Patent Application
20250034686
The present invention relates to a hot stamped body.
In recent years, in the automobile industry, lighter weight of car bodies has been sought from the viewpoint of improvement of fuel economy. To achieve both lighter weight of car bodies and collision safety, one effective method is to increase the strength of the steel sheet used. A high strength steel sheet is being developed due to such a background.
If making a steel sheet high in strength, the formability falls, and therefore it is generally difficult to achieve both strength and formability in the steel sheet. Hot stamping (hot pressing) is known as a technique for press-forming a material, which is difficult to form, such as a high strength steel sheet. Hot stamping is a technique of hot forming which heats then forms a material to be formed. This technique heats then forms the material, and therefore at the time of forming, the steel material is soft and has good formability. Therefore, even a high strength steel material can be formed into a complex shape with a good precision. Further, it is hardened at the same time as being formed by the press dies, and therefore a formed steel material is known to have sufficient strength.
In relation to this, PTL 1 describes a hot stamped body characterized by having a predetermined chemical composition and a microstructure containing former austenite with an average grain size of 3 μm or less and, further, containing at least one of lower bainite, martensite, and tempered martensite in an area ratio of 90% or more, wherein a grain boundary solid solution ratio Z, defined by Z=(mass % of one or both of Nb and Mo at the grain boundaries)/(mass % of one or both of Nb and Mo at time of melting), is 0.3 or more. Further, PTL 1 describes that the hot stamped body has a tensile strength of 2000 MPa or more.
If the strength of a steel material becomes higher, in general, the phenomenon of the steel material fracturing before reaching its maximum stress (early fracture) easily arises. For this reason, there is a high need for a steel material enabling suppression of such early fracture. In relation to this, PTL 1 teaches that in a hot stamped body having the above feature, there will be a high tensile strength of 2000 MPa or more and, in addition, early fracture will be suppressed. On the other hand, in the automobile industry, etc., further reduction of weight of the steel material is sought. To achieve such lighter weight, a need arises to raise the strength more than the past. Therefore, there is still a great need for a steel material, more specifically a hot stamped body, able to solve the problem of early fracture even if raising the strength equal to the past or more than the same.
Therefore, the present invention has as its object to provide a hot stamped body which is high in strength and able to suppress early fracture by a novel constitution.
The inventors discovered that, to achieve the above object, it is possible to reduce the variation in former austenite grain size in the microstructure of a hot stamped body so as to suppress the rise in local hardness acting as starting points of fracture and in addition possible to make specific elements segregate at the grain boundaries to reinforce the grain boundaries and further discovered that by the combination of such suppression of rise of local hardness and grain boundary strengthening, it is possible to sufficiently suppress early fracture regardless of the hot stamped body having a high tensile strength, and thereby completed the present invention.
The present invention able to achieve this object is as follows:
According to the present invention, it is possible to provide a hot stamped body which is high in strength and able to suppress early fracture.
The hot stamped body according to an embodiment of the present invention has a chemical composition comprising, by mass %,
As explained above, there is the problem that if the strength of a steel material becomes higher, in general, the phenomenon of the steel material fracturing before reaching the maximum stress (early fracture) easily occurs. Therefore, the inventors conducted studies focusing on the two viewpoints of the viewpoint of reducing hard regions able to become starting points for fracture and the viewpoint of strengthening the grain boundaries for preventing or suppressing fracture. More specifically, the inventors first discovered that if there is a large variation in former austenite grain size in the microstructure, the hardness becomes higher in a region with a smaller former austenite grain size and that such a local high hardness region can become a starting point of early fracture. As opposed to this, the inventors discovered that by controlling the standard deviation in a grain size distribution of former austenite grains to 5.0 μm or less, it is possible to reliably suppress such a rise in local hardness.
While not intending to be bound to any specific theory, it is believed that at the time of hot stamping, the starting temperature of martensite transformation changes in accordance with the size of the austenite grains. If explained in more detail, it is believed that austenite grains having larger size are higher in starting temperature of martensite transformation compared with austenite grains having smaller size, and therefore the hardness becomes lower. Austenite grains having smaller size rise in hardness since martensite transformation occurs at a lower temperature than large grains. Therefore, to suppress or reduce the rise in such local hardness, it becomes important to reduce the variation in the austenite grain size before martensite transformation. Due to such a reason, it is believed that by controlling the standard deviation in the grain size distribution of the former austenite grains to 5.0 μm or less to reduce the variation in former austenite grain size, it would become possible to remarkably suppress the rise in local hardness due to differences in timing of martensite transformation. If there is locally a region of a high hardness, it is believed that there will be a high possibility of such a region acting as a starting point triggering early fracture, and therefore reducing the variation in former austenite grain size would be extremely effective in suppressing early fracture.
In relation to this, as explained later in detail regarding the method of production of the hot stamped body, the inventors focused on the microstructure of a hot rolled steel sheet and discovered that by evenly dispersing the pearlite in the microstructure, it is possible to control the standard deviation of former austenite grains in the final microstructure of the hot stamped body to 5.0 μm or less. Along with the increasingly higher strength of steel materials, sometimes a relatively large amount of Mn is added so as to improve the hardenability of the steel material, but in the research by the inventors this time, it was learned that with such a high Mn content (for example, 0.50 mass % or more), pearlite is relatively easily formed and therefore compared with the case of a low Mn content, it is extremely difficult to evenly disperse the pearlite formed in large amounts in the microstructure of the hot rolled steel sheet. However, the inventors discovered that to deal with such a problem, it is possible to uniformly disperse the pearlite in the microstructure of the hot rolled steel sheet by relatively high reduction at the final stage of finish rolling and that as a result it is possible to remarkably reduce the variation in former austenite grain size in the final microstructure of the hot stamped body. On the other hand, with just reducing the variation in former austenite grain size and reducing the regions which can act as starting points of early fracture, there is a possibility that if a crack occurs, it will not be possible to reliably suppress its progression and prevent early fracture.
Therefore, next, the inventors engaged in further studies from the viewpoint of suppressing progression of cracks along the grain boundaries and discovered that by causing specific elements, more specifically at least one of Mo, W, Ta, Re, Os, Ir, and Tc, to segregate at the former austenite grain boundaries in a total amount of segregation of these of 0.10 atm % or more, it is possible to strengthen the former austenite grain boundaries at the microstructure of the hot stamped body. As a result, the inventors discovered that by the combination of suppression of the rise in local hardness explained above and the strengthening of former austenite grain boundaries by such specific grain boundary strengthening elements, despite the hot stamped body having a high tensile strength, it is possible to sufficiently suppress early fracture. In the past, the fact that, for example, from the viewpoint of improvement of the hardenability, etc., while it is known to add some of these elements to a steel sheet for hot stamping use, it is possible to add at least one of Mo, W, Ta, Re, Os, Ir, and Tc in a predetermined total amount and, further, control the heat treatment conditions to make them segregate at the former austenite grain boundaries and possible to strengthen the former austenite grain boundaries in a superhigh strength steel material containing 0.40 mas % or more of carbon and thereby suppress the progression of cracks along the grain boundaries was clarified the first time by the inventors this time. In the hot stamped body according to an embodiment of the present invention, by combining the suppression of rise of local hardness and grain boundary strengthening by grain boundary segregation of such specific elements, it is possible to reduce the regions which act as starting points of early fracture and reliably suppress the progression of cracks along the grain boundaries even in the case of occurrence of cracks at the hot stamped body. For this reason, according to the hot stamped body according to an embodiment of the present invention, early fracture can be suppressed regardless of having a high tensile strength, for example, a high tensile strength of 2200 MPa or more.
Below, the hot stamped body according to the embodiment of the present invention will be explained in more detail. In the following explanation, the “%” of the units of content of the elements, unless otherwise indicated, means “mass %”. Further, in this Description, “to” showing a numerical range, unless otherwise indicated, is used in the sense including the numerical values described before and after it as the upper limit value and lower limit value.
C is an element improving the strength of a hot stamped body. If the C content is less than 0.40%, it is not possible to obtain the desired strength at the hot stamped body. For this reason, the C content is 0.40% or more. The C content is preferably 0.42% or more, 0.44% or more, or 0.45% or more.
On the other hand, if the C content is more than 0.70%, the toughness of the martensite is too low and an excellent early fracture resistance cannot be obtained. For this reason, the C content is 0.70% or less. Preferably, the C content is 0.67% or less, 0.65% or less, or 0.60% or less.
Si is an element improving the strength of a hot stamped body by solid solution strengthening. If the Si content is less than 0.010%, it is not possible to obtain the desired strength. For this reason, the Si content is 0.010% or more. The Si content is preferably 0.05% or more, 0.10% or more, 0.15% or more, or 0.20% or more.
On the other hand, if the Si content is more than 3.00%, the amount of ferrite increases and the desired metallographic structure cannot be obtained. For this reason, the Si content is 3.00% or less. The Si content is preferably 2.50% or less, 2.00% or less, 1.00% or less, or 0.70% or less.
Mn is an element which promotes the transformation from austenite to pearlite in a hot rolled steel sheet in the process of production of a hot stamped body according to the present embodiment and contributes to control of the former austenite grain size distribution of a hot stamped body. To make the standard deviation in the grain size distribution of former austenite grains the desired range, the Mn content is 0.50% or more. The Mn content is preferably 0.70% or more, 1.00% or more, or 1.30% or more.
On the other hand, if the Mn content is more than 3.00%, transformation from austenite to pearlite in a hot rolled steel sheet is promoted too much and the standard deviation in the grain size distribution of former austenite grains in a hot stamped body cannot be controlled to a desired range. For this reason, the Mn content is 3.00% or less. Preferably, the Mn content is 2.70% or less, 2.50% or less, 2.30% or less, or 2.00% or less.
P is an impurity element and segregates at the grain boundaries to form starting points of fracture and cause the early fracture resistance to deteriorate. For this reason, the P content is 0.100% or less. The P content is preferably 0.050% or less or 0.010% or less. The lower limit of the P content is not particularly prescribed, but if less than 0.0001%, the dephosphorization cost greatly rises making this not preferable economically. For this reason, the P content may also be 0.0001% or more.
S is an impurity element and forms inclusions in the steel. The inclusions become starting points of fracture and cause the early fracture resistance to deteriorate, therefore the S content is 0.0100% or less. The S content is preferably 0.0080% or less, 0.0050% or less, or 0.0030% or less.
The lower limit of the S content is not particularly prescribed, but if less than 0.0001%, the desulfurization cost greatly rises making this not preferable economically. For this reason, the S content may also be 0.0001% or more.
N is an impurity element and forms inclusions in the steel. The inclusions become starting points of fracture and cause the early fracture resistance to deteriorate, therefore the N content is 0.0200% or less. The N content is preferably 0.0150% or less, 0.0100% or less, 0.0060% or less, or 0.0040% or less.
The lower limit of the N content is not particularly prescribed, but if less than 0.0001%, the denitridation cost greatly rises making this not preferable economically. For this reason, the N content may also be 0.0001% or more.
O, if contained in a large amount in the steel, forms coarse oxides acting as starting points of fracture and causes the early fracture resistance of a hot stamped body to deteriorate. For this reason, the O content is 0.0200% or less. The O content is preferably 0.0100% or less, 0.0070% or less, or 0.0040% or less.
From the viewpoint of reducing the refining costs, the O content may also be 0.0001% or more. To make a large number of fine oxides disperse at the time of deoxidation of the molten steel, the O content may be 0.0005% or more.
Al is an element having the action of deoxidizing the molten steel and making the steel sounder. If the Al content is less than 0.0010%, deoxidation will not sufficiently proceed and coarse oxides will be formed causing the early fracture resistance to deteriorate. For this reason, the Al content is 0.0010% or more. The Al content is preferably 0.005% or more, 0.010% or more, or 0.030% or more.
On the other hand, if the Al content is more than 0.500%, coarse oxides will form in the steel causing the early fracture resistance of a hot stamped body to fall. For this reason, the Al content is 0.500% or less. The Al content is preferably 0.400% or less, 0.300% or less, 0.200% or less, or 0.100% or less.
Nb is an element forming carbonitrides in steel and improving the strength of a hot stamped body by precipitation strengthening. If the Nb content is less than 0.0010%, the desired strength cannot be obtained. For this reason, the Nb content is 0.0010% or more. The Nb content is preferably 0.005% or more, 0.009% or more, or 0.015% or more.
On the other hand, if the Nb content is more than 0.100%, a large amount of coarse carbonitrides are formed in the steel and the early fracture resistance of a hot stamped body deteriorates. For this reason, the Nb content is 0.100% or less. The Nb content is preferably 0.080% or less, 0.060% or less, or 0.050% or less.
Ti is an element forming carbonitrides in steel and improving the strength of a hot stamped body by precipitation strengthening. If the Ti content is less than 0.010%, the desired strength cannot be obtained. For this reason, the Ti content is 0.010% or more. The Ti content is preferably 0.015% or more, 0.020% or more, or 0.025% or more. On the other hand, if the Ti content is more than 0.200%, a large amount of coarse carbonitrides are formed in the steel and the early fracture resistance of a hot stamped body deteriorates. For this reason, the Ti content is 0.200% or less. The Ti content is preferably 0.150% or less, 0.100% or less, 0.080% or less, 0.060% or less, or 0.050% or less.
Cr is an element dissolving in the former austenite grains at the time of heating before hot stamping and thereby raising the strength of a hot stamped body. If the Cr content is less than 0.010%, it is not possible to obtain the desired strength. For this reason, the Cr content is 0.010% or more. The Cr content is preferably 0.10% or more, 0.15% or more, or 0.20% or more. On the other hand, if the Cr content is more than 1.00%, in a hot stamped body, coarse intermetallic compounds are formed and the early fracture resistance deteriorates. For this reason, the Cr content is 1.00% or less. The Cr content is preferably 0.80% or less, 0.70% or less, 0.50% or less, or 0.40% or less.
Mo is an element segregating at the austenite grain boundaries at the time of heating in the hot stamping step to thereby make the strength of the former austenite grain boundaries rise and raise the early fracture resistance in the hot stamped body. If the Mo content is less than 0.0010%, the desired early fracture resistance cannot be obtained. For this reason, the Mo content is 0.0010% or more. The Mo content is preferably 0.010% or more, 0.050% or more, or 0.100% or more.
On the other hand, if the Mo content is more than 1.000%, in a hot stamped body, coarse intermetallic compounds are formed and the early fracture resistance deteriorates. For this reason, the Mo content is 1.000% or less. The Mo content is preferably 0.800% or less, 0.600% or less, or 0.400% or less.
B is an element improving the hardenability of steel. If the B content is less than 0.0005%, the desired strength cannot be obtained. For this reason, the B content is 0.0005% or more. The B content is preferably 0.0010% or more, 0.0015% or more, or 0.0020% or more. On the other hand, if the B content is more than 0.0200%, coarse intermetallic compounds are formed at a hot stamped body and the early fracture resistance of the hot stamped body falls. For this reason, the B content is 0.0200% or less. The B content is preferably 0.0150% or less, 0.0100% or less, 0.0080% or less, 0.0060% or less, or 0.0040% or less.
The basic chemical composition of a hot stamped body according to an embodiment of the present invention is as explained above. Furthermore, the hot stamped body may, if necessary, contain at least one of the following optional elements in place of part of the Fe of the balance. For example, the hot stamped body may contain at least one element selected from the group comprising Co: 0 to 4.00%, Ni: 0 to 3.00%, Cu: 0 to 3.00%, and V: 0 to 3.00%. Further, the hot stamped body may contain at least one element selected from the group comprising Ca: 0 to 1.000%, Mg: 0 to 1.000%, and REM: 0 to 1.000%. Further, the hot stamped body may also have at least one element selected from the group comprising Sb: 0 to 1.00%, Sn: 0 to 1.00%, and Zr: 0 to 1.00%. Further, the hot stamped body may contain As: 0 to 0.100%. Further, the hot stamped body may contain at least one element of W, Ta, Re, Os, Ir, and Tc in a total of 0 to 1.00%. Below, these optional elements will be explained in detail.
Co is an element improving the strength of a hot stamped body by solid solution strengthening. The Co content may be 0.001% or more, but to reliably obtain this effect, the Co content is preferably 0.01% or more or 0.05% or more.
On the other hand, even if made to be contained in a large amount, the effect becomes saturated, therefore the Co content is preferably 4.00% or less. The Co content may also be 3.00% or less, 2.50% or less, 2.00% or less, or 1.50%% or less.
Ni has the action of dissolving in the austenite grains at the time of heating in the hot stamping step and thereby raising the strength of a hot stamped body. The Ni content may be 0.001% or more, but to reliably obtain this effect, the Ni content is preferably 0.01% or more.
On the other hand, even if made to be contained in a large amount, the effect becomes saturated, therefore the Ni content is preferably 3.00% or less. The Ni content may also be 2.50% or less, 2.00% or less, 1.50% or less, 1.00% or less, or 0.80% or less.
Cu has the action of dissolving in the austenite grains at the time of heating in the hot stamping step and thereby raising the strength of a hot stamped body. The Cu content may be 0.001% or more, but to reliably obtain this effect, the Cu content is preferably 0.01% or more or 0.05% or more.
On the other hand, even if made to be contained in a large amount, the effect becomes saturated, therefore the Cu content is preferably 3.00% or less. The Cu content may also be 2.50% or less, 2.00% or less, 1.50% or less, 1.00% or less, or 0.80% or less.
V has the effect of forming carbonitrides in the steel to thereby improve the strength of the hot stamped body by precipitation strengthening. The V content may be 0.001% or more, but to reliably obtain this effect, the V content is preferably 0.01% or more or 0.05% or more. On the other hand, the V content is more than 3.00%, sometimes a large amount of carbonitrides are formed in the steel and the early fracture resistance of the hot stamped body deteriorates. For this reason, the V content is preferably 3.00% or less. The V content may also be 2.50% or less, 2.00% or less, 1.50% or less, 1.00% or less, or 0.80% or less.
Ca suppresses the formation of oxides acting as starting points of fracture and contributes to improvement of the early fracture resistance. The Ca content may be 0.0001% or more, but to reliably obtain this effect, the Ca content is preferably 0.0005% or more or 0.001% or more. On the other hand, even if made to be contained in a large amount, the effect becomes saturated, therefore the Ca content is preferably 1.000% or less. The Ca content may also be 0.100% or less, 0.050% or less, 0.010% or less, 0.005% or less, or 0.002% or less.
Mg forms oxides and sulfides in the molten steel to suppress the formation of coarse MnS, causes dispersion of large number of fine oxides, and contributes to increased fineness of the metallographic structure and improvement of the early fracture resistance. The Mg content may be 0.0001% or more, but to reliably obtain this effect, the Mg content is preferably 0.0005% or more or 0.001% or more.
On the other hand, even if made to be contained in a large amount, the effect becomes saturated, therefore the Mg content is preferably 1.000% or less. The Mg content may also be 0.100% or less, 0.050% or less, 0.010% or less, 0.005% or less, or 0.002% or less.
REM suppresses the formation of oxides acting as starting points of fracture and contributes to improvement of the early fracture resistance. The REM content may be 0.0001% or more, but to reliably obtain this effect, the REM content is preferably 0.0005% or more or 0.001% or more.
On the other hand, even if made to be contained in a large amount, the effect becomes saturated, therefore the REM content is preferably 1.000% or less. The REM content may be 0.100% or less, 0.050% or less, 0.010% or less, 0.005% or less, or 0.002% or less.
In the present embodiment, “REM” is the general term for the 17 elements of atomic number 21 scandium (Sc), atomic number 39 yttrium (Y), and the lanthanoids of atomic number 57 lanthanum (La) to atomic number 71 lutetium (Lu). The REM content is the total content of these elements.
Sb suppresses the formation of oxides acting as starting points of fracture and contributes to improvement of the early fracture resistance. To reliably obtain this effect, the Sb content is preferably 0.001% or more.
On the other hand, even if made to be contained in a large amount, the effect becomes saturated, therefore the Sb content is preferably 1.00% or less. The Sb content may also be 0.80% or less, 0.50% or less, 0.20% or less, or 0.10% or less.
Sn suppresses the formation of oxides which act as starting points of fracture and contributes to the improvement of the early fracture resistance. To reliably obtain this effect, the Sn content is preferably 0.001% or more.
On the other hand, even if contained in a large amount, the above effect is saturated, therefore the Sn content is preferably 1.00% or less. The Sn content may also be 0.80% or less, 0.50% or less, 0.20% or less, or 0.10% or less.
Zr suppresses the formation of oxides which act as starting points of fracture and contributes to the improvement of the early fracture resistance. To reliably obtain this effect, the Zr content is preferably 0.001% or more.
On the other hand, even if contained in a large amount, the above effect is saturated, therefore the Zr content is preferably 1.00% or less. The Zr content may also be 0.80% or less, 0.50% or less, 0.20% or less, or 0.10% or less.
As causes the temperature for forming an austenite single phase to fall and thereby refines the former austenite grains and contributes to improvement of the early fracture resistance. If reliably obtaining this effect, the As content is preferably 0.001% or more. On the other hand, even if contained in a large amount, the above effect is saturated, therefore the As content is preferably 0.100% or less. The As content may be 0.080% or less, 0.050% or less, 0.020% or less, or 0.010% or less.
W, Ta, Re, Os, Ir, and Tc are elements segregating at the former austenite grain boundaries in the same way as Mo to raise the strength of the grain boundaries. The total of the content of the at least one element of W, Ta, Re, Os, Ir, and Tc may be 0%, but to obtain such an effect, is preferably 0.001% or more. The total of the content of the at least one element of W, Ta, Re, Os, Ir, and Tc is preferably 0.01% or more, more preferably 0.10% or more, still more preferably 0.15% or more. On the other hand, even if excessively containing these elements, the effect becomes saturated. Therefore, including these elements in the steel material more than necessary is liable to invite a rise in the production costs. Therefore, the total of the contents of the at least one of W, Ta, Re, Os, Ir, and Tc is preferably 1.00% or less and may also be 0.80% or less, 0.60% or less, or 0.40% or less.
The chemical composition of the above hot stamped body may be measured by a general analysis method. For example, it may be measured using ICP-AES (Inductively Coupled Plasma-Atomic Emission Spectrometry). C and S may be measured using the combustion-infrared absorption method, N may be measured using the inert gas melting-thermal conductivity method, and O may be measured by the inert gas melting-nondispersion type infrared absorption method.
If the surface of the hot stamped body is provided with a plating layer, mechanical polishing may be used to remove the plating layer, then the chemical composition may be analyzed.
In a hot stamped body according to an embodiment of the present invention, the balance besides the above elements is comprised of Fe and impurities. The “impurities” are constituents, etc., entering due to various factors in the production process starting from materials such as ore and scrap, etc., when industrially producing hot stamped bodies.
The microstructure of the hot stamped body preferably includes, by area ratio, at least one of martensite, bainite, and tempered martensite in a total of 90% or more. The remaining structure is not particularly limited, but may also be comprised of at least one of 10% or less of ferrite, retained austenite, and pearlite. Martensite, bainite, and tempered martensite are extremely hard structures, therefore by the hot stamped body containing at least one of martensite, bainite, and tempered martensite in an area ratio of a total of 90% or more, a high tensile strength, specifically a tensile strength of 2200 MPa or more, can be achieved. The total of the area ratios of the least one of martensite, bainite, and tempered martensite is preferably 94% or more, more preferably 97% or more. The upper limit of the total of the area ratios of the at least one of martensite, bainite, and tempered martensite is not particularly prescribed and may also be 100%.
The microstructure in a hot stamped body is identified and the area ratios are calculated in the following way.
A sample is cut out from any position 50 mm or more away from the ends of a hot stamped body (if not possible to obtain a sample from this position, a position away from the ends) so as to enable a cross-section of thickness vertical to the surface to be examined. The size of the sample depends on the measurement device, but is a size enabling 10 mm or so to be examined in a direction vertical to the thickness direction.
The cross-section of the sample is polished using #600 to #1500 silicon carbide paper, then a liquid comprised of particle size 1 to 6 μm diamond powder dispersed in alcohol or other diluent or pure water is used to polish the surface to a mirror finish. Next, the examined surface is finished by electrolytic polishing. An area of a length 50 μm and 50 μm in the sheet thickness direction centered at a ¼ depth position of the sheet thickness at any position in the long direction of the sample cross-section is measured at 0.1 μm measurement intervals by electron backscatter diffraction to obtain crystal orientation information. For the measurement, an EBSD analysis apparatus comprised of a thermal field emission type scan electron microscope and EBSD detector may be used. For example, an EBSD analysis apparatus comprised of a JSM-7001F made by JEOL and a DVC5 model detector made by TSL may be used. At that time, the vacuum degree inside the EBSD analysis apparatus may be 9.6×10−5 Pa or less, the acceleration voltage may be 15 kV, and the beam current level may be made 13.
The obtained crystal orientation information is analyzed using the “Phase Map” function included in the software “OIM Analysis®” attached to the EBSD analysis apparatus. Structures with fcc crystal structures are judged to be retained austenite. The area ratio of the retained austenite is obtained by calculating the area ratio of this retained austenite. Next, regions with bcc crystal structures are judged to be bainite, tempered martensite, martensite, and ferrite. In these regions, using the “Grain Average Misorientation” function included in the software “OIM Analysis®” attached to the EBSD analysis apparatus, under conditions deeming a 5° grain boundary as a crystal grain boundary, a region having a “Grain Average Misorientation” of 0.5° or less is extracted as ferrite. The area ratio of ferrite is obtained by calculating the area ratio of the extracted ferrite.
Next, the remaining region (region with “Grain Average Misorientation” of more than 0.5°) is made the area ratio of the total of martensite, tempered martensite, and bainite. The area ratio of pearlite is calculated by subtracting from 100% the area ratio of the retained austenite and the area ratios of the bainite, tempered martensite, martensite, and ferrite.
In an embodiment of the present invention, the standard deviation at the grain size distribution of the former austenite grains is 5.0 μm or less. If the variation of the former austenite grain size is large, sometimes a rise in the local hardness is invited and early fracture is caused. According to an embodiment of the present invention, by controlling the standard deviation at the grain size distribution of the former austenite grains to 5.0 μm or less to reduce the variation in former austenite grain size, it is possible to reliably suppress the rise in local hardness acting as the starting points of early fracture. Preferably, the standard deviation is 4.0 μm or less, 3.0 μm or less, or 2.5 μm or less. From the viewpoint of suppressing the rise in local hardness acting as the starting points of early fracture, the smaller the standard deviation at the grain size distribution of the former austenite grains, the better. For this reason, the lower limit does not particularly have to be prescribed. For example, the standard deviation at the grain size distribution of the former austenite grains may be 0.1 μm or more and may be 0.5 μm or more, 1.0 μm or more, 1.2 μm or more, 1.5 μm or more, or 1.7 μm or more.
In an embodiment of the present invention, as explained above, it is important to control the standard deviation at the grain size distribution of the former austenite grains to 5.0 μm or less to reduce the variation of the former austenite grain size. For this reason, it is not necessary to control the former austenite grain size itself to a specific range. Therefore, the former austenite grain size is not particularly limited, but for example may be 10 μm or less. The “former austenite grain size” means the average crystal grain size of former austenite grains in the measurement of the standard deviation explained below.
The standard deviation at the grain size distribution of the former austenite grains can be obtained by the following method.
A sample is cut out from any position 50 mm or more away from the ends of the hot stamped body (if not possible to obtain a sample from this position, a position away from the ends) so as to enable a cross-section of thickness vertical to the surface to be examined. The size of the sample depends on the measurement device, but is a size enabling 10 mm or so to be examined in a direction vertical to the thickness direction.
The cross-section of the sample is polished using #600 to #1500 silicon carbide paper, then a liquid comprised of particle size 1 to 6 μm diamond powder dispersed in alcohol or other diluent or pure water is used to polish the surface to a mirror finish. Next, the examined surface is finished by electrolytic polishing. An area of a length 50 μm and 50 μm in the sheet thickness direction centered at a ¼ depth position of the sheet thickness at any position in the long direction of the sample cross-section is measured at 0.1 μm measurement intervals by electron backscatter diffraction to obtain crystal orientation information. For the measurement, an EBSD analysis apparatus comprised of a thermal field emission type scan electron microscope and EBSD detector may be used. For example, an EBSD analysis apparatus comprised of a JSM-7001F made by JEOL and a DVC5 model detector made by TSL may be used. At that time, the vacuum degree inside the EBSD analysis apparatus may be 9.6×10−5 Pa or less, the acceleration voltage may be 15 kV, and the beam current level may be made 13.
The obtained crystal orientation information is used to calculate the crystal orientation of the former austenite grains from the relationship of crystal orientation between general former austenite grains and crystal grains having bcc structures after transformation. This is used to calculate the average crystal grain size of the former austenite grains.
For the method of calculating the crystal orientations of the former austenite grains, the following method is used. First, a crystal orientation map of the former austenite grains is prepared by the method described in Acta Materialia, 58(2010), 6393-6403. The average value between the shortest diameter and the longest diameter of one former austenite grain included in the examined field is calculated. That average value is made the size of the former austenite grain. The above operation is performed for all of the former austenite grains except for the former austenite grains where the crystal grains as a whole are not included in the captured field, such as at the end parts of the captured field, to find the sizes of all of the former austenite grains in the captured field. From the obtained sizes of all former austenite grains, the average size is calculated whereupon the standard deviation at the grain size distribution of the former austenite grains is obtained.
In an embodiment of the present invention, the total amount of segregation of at least one of Mo, W, Ta, Re, Os, Ir, and Tc at the former austenite grain boundaries is 0.10 atm % or more. By making at least one of Mo, W, Ta, Re, Os, Ir, and Tc segregate at the former austenite grain boundaries to give a total amount of segregation of 0.10 atm % or more, it is possible to strengthen the former austenite grain boundaries at the microstructure of the hot stamped body. It is possible to not only suppress variation of the former austenite grain size to reduce the regions able to act as the starting points of early fracture, it is also possible to suppress the progression of cracks when they occur and thereby prevent early fracture. According to an embodiment of the present invention, by combining the suppression of the rise of local hardness and grain boundary strengthening by grain boundary segregation of such specific elements, it is possible to reduce the regions able to act as the starting points of early fracture and suppress the progression of cracks along the grain boundary even in the case of occurrence of cracks at the hot stamped bodies. For this reason, despite having a high tensile strength, for example, a high tensile strength of 2200 MPa or more, it is possible to suppress early fracture. The total amount of segregation of the at least one element of Mo, W, Ta, Re, Os, Ir, and Tc at the former austenite grain boundaries is preferably 0.13 atm % or more, more preferably 0.15 atm % or more, still more preferably 0.18 atm % or more, or 0.20 atm % or more. From the viewpoint of grain boundary strengthening, the higher the total amount of segregation of the at least one of Mo, W, Ta, Re, Os, Ir, and Tc at the former austenite grain boundaries, the more preferable. For this reason, the upper limit of the above total content is not particularly limited, but for example the total amount of segregation may be 3.00 atm % or less and may also be 2.00 atm % or less, 1.50 atm % or less, 1.00 atm % or less, 0.80 atm % or less, 0.60 atm % or less, 0.40 atm % or less, or 0.30 atm % or less.
In one embodiment, the amount of segregation of Mo at the former austenite grain boundaries may be 0.10 atm % or more, 0.13 atm % or more, 0.15 atm % or more, 0.18 atm % or more, or 0.20 atm % or more. Similarly, the amount of segregation of Mo at the former austenite grain boundaries may be 3.00 atm % or less, 2.00 atm % or less, 1.50 atm % or less, 1.00 atm % or less, 0.80 atm % or less, 0.60 atm % or less, 0.40 atm % or less, or 0.30 atm % or less. In another embodiment, the total amount of segregation of the amount of segregation of Mo and the amount of segregation of at least one of W, Ta, Re, Os, Ir and Tc at the former austenite grain boundaries may be 0.10 atm % or more, 0.13 atm % or more, 0.15 atm % or more, 0.18 atm % or more, or 0.20 atm % or more and/or may be 3.00 atm % or less, 2.00 atm % or less, 1.50 atm % or less, 1.00 atm % or less, 0.80 atm % or less, 0.60 atm % or less, 0.40 atm % or less, or 0.30 atm % or less.
The total amount of segregation of the at least one of Mo, W, Ta, Re, Os, Ir, and Tc at the former austenite grain boundaries is determined as follows: First, a test piece is taken from a position 50 mm or more away from the end faces of the hot stamped body. At that time, the front and back surfaces of the test piece are finished by machine polishing. Further, if there is a plating layer at the steel sheet surface, the plating layer is removed and then the front and back surfaces of the test piece of the steel sheet are finished by machine polishing. At that time, the sheet thickness is not particularly designated if the ¼ depth position of the sheet thickness can be measured, but the same amounts of the front and back surfaces of the test piece may also be removed by machine grinding so that the sheet thickness becomes 1.2 mm. The test piece is worked to a length of 20 mm and a width of 3.2 mm and formed with a V-notch of an angle of 450 at a position of a length of 11.5 mm. The test piece is dipped in a 20%-ammonium thiocyanate solution. At this time, the dipping time is not particularly limited. It is sufficient that the former austenite grain boundaries are exposed when set inside an Auger electron emission spectrometer and fracturing. For example, it may be 48 hours. The front and back surfaces of the test piece are galvanized within 10 minutes after ending the dipping. After plating, the test piece is quickly subjected to Auger electron emission spectrometry and fractured. At that time, the time after plating to fracture of the test piece is preferably within 1.5 hours, more preferably within 0.5 hour. The test piece is set within the Auger electron emission spectrometer and fractures from the notch portion of the test piece to expose the former austenite grain boundaries. At this time, the apparatus may be an Auger electron emission spectrometer. The model is not particularly limited, but a PHI680 made by ULVAC-PHI may be used. As the measurement conditions, the accelerating voltage may be 10 keV and the beam current may be 10 nA. An electron beam is fired at the exposed former austenite grain boundaries by a 1 to 30 kV accelerating voltage and the atm % of specific elements at the grain boundaries (specifically at least one of Mo, W, Ta, Re, Os, Ir, and Tc) are measured. The measurement is performed at the former austenite grain boundaries at 10 locations at a position of ¼ depth of the sheet thickness from the surface. To prevent contamination of the grain boundaries, quickly ending the measurement after fracture is preferable. The measurement should be ended within 30 minutes. The average value of the atm % of the obtained specific elements is calculated and determined as the total value of segregation of the at least one of Mo, W, Ta, Re, Os, Ir, and Tc.
The hot stamped body according to the present embodiment may having a plating layer at its surface. By having a plating layer at the surface, after the hot stamping, the corrosion resistance can be improved. As the plating layer, an aluminum plating layer, aluminum-zinc plating layer, aluminum-silicon plating layer, hot dip galvanized layer, electrogalvanized layer, hot dip galvannealed layer, zinc-nickel plating layer, aluminum-magnesium-zinc based plating layer, etc., may be illustrated.
According to the hot stamped body of an embodiment of the present invention, excellent mechanical properties, for example, a tensile strength of 2200 MPa or more, can be achieved. The tensile strength is preferably 2300 MPa or more, more preferably 2400 MPa or more, most preferably 2500 MPa or more. The upper limit is not particularly prescribed, but, for example, the tensile strength may be 3500 MPa or less, 3300 MPa or less, or 3000 MPa or less. The tensile strength of the hot stamped body is measured by preparing a No. 5 test piece and conducting a tensile test based on JIS Z 2241: 2011. At this time, for the purpose of removing roughness at the surface of the test piece, the surface layer parts of the front and back surfaces may be removed by machining or chemical polishing.
The hot stamped body according to an embodiment of the present invention, despite as explained above having, for example, a high tensile strength of 2200 MPa or more, can reliably suppress early fracture, and therefore is extremely useful for use as, for example, a frame member or bumper of an automobile or other structural member and reinforcing member where strength is required.
Next, a preferable method of production of the hot stamped body according to an embodiment of the present invention will be explained. The following explanation is intended to illustrate the characteristic method for producing the hot stamped body according to the embodiment of the present invention and is not intended to limit the hot stamped body to one produced by the method of production such as explained below.
In the method of production of the hot stamped body according to an embodiment of the present invention, in particular controlling the finish rolling conditions and preheating conditions is effective. Specifically, the method of production of the hot stamped body according to an embodiment of the present invention comprises:
First, a slab having the chemical composition explained above in relation to the hot stamped body is heated. The method of casting the molten steel is not particularly limited. The slab may be produced by continuous casting, ingot forming, or thin slab casting. The heating before the hot rolling is not particularly limited, but the slab used contains a relatively large amount of alloying elements for obtaining a high strength steel sheet. For this reason, the slab may also be heated before being sent on for hot rolling. For the purpose of making the alloying elements dissolve in the slab, the heating temperature may be 1100° C. or more.
In the present method, for example, the heated slab may be adjusted in thickness, etc., by rough rolling before finish rolling. The rough rolling need only secure the desired sheet bar dimensions. The conditions are not particularly limited.
The heated slab or the slab additionally rough rolled as needed is next finish rolled. In the present method, it is important that the rolling reduction of the final stage in the finish rolling be 40% or more. By making the rolling reduction of the final stage in the finish rolling 40% or more, the pearlite is evenly dispersed in the hot rolled steel sheet after rolling. This pearlite becomes starting points of austenite at the time of heating in the preheating step explained later in detail. For this reason, if the pearlite is evenly dispersed, at the hot stamped body, the standard deviation at the grain size distribution of the former austenite grains becomes smaller. As a result, the early fracture resistance of the hot stamped body can be improved. More preferably, the rolling reduction of the final stage in the finish rolling is 43% or more or 50% or more.
In a hot stamped body, for the purpose of securing a high hardenability, the amount of Mn added tends to be raised. For example, 0.50% or more of Mn is added. In relation to this, in research by the inventors, it was learned that with such a high Mn content, pearlite tends to be arranged relatively connected in a hot rolled steel sheet, and therefore compared with the case of low Mn content, it is extremely difficult to evenly disperse pearlite in the microstructure of a hot rolled steel sheet. Therefore, if finish rolling a steel material with such a high Mn content by a relatively low rolling reduction of less than 40%, it is believed that the presence of a part of connected pearlite at the microstructure becomes particularly prominent.
However, by making the rolling reduction of the final stage of finish rolling 40% or more, despite a high Mn content of 0.50% or more, it is possible to arrange pearlite sufficiently dispersed at the microstructure of the hot rolled steel sheet after the hot rolling step and coiling step. Therefore, in the microstructure of the hot rolled steel sheet rolled in this way, either there are no parts where pearlite is present connected or they are sufficiently reduced, therefore in the structure after the preheating step and hot stamping step, it is possible to reduce variation in the former austenite grain size. The upper limit of the rolling reduction in the final stage of the finish rolling is not particularly prescribed. Even in a steel material having such a high Mn content, in particular by suitable controlling the rolling reduction in the final stage of the finish rolling, it is possible to arrange the pearlite sufficiently dispersed in the microstructure of the hot rolled steel sheet and in turn possible to reduce the variation in the former austenite grain size and suppress a local rise in hardness.
The dominating factors in the morphology of this microstructure are the rolling reduction in the final stage of the finish rolling and the heating in the preheating step. The morphology is not particularly greatly affected by for example the heating in the hot stamping step after the preheating step, the optional cold rolling before the preheating step, the subsequent annealing, etc. This is because by forming the hot rolled steel sheet by a rolling reduction in the final stage of the finish rolling of 40% or more, even if the hot rolled steel sheet is cold rolled and then annealed at a relatively high temperature, there is a high tendency for formation of a microstructure where the carbides, grain boundaries, and retained austenite forming starting points of austenite after cooling are arranged dispersed. In general, if the rolling reduction in the final stage of the finish rolling is too high, fracture of the steel sheet at the time of rolling becomes a concern. Further, the preheating step explained in detail later is performed at an extremely high temperature compared with the hot stamping step, more specifically is performed at a temperature of more than 1200° C., and therefore generally this becomes a factor behind increased costs. Therefore, in particular the technical idea that by making the rolling reduction in the final stage of the finish rolling 40% or more and further by combining more than 1200° C. preheating before the stamping step, the variation in former austenite grain size is reduced and a local rise in hardness is suppressed has not existed up to now and was first discovered this time by the inventors.
Next, the finish rolled hot rolled steel sheet is coiled at a temperature of 750° C. or less. By making the coiling temperature 750° C. or less, it is possible to keep the ferrite from being arranged connected at the hot rolled steel sheet after rolling and the pearlite evenly disperses. This pearlite acts as starting points for austenite at the time of heating in the preheating step. For this reason, if the pearlite evenly disperses, in the hot stamped body, the standard deviation at the grain size distribution of the former austenite grains becomes smaller. As a result, the early fracture resistance of the hot stamped body can be improved.
Further, for the purpose of softening the hot rolled steel sheet, the coil after coiling may be heat treated to soften it. The method of heat treatment for softening is not particularly limited and may be made general conditions.
After the coiling step and before the cold rolling step, optionally, pickling may be performed for removing the oxide scale formed on the surface of the hot rolled steel sheet. The pickling may be formed under conditions suitable for removing oxide scale. It may be performed at one time or may be performed divided into several times so as to reliably remove the oxide scale.
After the coiling step, the steel sheet may be optionally cold rolled. The cold rolling is not particularly limited and may be performed under any suitable conditions. For example, the rolling reduction of the cold rolling may be 30 to 80%. The number of rolling passes and the rolling reduction per pass are not particularly limited and may be suitable set so that the rolling reduction of the cold rolling as a whole becomes the above range.
For example, after the cold rolling step, annealing may optionally be performed to adjust the microstructure and/or properties. The heating temperature of the annealing step is not particularly limited, but may for example be 800° C. or less.
For the purpose of improving the corrosion resistance, etc., the surface of the hot rolled steel sheet or cold rolled steel sheet may also be plated. The plating may be hot dip coating, alloyed hot dip coating, electroplating, or other treatment. For example, as the plating, the steel sheet may be hot dip galvanized. After hot dip galvanization, alloying treatment may be performed. As the plating layer, an aluminum plating layer, aluminum-zinc plating layer, aluminum-silicon plating layer, hot dip galvanized layer, electrogalvanized layer, hot dip galvannealed layer, zinc-nickel plating layer, aluminum-magnesium-zinc based plating layer, etc., may be illustrated. The specific conditions of the plating and alloying treatment are not particularly limited and may be any suitable conditions known to persons skilled in the art.
To correct the shape of the steel sheet or adjust the surface roughness, etc., it is possible, for example, to temper roll the steel sheet after the annealing step, or after the plating step.
In the present method, the obtained hot rolled steel sheet or cold rolled steel sheet is preheated to a temperature of more than 1200° C. before the hot stamping step, then is cooled by an average cooling speed of 10° C./s or more down to less than 350° C. In the hot stamped body according to an embodiment of the present invention, it is extremely important to make specific grain boundary strengthening elements, more specifically at least one type of Mo, W, Ta, Re, Os, Ir, and Tc, segregate at the former austenite grain boundaries in predetermined amounts. However, in the hot rolled steel sheet after the hot rolling step or in the cold rolled steel sheet after the optional cold rolling step or annealing step, these grain boundary strengthening elements are present as carbides and/or intermetallic compounds. Therefore, even if subjecting such steel sheet to the hot stamping step for usual heating and shaping without the preheating step, these grain boundary strengthening elements cannot be made to sufficiently segregate at the former austenite grain boundaries. In this case, it is no longer possible to sufficiently manifest the grain boundary strengthening action based on the grain boundary segregation of these elements. For this reason, in this method, it is extremely important to preheat the steel sheet before the hot stamping step to a relatively high temperature of more than 1200° C. to thereby make the carbides and/or intermetallic compounds of the grain boundary strengthening elements sufficiently melt and make the grain boundary strengthening elements dissolve in the steel sheet. The upper limit of the heating temperature of the preheating is not particularly prescribed, but the heating temperature may for example be 1400° C. or less. Further, after heating, the steel sheet is cooled by an average cooling speed of 10° C./s or more down to less than 350° C. By cooling by an average cooling speed of 10° C./s or more down to less than 350° C., it is possible to keep the grain boundary strengthening elements dissolved in the steel sheet from precipitating as compounds. The upper limit of the average cooling speed is not particularly prescribed, but for example the average cooling speed may be 3000° C./s or less, 1500° C./s or less, or 1200° C./s or less. The upper limit of the cooling speed is not particularly prescribed. The cooling method is also not particularly limited and may be die cooling, water cooling, oil cooling, or gas cooling. In particular, even with an extremely high average cooling speed, cooling can be relatively easily realized by utilizing die cooling or water cooled die cooling.
Finally, the preheated steel sheet is hot stamped in the hot stamping step to produce a hot stamped body having the desired chemical composition and microstructure. In particular, the grain boundary strengthening elements dissolved in the steel sheet in the previous preheating step disperse to the austenite grain boundaries and segregate there at the time of heating in the hot stamping step. For this reason, due to the following shaping and cooling operation, it is possible to achieve the desired total amount of segregation of the grain boundary strengthening elements at the former austenite grain boundaries after the martensite transformation. The dispersion and segregation of the grain boundary strengthening elements can be achieved by the usual heating, shaping, and cooling operations in the hot stamping step. Therefore, from the viewpoint of the dispersion and segregation of the grain boundary strengthening elements, these operations may be performed under suitable conditions known to persons skilled in the art. However, in particular from the viewpoint of obtaining the desired area ratio of the hard structures and former austenite grain size distribution, the steel sheet for hot stamping use is preferably heated to a temperature region of 800° C. to 1000° C. and held at that temperature region for 60 to 600 seconds. If the heating temperature is less than 800° C., sometimes the austenization becomes insufficient and the area ratio of the desired hard structures (at least one of martensite, bainite, and tempered martensite) and former austenite grain size distribution cannot be obtained and the tensile strength and early fracture resistance deteriorate. On the other hand, if the heating temperature is more than 1000° C., sometimes the austenite grains excessively grow, the desired former austenite grain size distribution cannot be obtained, and the early fracture resistance deteriorates. If the holding time is less than 60 seconds, sometimes the austenization becomes insufficient, the desired former austenite grain size distribution cannot be obtained, and the early fracture resistance deteriorates. If the holding time is more than 600 seconds, sometimes the austenite grains excessively grow, the desired former austenite grain size distribution cannot be obtained, and the early fracture resistance deteriorates.
The heating atmosphere is not particularly limited. Usual conditions are enough. For example, it may be an air atmosphere, a gas combustion atmosphere controlled in ratio of air and fuel, and a nitrogen atmosphere. The dew points may also be controlled in these gases.
The steel sheet is held at a temperature region of 800° C. to 1000° C., then hot stamped. After hot stamping, it may be cooled down to a temperature region of 250° C. or less by an average cooling speed of 20° C./s or more.
As the heating method, for example, furnace heating by an electric furnace, gas furnace, etc., flame heating, ohmic heating, high frequency heating, induction heating, etc., may be mentioned.
The hot stamped body according to the present embodiment is obtained by the above method. After hot stamping, it may be tempered at 130 to 600° C. or coated, then bake hardened (BH). Further, part of the hot stamped body may be tempered by being irradiated by a laser, etc., to partially form softened regions.
Below, examples will be used to explain the present invention in more detail, but the present invention is not limited to these examples in any way.
In the following examples, hot stamped bodies according to an embodiment of the present invention were produced under various conditions and the obtained tensile strengths and early fracture resistances of the hot stamped bodies were investigated.
First, molten steels having the chemical compositions shown in Table 1 were cast by continuous casting to produce slabs. The balances besides the constituents shown in Table 1 were Fe and impurities. These slabs were heated to a 1100° C. or more temperature and rough rolled under predetermined conditions, then were finish rolled and coiled under the conditions shown in Table 2. After the coiling, some of the hot rolled steel sheets were subjected to predetermined heat treatment for softening. Next, the obtained hot rolled steel sheets were pickled to remove the oxide scale cold rolled by 30 to 80% predetermined rolling reductions. (In Invention Example 337, cold rolling was not performed.) Next, some of the steel sheets were subjected to annealing, plating, or temper rolling under predetermined conditions. Next, the obtained steel sheets were preheated under the conditions shown in Table 2, then were cooled and finally were similarly hot stamped under the conditions shown in Table 2. The heating atmosphere and heating method in the hot stamping step, except when clearly indicated otherwise, were a gas combustion atmosphere (air-fuel ratio 0.85) and furnace heating. After the hot stamping, some of the hot stamped bodies were tempered or partially softened.
0.37
0.73
0.006
3.20
0.45
3.12
0.120
0.0151
0.0211
0.0237
0.0007
0.511
0.0006
0.126
0.007
0.211
0.008
1.12
0.0009
1.236
0.0221
16
769
1160
752
1021
631
The properties of the obtained hot stamped bodies were measured and evaluated by the following methods:
The tensile strength (TS) of a hot stamped body was obtained from any position of the hot stamped body by preparing a No. 5 test piece, removing the surface layer parts of the front and back surfaces of the test piece by machine grinding, and conducting a tensile test based on JIS Z 2241: 2011. The crosshead speed was 1 mm/min. If early fracture occurred at the time of the tensile test, i.e., if fracture occurred before reaching the maximum stress, the tensile strength of the hot stamped body was made the stress at the time of fracture.
The early fracture resistance was evaluated by the value of the tensile strength of the hot stamped body obtained by the above method divided by the value of the Vickers hardness obtained by the following method times 3.3 (tensile strength/(Vickers hardness×3.3)). If this value was 0.80 or more, the hot stamped body was deemed excellent in early fracture resistance and judged as passing while if it was less than 0.80, it was judged as failing. The “value of the Vickers hardness times 3.3” is the tensile strength estimated from the hardness. If the measured value of the tensile strength is 0.80 time or more of the estimated tensile strength, the hot stamped body can be judged to be excellent in early fracture resistance.
The Vickers hardness used for evaluation of early fracture resistance was obtained by the following method. First, a sample was cut out from any position 50 mm or more from the end faces of the hot stamped body so as to enable a cross-section vertical to the surface (sheet thickness cross-section) to be examined. The size of the sample, while depending also on the measuring device, was made a size enabling 10 mm or so to be examined in a direction vertical to the sheet thickness direction. The cross-section of the sample was polished using #600 to #1500 silicon carbide paper, then a liquid comprised of particle size 1 to 6 μm diamond powder dispersed in alcohol or other diluent or pure water was used to polish the surface to a mirror finish. The mirror finished cross-section was measured for hardness at a ¼ depth position of the sheet thickness from the surface in a direction parallel to the sheet surface by a load of 1 kgf at intervals of 3 times or more the indents using a microVickers hardness tester. A total of 20 points were measured and the average value was calculated to thereby obtain the Vickers hardness.
A case where the tensile strength was 2200 MPa or more and the above numerical value relating to the early fracture resistance was 0.80 or more was evaluated as a hot stamped body which is high in strength and able to suppress early fracture. The results are shown in Table 3. The “area ratio of hard structures” in Table 3 means the total of the area ratios of the martensite, bainite, and tempered martensite. Further, the balance of the structures other than the hard structures was comprised of ferrite, retained austenite, and/or pearlite. While not shown in Table 3, when measuring the standard deviation in the grain size distribution of the former austenite grains, the grain size of the former austenite grains of the hot stamped bodies in the invention examples in Table 3 was 10 tm or less in all cases.
2173
0.48
2170
74
2172
6.1
0.73
5.8
0.51
0.42
0.51
0.66
0.50
2093
0.54
2005
0.51
2008
0.52
0.05
0.50
0.03
0.73
2098
0.66
5.7
0.50
6.4
0.65
0.04
0.74
0.08
0.65
77
5.2
2078
0.42
5.6
0.41
6.2
0.59
5.6
0.48
Referring to Table 3, in Comparative Example 1, the C content was low, therefore the tensile strength fell. In Comparative Example 14, the C content was high, therefore the early fracture resistance fell. In Comparative Example 15, the Si content was low, therefore the tensile strength fell. In Comparative Example 29, the Si content was high, therefore the amount of ferrite increased, the desired metallographic structure was not obtained, and as a result the tensile strength similarly fell. In Comparative Example 30, the Mn content was low, therefore the standard deviation at the grain size distribution of the former austenite grains became greater and the early fracture resistance fell. In Comparative Example 43, the Mn content was high, therefore it is believed that in the hot rolled steel sheet, transformation from austenite to pearlite was promoted too much. As a result, in the hot stamped body, the standard deviation in the grain size distribution of the former austenite grains could not be made within the desired range and the early fracture resistance fell. In Comparative Examples 52, 61, 70, 78, 79, and 92, the respective P, S, N, O, or Al contents were not suitable, therefore the early fracture resistances fell. In Comparative Examples 93, 107, and 119, the respective Nb, Ti, and Cr contents were low, therefore the strengths could not be sufficiently improved by precipitation strengthening or solid solution strengthening and the tensile strengths fell. In Comparative Examples 106, 118, and 132, the respective Nb, Ti, and Cr contents were high, therefore it is believed large amounts of carbonitrides were formed or coarse intermetallic compounds were formed and as a result the early fracture resistances fell. In Comparative Example 133, the Mo content was low, therefore the total amount of segregation of the grain boundary strengthening elements at the former austenite grain boundaries become lower and the early fracture resistance fell. In Comparative Examples 145 and 156, the respective Mo and B contents were high, therefore it is believed that coarse intermetallic compounds were formed at the hot stamped bodies and as a result the early fracture resistance fell. In Comparative Example 146, the B content was low, therefore the tensile strength fell.
In Comparative Example 325, the rolling reduction of the final stage in the finish rolling of the hot rolling step was low, therefore it is believed the pearlite could not be evenly dispersed at the hot rolled steel sheet after rolling. As a result, in the hot stamped body, the standard deviation at the grain size distribution of the former austenite grains became greater and the early fracture resistance fell. In Comparative Example 333, the coiling temperature was high, therefore it is believed the ferrite was arranged connected and the pearlite could not be evenly dispersed. As a result, in the hot stamped body, the standard deviation at the grain size distribution of the former austenite grains became larger and the early fracture resistance fell. In Comparative Example 348, the heating temperature at the preheating step was low, therefore it is believed the grain boundary strengthening elements could not sufficiently dissolve in the steel sheet. As a result, the total amount of segregation of the grain boundary strengthening elements at the former austenite grain boundaries became low and the early fracture resistance fell. In Comparative Example 354, the average cooling speed at the preheating step was slow, therefore it is believed the grain boundary strengthening elements dissolved in the steel sheet due to the preheating precipitated as compounds. As a result, the total amount of segregation of the grain boundary strengthening elements at the former austenite grain boundaries became low and the early fracture resistance fell. In Comparative Example 361, the heating temperature at the hot stamping step was low, therefore the austenization became insufficient, the area ratio of the hard structures and the standard deviation at the grain size distribution of the former austenite grains could not be controlled to within the desired ranges, and the tensile strength and early fracture resistance fell. In Comparative Example 375, the heating temperature at the hot stamping step was high, therefore austenite grains excessively grew, the standard deviation at the grain size distribution of the former austenite grains could not be controlled to within the desired range, and the early fracture resistance fell. In Comparative Example 376, the holding time at the hot stamping step was long, therefore austenization became insufficient, the standard deviation at the grain size distribution of the former austenite grains could not be controlled to within the desired range, and the early fracture resistance fell. In Comparative Example 392, the holding time at the hot stamping step was long, therefore austenite grains excessively grew, the standard deviation at the grain size distribution of the former austenite grains could not be controlled to within the desired range, and the early fracture resistance fell.
In contrast to this, the hot stamped bodies according to all of the invention examples have the predetermined chemical compositions and metallographic structures, have standard deviations at the grain size distributions of the former austenite grains controlled to 5.0 μm or less, and have total amounts of segregation of the grain boundary strengthening elements at the former austenite grain boundaries, i.e., at least one of Mo, W, Ta, Re, Os, Ir, and Tc, controlled to 0.10 atm % or more, whereby early fracture can be reliably suppressed regardless of having high tensile strengths of 2200 MPa or more.
| Number | Date | Country | Kind |
|---|---|---|---|
| 2022-060624 | Mar 2022 | JP | national |
| Filing Document | Filing Date | Country | Kind |
|---|---|---|---|
| PCT/JP2023/007869 | 3/2/2023 | WO |
