Patent Application
20250197959
The present invention relates to a high strength steel sheet excellent in tensile strength, flatness in the width direction, and working embrittlement resistance, and to a method for manufacturing the same. The high strength steel sheet according to aspects of the present invention may be suitably used as structural members, such as automobile parts.
Steel sheets for automobiles are being increased in strength in order to reduce CO2 emissions by weight reduction of vehicles and to enhance crashworthiness by weight reduction of automobile bodies at the same time, with introduction of new laws and regulations one after another. To increase the strength of automobile bodies, high strength steel sheets having a tensile strength (TS) of 1180 MPa or higher grade are increasingly applied to principal structural parts of automobiles.
From the point of view of the performance of parts, high strength steel sheets used in automobiles require high working embrittlement resistance and excellent yield ratio. For example, high strength steel sheets applied to automobile frame parts, such as bumpers, are suitably those that excel in working embrittlement resistance and are not embrittled upon being press-formed, and have excellent collision impact absorption properties which are correlated with YR.
Furthermore, high strength steel sheets used in automobiles require high flatness. Patent Literature 1 describes that warpage of a steel sheet causes operational troubles in forming lines and adversely affects the dimensional accuracy of products. The present inventors carried out extensive studies and have found that the dimensional accuracy of products is affected not only by the warpage of steel sheets but also by the flatness in the width direction that is evaluated as steepness. For example, the steepness in the width direction is suitably 0.02 or less in order to achieve excellent dimensional accuracy.
To meet the above demands, for example, Patent Literature 2 provides a high strength steel sheet having a tensile strength of 1100 MPa or more and being excellent in YR, surface quality, and weldability, and a method for manufacturing the same. However, the technique described in Patent Literature 2 does not take into consideration flatness in the width direction and working embrittlement resistance.
Patent Literature 3 provides a hot-dip galvanized steel sheet with excellent press formability and low-temperature toughness that has a tensile strength of 980 MPa or more, and a method for manufacturing the same. While the steel sheet of Patent Literature 3 is improved in embrittlement at low temperatures, the technique does not take into consideration the working embrittlement of the steel sheet or the flatness in the width direction.
Aspects of the present invention have been developed in view of the circumstances discussed above. Objects of aspects of the present invention are therefore to provide a high strength steel sheet having 1180 MPa or higher TS and 85% or more YR and being excellent in flatness in the width direction and working embrittlement resistance; and to provide a method for manufacturing the same.
The present inventors carried out extensive studies directed to solving the problems described above and have consequently found the following facts.
Aspects of the present invention have been made based on the above findings. Specifically, a summary of aspects of the present invention is as follows.
[1] A high strength steel sheet having a chemical composition including, in mass %, C: 0.030% or more and 0.500% or less, Si: 0.01% or more and 2.50% or less, Mn: 0.10% or more and 5.00% or less, P: 0.100% or less, S: 0.0200% or less, Al: 1.000% or less, N: 0.0100% or less, and O: 0.0100% or less, a balance being Fe and incidental impurities, the high strength steel sheet being such that in a region at ¼ sheet thickness, an area fraction of tempered martensite is 90% or more, a volume fraction of retained austenite is less than 3%, an area fraction of the total of ferrite and bainitic ferrite is less than 10%, an average grain size of prior austenite is 20 μm or less, and an average of the proportions of packets having the largest area in prior austenite grains is 70% by area or less of the prior austenite grain.
[2] The high strength steel sheet according to [1], wherein the chemical composition further includes at least one element selected from, in mass %, Ti: 0.200% or less, Nb: 0.200% or less, V: 0.200% or less, Ta: 0.10% or less, W: 0.10% or less, B: 0.0100% or less, Cr: 1.00% or less, Mo: 1.00% or less, Co: 0.010% or less, Ni: 1.00% or less, Cu: 1.00% or less, Sn: 0.200% or less, Sb: 0.200% or less, Ca: 0.0100% or less, Mg: 0.0100% or less, REM: 0.0100% or less, Zr: 0.100% or less, Te: 0.100% or less, Hf: 0.10% or less, and Bi: 0.200% or less.
[3] The high strength steel sheet according to [1] or [2], which has a coated layer on a surface of the steel sheet.
[4] A method for manufacturing the high strength steel sheet described in [1] or [2], the method including providing a cold rolled steel sheet produced by subjecting a steel having the chemical composition described above to hot rolling, pickling, and cold rolling; heating the steel sheet at an annealing temperature T1 of 750° C. or above and 950° C. or below for a holding time t1 at the annealing temperature T1 of 10 seconds or more and 1000 seconds or less; cooling the steel sheet in such a manner that an average cooling rate from 750° C. to 600° C. is 20° C./s or more, an average cooling rate from (Ms+50° C.) to a quench start temperature T2 is 5° C./s or more and 30° C./s or less wherein the quench start temperature T2 is (Ms−80° C.) or above and is below Ms where Ms is martensite start temperature (° C.) defined by formula (1), and an average cooling rate from the quench start temperature T2 to 80° C. is 300° C./s or more; and heating the steel sheet at a tempering temperature T3 of 100° C. or above and 400° C. or below for a holding time t3 at the tempering temperature T3 of 10 seconds or more and 10000 seconds or less,
wherein [% C], [% Mn], [% Cr], [% Mo], and [% Ni] indicate the contents (mass %) of C, Mn, Cr, Mo, and Ni, respectively, and are zero when the element is absent.
[5] The method for manufacturing the high strength steel sheet according to [4], further including performing a coating treatment.
According to aspects of the present invention, a high strength steel sheet can be obtained that has 1180 MPa or higher TS and 85% or more YR and excels in flatness in the width direction and working embrittlement resistance. Furthermore, for example, the high strength steel sheet according to aspects of the present invention may be applied to automobile structural members to reduce the weight of automobile bodies and thereby to enhance fuel efficiency. Thus, aspects of the present invention are highly valuable in industry.
Embodiments of the present invention will be described below.
First, appropriate ranges of the chemical composition of the high strength steel sheet and the reasons why the chemical composition is thus limited will be described. In the following description, “%” indicating the contents of constituent elements of steel means “mass %” unless otherwise specified.
Carbon is one of the important basic components of steel. Particularly in accordance with aspects of the present invention, carbon is an important element that affects the fraction of tempered martensite and the working embrittlement resistance. When the C content is less than 0.030%, the fraction of tempered martensite is so small that realizing 1180 MPa or higher TS is difficult. When, on the other hand, the C content is more than 0.500%, tempered martensite becomes brittle to cause deterioration in working embrittlement resistance. Thus, the C content is limited to 0.030% or more and 0.500% or less. The C content is preferably 0.050% or more. The C content is preferably 0.400% or less. The C content is more preferably 0.100% or more. The C content is more preferably 0.350% or less.
Silicon is one of the important basic components of steel. Silicon suppresses the occurrence of carbides during continuous annealing and promotes the formation of retained austenite. Thus, particularly in accordance with aspects of the present invention, silicon is an important element that affects TS and the amount of retained austenite. When the Si content is less than 0.01%, realizing 1180 MPa or higher TS is difficult. When, on the other hand, the Si content is more than 2.50%, the amount of retained austenite is increased excessively to make it difficult to achieve 85% or more YR. Thus, the Si content is limited to 0.01% or more and 2.50% or less. The Si content is preferably 0.05% or more. The Si content is preferably 2.00% or less. The Si content is more preferably 0.10% or more. The Si content is more preferably 1.20% or less.
Manganese is one of the important basic components of steel. Particularly in accordance with aspects of the present invention, manganese is an important element that affects the fraction of tempered martensite and the working embrittlement resistance. When the Mn content is less than 0.10%, the fraction of tempered martensite is so small that realizing 1180 MPa or higher TS is difficult. When, on the other hand, the Mn content is more than 5.00%, tempered martensite becomes brittle to cause deterioration in working embrittlement resistance. Thus, the Mn content is limited to 0.10% or more and 5.00% or less. The Mn content is preferably 0.50% or more. The Mn content is preferably 4.50% or less. The Mn content is more preferably 0.80% or more. The Mn content is more preferably 4.00% or less.
Phosphorus is segregated at prior austenite grain boundaries and makes the grain boundaries brittle, thereby lowering the ultimate deformability of steel sheets and causing deterioration in working embrittlement resistance. Thus, the P content needs to be 0.100% or less. The lower limit of the P content is not particularly specified. In view of the fact that phosphorus is a solid solution strengthening element and can increase the strength of steel sheets, the lower limit is preferably 0.001% or more. For the reasons above, the P content is limited to 0.100% or less. The P content is preferably 0.001% or more. The P content is preferably 0.070% or less.
Sulfur forms sulfides and lowers the ultimate deformability of steel sheets to cause deterioration in working embrittlement resistance. Thus, the S content needs to be 0.0200% or less. The lower limit of the S content is not particularly specified but is preferably 0.0001% or more due to production technique limitations. For the reasons above, the S content is limited to 0.0200% or less. The S content is preferably 0.0001% or more. The S content is preferably 0.0050% or less.
Aluminum raises the A3 transformation temperature to allow more ferrite to be contained in the microstructure. The fraction of tempered martensite is correspondingly lowered to make it difficult to realize 1180 MPa or higher TS. Thus, the Al content needs to be 1.000% or less. The lower limit of the Al content is not particularly specified. In view of the fact that aluminum suppresses the occurrence of carbides during continuous annealing and promotes the formation of retained austenite, the Al content is preferably 0.001% or more. For the reasons above, the Al content is limited to 1.000% or less. The Al content is preferably 0.001% or more. The Al content is preferably 0.500% or less.
Nitrogen forms nitrides and lowers the ultimate deformability of steel sheets to cause deterioration in working embrittlement resistance. Thus, the N content needs to be 0.0100% or less. The lower limit of the N content is not particularly specified but the N content is preferably 0.0001% or more due to production technique limitations. For the reasons above, the N content is limited to 0.0100% or less. The N content is preferably 0.0001% or more. The N content is preferably 0.0050% or less.
Oxygen forms oxides and lowers the ultimate deformability of steel sheets to cause deterioration in working embrittlement resistance. Thus, the O content needs to be 0.0100% or less. The lower limit of the O content is not particularly specified but the O content is preferably 0.0001% or more due to production technique limitations. For the reasons above, the O content is limited to 0.0100% or less. The O content is preferably 0.0001% or more. The O content is preferably 0.0050% or less.
The chemical composition of the high strength steel sheet according to an embodiment of the present invention includes the components described above, and the balance is Fe and incidental impurities. Here, the incidental impurities include Zn, Pb, As, Ge, Sr, and Cs. A total of 0.100% or less of these impurities is acceptable.
In addition to the components in the proportions described above, the high strength steel sheet according to aspects of the present invention may further include at least one element selected from, in mass %, Ti: 0.200% or less, Nb: 0.200% or less, V: 0.200% or less, Ta: 0.10% or less, W: 0.10% or less, B: 0.0100% or less, Cr: 1.00% or less, Mo: 1.00% or less, Ni: 1.00% or less, Co: 0.010% or less, Cu: 1.00% or less, Sn: 0.200% or less, Sb: 0.200% or less, Ca: 0.0100% or less, Mg: 0.0100% or less, REM: 0.0100% or less, Zr: 0.100% or less, Te: 0.100% or less, Hf: 0.10% or less, and Bi: 0.200% or less. These elements may be contained singly or in combination.
When the contents of Ti, Nb, and V are each 0.200% or less, coarse precipitates and inclusions will not occur in large amounts and thus will not cause lowering of the ultimate deformability of steel sheets; hence there will be no deterioration in working embrittlement resistance. Thus, the contents of Ti, Nb, and V are each preferably 0.200% or less. The lower limits of the contents of Ti, Nb, and V are not particularly specified. These elements form fine carbides, nitrides, or carbonitrides during hot rolling or continuous annealing to increase the strength of steel sheets. In view of this fact, the contents of Ti, Nb, and V are each more preferably 0.001% or more. When titanium, niobium, and vanadium are added, the contents thereof are each limited to 0.200% or less for the reasons above. The contents are each more preferably 0.001% or more. The contents are each more preferably 0.100% or less.
When the contents of Ta and W are each 0.10% or less, coarse precipitates and inclusions will not occur in large amounts and thus will not cause lowering of the ultimate deformability of steel sheets; hence there will be no deterioration in working embrittlement resistance. Thus, the contents of Ta and W are each preferably 0.10% or less. The lower limits of the contents of Ta and W are not particularly specified. These elements form fine carbides, nitrides, or carbonitrides during hot rolling or continuous annealing to increase the strength of steel sheets. In view of this fact, the contents of Ta and W are each more preferably 0.01% or more. When tantalum and tungsten are added, the contents thereof are each limited to 0.10% or less for the reasons above. The contents are each more preferably 0.01% or more. The contents are each more preferably 0.08% or less.
When the B content is 0.0100% or less, inner cracks that lower the ultimate deformability of steel sheets will not form during casting or hot rolling and thus there will be no deterioration in working embrittlement resistance. Thus, the B content is preferably 0.0100% or less. The lower limit of the B content is not particularly specified. The B content is more preferably 0.0003% or more in view of the fact that this element is segregated at austenite grain boundaries during annealing and enhances hardenability. When boron is added, the content thereof is limited to 0.0100% or less for the reasons above. The content is more preferably 0.0003% or more. The content is more preferably 0.0080% or less.
When the contents of Cr, Mo, and Ni are each 1.00% or less, coarse precipitates and inclusions will not occur in increased amounts and thus will not cause lowering of the ultimate deformability of steel sheets; hence there will be no deterioration in working embrittlement resistance. Thus, the contents of Cr, Mo, and Ni are each preferably 1.00% or less. The lower limits of the contents of Cr, Mo, and Ni are not particularly specified. In view of the fact that these elements enhance hardenability, the contents of Cr, Mo, and Ni are each more preferably 0.01% or more. When chromium, molybdenum, and nickel are added, the contents thereof are each limited to 1.00% or less for the reasons above. The contents are each more preferably 0.01% or more. The contents are each more preferably 0.80% or less.
When the Co content is 0.010% or less, coarse precipitates and inclusions will not occur in increased amounts and thus will not cause lowering of the ultimate deformability of steel sheets; hence there will be no deterioration in working embrittlement resistance. Thus, the Co content is preferably 0.010% or less. The lower limit of the Co content is not particularly specified. In view of the fact that this element enhances hardenability, the Co content is more preferably 0.001% or more. When cobalt is added, the content thereof is limited to 0.010% or less for the reasons above. The content is more preferably 0.001% or more. The content is more preferably 0.008% or less.
When the Cu content is 1.00% or less, coarse precipitates and inclusions will not occur in increased amounts and thus will not cause lowering of the ultimate deformability of steel sheets; hence there will be no deterioration in working embrittlement resistance. Thus, the Cu content is preferably 1.00% or less. The lower limit of the Cu content is not particularly specified. In view of the fact that this element enhances hardenability, the Cu content is preferably 0.01% or more. When copper is added, the content thereof is limited to 1.00% or less for the reasons above. The content is more preferably 0.01% or more. The content is more preferably 0.80% or less.
When the Sn content is 0.200% or less, inner cracks that lower the ultimate deformability of steel sheets will not form during casting or hot rolling and thus there will be no deterioration in working embrittlement resistance. Thus, the Sn content is preferably 0.200% or less. The lower limit of the Sn content is not particularly specified. The Sn content is more preferably 0.001% or more in view of the fact that tin enhances hardenability (in general, is an element that enhances corrosion resistance). When tin is added, the content thereof is limited to 0.200% or less for the reasons above. The content is more preferably 0.001% or more. The content is more preferably 0.100% or less.
When the Sb content is 0.200% or less, coarse precipitates and inclusions will not occur in increased amounts and thus will not cause lowering of the ultimate deformability of steel sheets; hence there will be no deterioration in working embrittlement resistance. Thus, the Sb content is preferably 0.200% or less. The lower limit of the Sb content is not particularly specified. In view of the fact that this element enables control of the thickness of surface layer softening and the strength, the Sb content is more preferably 0.001% or more. When antimony is added, the content thereof is limited to 0.200% or less for the reasons above. The content is more preferably 0.001% or more. The content is more preferably 0.100% or less.
When the contents of Ca, Mg, and REM are each 0.0100% or less, coarse precipitates and inclusions will not occur in increased amounts and thus will not cause lowering of the ultimate deformability of steel sheets; hence there will be no deterioration in working embrittlement resistance. Thus, the contents of Ca, Mg, and REM are each preferably 0.0100% or less. The lower limits of the contents of Ca, Mg, and REM are not particularly specified. In view of the fact that these elements change the shapes of nitrides and sulfides into spheroidal and enhance the ultimate deformability of steel sheets, the contents of Ca, Mg, and REM are each more preferably 0.0005% or more. When calcium, magnesium, and rare earth metal(s) are added, the contents thereof are each limited to 0.0100% or less for the reasons above. The contents are each more preferably 0.0005% or more. The contents are each more preferably 0.0050% or less.
When the contents of Zr and Te are each 0.100% or less, coarse precipitates and inclusions will not occur in increased amounts and thus will not cause lowering of the ultimate deformability of steel sheets; hence there will be no deterioration in working embrittlement resistance. Thus, the contents of Zr and Te are each preferably 0.100% or less. The lower limits of the contents of Zr and Te are not particularly specified. In view of the fact that these elements change the shapes of nitrides and sulfides into spheroidal and enhance the ultimate deformability of steel sheets, the contents of Zr and Te are each more preferably 0.001% or more. When zirconium and tellurium are added, the contents thereof are each limited to 0.100% or less for the reasons above. The contents are each more preferably 0.001% or more. The contents are each more preferably 0.080% or less.
When the Hf content is 0.10% or less, coarse precipitates and inclusions will not occur in increased amounts and thus will not cause lowering of the ultimate deformability of steel sheets; hence there will be no deterioration in working embrittlement resistance. Thus, the Hf content is preferably 0.10% or less. The lower limit of the Hf content is not particularly specified. In view of the fact that this element changes the shapes of nitrides and sulfides into spheroidal and enhances the ultimate deformability of steel sheets, the Hf content is more preferably 0.01% or more. When hafnium is added, the content thereof is limited to 0.10% or less for the reasons above. The content is more preferably 0.01% or more. The content is more preferably 0.08% or less.
When the Bi content is 0.200% or less, coarse precipitates and inclusions will not occur in increased amounts and thus will not cause lowering of the ultimate deformability of steel sheets; hence there will be no deterioration in working embrittlement resistance. Thus, the Bi content is preferably 0.200% or less. The lower limit of the Bi content is not particularly specified. In view of the fact that this element reduces the occurrence of segregation, the Bi content is more preferably 0.001% or more. When bismuth is added, the content thereof is limited to 0.200% or less for the reasons above. The content is more preferably 0.001% or more. The content is more preferably 0.100% or less.
When the content of any of Ti, Nb, V, Ta, W, B, Cr, Mo, Ni, Co, Cu, Sn, Sb, Ca, Mg, REM, Zr, Te, Hf, and Bi is below the preferred lower limit, the element does not impair the advantageous effects according to aspects of the present invention and is regarded as an incidental impurity.
Next, the steel microstructure of the high strength steel sheet according to aspects of the present invention will be described.
This configuration is a very important requirement that constitutes an aspect of the present invention. 1180 MPa or higher TS can be achieved when tempered martensite is the principal phase. In order to obtain this effect, the area fraction of tempered martensite needs to be 90% or more. Thus, the area fraction of tempered martensite is limited to 90% or more. The area fraction is preferably 94% or more, and more preferably 96% or more.
Here, tempered martensite is measured as follows. A longitudinal cross section of the steel sheet is polished and is etched with 3 vol % Nital. A portion at ¼ sheet thickness (a location corresponding to ¼ of the sheet thickness in the depth direction from the steel sheet surface) is observed using SEM in 10 fields of view at a magnification of ×2000. In the microstructure images, tempered martensite is structures that have fine irregularities inside the structures and contain inner carbides. The values thus obtained are averaged to determine the tempered martensite.
[Amount of Retained Austenite: Less than 3%]
This configuration is a very important requirement that constitutes an aspect of the present invention. When the volume fraction of retained austenite is 3% or more, it is difficult to realize 85% or more YR. The reason for low YR is that retained austenite with a high fraction gives rise to a lowering in YS by undergoing strain-induced transformation. Thus, the retained austenite is limited to less than 3%. The amount of retained austenite is preferably 1% or less. The lower limit of retained austenite is not particularly limited and may be 0%.
Here, retained austenite is measured as follows. The steel sheet is polished to expose a face 0.1 mm below ¼ sheet thickness and is thereafter further chemically polished to expose a face 0.1 mm below the face exposed above. The face is analyzed with an X-ray diffractometer using CoKα radiation to determine the integral intensity ratios of the diffraction peaks of {200}, {220}, and {311} planes of fcc iron and {200}, {211}, and {220} planes of bcc iron. Nine integral intensity ratios thus obtained are averaged to determine retained austenite.
[Area Fraction of the Total of Ferrite and Bainitic Ferrite: Less than 10%]
This configuration is a very important requirement that constitutes an aspect of the present invention. When the total of ferrite and bainitic ferrite is 10% or more, it is difficult to realize 1180 MPa or higher TS and it is also difficult to achieve 85% or more YR. The reason for low YR is that ferrite and bainitic ferrite are soft microstructures and hasten the occurrence of yielding. Thus, the total of ferrite and bainitic ferrite is limited to less than 10%. The total amount is preferably 8% or less, and more preferably 5% or less. The lower limit of the total of ferrite and bainitic ferrite is not particularly limited and may be 0%.
Here, the total of ferrite and bainitic ferrite is measured as follows. A longitudinal cross section of the steel sheet is polished and is etched with 3 vol % Nital. A portion at ¼ sheet thickness (a location corresponding to ¼ of the sheet thickness in the depth direction from the steel sheet surface) is observed using SEM in 10 fields of view at a magnification of ×2000. In the microstructure images, ferrite and bainitic ferrite are recessed structures having a flat interior and containing no inner carbides. The values thus obtained are averaged to determine the total of ferrite and bainitic ferrite.
Possible microstructures other than those described above include pearlite, fresh martensite, and acicular ferrite. These microstructures do not affect characteristics as long as their fractions are 5% or less, and thus may be present within that range.
This configuration is a very important requirement that constitutes an aspect of the present invention. Reducing the average grain size of prior austenite can suppress crack propagation and thereby enhances the working embrittlement resistance of steel sheets. In order to obtain these effects, the average grain size of prior austenite needs to be 20 μm or less. The lower limit of the average grain size of prior austenite is not particularly specified. When, however, the average grain size of prior austenite is less than 2 μm, more retained austenite may form. Thus, the average grain size is preferably 2 μm or more. For the reasons above, the average grain size of prior austenite is limited to 20 μm or less. The average grain size is preferably 2 μm or more. The average grain size is preferably 15 μm or less. The average grain size is more preferably 3 μm or more. The average grain size is more preferably 10 μm or less.
Here, the average grain size of prior austenite is measured as follows. A longitudinal cross section of the steel sheet is polished and is etched with, for example, a mixed solution of picric acid and ferric chloride to expose prior austenite grain boundaries. Portions at ¼ sheet thickness (locations corresponding to ¼ of the sheet thickness in the depth direction from the steel sheet surface) are photographed with an optical microscope each in 3 to 10 fields of view at a magnification of ×400. Twenty straight lines including 10 vertical lines and 10 horizontal lines are drawn at regular intervals on the image data obtained, and the grain size is determined by a linear intercept method.
This configuration is a very important requirement that constitutes an aspect of the present invention. The proportion of a packet having the largest area in a prior austenite grain affects the flatness in the width direction and the working embrittlement resistance. As illustrated in
Here, the average proportion of packets having the largest area in prior austenite grains is measured as follows. First, a test specimen for microstructure observation is sampled from the cold rolled steel sheet. Next, the sampled test specimen is polished by vibration polishing with colloidal silica to expose a cross section in the rolling direction (a longitudinal cross section) for use as observation surface. The observation surface is specular. Next, electron backscatter diffraction (EBSD) measurement is performed with respect to a portion at ¼ sheet thickness (a location corresponding to ¼ of the sheet thickness in the depth direction from the steel sheet surface) to obtain local crystal orientation data. Here, the SEM magnification is ×1000, the step size is 0.2 μm, the measured region is 80 μm square, and the WD is 15 mm. The local orientation data obtained is analyzed with OIM Analysis 7 (OIM), and a map (a CP map) that shows close-packed plane groups (CP groups) with different colors is created using the method described in Non Patent Literature 1. In accordance with aspects of the present invention, a packet is defined as a region or regions belonging to the same CP group. From the CP map obtained, the area of the packet having the largest area is determined and is divided by the area of the whole prior austenite grain to give the proportion of the packet having the largest area in the prior austenite grain. This analysis is performed with respect to 10 or more adjacent prior austenite grains, and the results are averaged to give the average proportion of packets having the largest area in prior austenite grains.
Next, a manufacturing method according to aspects of the present invention will be described.
In accordance with aspects of the present invention, a steel material (a steel slab) may be obtained by any known steelmaking method without limitation, such as a converter or an electric arc furnace. To prevent macro-segregation, the steel slab (the slab) is preferably produced by a continuous casting method.
In accordance with aspects of the present invention, the slab heating temperature, the slab soaking holding time, and the coiling temperature in hot rolling are not particularly limited. For example, the steel slab may be hot rolled in such a manner that the slab is heated and is then rolled, that the slab is subjected to hot direct rolling after continuous casting without being heated, or that the slab is subjected to a short heat treatment after continuous casting and is then rolled. The slab heating temperature, the slab soaking holding time, the finish rolling temperature, and the coiling temperature in hot rolling are not particularly limited. The lower limit of the slab heating temperature is preferably 1100° C. or above. The upper limit of the slab heating temperature is preferably 1300° C. or below. The lower limit of the slab soaking holding time is preferably 30 minutes or more. The upper limit of the slab soaking holding time is preferably 250 minutes or less. The lower limit of the finish rolling temperature is preferably Ar3 transformation temperature or above. Furthermore, the lower limit of the coiling temperature is preferably 350° C. or above. The upper limit of the coiling temperature is preferably 650° C. or below.
The hot rolled steel sheet thus produced is pickled. Pickling can remove oxides on the steel sheet surface and is thus important to ensure good chemical convertibility and a high quality of coating in the final high strength steel sheet. Pickling may be performed at a time or several. The hot rolled sheet that has been pickled may be cold rolled directly or may be subjected to heat treatment before cold rolling.
The rolling reduction in cold rolling and the sheet thickness after rolling are not particularly limited. The lower limit of the rolling reduction is preferably 30% or more. The upper limit of the rolling reduction is preferably 80% or less. The advantageous effects according to aspects of the present invention may be obtained without any limitations on the number of rolling passes and the rolling reduction in each pass.
The cold rolled steel sheet obtained as described above is annealed. Annealing conditions are as follows.
When the annealing temperature T1 is below 750° C., the area fraction of the total of ferrite and bainitic ferrite is 10% or more to make it difficult to realize 1180 MPa or higher TS and 85% or more YR. When, on the other hand, the annealing temperature T1 is above 950° C., prior austenite grains are excessively increased in size and the prior austenite grain size exceeds 20 μm to give rise to a decrease in working embrittlement resistance. Thus, the annealing temperature T1 is limited to 750° C. or above and 950° C. or below. The annealing temperature T1 is preferably 800° C. or above. The annealing temperature T1 is preferably 900° C. or below.
When the holding time t1 at the annealing temperature T1 is less than 10 seconds, austenitization is insufficient and the area fraction of the total of ferrite and bainitic ferrite is 10% or more. As a result, it is difficult to achieve 1180 MPa or higher TS and it is also difficult to realize 85% or more YR. When, on the other hand, the holding time at the annealing temperature T1 is more than 1000 seconds, the prior austenite grain size is excessively increased, and the working embrittlement resistance is lowered. For the reasons above, the holding time t1 at the annealing temperature T1 is limited to 10 seconds or more and 1000 seconds or less. The holding time t1 is preferably 50 seconds or more. The holding time t1 is preferably 500 seconds or less.
[Average Cooling Rate from 750° C. to 600° C.: 20° C./s or More]
When the average cooling rate from 750° C. to 600° C. is less than 20° C./s, the area fraction of the total of ferrite and bainitic ferrite is 10% or more to make it difficult to achieve 1180 MPa or higher TS and to realize 85% or more YR. For the reasons above, the average cooling rate from 750° C. to 600° C. is limited to 20° C./s or more. The average cooling rate is preferably 30° C./s or more. The upper limit is not necessarily specified but is preferably 2000° C./s or less.
[Average Cooling Rate from (Ms+50° C.) to a Quench Start Temperature T2: 5° C./s or More and 30° C./s or Less]
This configuration is a very important requirement that constitutes an aspect of the present invention. The average cooling rate from (Ms+50° C.) to a quench start temperature T2 affects the area fraction of the total of ferrite and bainitic ferrite and the average proportion of packets having the largest area in prior austenite grains. When the average cooling rate from (Ms+50° C.) to a quench start temperature T2 is less than 5° C./s, the area fraction of the total of ferrite and bainitic ferrite is 10% or more to make it difficult to achieve 1180 MPa or higher TS and to realize 85% or more YR. When, on the other hand, the average cooling rate from (Ms+50° C.) to a quench start temperature T2 is more than 30° C./s, the average proportion of packets having the largest area in prior austenite grains exceeds 70% to cause deterioration in flatness in the width direction and working embrittlement resistance. For the reasons above, the average cooling rate from (Ms+50° C.) to a quench start temperature T2 is limited to 5° C./s or more and 30° C./s or less. The average cooling rate is preferably 10° C./s or more. The average cooling rate is preferably 20° C./s or less.
This configuration is a very important requirement that constitutes an aspect of the present invention. The quench start temperature T2 is controlled to (Ms−80° C.) or above and below Ms to ensure that the martensite transformation rate before the start of quenching is 1% or more and 80% or less. In this manner, quenching can give microstructures in which the average proportion of packets having the largest area in prior austenite grains is 70% or less and the volume fraction of retained austenite is less than 3%. When the quench start temperature T2 is below (Ms−80° C.), the martensite transformation rate before the start of quenching exceeds 80% and consequently the volume fraction of retained austenite is 3% or more to make it difficult to achieve 85% or more YR. When, on the other hand, the quench start temperature T2 is above Ms, the martensite transformation rate before the start of quenching is less than 1% and the average proportion of packets having the largest area in prior austenite grains exceeds 70% to cause deterioration in flatness in the width direction and working embrittlement resistance. Thus, the quench start temperature T2 is limited to (Ms−80° C.) or above and below Ms. The quench start temperature T2 is preferably (Ms−50° C.) or above. The quench start temperature T2 is preferably (Ms−5° C.) or below. The martensite start temperature Ms (° C.) is defined by the following formula (1):
wherein [% C], [% Mn], [% Cr], [% Mo], and [% Ni] indicate the contents (mass %) of C, Mn, Cr, Mo, and Ni, respectively, and are zero when the element is absent.
[Average Cooling Rate from the Quench Start Temperature T2 to 80° C.: 300° C./s or More]
When the average cooling rate from the quench start temperature T2 to 80° C. is less than 300° C./s, the volume fraction of retained austenite is 3% or more to make it difficult to achieve 85% or more YR. Thus, the average cooling rate from the quench start temperature T2 to 80° C. is limited to 300° C./s or more. The average cooling rate is preferably 800° C./s or more. The upper limit is not necessarily specified but is preferably 2000° C./s or less.
In accordance with aspects of the present invention, tempered martensite is a microstructure that is formed when martensite at 80° C. or below is heat-treated at a tempering temperature of 100° C. or above for a holding time of 10 seconds or more. Thus, martensite is not sufficiently tempered when the tempering temperature T3 is below 100° C. The resultant microstructures will be based on as-quenched martensite, which deteriorates the working embrittlement resistance. When, on the other hand, the tempering temperature T3 is above 400° C., tempered martensite is decomposed into ferrite and the area fraction of tempered martensite is less than 90% to make it difficult to achieve 1180 MPa or higher TS. For the reasons above, the tempering temperature T3 is limited to 100° C. or above and 400° C. or below. The tempering temperature T3 is preferably 150° C. or above. The tempering temperature T3 is preferably 350° C. or below.
In accordance with aspects of the present invention, tempered martensite is a microstructure that is formed when martensite at 80° C. or below is heat-treated at a tempering temperature of 100° C. or above for a holding time of 10 seconds or more. Thus, martensite is not sufficiently tempered when the holding time t3 at the tempering temperature T3 is less than 10 seconds. The resultant microstructures will be based on as-quenched martensite, which deteriorates the working embrittlement resistance. When, on the other hand, the tempering temperature T3 is more than 10000 seconds, tempered martensite is decomposed into ferrite and the area fraction of tempered martensite is less than 90% to make it difficult to achieve 1180 MPa or higher TS. For the reasons above, the holding time t3 at the tempering temperature T3 is limited to 10 seconds or more and 10000 seconds or less. The holding time t3 is preferably 50 seconds or more. The holding time t3 is preferably 5000 seconds or less.
Post-temper cooling is not particularly limited and the steel sheet may be cooled to a desired temperature in an appropriate manner. Incidentally, the desired temperature is preferably about room temperature.
Furthermore, the high strength steel sheet described above may be worked under conditions where the amount of equivalent plastic strain is 0.10% or more and 5.00% or less. The working may be followed by reheating at 100° C. or above and 400° C. or below.
When the high strength steel sheet is a product that is traded, the steel sheet is usually traded after being cooled to room temperature.
The high strength steel sheet may be subjected to coating treatment during annealing or after annealing.
For example, the coating treatment during annealing may be hot-dip galvanizing treatment performed when the steel sheet is being cooled or has been cooled from 750° C. to 600° C. at an average cooling rate of 20° C./s or more. The hot-dip galvanizing treatment may be followed by alloying. For example, the coating treatment after annealing may be Zn—Ni electrical alloy coating treatment or pure Zn electroplated coating treatment performed after tempering. A coated layer may be formed by electroplated coating, or hot-dip zinc-aluminum-magnesium alloy coating may be applied. While the coating treatment has been described above focusing on zinc coating, the types of coating metals, such as Zn coating and Al coating, are not particularly limited. Other conditions in the manufacturing method are not particularly limited. From the point of view of productivity, the series of treatments including annealing, hot-dip galvanizing, and alloying treatment of the coated zinc layer is preferably performed on hot-dip galvanizing line CGL (continuous galvanizing line). To control the coating weight of the coated layer, the hot-dip galvanizing treatment may be followed by wiping. Conditions for operations, such as coating, other than those conditions described above may be determined in accordance with the usual hot-dip galvanizing technique.
After the coating treatment after annealing, the steel sheet may be worked again under conditions where the amount of equivalent plastic strain is 0.10% or more and 5.00 or less. The working may be followed by reheating at 100° C. or above and 400° C. or below.
Steels having a chemical composition described in Table 1 and 2, with the balance being Fe and incidental impurities, were smelted in a converter and were continuously cast into slabs. Next, the slabs obtained were heated, hot rolled, pickled, cold rolled, and subjected to annealing treatment and tempering treatment described in Tables 3 to 6. High strength cold rolled steel sheets having a sheet thickness of 0.6 to 2.2 mm were thus obtained. Incidentally, some of the steel sheets were subjected to coating treatment during or after annealing.
0.021
0.522
0.003
2.532
0.08
5.12
0.121
0.0222
1.135
0.0112
0.213
0.0123
0.204
1.12
743
965
1015
15
38
349
496
284
554
529
420
The high strength cold rolled steel sheets obtained as described above were used as test steels. Tensile characteristics, flatness in the width direction, and working embrittlement resistance were evaluated in accordance with the following test methods.
The amount of tempered martensite, the amount of retained austenite, the total amount of ferrite and bainitic ferrite, and the average grain size of prior austenite were determined by the methods described hereinabove.
The average proportion of packets having the largest area in prior austenite grains was determined by the method described hereinabove.
A JIS No. 5 test specimen (gauge length: 50 mm, parallel section width: 25 mm) was sampled so that the longitudinal direction of the test specimen would be perpendicular to the rolling direction. A tensile test was performed in accordance with JIS Z 2241 under conditions where the crosshead speed was 1.67×10−1 mm/sec. YS and TS were thus measured. In accordance with aspects of the present invention, 1180 MPa or higher TS was determined to be acceptable, and 85% or more yield ratio (YR) was determined to be acceptable. YR is determined from the formula (2) below:
The cold rolled steel sheets obtained as described above were analyzed to measure the flatness in the width direction. The measurement is illustrated in
The working embrittlement resistance was evaluated by Charpy test. A Charpy test specimen was a 2 mm deep V-notched test piece that was a stack of steel sheets fastened together with bolts to eliminate any gaps between the steel sheets. The number of steel sheets that were stacked was controlled so that the thickness of the stack as the test piece would be closer to 10 mm. When, for example, the sheet thickness was 1.2 mm, eight sheets were stacked to give a 9.6 mm thick test piece. The sheets for stacking into the Charpy test specimen were sampled so that the width direction would be the longitudinal direction. As an index of the working embrittlement resistance, the ratio vE0%/vE10% of the absorbed impact energy at room temperature of the as-produced (unworked) steel sheet to that of the steel sheet after 10% rolling was measured. The working embrittlement resistance was rated as “x” when vE0%/vE10% was less than 0.6, as “◯” when VE0%/vE10% was 0.6 or more and less than 0.7, and as “⊚” when vE0%/VE10% was 0.7 or more. The Charpy test specimen was evaluated as “excellent in working embrittlement resistance” when vE0%/vE10% was 0.6 or more. Conditions other than those described above conformed to JIS Z 2242:2018.
The results are described in Tables 7 to 10. As shown in the tables, INVENTIVE EXAMPLES achieved 1180 MPa or higher TS, 85% or more YR, excellent flatness in the width direction, and excellent working embrittlement resistance. In contrast, COMPARATIVE EXAMPLES were unsatisfactory in one or more of TS, YR, flatness in the width direction, and working embrittlement resistance.
72
28
1020
73
24
72
17
82
24
72
14
82
72
12
81
89
7
78
95
79
6
82
5
83
6
83
79
7
74
93
72
20
76
1154
8
78
72
19
81
72
22
1002
78
88
X
X
| Number | Date | Country | Kind |
|---|---|---|---|
| 2022-049756 | Mar 2022 | JP | national |
This is the U.S. National Phase application of PCT/JP2023/002913 filed Jan. 30, 2023 which claims priority to Japanese Patent Application No. 2022-049756, filed Mar. 25, 2022, the disclosures of these applications being incorporated herein by reference in their entireties for all purposes.
| Filing Document | Filing Date | Country | Kind |
|---|---|---|---|
| PCT/JP2023/002913 | 1/30/2023 | WO |
