Field

The present disclosure relates to an alloyed steel and a method for producing the same. More particularly to a vanadium alloyed steel having an enhanced elastic modulus.

Background

Given decarbonisation trends towards using stronger and lighter-gauge structural steel plate, there exists the need to develop steels that offer improved properties, such as fire-resistance, buckling resistance and deflection resistance.

The development and use of fire-resistant structural steels in civil engineering is a particular problem in the construction industry. In addition, there is a strong trend towards reducing the use of embodied carbon in the steel industry.

The development of a lean-alloyed high-strength steel with superior fire-resistance would be of great economic interest, enabling the use of more slender columns and beam designs, and accelerating progress towards global decarbonisation goals. However, the fire-resistance performance requirements and decarbonisation goals are often incompatible.

Furthermore, provision of steels that offer enhanced buckling and deflection resistance would be beneficial in both the automotive industry and in the construction of wind towers, for example.

It is an aim of the current disclosure to address or at least mitigate the problems associated with the prior art.

Summary

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

In some embodiments a vanadium alloyed steel is provided comprising, preferably consisting essentially of, more preferably consisting of carbon, vanadium and balance iron; wherein the steel microstructure comprises coherent interphase precipitates comprising vanadium.

In some embodiments, the vanadium alloyed steel comprises, preferably consists essentially of, more preferably consists of (in wt%):

Carbon in the range of from about 0.06 to about 1.1; and

Vanadium in the range of from about 0.1 to about 1.5; wherein the steel microstructure comprises coherent interphase precipitates comprising vanadium.

In some embodiments, the vanadium alloyed steel comprises, preferably consists essentially of, more preferably consists of (in wt%):

Carbon in the range of from about 0.06 to about 1.1;

Vanadium in the range of from about 0.1 to about 1.5; and balance iron, wherein the steel microstructure comprises coherent interphase precipitates comprising vanadium.

In some embodiments, the vanadium alloyed 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 balance iron; wherein the steel microstructure comprises coherent interphase precipitates comprising vanadium.

In some embodiments, the vanadium alloyed steel comprises, preferably consists essentially of, more preferably consists of (in wt%):

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

Vanadium in the range of from about 0.1 to about 1.5; optionally one or more of: about 0.015 or less Nitrogen, about 1.6 or less Molybdenum, about 1 or less Copper, about 1.2 or less Silicon, about 0.3 or less Chromium, and/or about 1.6 or less Manganese; and balance iron; wherein the steel microstructure comprises coherent interphase precipitates comprising vanadium. As will be described below, in the steels disclosed herein, the average elastic modulus (i.e. Young's modulus) of the steel is enhanced. It is believed that the coherent vanadium precipitates generate large strain fields in the steel, resulting in latticeparameter changes which can modify the bulk elastic modulus of the steel.

The enhancement of the elastic modulus is believed to be applicable at high- temperatures, due to the greater stability of coherent precipitates against tempering and coarsening, thereby enhancing the buckling resistance of the steel. In this way, improved fire-resistant performance is achieved, for example using lower alloying embodiments.

The provision of higher modulus steels as disclosed herein may also be beneficial in providing enhanced structural benefits such as greater buckling resistance, for example using higher alloyed embodiments. This may enable construction of taller wind-turbines.

The provision of higher modulus steels as disclosed herein may also be beneficial in providing enhanced deflection resistance, for example using higher alloyed embodiments. This may be beneficial in creating lighter, stiffer automotive structures.

As illustrated in Figure 1, it is believed that the lattice-parameter change effect is a direct consequence of the matrix dilatation and contraction caused upon elastic accommodation of the precipitates in the surrounding steel matrix. The precipitates comprising vanadium 2 have a different lattice parameter to that of the surrounding steel (e.g. ferrite) matrix 4. In other words, there is a lattice-parameter misfit between the precipitates comprising vanadium and the surrounding steel, which results in the generation of large strain fields (shown by the shaded region indicated with reference numeral 6 in Figure 1) when coherent precipitates comprising vanadium are formed. Consequently, the average lattice-parameter of the overall steel is altered, modifying the bulk elastic modulus of the steel. In contrast, where incoherent precipitates are formed, no such strain field is observed.

It is believed that this change in lattice parameter is anisotropic. However, as would be expected in a polycrystalline material such as steel, random grain orientations can neutralise this anisotropy, which facilitates a modulus enhancement on a macro scale. Combined with this physical effect is a modification of the electron distribution within vicinity of the coherent interfaces, which further contributes to enhancing the bulk elastic modulus.

Figure 2 demonstrates the uniqueness of vanadium in having relatively low misfits between the lattice parameter of precipitates comprising vanadium as compared to ferrite and austenite, for example. The data points relating to the misfit in ferrite are shown by the solid circles and triangles, and the data points relating to the misfit in austenite are shown by the outlined circles and triangles. Further, Figure 2 illustrates the high limiting sizes for coherent precipitates comprising vanadium, in this case in ferrite, as compared to precipitates including other elements commonly used in steel. It is believed that this promotes the generation of coherent precipitate formation, since the greater coherency limit has the effect that it is easier to maintain coherency via established steel production methods.

Turning to Figure 3, this illustrates a relationship between strengthening effects and the size of precipitates. As can be seen, the shear strength increases with increasing particle size up to a critical particle size. This is due to coherency strain hardening, as described above. When the size of the particles exceeds the critical particle size, incoherent precipitates are formed and the modulus strengthening effect is no longer seen.

It is believed that precipitates comprising vanadium are very efficient in increasing the elastic modulus of steel by function of their high coherency limit and high shear modulus, as compared to other precipitates.

When the precipitates are small and coherent with the surrounding matrix, then the dislocation mechanism shifts from the well-established ashby-orowan looping mechanism (which is applicable for incoherent precipitates) to a particle shearing mechanism. The strengthening mechanisms attributed to dislocations shearing the ordered coherent particles results in the operation of order and modulus strengthening mechanisms. The resulting modulus strengthening increment can be described by the Knowles-Kelly equation (1).

Where, b « 0.248 nm is the Burgers vector of the dislocation, r is the mean radius of the precipitates comprising vanadium, e.g. VC particles (~ 3.1±1.0nm), and G « 81.6 GPa is the shear modulus of the ferrite matrix. AG is the difference in the shear modulus between the matrix and precipitates (175.7GPa-81.6GPa=94.1GPa). The f = (4/3)nnr ³ is the volume fraction of the VC nano-precipitates, where n is the number density of the precipitates (~ 1.18±0.03xl0' ⁴ nm' ³ ).

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

By way of example, Figure 5a shows a simulation of a ferritic steel as disclosed herein and illustrates the elastic strain fields associated with coherent interphase precipitates comprising vanadium, which are formed as nano-sized platelets in this embodiment. The platelets are arranged in rows with the spacing between rows indicated by the reference numeral 22. Only two rows of interphase precipitates are shown for the purposes of clarity.

In precipitates containing vanadium, such as vanadium carbides and/or vanadium carbonitrides, there is a large relative difference in atomic misfit parallel and perpendicular components relative to the precipitate's orientation. Parallel to the platelet precipitate's surface there exists a very low misfit, which translates into a large coherency critical size limit and a dilation effect. Perpendicular to the precipitate, we observe a relatively larger misfit, and lower coherency size limit, which results in a contraction effect on the lattice spacing perpendicular to the broad platelet surface of the precipitate. This contraction reduces the average lattice spacing i.e. the burgers vector, which in turn increases the elastic modulus of the unit cell, which when applied over multiple grains on a macro scale, results in a noticeable increase in the Young's modulus on the macro level.

The change in lattice-parameter is shown along line A-A. This illustrates that the average lattice spacing is contracted as a result of anisotropic misfit strains, which modifies the bulk elastic modulus.

In some embodiments, the steel is a fire-resistant steel. The enhancement of the elastic modulus as described above is applicable at high-temperatures, due to the greater stability of coherent precipitates against tempering and coarsening, thereby enhancing the buckling resistance of fire-resistant steels. The stability of buildings and other structures in a fire depends on the extent to which the steel structures soften when heated to the temperatures created by the fire. A steel is generally regarded as being fire-resistant if its short-term strength when heated to such temperatures is approximately 0.6-0.7 of its strength at room temperature.

In exemplary embodiments, the average elastic modulus is 120GPa or higher at temperatures in the range from about 600°C about 700°C.

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

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

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

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

The use of vanadium to form interphase precipitates is considered to be advantageous as compared to other alloying elements such at niobium and titanium, since vanadium exerts a relatively low influence on the transformation temperature (i.e. the temperature at which the steel transforms from austenite to ferrite). Lowering the transformation temperature has the disadvantageous effect of increasing the probability of random precipitation and not interphase precipitation. Accordingly, use of vanadium is beneficial in maintaining the transformation temperature in a desired range.

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

In some embodiments, the steel comprises carbon in the range of from about 0.06 to about 0.2 (wt%). For example, the steel is a ferritic steel (e.g. fire-resistant alloy HSLA (high strength low alloy)). In some embodiments, the steel comprises carbon in the range of from 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 from about 0.06 to about 0.3 (wt%). For example, the steel is a ferritic steel (e.g. a hydrogen resistant steel).

In some embodiments, the steel comprises carbon in the range of from about 0.6 to about 1.1 (wt%). For example, the steel is a pearlitic steel.

In some embodiments, the steel is a modulus-enhanced steel comprising an average elastic modulus (i.e. a measured bulk elastic modulus) of above about 210 Pa, for example about 220GPa or higher, for example in the range of from about 210GPa to about 300GPa, e.g. in the range of from about 220GPa to about 300GPa.

The elastic modulus may be measured using dynamic resonance methods. A test-piece of steel with a uniform cross-section (e.g. round, square or rectangle) is prepared and the characteristic vibration frequencies of the steel are determined and related to the elastic modulus using known equations. ASTM standard C1259 sets out the standard test method for Dynamic Young's Modulus, Shear Modulus, and Poisson's Ratio for Advanced Ceramics by Impulse Excitation of Vibration, for rectangular and round cross section bars.

Optionally, the steel is a modulus-enhanced AHSS steel. Optionally, the steel is a fire- resistant steel (e.g. a HSLA steel) comprising an average elastic modulus of about 120 GPa or higher at temperatures in the range of from about 600°C to about 700°C.

Optionally, the coherent interphase precipitates have a particle size of about 9nm or less, optionally in the range of from about 5nm to about 9nm, optionally greater than about 5nm This can be measured with high-resolution transmission electron microscopy, X-ray diffraction, or any other suitable method.

The coherent interphase precipitate particle size is taken to be the mean value of the largest particle dimension of a representative sample of particles.

For coherent interphase precipitation in ferrite, lamellae specimens are typically prepared using a focused ion beam (FIB) microscope. Analysis of thin lamellae is also typically carried out using a cold field emission guns in-situ with probe-corrected high- resolution scanning transmission electron microscope (HR STEM). Through this method, high angle annular dark-field (HAADF) images can be obtained at suitable collection angles. The sizes, morphology and number density of coherent precipitates can be obtained given such HAADF and STEM micrographs, whereby the average number density is derived from averages taken from different grains. Concentration profiles can also be obtained via such HAADF data, which can verify the nature of our combined nano-particles. Atomic probe tomography, can also be used to verify and offer improved 3D spatial resolution of the atomic concentration profiles (from HAADF micrographs) in the core-shell nanoparticles.

As shown in Figure 2 and as is described above, this is relatively large as compared to precipitates formed from other commonly used elements and is believed to contribute to an increased elastic modulus of the bulk steel (as shown in Figure 3).

Optionally the coherent interphase precipitates comprise vanadium carbides, optionally wherein the coherent interphase precipitates comprise precipitates having enhanced levels of vanadium with the chemical formula V _x C _y , where x>y (e.g. the coherent interphase precipitates comprise V4C3, VeCs, and/or V5C3). During the transition from coherent to incoherent particles, the precipitate-matrix interfacial energy increases thus a larger fraction of vacancies is necessary for the precipitate stability, causing a slight gradual decrease in the carbon-to-vanadium ratio, and thus, increases their stability.

It is also believed that small increases in the vanadium content within ferritic steels can reduce the carbon to vanadium ratio within the precipitates, which serves to bolster the precipitate shear modulus, resulting in modulus strengthening.

In some embodiments the steel is a hydrogen-resistant steel, for example using higher alloyed embodiments. Vanadium carbides within the steel are believed to act as hydrogen traps, which can irreversibly trap hydrogen.

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

This is illustrated in Figures 4a and 4b. Figure 4b shows a 3D illustration of hydrogen atoms 8 occupying carbon vacancies in a V4C3 coherent interphase precipitate 2. Figure 4a illustrates the potential energy wells associated with a TiC precipitate 14, a NbC precipitate 12 and a V4C3 precipitate 10. As can be seen, the potential energy well associated with V4C3 precipitates 10 is deeper than that of the other precipitates, and hence acts as a more effective hydrogen trap.

Further, it is believed that vanadium carbide precipitates containing enhanced vanadium contain carbon vacancies which are inherently more thermally stable under fire exposure conditions.

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

In some embodiments, the coherent interphase precipitates are formed as an austenite phase transitions to a ferrite phase, in other words at conditions corresponding to a migrating phase boundary between austenite and ferrite.

Upon cooling, in order to accommodate accumulated elastic strain caused by formation of coherent interphase precipitates, it is believed that nodular ferrite (NF) can be formed, which changes its orientation during growth. Nodular Ferrite has only previously been identified in eutectic cast steels, caused due to lattice parameter misfit resulting from iron carbides (i.e. in cementite). It has been found that nodular ferrite can also be created due to the misfit generated by coherent alloy carbides. Nodular ferrite is characterised by chaotic grain orientations and chaotic grain boundary misorientations. The grains are not isomorphs and resemble a knotted appearance. Ferrite transformation with interphase precipitation is a kind of eutectoid transformation analogous to degenerate pearlite in high carbon steels, where alloy carbide is formed instead of ironcarbide, cementite. Therefore, it is believed that this nodular or knotted ferrite can be considered analogous to that of pearlite observed in high carbon steels.

This branching of NF, and thus variation of ferrite/austenite interface planes for interphase precipitation, is believed to optimise the texture of the microstructure towards greater hydrogen embrittlement resistance. It is believed that low-angle grain boundaries and Coincidence Site Lattices (CSLs), which exhibit lower relative grain boundary energies, can significantly bolster the hydrogen embrittlement resistance of ferritic steels. One possible mechanism for this is that the number of atoms arranged on the CSL grain boundaries (e.g. between Z5 and £13) is increased, which could decrease the vacancy density and further mitigate hydrogen atom segregation at the grain boundary. 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 nano-particles, which exhibit a higher capillary-driven coarsening resistance and a lattice parameter contraction. With reference to Figure 7b, in some exemplary embodiments, core-shell nano-particles 16 may comprise a core 18 of VN. For example, the core-shell nano-particles may comprise a shell 20 comprising Mo, Cr and/or Cu.

Capillary-driven particle coarsening is impeded due to the interfacial segregation of Mo, Cr and/or Cu, which reduces the driving force for further particle coarsening due to a reduction in interface energy. The coarsening resistance is conducive to grain refinement and prior austenite grain size control during austenisation. Developing core-shell nanoparticles are inherently resistant to coarsening and exert a zener pinning mechanism on the austenite grain boundary to limit grain size.

Accordingly, improved toughness is achieved.

Copper is thought to have low misfits with the ferrite lattice, which enable it to contribute to elastic misfit strains by forming a high density of coherent interphase particles. It is thought that the low magnetic moment of Cu plays a positive role in increasing the number density of complex nano-particles. 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, precipitation is not hindered by the magnetic field.

In some embodiments, the steel comprises copper in the range of from greater than 0 to about 0.5 (wt%). For example, the steel is a ferritic steel (e.g. fire-resistant alloy HSLA (high strength low alloy)).

In some embodiments, the steel comprises copper in the range of from 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 from greater than 0 to about 1 (wt%). For example, the steel is a ferritic steel (e.g. a hydrogen resistant steel). In some embodiments, the steel comprises copper in the range of from greater than 0 to about 0.5 (wt%). For example, the steel is a pearlitic steel.

The presence of manganese is thought to contribute to refinement of ferrite gains, for example when higher coiling temperatures are used.

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

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

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

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

In some embodiments, the steel comprises silicon in the range of from greater than 0 to about 0.5 (wt%). For example, the steel is a ferritic steel (e.g. fire-resistant alloy HSLA (high strength low alloy)).

In some embodiments, the steel comprises silicon in the range of from greater than 0 to about 0.5 (wt%). 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 from greater than 0 to about 1.5 (wt%). For example, the steel is a ferritic steel (e.g. a hydrogen resistant steel). In some embodiments, the steel comprises silicon in the range of from greater than 0 to about 1.2 (wt%). For example, the steel is a pearlitic steel.

The presence of chromium is thought to increase the solubility of other microalloying elements and reduce the precipitate sizes. In this way, the mean precipitate size can be kept below the limiting size for coherency, resulting in a greater proportion of coherent precipitates.

In some embodiments, the steel comprises chromium in the range of from greater than 0 to about 0.3 (wt%). For example, the steel is a ferritic steel (e.g. fire-resistant alloy HSLA (high strength low alloy)).

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

In some embodiments, the steel comprises chromium in the range of from greater than 0 to about 0.1 (wt%). For example, the steel is a pearlitic steel.

The presence of molybdenum is thought to contribute to increasing the number density of the coherent interphase particles, by forming a complex with the precipitates which can improve diffusion kinetics of the precipitates.

Optionally, the coherent interphase precipitates comprise Mo, optionally wherein the coherent interphase precipitates comprise precipitates having the chemical formula (Mo,V) _x Cy, where x>y, e.g. (Mo,V) ₄ C3/(Mo,V)C

It is believed that this has the effect of reducing the lattice mismatch between the interphase precipitate and the ferritic phase, resulting in improved coherency.

Further, it is believed that the presence of Mo reduces the migration speed of the austenite to ferrite interphase during the transformation, which enables a higher number of interphase precipitates to form. In some embodiments, the steel comprises molybdenum in the range of from greater than 0 to about 0.2 (wt%). For example, the steel is a ferritic steel (e.g. fire-resistant alloy HSLA (high strength low alloy)).

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

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

Optionally, the ferrite phase comprises grains having a mean size of less than about 20pm. For example, having a mean grain size in the range of from about 5 pm to about 20 pm. Grain size can be measured by any suitable method, for example using a scanning electron microscope and electron-backscattering diffraction. The intercept method may then be used to calculate the average grain size. In this method, a straight line is drawn on a micrograph and the number of grain boundaries crossing the line are counted. The average grain size is found by dividing the number of intersections by the actual line length.

Optionally, the steel microstructure comprises VN precipitates, optionally wherein the microstructure comprises intra-granular VN nucleated acicular ferrite (e.g. V(C,N)). Optionally the VN precipitates form in austenite as intra-granular nucleants for acicular ferrite formation. It is believed that this is due to the lower solubility of VN in austenite, relative to Vanadium Carbides.

It is believed that the intragranular VN nucleated acicular ferrite structure improves the coherency of the interphase particles and refines the ferrite grains, resulting in the consequential benefits to the elastic modulus and toughness of the steel.

In this way, an acicular microstructure is formed simultaneously with a system of nanoscale interphase precipitates. This acicular microstructure is illustrated schematically in Figure 7a. In some embodiments, the steel comprises nitrogen in the range of from greater than 0 to about 0.015 (wt%). For example, the steel is a ferritic steel (e.g. fire-resistant alloy HSLA (high strength low alloy)).

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

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

In some embodiments, where not all the nitrogen is used to form VN precipitates as intra- granular nucleants for acicular ferrite formation, interphase vanadium carbonitrides may form as coherent interphase precipitates. The nitrogen will also further increase driving force for precipitation, which is beneficial for obtaining a high number density of interphase precipitates.

Optionally, the steel comprises a single phase ferritic steel e.g. HSLA, AHSS and/or hydrogen resistant steel.

Optionally, the steel comprises a pearlitic steel. Accordingly, the steel comprises a ferritic phase and a cementite phase. The modulus of the ferrite phase is enhanced via the presence of coherent interphase particles and the mechanisms described above.

Optionally, the pearlitic steel comprises a vanadium enhanced cementite phase, optionally, the vanadium enhanced cementite phase comprises Vanadium dissolved in cementite e.g. to form FezVC and/or FeVzC

The average lattice spacing in the ferrite phase is contracted as a result of misfit strains from both coherent interphase precipitates comprising vanadium, and cementite. In addition, vanadium can easily dissolve in cementite given an appropriate isothermal holding regime, which can further contribute to misfit-induced lattice strain due to the change in the cementite lattice parameter. Cementite is notorious for its anisotropy and this applies also to the anisotropy of its elastic modulus components as a function of crystallographic orientation. The dissolution of vanadium into cementite reduces this anisotropy. It is believed that this will evolve as a function of isothermal holding duration as the composition of the cementite changes with ageing.

The lattice strain associated with the elastic misfits of alloy carbides or V-enhanced cementite can be verified by the use of high-resolution scanning transmission electron microscopy (HR STEM) and X-ray diffraction (XRD), respectively. Atomic Probe Tomography could also be used as an additional method to quantify atomic concentrations within 3-dimensions.

For the lattice strain associated with iron carbides and the formation of the strain induced by V-enhanced cementite in higher carbon pearlite steels, the use of electron backscattered diffraction (XRD) is advised. The elastic strain is quantified from the full width at half maximum (FWHM) of the ferrite diffraction peak measured by XRD. The lattice strain in lamellar ferrite, which is quantified by XRD measurement, would also correlate with the proof stress. The modulus enhancement would contribute also to the proof stress, which can be conventionally obtained by tensile testing equipment.

With reference to Figure 5b, an exemplary embodiment is illustrated in which the steel comprises cementite colonies. The additional strain applied to the ferrite lattice is illustrated due to cementite is illustrated. This results in a reduction in the average lattice spacing of the ferrite matrix in regions proximal a cementite colony.

The change in lattice-parameter is shown along line A-A. This illustrates that the average lattice spacing is contracted as a result of misfit strains from both coherent vanadium- alloy carbides and cementite, which modifies the bulk elastic modulus.

Figure 5b also shows a schematic illustration of the microstructure of pearlite including V-enhanced cementite.

Optionally, the steel comprises a tensile strength in the range from about 690 MPa to about 2000 MPa, for example in the range from about 690 MPa to about 1800 MPa.

The tensile strength may be measured by gripping the ends of a suitably prepared standard test piece in a tensile test machine and then applying a continuously increasing uni-axial load until such time as failure occurs. ASTM E8/E8M-13: "Standard Test Methods for Tension Testing of Metallic Materials" may be used.

The net result of greater number density and optimized precipitate compositions can tangibly impact the applicability of the steels disclosed herein for fire-resistant or hydrogen transport/storage applications.

The structure of the steel alloy described herein can be determined by conventional microstructural characterisation 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 alloyed steel as disclosed herein is provided, the method comprising: a. providing a starting steel; b. heating the starting steel to at least partially austenitise the steel; c. holding the temperature of the steel at 650°C ± 200°C for about 25 mins or less; wherein the vanadium alloyed steel comprises coherent interphase precipitates comprising 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 wt%):

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

Vanadium in the range of from about 0.1 to about 1.5.

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

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

Vanadium in the range of from about 0.1 to about 1.5; and balance 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 balance iron.

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

Carbon in the range of from about 0.06 to about 1.1, and Vanadium in the range of from about 0.1 to about 1.5; one or more of: about 0.015 or less Nitrogen, about 1.6 or less Molybdenum, about 1 or less Copper, about 1.2 or less Silicon, about 0.3 or less Chromium, and/or about 1.6 or less Manganese; and balance iron.

Optionally, heating the starting steel to at least partially austenitise the steel involves raising the temperature to at least about 900°C, optionally to a temperature in the range of from about 920° to about 1300°C, for example to a temperature in the range of from about 1100°C to about 1300°C.

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, said heating may be carried out by inductive heating and/or a roller hearth furnace.

Holding the temperature of the steel at 650°C ± 200°C for about 25 mins or less (i.e. isothermal holding) enhances the number density of coherent interphase precipitates, which results in enhanced elastic modulus (in accordance with equation (1) above).

It has been found that precipitate formation at the ferrite/austenite interface is impeded when the interface advances rapidly. Accordingly, isothermal holding in this manner enhances interphase precipitation (e.g. precipitation of vanadium carbide and/or vanadium carbonitride), which consequently improves the properties of the steel. The sheet spacing between the interphase precipitate rows decreases with increasing transformation time. Accordingly, isothermal holding has the effect of decreasing the spacing between rows of interphase precipitates, thereby increasing the number density of interphase precipitates.

Furthermore, isothermal holding in this manner is believed to have the effect of increasing the ratio of V:C in the interphase precipitates, such that a large degree of carbon vacancies are created at the coherent interfaces of the steel (e.g. ferrite) lattice matrix. These vacancies act as hydrogen traps, leading to enhanced hydrogen embrittlement resistance (for example, see figure 4a and 4b).

In some embodiments, the temperature is held at 650°C ± 200°C for about 20 mins or less.

In some embodiments, the temperature is held by placing the steel in an insulated thermal box.

In some embodiments, batch annealing is carried out, for example using an insulated thermal box to control the temperature applied to the steel. For example, a bell furnace may be used to apply heat treatment to a full coil.

In some embodiments, the temperature is held by using an induction heating system, e.g. an induction furnace.

Optionally, the temperature is held at 650°C ± 200°C for about 15 minutes or less. In some embodiments, the temperature is held at 650°C ± 200°C for in the range of from about 5 minutes to about 20 minutes, e.g. from about 10 to about 20 minutes.

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

Optionally, the temperature is held at 650°C ± 100°C.

Optionally, the steel is cooled at a cooling rate in the range of from about 2°C/s to about 80°C/s, prior to the temperature being held at 650°C ± 200°C, optionally at a cooling rate in the range of from about 2°C/s to about 50°C/s.

It has been found that precipitate formation at the ferrite/austenite interface is impeded when the interface advances rapidly. Accordingly, controlling the cooling rate in this manner enhances interphase precipitation (e.g. vanadium carbide and/or vanadium carbonitride precipitation), which consequently improves the properties of the steel. Furthermore, rapidly cooling the steel in this way prevents or impedes the formation of undesirable phases and/or promotes the formation of acicular ferrite.

In some embodiments, the cooling is carried out by air cooling, forced air cooling, lamellar cooling, and/or any other suitable cooling method.

Where Mn is present in the steel composition, this may facilitate higher number density at faster cooling rates.

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

By carrying out hot rolling at temperatures above the RST, grain deformation is limited, thereby facilitating the creation of interphase precipitates.

Optionally, hot rolling the steel is carried out prior to cooling the steel to 650°C ± 200°C.

In some embodiments, hot rolling is carried out at a temperature(s) in the range of from about 820°C to about 1200°C.

In some embodiments, hot rolling comprises roughing, intermediate and/or finishing steps.

Optionally, for compositions including relatively high amounts of Nitrogen, V(C,N) precipitation controlled rolling may be used. This achieves low grain deformation levels.

In some embodiments, hot rolling the steel comprises hot rolling at substantially the V(C,N) precipitation temperature time nose.

For example, where V(C,N) precipitation controlled rolling is used, this may be carried out at substantially the V(C,N) precipitation temperature time nose.

At the V(C,N) precipitation temperature time nose, the system has optimal free energy for VN to precipitate. Accordingly, hot rolling at this temperature facilitates the formation of acicular ferrite. The V(C,N) precipitation temperature time nose depends on the particular composition of the steel, but is typically around 850°C.

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

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 variable magnetic field e.g. applied by an induction heating system. In some embodiments, the variable magnetic field is applied by an inductive heating system, which also controls the temperature of the steel (e.g. is used to hold the temperature at 650°C ± 200°C).

The application of a magnetic field facilitates increasing the number density of coherent precipitates and also optimises the composition of the constituent nano-sized precipitates.

In practice, long isothermal holding durations can be detrimental to productivity. The proposed environmentally-friendly method to accelerate this isothermal hold duration is to utilize a magnetic field, which can act to decelerate the austenite to ferrite transformation kinetics (thus refining the spacing between precipitates and maximize number density of interphase precipitates). This can also increase the extent of higher V:C ratio carbides due to paramagnetic effects.

Vanadium in solution, prior to precipitation, acts to increase the Curie temperature, Tc, (but exerts a small effect on the magnetic moment and small interaction energy with the external magnetic field) so that the range of temperature between Tc and a given transformation temperature below the Curie point is increased and therefore, the magnetisation is larger at the transformation point. The external magnetization energy serves to significantly increase the Gibbs free energy of the system, increasing the free energy of the total nucleation of the nanoparticles, which subsequently increases the nucleation rate and significantly increases the quantitative number density of coherent interphase vanadium precipitates.

In steel compositions comprising higher amounts of carbon (i.e. pearlitic steel), the application of a magnetic field also acts to refine the interlamellar spacing of the pearlite (due to an increase in the nucleation rate of cementite). This can further contribute to lattice strain and modulus enhancement, as a result of strain applied to the ferrite lattice in regions proximal cementite colonies, as previously described. Optionally, after the temperature has been held at 650°C ± 200°C for about 25 mins or less, the steel is reheated to at least partially austenitise the steel.

Optionally this reheating step comprises raising the temperature to in the range of from about 920°C to about 1150°C.

Optionally, after the steel has been reheated, again holding the temperature at 650°C ± 200°C for about 25 mins or less.

In some embodiments, the temperature is held at 650°C ± 200°C for about 20 mins or less. Optionally, the temperature is held at 650°C ± 200°C for about 15 minutes or less. In some embodiments, the temperature is held at 650°C ± 200°C for in the range of from about 5 minutes to about 20 minutes, e.g. in the range of from about 10 minutes to about 20 minutes. Optionally, following isothermal holding, the steel is air cooled.

Optionally, the temperature at is held at 650°C ± 100°C.

Isothermal holding in this manner enhances interphase precipitation (e.g. vanadium carbide and/or vanadium carbonitride precipitation), as previously described.

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

Optionally the cooling rate is in the range of from about 2°C/s to 50°C/s. Controlled cooling enhances interphase precipitation (e.g. vanadium carbide and/or vanadium carbonitride precipitation) and inhibits the formation of undesirable phases, as previously described.

On reheating, at least a portion of the precipitates previously generated in the steel will remain, therefore, the step of reheating, subsequent cooling and/or subsequent isothermal holding will result in an increased number density of precipitates in the final steel.

Figures

Figure 1 illustrates a schematic showing the interaction of a coherent precipitate and an incoherent precipitate with a surrounding ferrite lattice matrix; Figure 2 shows a comparison of lattice parameter misfits and limiting coherent precipitate sizes for precipitates including Ti, V, Zr or Nb;

Figure 3 shows a relationship between strength and precipitate particle size;

Figures 4a and 4b are schematic illustrations of the capacity of coherent interphase precipitates to act as a hydrogen trap;

Figure 5a illustrates the impact of a strain field in a ferrite matrix on the average lattice parameter in a ferrite steel;

Figure 5b illustrates the impact of a strain field in a ferrite matrix on the average lattice parameter in a pearlite steel;

Figure 6 illustrates three typical beam compositions (in accordance with Eurocode 3);

Figure 7a shows a schematic illustration of a VN-nucleated acicular ferrite structure;

Figure 7b shows a schematic illustration of a core-shell nano-particle;

Figure 8 shows a process diagram representative of a process for manufacturing a steel in accordance with the present disclosure;

Figure 9 shows a flow chart describing additional steps that may be applied to the process of Figure 8;

Figure 10 shows a process diagram representative of a process for manufacturing a modulus-enhanced hot-rolled high carbon wire steel in accordance with the present disclosure; and

Figure 11 shows a flow chart describing additional steps that may be applied to the process of Figure 10.

Examples

The vanadium alloyed steels and production processes disclosed herein will now be explained 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 includes a balance amount of iron.

Table 1 (all wt%)

It has been determined that the ferritic compositions set out in Table 1 will have the characteristics and properties set out in Table 2.

Table 2

Example 2 A typical composition for an enhanced-modulus AHSS steel is provided in Table 3. Such a composition also includes a balance amount of iron. In this case, the balance amount of iron is 97.6955 wt%.

Table 3 (all wt%)

Example 3

A conventional steel of grade S690MC (in accordance with EN 10051) was compared to a nano-structured enhanced modulus steel as disclosed herein. The conventional S690MC steel is taken as having a nominal elastic modulus of 210 GPa (in accordance with European Standard EN 1993-1-1: Eurocode 3: Design of steel structures, and European Standard EN 1993-1-12: General - High strength steels). The modulus enhanced steel is taken to have an average elastic modulus of 230 GPa, in other words, an increase of 20 GPa.

Figure 6 illustrates three typical beam compositions (in accordance with Eurocode 3). The consequences with respect to reduced deflections with an enhanced elastic modulus are set out in Table 4. These results were obtained via a computer simulation.

Table 4

Example 4

With reference to Figure 8, an exemplary method for producing the enhanced modulus steel disclosed herein will now be described. A slab (or alternatively a billet) of cast steel is heated to a reheating temperature in the range of from about 1100 to about 1300 °C (Treh). This is indicated by reference numeral 102 on Figure 8. This temperature is maintained until the steel slab is fully heated through its thickness (i.e. at Teq as shown in Figure 8). By heating the steel slab in this way, the steel is heated to form austenite (y), having a grain microstructure as illustrated in the schematic indicated by reference numeral 104.

The heated steel is then hot rolled at temperatures in the range of from about 820°C to about 1200°C, firstly in a roughing mill as indicated by reference numeral 106, followed by hot rolling in a finishing mill indicated by reference numeral 108. During hot rolling, the temperature of the steel cools under standard air cooling, for example at about 0.7°C/s.

Hot rolling in the finishing mill 108 can be carried out using conventional recrystallisation controlled rolling. Alternatively, in particular where the steel has a composition including a relatively high amount of nitrogen, V(C,N) precipitation controlled rolling can be carried out. This is carried out at the V(C,N) precipitation time temperature nose 110, at which there is optimal free energy for VN to precipitate. The V(C,N) precipitation time temperature nose varies depending to the particular steel composition, however is typically about 850°C.

As the steel is hot rolled, in particular in the finishing mill, VN precipitates in austenite as intra-granular nucleants for acicular ferrite formation. At this stage, the steel has a microstructure as illustrated schematically at reference numeral 112.

Prior to 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 from about 2°C/s to about 80°C/s until a temperature of 650 ±100°C is reached. This may be by forced air cooling, lamellar cooling or any other suitable cooling means. Rapidly cooling the steel in this way prevents or impedes the formation of undesirable phases and/or promotes the formation of acicular ferrite, illustrated schematically at reference numeral 116.

When the temperature of 650 ±100°C is reached, the steel is formed into a coil. The temperature of the steel is then held 118 at 650 ± 100°C for a time in the range of from about 10 minutes to about 20 minutes, after which the steel is air cooled. As described above, isothermal holding in this manner promotes the generation of interphase precipitates. In this way, a modulus enhanced nanostructured steel is produced.

Optionally, the method may further include the steps set out in Figure 9 to increase the number density of interphase precipitates.

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

The steel may then be austenitised 122 a second time by heating to a temperature in the range of from about 920°C to about 1150°C, for example using a roller hearth furnace and/or inductive heating equipment. The reheated steel is then cooled 124 at a rate in the range of from about 2°C/s to about 50°C/s until a temperature of 650 ± 40°C is reached. Cooling can be carried out using forced air cooling, lamellar cooling, or any other suitable method.

A static magnetic field in the range of from about 0.2T to about 16T can then be applied 26 to the steel. Whilst the magnetic field is applied, the temperature is held at 650 ± 40°C for about 15 minutes or less. For example, this may be carried out using an insulated thermal box.

Alternatively, after a blank has been stamped or cut from the decoiled strip 120, a variable magnetic field is applied 128 to the steel. For example, this may be applied using an induction heating system. Whilst the variable magnetic field is applied, the temperature is held at 650 ±100°C for 15 minutes or less.

In alternative embodiments, a magnetic field is applied to the steel during the step of isothermal holding 118, either by using a static magnetic field or a variable magnetic field.

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

Alternatively, instead of decoiling and stamping a blank from the steel, a full coil can be reheated (i.e. austenitised) 122, cooled 124 then held at 650 ± 100°C for the required time. This can be carried out as a batch annealing process using an insulated thermal box, e.g. a bell furnace suitable for applying heat treatment to full coils. This method is applicable to both ferritic and pearlitic steels.

Example 5

Figure 10 shows a method of making a modulus-enhanced hot-rolled high carbon wire steel. A billet of cast steel is heated 202 to a reheating temperature, as described in relation to example 4.

The heated steel is then hot rolled 206 at temperatures in the range of from about 840°C to about 1200°C using a roughing mill, intermediate mill, finishing mill and/or notwist V-block mill. Hot rolling is carried out above the recrystallization stop temperature, such that grain deformation is kept to a minimum.

After hot rolling has been completed, the steel is rapidly cooled 214 at a cooling rate of in the range of from about 2°C/s to about 80°C/s until a temperature of 650 ±100°C is reached.

The temperature is then held 218 at 650 ±100°C, for example, using an in-line induction heating coil. With reference to Figure 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 austenised, fan cooled to 600±50°C, then the temperature is held at 650±200°C for less than 5 minutes. These steps are applied in a continuous, slow moving fashion. Optionally, a static magnetic field may also be applied at this stage in order to refine the interlamellar spacing of the pearlite, which can further contribute to lattice strain and modulus enhancement.

Unless otherwise stated, each of the integers described herein may be used in combination with any other integer as would be understood by the person skilled in the art. Further, although all aspects of the invention preferably "comprise" the features described in relation to that aspect, it is specifically envisaged that they may "consist" or "consist essentially" of those features outlined in the claims. In addition, all terms, unless specifically defined herein, are intended to be given their commonly understood meaning in the art.

Further, in the discussion of the invention, unless stated to the contrary, the disclosure of alternative values for the upper or lower limit of the permitted range of a parameter, is to be construed as an implied statement that each intermediate value of said parameter, lying between the smaller and greater of the alternatives, is itself also disclosed as a possible value for the parameter.

In addition, unless otherwise stated, all numerical values appearing in this application are to be understood as being modified by the term "about".