integer as will be understood by one skilled in the art. Furthermore, although all aspects of the invention preferably "comprise" features described in relation to that aspect, it is specifically contemplated that they may be "composed" or "substantially composed" of those features outlined in the claims. Furthermore, unless specifically defined herein, all terms are intended to be given their commonly understood meanings in the art.

Furthermore, in the discussion of the present invention, disclosure of alternative values for the upper or lower limit of the allowable range of parameters should be construed as an implied statement as follows unless otherwise indicated: each intermediate value of the parameter between the smaller alternative value and the larger alternative value is itself also disclosed as a possible value of the parameter.

Furthermore, unless otherwise indicated, all numbers appearing in this application are to be understood as modified by the term "about".

Claims

1. A vanadium alloy steel comprising carbon and vanadium, wherein the microstructure of the steel comprises coherent interphase precipitates comprising vanadium.

2. Vanadium alloy steel according to claim 1, comprising (in weight-%):

carbon in the range of about 0.06 to about 1.1; and

vanadium in the range of about 0.1 to about 1.5.

3. The vanadium alloy steel according to claim 1 or 2, wherein the steel is a modulus-strengthened steel having an average elastic modulus of greater than about 210GPa, optionally wherein the steel has an average elastic modulus of up to about 300 GPa.

4. A vanadium alloy steel according to claim 1, 2 or 3, wherein the particle size of the coherent interphase precipitate is about 9nm or less.

5. The vanadium alloy steel according to claim 4, wherein the particle size of the coherent interphase precipitate is in the range of about 5nm to about 9 nm.

6. A vanadium alloy steel according to any preceding claim, wherein the coherent interphase precipitate comprises vanadium carbide.

7. The vanadium alloy steel according to claim 6 wherein the coherent interphase precipitate comprises a compound of formula V _x C _y Wherein x > y (e.g., the coherent interphase precipitate includes V) ₄ C ₃ )。

8. The vanadium alloy steel according to claim 7, wherein the coherent interphase precipitate comprises V ₆ C ₅ And/or V ₅ C ₃ 。

9. Vanadium alloy steel according to any one of the preceding claims, wherein the composition of the steel comprises at least one or more of the following (in weight%):

about 0.015 or less of nitrogen,

about 1.6 or less of molybdenum,

about 1 or less of the copper is present,

about 1.2 a or less of silicon,

about 0.3 or less of chromium, and/or

About 1.6 or less.

10. The vanadium alloy steel according to any one of the preceding claims, wherein the composition of the steel comprises:

carbon in the range of about 0.06 to about 1.1,

vanadium in the range of about 0.1 to about 1.5,

about 0.015 or less of nitrogen,

about 1.6 or less of molybdenum,

about 1 or less of the copper is present,

about 1.2 a or less of silicon,

about 0.3 or less, and

about 1.6 or less.

11. Vanadium alloy steel according to claim 9 or 10, wherein the interphase precipitate comprises Mo, optionally wherein the interphase precipitate comprises a compound of formula (Mo, V) _x C _y Wherein x > y, e.g. (Mo, V) ₄ C ₃ /(Mo,V)C。

12. The vanadium alloy steel according to claim 9, 10 or 11, wherein the microstructure of the steel comprises VN precipitates, optionally wherein the microstructure comprises intragranular VN nucleated acicular ferrite.

13. A vanadium alloy steel according to any preceding claim, wherein the steel comprises a ferrite phase and the coherent interphase precipitate is formed in the ferrite phase to form spherical or nodular ferrite.

14. The vanadium alloy steel according to claim 13, wherein the ferrite phase comprises grains having an average size of less than about 20 μιη, such as in the range of about 5 μιη to about 20 μιη.

15. A vanadium alloy steel according to any preceding claim wherein the steel comprises single phase ferritic steel, such as HSLA, AHSS.

16. The vanadium alloy steel according to any one of claims 1 to 14, wherein the steel comprises pearlite phase, optionally wherein the steel comprises vanadium enhanced cementite.

17. Vanadium alloy steel according to claim 16, wherein the vanadium enhanced cementite comprises vanadium dissolved in cementite, for example to form Fe ₂ VC and/or FeV ₂ C。

18. The vanadium alloy steel according to any one of the preceding claims, wherein the steel has a tensile strength in the range of about 360Mpa to about 2000 Mpa.

19. A method for preparing the vanadium alloy steel of any preceding claim, the method comprising:

a. providing a starting steel;

b. Heating the starting steel to at least partially austenitize the starting steel;

c. the temperature is maintained at 650 ℃ ± 200 ℃ for about 25 minutes or less.

20. The method of claim 18, wherein the starting steel has a composition (in weight%) comprising:

carbon in the range of about 0.06 to about 1.1, and

vanadium in the range of about 0.1 to about 1.5;

and optionally one or more of the following (in wt.%):

about 0.015 or less of nitrogen,

about 1.6 or less of molybdenum,

about 1 or less of the copper is present,

about 1.2 a or less of silicon,

about 0.3 or less of chromium, and/or

About 1.6 or less.

21. The method of claim 19 or 20, wherein the steel is cooled, optionally to a temperature of 650 ℃ ± 200 ℃ at a cooling rate in the range of about 2 ℃/sec to about 80 ℃/sec, before maintaining the temperature at 650 ℃ ± 200 ℃.

22. A method according to claim 19, 20 or 21, wherein after heating the starting steel, the steel is hot rolled at a temperature above the Recrystallization Stop Temperature (RST), optionally by recrystallization controlled rolling and/or V (C, N) precipitation controlled rolling.

23. The method of claim 22, wherein hot rolling the steel comprises hot rolling at substantially the V (C, N) precipitation temperature time inflection point.

24. A method according to any one of claims 19 to 23, wherein a magnetic field is applied to the composition of the steel during or after maintaining the temperature.

25. The method of any one of claims 19 to 24, wherein the steel is reheated to at least partially austenitize the steel after maintaining the temperature at 650 ℃ ± 200 ℃ for about 25 minutes or less.

26. The method of claim 25, wherein after reheating the steel, the temperature is again maintained at 650 ℃ ± 200 ℃ for about 25 minutes or less.

27. The method of claim 25 or 26, wherein after reheating the steel, the steel is cooled, optionally to a temperature of 650 ℃ ± 200 ℃ at a cooling rate in the range of about 2 ℃/sec to about 80 ℃/sec.