Patent Application
20240191319
The invention relates to a steel strip consisting of a high-strength multiphase steel which has a tensile strength of at least 780 MPa in the longitudinal direction.
The invention also relates to a method for producing a steel strip consisting of a high-strength multiphase steel which has a tensile strength of at least 780 MPa.
Multiphase steels have, by reason of their multiphase microstructure, an excellent combination of strength, deformability and ductility. In particular, phase proportions over 30 vol. % martensite and/or bainite are an essential microstructure constituent for achieving high tensile strengths (e.g. >600 MPa). With an increasing proportion e.g. of 50 or 70 vol. %, strength classes up to above 980 MPa are also possible depending upon the chemical composition. In particular, in the case of annealing treatments with slow cooling rates and/or a coarse microstructure, resulting from an annealing treatment at high annealing temperatures, the proportion of hard phase constituents (martensite or bainite, optionally also tempered) must be higher in order to achieve a higher level of strength. In order to ensure that a low-alloyed steel can form sufficient phase proportions of bainite and martensite, an annealing treatment significantly above the conversion temperature A1, which is characteristic of steel, with subsequent sufficiently high cooling rates is necessary. For large-scale production of low-alloyed multiphase steels, continuous annealing installations, such as so-called “continuous annealing” or hot-dip galvanizing lines are required, in which the cooling rates are considerably above 1 K/s, the steel strip is thus cooled within seconds or minutes from temperatures above the A1 temperature to room temperature.
Steel strip is understood hereinafter to be a hot-rolled or cold-rolled and annealed steel strip. Typical thicknesses of hot-rolled steel strip, also referred to as hot strip, are between 1.8 mm and 18 mm. Cold-rolled, annealed steel strips are referred to as cold strip or as fine sheet and typically have thicknesses in the range of 0.5 mm to 2.5 mm, wherein the strip thickness can also be adjusted in a targeted manner with varying flexibility by means of targeted processing even within a cold strip or fine sheet.
In addition to the large-scale heat treatment by continuous annealing in a continuous annealing furnace, strip sheets are heat-treated “as a whole” on an industrial scale also as coiled strip in furnaces, such as e.g. so-called batch-type annealing installations. Batch-type annealing treatments of low-alloyed strip sheets are effected either as recovery annealing or recrystallisation/soft annealing. In the case of recovery annealing, a generally cold-formed strip sheet is annealed at temperatures below 700° C. in order to achieve a high tensile strength with a simultaneously high yield strength and low ductility in the steel strip produced by the annealing. Typically, a recovery-annealed steel has a pronounced yield strength, moderate ductility and a high yield strength/tensile strength ratio >0.8, which can be critical for the further processing of the steel strip. The materials-science mechanism of recovery, which is causal for the technological characteristic values after the batch-type annealing treatment, is very greatly dependent upon the annealing temperature, the annealing duration and the previous cold forming (e.g. degree of cold forming) of the strip sheet in the annealing treatment. In the case of soft/recrystallisation annealing, the strip is annealed at temperatures around the A conversion temperature for several hours to days. The tensile strength after the previously described annealing treatment is at strengths under 600 MPa and is significantly lower compared to the strength prior to the annealing treatment. However, the ductility increases considerably on account of recrystallisation annealing in comparison with the non-annealed and cold-rolled material.
The fiercely competitive car market means that producers are constantly forced to find solutions for reducing fleet fuel consumption and CO2 exhaust emissions whilst maintaining the highest possible level of comfort and passenger protection. On the one hand, the weight saving of all of the vehicle components plays a decisive role as does, on the other hand, the most favourable possible behaviour of the individual components in the event of high static and dynamic loading during operation and also in the event of a crash. The steel producers contribute to the solution of this problem through the provision of high-strength steels. Furthermore, through the provision of high-strength steels having a smaller sheet thickness, the weight of the vehicle components can be reduced whilst the behaviour of components remains the same or is possibly even improved.
These newly developed multiphase steels must satisfy not only the required weight reduction but also the high material requirements in relation to elasticity limit, tensile strength and elongation at fracture.
Multiphase steels are known e.g. from laid-open documents DE 10 2017 131 247 A1, DE 2017 130 237 A1 and DE 10 2015 111 177 A1. The material properties disclosed therein result from a high phase proportion of bainite and/or martensite which require sufficiently rapid cooling conditions. Large-scale processing of such multiphase steels is effected with continuous annealing installations.
The present invention provided alternative measures for providing steel strips consisting of high-strength multiphase steel.
In accordance with an embodiment of the invention, a steel strip consisting of a high-strength multiphase steel which has a tensile strength of at least 780 MPa in the longitudinal direction, wherein the multiphase steel consists of the following elements in wt. %:
and optionally one or more of the following elements in wt. %:
with the remainder being iron, including typical steel-associated, melting-induced impurities, and has a carbon equivalent CEV which is greater than 0.49 and less than 0.9, wherein the carbon equivalent CEV is determined according to the formula
CEV=C+Mn/6+(Cu+Ni)/15+(Cr+Mo+V)/5
from the contents of the corresponding elements in wt. % (i.e. the mass proportions of these elements) and wherein the ratio of the carbon equivalent CEV and the sum of the contents of Si and Al in wt. % is less than 2.3 (CEV/(Si+Al)<2.3), wherein the multiphase steel has a microstructure, in which the sum of the volume proportions of the microstructure constituents of martensite, tempered martensite, residual austenite, upper bainite and/or lower bainite is at least 30 vol. % and the remaining microstructure consists of ferrite and pearlite. Such a steel strip can be produced from a rolled strip sheet consisting of steel of a corresponding composition by heat-treating this strip sheet—which is rolled up into a coil—“as a whole” by means of the method for producing a steel strip described hereinafter. Said heat treatment is also referred to as “annealing” and can be effected e.g. by the batch-type annealing installation mentioned in the introduction.
The invention thus makes it possible to provide a steel strip which has a high tensile strength >780 MPa, in particular with good ductility A80>5% and a low elasticity limit ratio Rp0.2/Rm<0.8, and in which these technological characteristic values are not influenced substantially by the microstructure after heat treatment or by cold deformation before heat treatment. In other words, provision is made in particular that the ratio of elasticity limit to tensile strength Rp0.2/Rm is less than 0.8 and the elongation at fracture A80 is >5%.
According to another embodiment of the invention, provision is made that the content in wt. % of the element C is between 0.09 and 0.2 and/or the content in wt. % of the element Mo is less than 0.4.
According to a further embodiment of the invention, provision is made that the content in wt. % of the element Mn is between 1.8 and 2.5 and/or that the content in wt. % of the sum of the elements Si+Al is between 0.25 and 1.
According to one particular embodiment of the invention, provision is made that the carbon equivalent CEV is less than 0.7. Such a steel can be processed in a particularly effective manner.
According to yet a further embodiment of the invention, provision is made that the volume proportion of the common microstructure constituents of martensite, tempered martensite, residual austenite, upper bainite and/or lower bainite in the microstructure of the multiphase steel is at least 50 vol. %, particularly preferably at least 70 vol. %, and the remaining microstructure consists of ferrite and pearlite.
The steel strip has in particular a constant thickness, wherein the term “constant thickness” is to be understood in terms of the conventional standard tolerance (e.g. corresponding to EN 10051). Alternatively, provision is made in an advantageous manner that the steel strip has a thickness which specifically varies in the longitudinal extension. The ratio between maximum thickness and minimum thickness is, in particular, between 1.16 and 3. For nominal strip thicknesses of 2.0 mm or more, this ratio is outside of the conventional standard tolerance. In particular, such a steel strip having a thickness which specifically varies in the longitudinal extension is a flexibly rolled steel strip for so-called “Tailor Rolled Blanks”. The flexibly rolled steel strip is flexibly rolled as a strip sheet prior to the heat treatment, wherein the rollers produce different sheet thicknesses by means of up and down movement. The homogeneous transition between different thicknesses is advantageous.
According to yet a further embodiment of the invention, provision is made that the steel strip has, in particular, a thickness between 4 mm and 18 mm, which is not readily possible in the case of industrial production in continuous annealing furnaces.
In another embodiment of the invention, a method for producing a steel strip consisting of a high-strength multiphase steel, in particular a steel strip stated above, which has a tensile strength of at least 780 MPa in the longitudinal direction, wherein a rolled strip sheet of steel consisting of the following elements in wt. %
optionally one or more of the following elements:
with the remainder being iron, including typical steel-associated, melting-induced impurities, and having a carbon equivalent CEV which is greater than 0.49 and less than 0.9, preferably greater than 0.49 and less than 0.75, wherein the carbon equivalent CEV is determined according to the formula
CEV=C+Mn/6+(Cu+Ni)/15+(Cr+Mo+V)/5
from the contents of the corresponding elements in wt. % and wherein the ratio of the carbon equivalent CEV and the sum of the contents of Si and Al in wt. % is less than 2.3, is heat-treated as a whole—in particular rolled up into a coil—such that it assumes a temperature above 750° C. and after this heat treatment is cooled to a temperature below 200° C., wherein the cooling between 750° C. and 200° C. is effected at an average cooling rate greater than 1 K/h and less than 300 K/h.
The method disclosed herein for producing a steel strip does not require a continuous annealing procedure for large-scale production and the final product, steel strip, still has a high tensile strength >780 MPa, with good ductility A80>5% and a low elasticity limit ratio Rp0.2/Rm<0.8, and the technological characteristic values after the heat treatment are not influenced substantially by the microstructure or the cold deformation before heat treatment.
The method may be carried out on a large scale e.g. using a batch-type annealing installation. In the case of low-alloy steels, this is only possible in terms of materials science by means of the production method in accordance with the invention because the method includes targeted phase conversions during the heat treatment. For this purpose, the annealing is performed within a temperature range above the A1 temperature, similar to a heat treatment in the case of continuous annealing, although the temperature does not necessarily have to be above the A3 temperature. The annealing temperatures can vary depending upon the chemical composition of the steel strip.
The corresponding product is a hot-rolled and/or cold-rolled steel strip having a multiphase microstructure and the associated characteristic technological properties of multiphase steels stated above.
The phase constituents of bainite and/or martensite characteristic of multiphase steels are formed from austenitic phase proportions during cooling of a steel from temperatures above the A1 temperature. In order to ensure that the austenite does not convert into the phases of ferrite and/or pearlite, it is necessary to have sufficient thorough hardenability of the material corresponding to the technically possible cooling rate. The thorough hardenability of a steel is dependent upon the chemical composition and can be described approximately by the following carbon equivalent CEV:
CEV=C+Mn/6+(Cu+Ni)/15+(Cr+Mo+V)/5
However, an excessively high proportion of alloy elements, such as manganese (Mn), chromium (Cr), carbon (C), vanadium (V), molybdenum (Mo), copper (Cu) and/or nickel (Ni) cannot be tolerated for the previous process steps, such as continuous casting, hot rolling or cold rolling and the subsequent joining operations, such as welding. In the case of steels having lower proportions of alloy elements, it is optional to select a low annealing temperature in order to locally enrich the alloy elements in the austenite and thus to achieve better thorough hardenability locally in the austenite. For the reasons stated above, the CEV is restricted to 0.49 to a maximum of 0.9, optionally to 0.75, and optionally to a maximum of 0.7. For cost reasons, the proportion of alloy elements and thus also the CEV should be kept low. For the reasons stated above, it is recommended for the average cooling rate in the critical temperature range of 750° C. to 200° C. to be between 1 K/h and 300 K/h.
The strength-building microstructure constituents of the multiphase steels, such as bainite and/or martensite, are formed from the austenitic phase proportions during cooling at temperatures below 570° C. At temperatures above room temperature, in particular above 200° C., the local high strengths of the phases of martensite and bainite are reduced by so-called tempering or self-tempering. In the case of this materials-science mechanism, the precipitation of forcibly dissolved carbon to carbides and the reduction of conversion-induced stresses result in particular in a reduction in strength of the hard phases of bainite and martensite and thus also to a decrease in strength of the annealed steel strip. This tempering mechanism is thermally activated. The decrease in strength as a result of the tempering/self-tempering increases accordingly during longer dwell times at higher temperatures, in particular at temperatures above 200° C. In order to keep the decrease in strength to a minimum, it is absolutely necessary, after formation of the hard phases, to cool to temperatures below 200° C. in an accelerated manner and to counteract the tempering mechanisms with alloying concepts.
In particular, the elements Si and Al are useful for delaying the kinetics of the carbide formation and thus for stabilising the hard phases. Lowering the conversion temperatures of bainite and martensite also results in fewer tempering effects with the same process control. The elements described in the CEV likewise serve to lower the conversion temperatures of bainite and, in particular, martensite, which is positive for the invention.
However, high proportions of Cr or Mn can result, even during the annealing treatment, in the formation of additional carbides which likewise can result in a lower maximum strength. For the reasons stated above, it is necessary to restrict the sum content of the alloy elements Si and Al to greater than 0.25 wt. % and likewise to restrict the ratio of the CEV to the aforementioned sum content of Si and Al to a maximum value of 2.3.
To ensure that sufficient proportions of austenite are formed during the annealing treatment, temperatures above 750° C., preferably above 780° C., even more preferably above 790° C., are to be maintained. However, excessively high annealing temperatures result in undesired grain growth such that the maximum annealing temperature is preferably not above 70° C.+Ar3 temperature. The Ar3 temperature is dependent upon the chemical composition and can be estimated using the following formula.
Ar3=910−203√{square root over (C)}−30Mn+44.7Si−11Cr+31.5Mo−15.2Ni
Correspondingly, according to one embodiment of the method in accordance with the invention, provision is made that the strip sheet is heated from 100° ° C. to a temperature of 750° C. during the heat treatment at an average heating rate between 1 K/h and 300 K/h, and in which the strip sheet remains in the temperature range from 750° C. to Ar3+70° C. for at least 1 h, wherein the numerical value of the temperature Ar3 is calculated by means of the above formula from the contents of the corresponding elements in wt. %.
In order to save energy and with regard to the temperature resistance of furnace linings, it is recommended to restrict the temperatures to a maximum of 900° C. or even more so to 850° C. Accordingly, the strip sheet consisting of steel optionally reaches a maximum temperature of at least 780° C. and at most 900° C., preferably of at least 790° C. and at most 850° C. during the heat treatment.
In order to ensure homogeneous thorough heating of the coiled steel strip, provision is optionally made for thorough heating of at least 1 hour. Longer retention times are conducive to more homogeneous thorough heating, but are not recommended by reason of grain growth associated therewith, which in turn causes a decrease in strength.
By processing a coiled steel strip e.g. in a batch-type annealing furnace, it is technically possible to achieve very low heating rates than in a continuous annealing procedure. However, it is still recommended to heat the entire annealing cycle at the highest possible heating rates in order to avoid undesired thermodynamically stable precipitations as well as undesired grain growth during heating. Again, excessively rapid heating may be detrimental to uniform thorough heating of the steel strip, and so it is recommended for the average heating rates in the critical temperature range from 100° C. to the minimum annealing temperature of 750° ° C. to be between 1 K/h and 300 K/h.
Furthermore, provision is preferably made that the steel strip is provided with a surface coating in the form of a metallic coating, organic coating or lacquer after cooling.
The steel strip which is provided for the heat treatment has, in particular, a constant thickness, wherein the term “constant thickness” is to be understood in terms of the conventional standard tolerance (e.g. corresponding to EN 10051). Alternatively, the strip sheet provided for the heat treatment advantageously has thicknesses which specifically vary in the longitudinal extension, wherein the ratio between maximum thickness and minimum thickness is, in particular, between 1.16 and 3. For nominal strip thicknesses of 2.0 mm or more this ratio is outside of the conventional standard tolerance. In particular, such a strip sheet having a thickness which specifically varies in the longitudinal extension is a flexibly rolled steel strip for so-called “Tailor Rolled Blanks”. The flexibly rolled steel strip is rolled once again prior to the heat treatment, wherein the rollers produce different sheet thicknesses by means of up and down movement. The homogeneous transition between different thicknesses is advantageous. The resulting steel strip is then a flexibly rolled steel strip consisting of a high-strength multiphase steel.
The effect of the elements in the steel strip in accordance with the invention having a multiphase microstructure will be described in greater detail hereinafter. Typically, the multiphase steels are chemically structured in such a way that alloy elements with and also without microalloy elements are combined. Associated elements are unavoidable and, if necessary, are taken into consideration in the analysis concept in terms of their effect.
Associated elements are elements which are already present in the iron ore or get into the steel as a result of the production process. They are generally undesired by reason of their predominantly negative influences. The attempt is made to remove them to a tolerable content level or to convert them into less damaging forms.
Hydrogen (H) can diffuse as a single element through the iron lattice, without producing lattice tensions. As a result, the hydrogen in the iron lattice is relatively mobile and can be relatively easily absorbed during the manufacture of the steel. Hydrogen can be absorbed into the iron lattice only in atomic (ionic) form. Hydrogen exerts a significant embrittling effect and diffuses preferably to locations which are favourable in terms of energy (flaws, grain boundaries etc.). Flaws thus function as hydrogen traps and can considerably increase the dwell time of the hydrogen in the material. Cold cracks can be produced by means of a recombination to molecular hydrogen. This behaviour occurs in the event of hydrogen embrittlement or in the event of hydrogen-induced stress crack corrosion. Even in the case of delayed cracking, so-called delayed fracture, which occurs without external stresses, hydrogen is often stated to be the reason. Therefore, the hydrogen content in the steel should be kept as small as possible.
Oxygen (O): In the molten state, the steel has a relatively large absorbency for gases, however at room temperature oxygen is soluble only in very small quantities. In a similar manner to hydrogen, oxygen can diffuse only in atomic form into the material. Owing to the highly embrittling effect and the negative effects upon the ageing resistance, every attempt is made during production to reduce the oxygen content. On the one hand, procedural approaches such as vacuum treatment and, on the other hand, analytical approaches are provided in order to reduce the oxygen. By adding specific alloy elements, the oxygen can be converted into less dangerous states. For instance, it is generally conventional to bind the oxygen via manganese, silicon and/or aluminium. However, the resulting oxides can produce negative properties as flaws in the material. In contrast, in the case of fine precipitation, specifically of aluminium oxides, grain refinement can also take place. Therefore, for the reasons stated above the oxygen content in the steel should be kept as small as possible.
Nitrogen (N) is likewise an associated element from the production of steel. Steels with free nitrogen tend to have a strong ageing effect. The nitrogen diffuses even at low temperatures to dislocations and blocks same. It thus produces an increase in strength associated with a rapid loss of toughness. Binding of the nitrogen in the form of nitrides is possible by addition by alloying of e.g. aluminium or titanium. For the reasons stated above, the sulphur content is limited to ≤ 0.016 wt. % or to quantities which are unavoidable in the production of steel.
Sulphur (S), like phosphorous, is bound as a trace element in the iron ore. It is not desirable in steel (the exception being machining steels) because it exhibits a strong tendency towards segregation and has a greatly embrittling effect. An attempt is therefore made to achieve amounts of sulphur in the melt which are as low as possible (e.g. by deep vacuum treatment). Furthermore, the sulphur present is converted by the addition of manganese into the relatively innocuous compound manganese sulphide (MnS). The manganese sulphides are often rolled out in lines during the rolling process and function as nucleation sites for the conversion. Primarily in the case of a diffusion-controlled conversion this produces a microstructure of pronounced lines and, in the case of a highly pronounced line formation, can result in impaired mechanical properties (e.g. pronounced martensite lines instead of distributed martensite islands, anisotropic material behaviour, reduced elongation at fracture). For the reasons stated above, the sulphur content is limited to ≤ 0.01 wt. % or to quantities which are unavoidable in the production of steel.
Phosphorous (P) is a trace element from the iron ore and is dissolved in the iron lattice as a substitution atom. Phosphorus increases hardness by means of mixed crystal hardening and improves hardenability. However, attempts are generally made to lower the phosphorus content as much as possible because inter alia it exhibits a strong tendency towards segregation owing to its low diffusion rate and greatly reduces the level of toughness. The attachment of phosphorous to the grain boundaries causes grain boundary fractures. Moreover, phosphorous increases the transition temperature from tough to brittle behaviour up to 300° C. During hot rolling, near-surface phosphorous oxides at the grain boundaries can result in the formation of fractures. The addition by alloying of small quantities of boron can partially compensate for the negative effects of phosphorus. It is believed that boron increases grain boundary cohesion and reduces phosphorus segregation at grain boundaries. However, in some steels owing to the low costs and high increase in strength, P is used in small quantities (<0.1%) as a microalloy element, e.g. in higher-strength IF steels (interstitial free). For the reasons stated above, the phosphorous content is limited to ≤ 0.050% or to quantities which are unavoidable in the production of steel.
Alloy elements are generally added to the steel in order to influence specific properties in a targeted manner. An alloy element can thereby influence different properties in different steels. The correlations are varied and complex. The effect of the alloy elements will be discussed in greater detail hereinafter.
Carbon (C) is considered to be the most important alloy element in steel. Its targeted introduction at an amount up to 2.06% turns iron first into steel. The carbon proportion is often drastically reduced during the production of steel. In the case of the multiphase steel in accordance with the invention, its proportion is 0.08 wt. % to 0.23 wt. %. Carbon is interstitially dissolved in the iron lattice owing to its comparatively small atomic radius. The solubility is at most 0.02% in the α-iron and is at most 2.06% in the γ-iron. In dissolved form, carbon considerably increases the hardenability of steel. The different solubility makes pronounced diffusion procedures necessary during the phase conversion, which procedures can result in very different kinetic conditions. Moreover, carbon increases the thermodynamic stability of the austenite, which is demonstrated in the phase diagram in an extension of the austenite region at lower temperatures. As the forcibly dissolved carbon content in the martensite increases, the lattice distortions and, associated therewith, the strength of the phase produced without diffusion increase. Moreover, carbon is required for carbide formation. A representative which occurs almost in every steel is cementite (Fe3C). However, substantially harder special carbides can be formed with other metals, such as e.g. chromium, titanium, niobium and vanadium. Therefore, it is not only the type but also the distribution and extent of the precipitations which is of crucial significance for the resulting increase in strength. Therefore, in order to ensure, on the one hand, sufficient strength and, on the other hand, good weldability, the minimum C content is fixed to 0.08 wt. % and the maximum C content is fixed to 0.23 wt. %, preferably between 0.09 and 0.2 wt. %.
Aluminium (Al) is generally added to the steel by alloying in order to bind the oxygen and nitrogen dissolved in the iron. The oxygen and nitrogen are thus converted into aluminium oxides and aluminium nitrides. These precipitations can effect grain refinement by increasing the nucleation sites and can thus increase the toughness properties and strength values. Aluminium nitride is not precipitated if titanium is present in a sufficient quantity. Titanium nitrides have a lower enthalpy of formation and are formed at higher temperatures. In the dissolved state, aluminium, like silicon, shifts the formation of ferrite towards shorter times and thus permits the formation of sufficient ferrite. It also suppresses the formation of carbide and thus results in a delayed conversion of the austenite. For this reason, Al is also used as an alloy element in residual austenite steels in order to substitute a part of the silicon with aluminium. The reason for this approach resides in Al being slightly less critical for the galvanization reaction than Si.
During casting, silicon (Si) binds oxygen and therefore reduces segregations and impurities in the steel. Moreover, by means of mixed crystal hardening silicon increases the strength of the ferrite with the elongation at fracture only decreasing slightly. A further important effect is that silicon shifts the formation of ferrite towards shorter times and therefore permits the production of sufficient ferrite prior to quench hardening with a continuously annealed material. An effect which is particularly advantageous when using low-alloyed steels in the inventive batch-type annealing treatment of multiphase steels. The formation of ferrite causes the austenite to be enriched with carbon and stabilised. In the case of higher contents, silicon markedly stabilises the austenite in the low temperature range specifically in the region of bainite formation by preventing the formation of carbide. During hot rolling, highly adhesive scales which can impair further processing can form at high silicon contents.
By suppressing carbides (in particular M3C-carbides, where M is a metallic alloying element) in bainitic microstructure constituents, or tempered martensite, additions of both Si and Al prevent a reduction in strength from the aforementioned hard phases of martensite and/or bainite and cause the strength to decrease less significantly after an annealing treatment. For the reasons stated above, a sum content of Al and Si is fixed to 0.25 to 2 wt. %, preferably to a maximum of 1 wt. %.
Manganese (Mn) is added to almost all steels for the purpose of desulphurisation in order to convert the noxious sulphur into manganese sulphides. Moreover, by means of mixed crystal hardening manganese increases the strength of the ferrite and shifts the conversion towards lower temperatures. A main reason for adding manganese by alloying is the considerable improvement in the potential hardness increase. By reason of the inhibition of diffusion, the pearlite and bainite conversion is shifted towards longer times and the martensite starting temperature is decreased. Manganese, like silicon, tends to form oxides on the steel surface during the annealing treatment. In dependence upon the annealing parameters and the contents of other alloy elements (in particular Si and Al) manganese oxides (e.g. MnO) and/or Mn mixed oxides (e.g. Mn2SiO4) can occur.
However, manganese is to be considered to be less impactful in a small Si/Mn or Al/Mn ratio because globular oxides are more likely to form instead of oxide films. Therefore, the Mn content is selected to be in a range from 1.5 wt. % to 3.5 wt. %, optionally 1.8 to 2.5 wt. %.
Chromium (Cr): the addition of chromium mainly improves the potential hardness increase. Chromium in the dissolved state shifts the pearlite and bainite conversion towards longer times and at the same time lowers the martensite starting temperature. A further important effect is that chromium increases the tempering resistance considerably. Moreover, chromium is a carbide forming agent. Should chromium be present in carbide form, the austenitizing temperature must be selected, prior to hardening, to be high enough to dissolve chromium carbides. Otherwise, the increased number of nuclei can cause a deterioration in the potential hardness increase. Chromium likewise tends to form oxides on the steel surface during the annealing treatment, as a result of which the galvanizing quality can be impaired. Therefore, the optional Cr content is selected to be a range of values of 0.05 to 1.0 wt. %.
Molybdenum (Mo): The addition of molybdenum is effected, in a similar manner to the addition of chromium, to improve hardenability. The pearlite and bainite conversion is shifted towards longer times and the martensite starting temperature is decreased. Moreover, molybdenum considerably increases the tempering resistance and effects an increase in strength of the ferrite owing to mixed crystal hardening. The Mo content is added in dependence upon the dimension, the system configuration and the microstructure setting. For cost reasons, the optional Mo content is selected to be in a range from 0.05 to 1.0 wt. %, optionally to a maximum of 0.4 wt. %.
Copper (Cu): the addition of copper can increase the tensile strength and the potential hardness increase. In conjunction with nickel, chromium and phosphorous, copper can form a protective oxide layer on the surface which can considerably reduce the corrosion rate. In conjunction with oxygen, copper can form, at the grain boundaries, noxious oxides which can produce negative effects particularly for hot-deformation processes. Therefore, the maximum content of copper is limited to 0.2 wt. %.
Calcium (Ca): calcium is used in the production of high-strength steels for deoxidation, desulphurisation and to control the size and shape of oxides and sulphides. This produces improved ductility and toughness particularly in high-strength steels. Furthermore, steels with additions of calcium tend to a lesser extent to incur hot cracks, e.g. during hot rolling. For the reasons stated above and owing to the very low solubility of calcium in steel, the content of calcium—where required accordingly—is thus limited to 0.0005 to 0.0060 wt. %.
Nickel (Ni): In conjunction with oxygen, nickel can form, at the grain boundaries, noxious oxides which can produce negative effects particularly for hot-deformation processes. However, nickel likewise increases the hardenability and lowers the conversion temperature Ac3. For the reasons stated above and for cost reasons, the optional content of nickel is thus limited to 0.05 to 0.50 wt. %.
Microalloy elements are generally added only in very small amounts (<0.1%). Typical microalloy elements are aluminium, vanadium, titanium, niobium and boron. In contrast to the alloy elements, they mainly act by precipitate formation but can also influence the properties in the dissolved state. Despite the small amounts added, microalloy elements greatly influence the production conditions and the processing properties and final properties. In general, carbide and nitride forming agents which are soluble in the iron lattice are used as microalloy elements. Formation of carbonitrides is likewise possible by reason of the complete solubility of nitrides and carbides in one another. The tendency to form oxides and sulphides is generally most pronounced with the microalloy elements, but generally is specifically prevented by reason of other alloy elements. This property can be used positively by binding the generally harmful elements sulphur and oxygen. However, the binding can also have negative effects if, as a result, there are no longer sufficient microalloy elements available for the formation of carbides.
Titanium (Ti) forms very stable nitrides (TiN) and sulphides (TiS2) even at high temperatures. They only partly dissolve in the melt in dependence upon the nitrogen content. If the thus produced precipitations are not removed with the slag, they form coarse particles in the material owing to the high formation temperature and are generally not conducive to the mechanical properties. A positive effect on the toughness is produced by binding of the free nitrogen and oxygen. Therefore, titanium protects other dissolved microalloy elements, such as niobium, against being bound by nitrogen. These can then optimally demonstrate their effect. Nitrides which are produced only at lower temperatures by lowering the oxygen and nitrogen content can additionally ensure effective hindrance of the austenite grain growth. Non-bound titanium forms, at temperatures from 1150° C., titanium carbides and can thus effect grain refinement (inhibition of the austenite grain growth, grain refinement by delayed recrystallisation and/or increase in the number of nuclei in α/γ conversion) and precipitation hardening. The optional Ti content thus has values of 0.005 to 0.150 wt. %.
Niobium (Nb) effects considerable grain refinement because it effects a delay in the recrystallisation most effectively among all micro-alloy elements and additionally impedes the austenite grain growth. However, the strength-increasing effect is to be estimated to be qualitatively higher than that of titanium, as can be seen by the increased grain refinement effect and the larger number of strength-increasing particles (binding of the titanium to coarse TiN at high temperatures). Niobium carbides form at temperatures below 1200° C. In the case of nitrogen binding with titanium, niobium can increase its strength-increasing effect by forming small carbides which are effective in terms of their effect in the lower temperature range (smaller carbide sizes). A further effect of the niobium is the delay of the α/γ conversion and the reduction of the martensite starting temperature in the dissolved state. On the one hand, this occurs by the solute-drag effect and on the other hand by the grain refinement. This effects an increase in strength of the microstructure and thus also a higher resistance to the increase in volume upon martensite formation. In principle, the addition of niobium by alloying is limited until its solubility limit is reached. Although this limits the amount of precipitations, it primarily effects an early formation of precipitation with quite coarse particles when said limit is exceeded. The precipitation hardening can thus become effective in real terms primarily in steels with a low C content (higher supersaturation possible) and in hot deformation processes (deformation-induced precipitation). Therefore, Nb content is selected to be in a range from 0.005 to 0.150 wt. %.
Vanadium (V): The carbide and nitride formation by vanadium first begins at temperatures from about 1000ºC or even after the α/γ conversion, i.e. substantially later than for titanium and niobium. Vanadium thus barely has a grain-refining effect owing to the low number of precipitations provided in the austenite. The austenite grain growth is also not hindered by the late precipitation of the vanadium carbides. Therefore, the strength-increasing effect is based virtually exclusively on the precipitation hardening. One advantage of the vanadium is the high solubility in the austenite and the high volume proportion of fine precipitations caused by the low precipitation temperature. Therefore, the optional V content is selected to be in a range of 0.001 to 0.300 wt. %.
Boron (B) forms nitrides and carbides with nitrogen and with carbon respectively; however, this is generally not desired. On the one hand, only a low amount of precipitations are formed owing to the low solubility and on the other hand these are mostly precipitated at the grain boundaries. An increase in hardness at the surface is not achieved (the exception being boronising with formation of FeB and Fe2B in the edge zone of a workpiece). To prevent nitride formation, an attempt is generally made to bind the nitrogen by means of more affine elements. In particular, titanium can ensure the binding of all of the nitrogen. In the dissolved state, in very small amounts, boron results in a considerable improvement in the potential hardness increase. The mechanism of action of boron can be described in such a way that boron atoms accumulate at the grain boundaries under suitable temperature control and at that location, by lowering the grain boundary energy, significantly hamper the formation of ferrite nuclei capable of growth. When controlling the temperature, care must be taken to ensure that boron is predominantly distributed atomically in the grain boundary and is not present in the form of precipitations by reason of excessively high temperatures. The efficacy of boron is decreased as the grain size increases and the carbon content increases (>0.8%). An amount over 60 ppm additionally causes decreasing hardenability because boron carbides act as nuclei on the grain boundaries. Boron diffuses extraordinarily well by reason of the small atomic diameter and has an extremely high affinity to oxygen which can lead to a reduction in the boron content in regions near to the surface (up to 0.5 mm). In this connection, annealing at over 1000° C. is discouraged. This is also to be recommended because boron can result in an excessive coarse grain formation at annealing temperatures above 1000° C. For the reasons stated above, the B content is selected to be in a range of 0.0005 to 0.0050 wt. %.
Finally, the invention also relates to a use of an aforementioned steel strip for producing a motor vehicle component.
Embodiments of the invention will be explained hereinafter with reference to examples and by means of figures and tables.
Basically, the annealing treatments in accordance with the invention can be multi-stage or additional annealing treatments can also be provided in relation to the entire process. An exemplary time-temperature cycle which represents the characteristic temperature ranges for retention times, cooling rates and heating rates is specified in
For this purpose, a rolled strip sheet consisting of steel of a corresponding composition is put into a compact form, in particular rolled into a coil, which makes it possible to transport the strip sheet as a whole to an apparatus for heat treatment. At this location, in a first step S1 the sheet strip is heated to a temperature T≥ 750° C. within about 3 h. Subsequently, in a second step S2 the strip sheet is held at a temperature above 750° C. for about 8 h by means of the apparatus. The following applies to the maximum temperature reached in the second step S2: Tmax<Ar3+70 K. The strip sheet is then cooled. During this cooling, the temperature passes through the temperature range of 750° C. to 200° C. in a time period of about 14 h. This gives rise to a third step S3 of cooling from 750° C. to 200° C. at an average cooling rate of about 40 K/h. When the strip sheet consisting of steel of a corresponding steel concept, i.e. a suitable composition, is cooled, the desired microstructure is produced and the steel strip consisting of high-strength multiphase steel is produced. The cooling is effected to a specific temperature preferably in the apparatus for heat treatment. This is e.g. a batch-type annealing installation. The example shown, at about 40 K/h, is in a preferred cooling range of 20 K/h to 80 K/h.
Material concepts, more specifically steel concepts, and their chemical composition in wt. % are listed by way of example in the following table 1. Steel concepts in accordance with the invention are characterised accordingly. In addition to the steel concepts in accordance with the invention which, in the form of a hot-rolled or cold-rolled strip sheet, are used as input material for production in accordance with the invention of a product in accordance with the invention, steel concepts are likewise indicated as a comparison but are not in accordance with the invention.
The parameters of a production method in accordance with the invention and the characteristic values of the inventive product of this production method, i.e. the steel strip consisting of high-strength multiphase steel, are listed in Table 2. The exemplary material concepts are explained hereinafter. They are designated as “Steel A”, “Steel B”, “Steel C” and “Steel D”.
Steel A is not in accordance with the invention because the sum of alloy elements which increase thorough hardenability, as described by the CEV, is below the required value of 0.49. After a heat treatment involving the process parameters in accordance with the invention, steel A has a microstructure consisting of ferrite and pearlite and no proportions of bainite and/or martensite are formed. The associated stress-strain curve can be seen in
Steel B from Table 2 is likewise not in accordance with the invention, although the steel B, with a CEV value of 2.34, has sufficient thorough hardenability. However, the ratio of CEV/(Si+Al) is >2.34 and thus the content of Si and Al in relation to the use of elements which increase thorough hardenability is not sufficient. This is also apparent from the achievable maximum tensile strengths of 762 MPa.
By way of example, steels C and D are material concepts which are suitable for an annealing treatment in accordance with the invention and for the production of steel strips produced in accordance with the invention. After a heat treatment involving the process parameters in accordance with the invention, the steels C and D have a martensite and/or bainite proportion of over 30%. By reason of the microstructure set in this way, the steels also have the material properties characteristic of multiphase steels, such as an elasticity limit-tensile strength ratio (Rp0.2/Rm) between 0.45 and 0.6, a high tensile strength Rm above 780 MPa and, at the same time, a high elongation at fracture A80>8%.
| Number | Date | Country | Kind |
|---|---|---|---|
| 10 2021 108 448.2 | Apr 2021 | DE | national |
The present application is a national stage application of PCT/EP 2022/058767 filed on Apr. 1, 2022, which claims the benefit of German Application 10 2021 108 448.2, filed on Apr. 1, 2021.
| Filing Document | Filing Date | Country | Kind |
|---|---|---|---|
| PCT/EP2022/058767 | 4/1/2022 | WO |
