Patent Application
20250075283
The present application discloses a steel sheet and a method of production of the same.
In recent years, to realize improved fuel efficiency of automobiles, the weight of automobile bodies has been increasingly lightened by use of high strength steel sheet. Further, to secure the safety of passengers as well, much use has been made of high strength steel sheet for automobile bodies in place of mild steel sheet. In the future, to further lighten the weight of automobile bodies, it will be necessary to increase the level of strength of high strength steel sheet to higher than that of the past.
Further, auto parts are being asked to function to suppress deformation at the time of collision of automobiles. To increase the resistance of auto parts to deformation at the time of collision of automobiles, the bending strength of auto parts is preferably increased. In addition, to increase the bending strength by rectification of the shape of parts from the viewpoint of structure, a high shapeability is sought from steel sheet. For this reason, steel sheet applied to auto parts should be high strength, be provided with excellent bending strength, and further exhibit high elongation. However, in the prior art, while the workability of high strength steel sheet etc. have been studied (for example, the following PTLs 1 to 3), securing bending strength against bending deformation from both the front and back surfaces has not been sufficiently studied.
PTL 1 discloses high strength steel sheet excellent in workability comprising steel sheet having ferrite as its main phase, containing retained austenite in an average of 5 vol % or more, and having a difference ΔVγ of 3.0 vol % or less between the maximum and minimum of content of retained austenite at different positions in the sheet thickness direction between 0.1 mm from the front surface of the steel sheet and 0.1 mm from the back surface of the steel sheet.
PTL 2 discloses hull-use steel plate excellent in shock absorbing ability able to minimize fracture of a hull at the time of collision of tankers, which comprises plate thickness 8 mm or more steel plate containing C, Si, Mn, and Al, further containing, as needed, strengthening elements, and having a balance of Fe and impurities, which contains, by area ratio, 1.0 to 20% of retained γ at least at ⅛ or more of the plate thickness at the front and back layers of the steel plate.
PTL 3 discloses structural use thick gauge steel plate able to simultaneously strikingly improve the brittle crack arresting property and Charpy property without relying on addition of Ni and other expensive alloy elements, which comprises, by wt %, C: 0.04 to 0.30%, Si: 50.5%, Mn: ≤2.0%, Al: ≤0.1%, Ti: 0.001 to 0.10%, N: 0.001 to 0.01%, and a balance of Fe and unavoidable impurities, has an average crystal grain size “d” of the microstructure at predetermined regions of the front and back layer parts of plate thickness of 3 μm or less, and has a Vickers hardness of the microstructure satisfying a predetermined requirement.
The present application, in consideration of the above situation, discloses a steel sheet excellent in mechanical properties such as strength and elongation and also excellent in bending strength, and a method of production of the same.
The inventors intensively researched means for solving the above problem and clarified that by optimizing the ratios of the steel sheet structures such as retained austenite and by decreasing the difference in number densities of precipitates at the front and back surfaces, steel sheet excellent in strength, elongation, and other mechanical properties and exhibiting a high bending strength is obtained. Along with this, they confirmed that in steel sheet with a difference in number densities of precipitates at the front and back surfaces of more than 10%, the bending strength changes depending on the direction of the bending and the deformation resistance of parts at the time of collision occasionally declines.
Further, the inventors discovered that by using an integrated production process characterized by performing annealing twice on cold rolled steel sheet and coiling or tempering the steel sheet for predetermined aging treatment between the two annealing steps, the microstructure is optimized and steel sheet with a small difference in number densities of precipitates at the front and back surfaces can be produced.
Further, the inventors discovered by various research conducted over time that steel sheet increased in bending strength by reducing the difference in number densities of precipitates at the front and back surfaces as explained above is difficult to produce just by modifying the hot rolling conditions, annealing conditions, etc. alone and can only be produced by optimization of the same in a so-called integrated process of hot rolling and annealing steps etc.
The gist of the present invention is as follows:
The steel sheet of the present disclosure is excellent in strength, elongation, and other mechanical properties and excellent in bending strength.
Below, embodiments of the present invention will be explained. Note that the explanation of these is intended to simply illustrate embodiments of the present invention. The present invention is not limited to the following embodiments.
The steel sheet according to the present embodiment is characterized by
First, the reasons for limiting the chemical composition according to embodiments of the present invention will be explained. Here, the “%” relating to the constituents means mass %. Furthermore, in this Description, the “to” showing numerical ranges, unless otherwise indicated, is used in the sense including the numerical values described before and after it as a lower limit value and upper limit value.
C is an element inexpensively making the tensile strength increase and is an element extremely important for controlling the strength of steel. If the C content is 0.10% or more, such an effect is easy to obtain. The C content may also be 0.12% or more. On the other hand, if excessively including C, the elongation falls, brittle fracture of steel is invited, and sometimes a drop in the bending strength at the time of deformation of a part is promoted. If the C content is 0.30% or less, such a problem is easily avoided. The C content may also be 0.28% or less.
Si is an element which acts as a deoxidizer, increases the stability with respect to working of retained austenite structures, and suppresses the precipitation of carbides in martensite structures at the time of aging. If the Si content is 0.60% or more, such an effect is easy to obtain. The Si content may also be 0.70% or more. On the other hand, if excessively including Si, in the aging treatment, formation of E carbides may be suppressed, and the bending strength may be dropped. If the Si content is 1.20% or less, such a problem is easily avoided. The Si content may also be 1.00% or less.
Mn is a factor having an effect on the ferrite transformation of steel and an element which suppresses ferrite transformation in the cooling process of the later explained Q-annealing, increases the ratio of martensite structures after Q-annealing, and is effective for raising strength. If the Mn content is 1.00% or more, such an effect is easy to obtain. The Mn content may also be 1.30% or more. On the other hand, if excessively including Mn, a concentrated Mn layer formed due to microsegregation and center segregation remarkably appears in the steel sheet and, due to the difference in solidification speeds at the front and back surfaces of the slab, differences are formed in the state of distribution of the concentrated Mn layer at the front and back surfaces of the steel sheet, so differences may be invited in the bending strength of the front and back surfaces due to differences in formation of segregated Mn bands. If the Mn content is 3.50% or less, such a problem is easily avoided. The Mn content may also be 3.00% or less.
P is an element strongly segregating at the ferrite grain boundaries and promoting embrittlement of the grain boundaries. The less the better. Further, if excessively including P, brittle fracture of steel is invited and sometimes a drop in the bending strength at the time of deformation of a part is promoted. On this point, the P content is 0.0200% or less. The P content may also be 0.0180% or less. On the other hand, the lower limit of the P content is not particularly prescribed. The P content is 0% or more, may be 0.0001% or more, and may be 0.0010% or more.
S is an element which forms MnS and other nonmetallic inclusions in steel and invites a drop in ductility of steel parts. The less the better. Further, if excessively including S, formation of voids starting from nonmetallic inclusions at the time of deformation of a part is invited and sometimes the bending strength is made to fall. On this point, the S content is 0.0200% or less. The S content may also be 0.0180% or less. On the other hand, the lower limit of the S content is not particularly prescribed. The S content is 0% or more, may be 0.0001% or more, or may be 0.0005% or more.
Al is an element acting as a deoxidizer of steel and stabilizes ferrite and is added according to need. If the Al content is 0.001% or more, such an effect is easy to obtain. The Al content may also be 0.010% or more. On the other hand, if excessively including Al, ferrite transformation and bainite transformation in the cooling process at annealing may be excessively promoted and the strength of the steel sheet may fall. If the Al content is 1.000% or less, such a problem is easily avoided. The Al content may also be 0.800% or less.
N is an element which forms coarse nitrides in steel sheet and lowers the workability of steel sheet. Further, N is an element becoming a cause of formation of blowholes at the time of welding. Further, if excessively including N, it bonds with Al or Ti to form a large amount of AlN or TiN. These nitrides become starting points for formation of voids at the time of deformation of parts and sometimes invite a drop in the bending strength. On this point, the N content is 0.0200% or less. The N content may also be 0.0160% or less. On the other hand, the lower limit of the N content is not particularly prescribed. The N content is 0% or more, may also be 0.0001% or more, or may also be 0.0010% or more.
The basic chemical composition of the steel sheet in the present embodiment is as explained above. Furthermore, the steel sheet in the present embodiment may if necessary include at least one of the following optional elements. These elements need not be included, so the lower limits are 0%.
Ti is a strengthening element. It contributes to raising strengthen of steel sheet through precipitation strengthening, fine grain strengthening through suppression of growth of crystal grains, and dislocation strengthening through suppression of recrystallization. On the other hand, if excessively including Ti, coarse carbides precipitate more, these carbides become starting points for formation of voids at the time of deformation of parts, and sometimes a drop in bending strength is invited. The Ti content is 0% or more, may be 0.0010% or more, or may be 0.005% or more. Further, it is 0.500% or less and may be 0.400% or less.
Co is an element effective for control of the form of carbides and increase of strength and is added as needed for control of the strength. On the other hand, if excessively including Co, fine Co carbides precipitate in large numbers. These carbides become starting points for formation of voids at the time of deformation of parts and sometimes invite a drop in the bending strength. The Co content is 0% or more and may also be 0.001% or more. Further, it is 0.500% or less and may be 0.400% or less.
Ni is a strengthening element and is effective for improving quenchability. In addition, it may be added since it improves the wettability of the steel sheet and plating and promotes an alloying reaction. On the other hand, if excessively including Ni, the peelability of oxide scale at the time of hot rolling may be affected and formation of defects at the surfaces of the steel sheet may be promoted, so the yield strength at the time of bending deformation may fall. The Ni content is 0% or more and may be 0.001% or more. Further, it is 0.500% or less and may be 0.400% or less.
Mo is an element effective for raising the strength of steel sheet. Further, Mo is an element having the effect of suppressing ferrite transformation occurring at the time of heat treatment at continuous annealing facilities or continuous hot dip galvanization facilities. On the other hand, if excessively including Mo, fine Mo carbides may precipitate in large numbers. These carbides may become starting points for formation of voids at the time of deformation of parts and may invite a drop in the bending strength. The Mo content is 0% or more and may be 0.001% or more. Further, it is 0.500% or less and may be 0.400% or less.
Cr, like Mn, is an element which suppresses pearlite transformation and is effective for making steel high strength. It is added according to need. On the other hand, if excessively including Cr, formation of retained austenite may be promoted and a drop in bending strength may be invited due to the presence of excessive retained austenite. The Cr content is 0% or more and may be 0.001% or more. Further, it is 2.000% or less and may be 1.500% or less.
O forms oxides and causes the workability to deteriorate, so it is necessary to suppress the amount of addition. In particular, oxides in many cases are present as inclusions and, if present at punched edges or cut surfaces, notches or coarse dimples are formed at the end faces, so at the time of bending deformation, stress concentration may be invited and may become starting points of crack formation and sometimes a drop in bending strength is invited. The O content is 0.0100% or less and may also be 0.0080% or less. Further, the O content is 0% or more, but controlling the O content to less than 0.00010% increases the refining time and is liable to invite an increase in the production costs. Due to the aim of preventing a rise in the production costs, the O content may also be 0.0001% or more and may also be 0.0010% or more.
B is an element suppressing the formation of ferrite and pearlite in the cooling process from austenite and promoting formation of bainite, martensite, and other low temperature transformed structures. Further, B is an element advantageous for making the steel high strength and is added in accordance with need. On the other hand, if excessively including B, formation of coarse B inclusions in the steel may be invited and these inclusions may become starting points for formation of voids, so a drop in bending strength at the time of deformation of a part may be invited. The B content is 0% or more, may be 0.0001% or more, and may be 0.0010% or more. Further, it is 0.0100% or less and may be 0.0080% or less.
Nb is an element effective for control of the form of carbides and is an element effective for raising the toughness since it refines the microstructure by its addition. On the other hand, if excessively including Nb, fine, hard Nb carbides may precipitate in large numbers. These carbides may become starting points for formation of voids, so may invite a decline in the bending strength at the time of deformation of a part. The Nb content is 0% or more and may be 0.001% or more. Further, it is 0.500% or less and may be 0.400% or less.
V is a strengthening element. It contributes to a rise of strength of the steel sheet by precipitation strengthening, fine grain strengthening through suppression of growth of ferrite crystal grains, and dislocation strengthening through suppression of recrystallization. On the other hand, if excessively including V, carbonitrides may precipitate in large amounts. These carbonitrides may become starting points of formation of voids, so a drop in bending strength at the time of deformation of a part may be invited. The V content is 0% or more and may be 0.001% or more. Further, it is 0.500% or less and may be 0.400% or less.
Cu is an element effective for raising the strength of steel sheet. On the other hand, if excessively including Cu, the steel material may become brittle during hot rolling and hot rolling may become difficult. Further, along with the rise in strength of the steel, the ductility may fall so a drop in bending strength at the time of deformation of a part may be invited. The Cu content is 0% or more and may be 0.001% or more. Further, it is 0.500% or less and may be 0.400% or less.
W is effective for raising the strength of steel sheet. Further the precipitates and crystallized substances containing W become hydrogen trap sites. On the other hand, if excessively including W, voids may be easily formed and proceeded starting from the coarse carbides, so a drop of bending strength at the time of deformation of a part may be invited. The W content is 0% or more, may be 0.0001% or more, and may be 0.0010% or more. Further, it is 0.1000% or less and may be 0.0800% or less.
Ta, like Nb, V, and W, is an element effective for controlling the form of carbides and increasing the strength and is added in accordance with need. On the other hand, if excessively including Ta, fine Ta carbides may precipitate in large numbers. Formation of voids may become easy starting from these carbides, so a drop in the bending strength at the time of deformation of a part may be invited. The Ta content is 0% or more, may be 0.00010% or more, and may be 0.0010% or more. Further, it is 0.1000% or less and may be 0.0800% or less.
Sn is an element contained in the steel when using scrap as the raw material. The less the better. If excessively including Sn, a drop in bending strength at the time of deformation of a part due to embrittlement of the steel sheet may be invited. The Sn content is 0.0500% or less and may be 0.0400% or less. Further, the Sn content may be 0%, but controlling the Sn content to less than 0.0001% increases the refining time and is liable to invite an increase in the production costs. Due to the aim of preventing a rise in the production costs, the Sn content may also be 0.0001% or more and may also be 0.0010% or more.
Sb, like Sn, is an element contained in steel when using scrap as the raw material. Sb strongly segregates at the grain boundaries and invites embrittlement of the grain boundaries and a drop in ductility, so the less the better. If excessively including Sb, a drop in bending strength at the time of deformation of a part due to embrittlement of the steel sheet may be invited. The Sb content is 0.0500% or less and may be 0.0400% or less. Further, the Sb content may be 0%, but controlling the Sb content to less than 0.00010% increases the refining time and is liable to invite an increase in the production costs. Due to the aim of preventing a rise in the production costs, the Sb content may also be 0.00010% or more and may also be 0.0010% or more.
As, like Sn and Sb, is an element contained in the steel when using scrap as the raw material. It strongly segregates at the grain boundaries. The less the better. If excessively including As, a drop in bending strength at the time of deformation of a part due to embrittlement of the steel sheet may be invited. The As content is 0.0500% or less and may be 0.0400% or less. Further, the As content may be 0%, but controlling the As content to less than 0.0001% increases the refining time and is liable to invite an increase in the production costs. Due to the aim of preventing a rise in the production costs, the As content may also be 0.00010% or more and may also be 0.0010% or more.
Mg is an element able to control the form of sulfides by trace addition and is added according to need. On the other hand, if excessively including Mg, coarse inclusions may be formed. These inclusions may become starting points for formation of voids, so a drop in the bending strength at the time of deformation of a part may be invited. The Mg content is 0% or more, may be 0.00010% or more, and may be 0.0010% or more. Further, it is 0.0500% or less and may be 0.0400% or less.
Ca is useful as a deoxidizing element and also has an effect on the control of the form of sulfides. On the other hand, if excessively including Ca, a drop in bending strength at the time of deformation of a part due to embrittlement of the steel sheet may be invited The Ca content is 0% or more, may be 0.0001% or more, and may be 0.0010% or more. Further, it is 0.0500% or less and may be 0.0400% or less.
Y, like Mg and Ca, is an element enabling the control of sulfides by trace addition and is added according to need. On the other hand, if excessively including Y, coarse Y inclusions may be formed. These inclusions may become starting points for formation of voids, so a drop in the bending strength at the time of deformation of a part may be invited. The Y content is 0% or more, may be 0.00010% or more, and may be 0.0010% or more. Further, it is 0.0500% or less and may be 0.0400% or less.
Zr, like Mg, Ca, and Y, is an element enabling the control of sulfides by trace addition and is added in according to need. On the other hand, if excessively including Zr, coarse Zr inclusions may be formed. These inclusions may become starting points for formation of voids, so a drop in the bending strength at the time of deformation of a part may be invited. The Zr content is 0% or more, may be 0.0001% or more, and may be 0.0010% or more. Further, it is 0.0500% or less and may be 0.0400% or less.
La is an element enabling the control of sulfides by trace addition and is added according to need. On the other hand, if excessively including La, La inclusions may be formed. These inclusions may become starting points for formation of voids, so a drop in the bending strength at the time of deformation of a part may be invited. The La content is 0% or more, may be 0.0001% or more, and may be 0.0010% or more. Further, it is 0.0500% or less and may be 0.0400% or less.
Ce, like La, is an element enabling the control of sulfides by trace addition and is added according to need. On the other hand, if excessively including Ce, Ce inclusions may be formed. These inclusions may become starting points for formation of voids, so a drop in the bending strength at the time of deformation of a part may be invited. The Ce content is 0% or more, may be 0.0001% or more, and may be 0.0010% or more. Further, it is 0.0500% or less and may be 0.0400% or less.
In the steel sheet in the present embodiment, the balance of the constituents explained above is comprised of Fe and impurities. “Impurities” are constituents entering due to various factors in the production process, such as the ore, scraps, and other such raw materials, when industrially producing steel sheet or the present embodiment.
Next, the microstructure and properties of the steel sheet according to an embodiment of the present invention will be explained.
Ferrite, pearlite, and bainite are structures effective for improving the strength-ductility balance of steel sheet, but with inclusion of large amounts, a drop in the local ductility may be invited. Further, from the viewpoint of efficiently raising the strength of steel as well, the smaller the area ratios of ferrite, pearlite, and bainite, the better. The total of the area ratios of the ferrite, pearlite, and bainite may be 0% and may be 1.0% or more. Further, it may be 30.0% or less, may be 25.0% or less, and may be 20.0% or less. Further, the productivity falls somewhat, but by controlling the integrated production conditions to a high precision, the total of the area ratios of the ferrite, pearlite, and bainite can be made 0%.
Retained austenite are effective for improving the strength-ductility balance of steel sheet. If the area ratio of the retained austenite is too small, at the time bending deformation is applied to the steel sheet, the effect of raising the strength by work induced transformation from retained austenite to martensite cannot be obtained, so a drop in the bending strength may be invited. On the other hand, if the area ratio of the retained austenite is too large, a drop in the yield strength along with a drop in the bending strength may be invited. The area ratio of the retained austenite is 10.0% or more and may be 13.0% or more. Further, it is 30.0% or less and may be 25.0% or less.
In the steel sheet according to the present embodiment, the microstructure of the steel sheet preferably includes the retained austenite in an acicular form. By the retained austenite being “acicular” in form, the following effect can be expected. That is, even if the retained austenite is spherical (lumpy) in form, work induced transformation easily occurs along with deformation of steel sheet, so bending deformation may start due to the low stress. As opposed to this, if the retained austenite is acicular in form, work induced transformation becomes hard to occur and the bending strength becomes much higher. In the steel sheet according to the present embodiment, by combining the effect due to the acicular retained austenite and the effect due to the difference in number densities of precipitates, the bending strength of the steel sheet is remarkably improved. The area ratio of acicular retained austenite may be 30% or more or 50% or more and may be 95% or less or 90% or less in the case the area ratio of retained austenite as a whole is 100%. Further, in the present application, “acicular retained austenite” means a ratio of a long axis and short axis (long axis/short axis) of 3.0 or more. The “long axis” and “short axis” of the retained austenite can be identified by examination of the microstructure by EBSD. Specifically, in examination of the microstructure, one retained austenite crystal grain is specified and minimum ferret diameter of that crystal grain is identified as the short axis and the maximum ferret diameter is identified as the long axis.
Fresh martensite and tempered martensite are structures extremely effective for raising the strength of steel sheet. The higher the area ratios, the better. In the steel sheet according to the present embodiment, the balance other than the above ferrite, pearlite, bainite, and retained austenite is comprised of fresh martensite and tempered martensite. The total of the area ratios of the fresh martensite and tempered martensite may be 40.0% or more, may be 45.0% or more, may be 50.0% or more and, further, may be 90.0% or less and may be 85.0% or less. Further, the area ratio of the fresh martensite may be 5% or more, 10% or more, 20% or more, 30% or more, or 40% or more and may be 80% or less, 70% or less, 60% or less, 50% or less, or 40% or less. Further, the area ratio of the tempered martensite may be 5% or more, 10% or more, 20% or more, 30% or more, or 40% or more and may be 80% or less, 70% or less, 60% or less, 50% or less, or 40% or less.
The number densities of precipitates in the tempered martensite at the first surface of the front side and second surface of the back side of the steel sheet are an important factor in raising the resistance to bending deformation. The higher both the number densities of the first surface of the front side and second surface of the back side of the steel sheet, the higher the bending strength. If the number density of either one of the surfaces is high, uneven strain is caused at the time of bending deformation. At the other low number density surface, yield occurs and the bending strength falls. For this reason, the smaller the difference between the number density of precipitates in the tempered martensite at the first surface at the front side of the steel sheet and the number density of precipitates in the tempered martensite at the second surface at the back side of the steel sheet, the better. Specifically, it is important that the difference in number densities of precipitates be 10.0% or less. The difference of the number densities may be 8.0% or less, may be 6.0% or less, may be 4.0% or less, or may be 2.0% or less. In other words, in the present embodiment, the ratio A1/A2 of the number density of precipitates A1 at the first surface at the front side of the steel sheet and the number density of precipitates A2 at the second surface of the back side of the steel sheet is 0.90 or more and 1.10 or less and may also be 0.92 or more, 0.94 or more, 0.96 or more, or 0.98 or more and may be 1.08 or less, 1.06 or less, 1.04 or less, or 1.02 or less. Note that, making the difference of the number densities 0% becomes a factor behind a rise of the production load and increasing the production costs for fine control of the steel sheet microstructure. On this point, the difference of the number densities may also be 0.10% or more. Further, the precipitates are mainly comprised of carbides formed by tempering of martensite. The carbides may be cementite and may also be Fe-based carbides or alloy carbides obtained by Cr, Ti, V, and other alloy elements bonding with carbon in place of iron. Further, the specific value of the number density of precipitates in the tempered martensite at the first surface of the front side and second surface of the back side of the steel sheet may be, for example, 1/μm2 or more, 5/μm2 or more, or 10/μm2 or more and may be 300/μm2 or less, 100/μm2 or less, or 30/μm2 or less.
Further, in the present application, for convenience of the explanation, the “front” and “back” of the steel sheet are differentiated, but which side of the steel sheet is the front and which is the back are not particularly limited.
To lighten structures using steel as their materials and raising the yield strength at which plastic deformation starts, the higher the yield strength of the steel material, the better. On the other hand, if the yield strength becomes too high, the change of shape by elastic deformation after plastic working and effect of so-called springback may become greater. The yield strength of the steel sheet according to the present embodiment may be 600 MPa or more or more and may be 650 MPa or more. The upper limit of the yield strength is not particularly prescribed, but from the viewpoint of suppressing the effect of springback, it may be 1100 MPa or less and may be 1050 MPa or less.
To lighten structures using steel as their materials and raising the resistance of structures at plastic deformation, it is preferable that the steel material have a large work hardening ability and exhibit the maximum strength. On the other hand, if the tensile strength becomes too large, fracture may easily occur by low energy during plastic deformation and the shapeability may fall. The tensile strength of the steel sheet is not particularly limited, but may be 900 MPa or more and may be 980 MPa or more. Further, it may be 2000 MPa or less and may be 1800 MPa or less.
(Total Elongation t-El)
When cold forming a steel sheet material to produce a structural member, elongation is required for finishing it into a complicated shape. If the total elongation becomes too low, the material may crack in the cold forming process. On the other hand, the higher the total elongation the better, but if trying to excessively increase the total elongation, a large amount of retained austenite becomes required in the microstructure. Due to this, the yield strength at the time of bending deformation may fall. The total elongation of the steel sheet is not particularly limited, but may be 13% or more and may be 20% or more. Further, it may be 35% or less and may be 30% or less.
When cold forming a steel sheet material to produce a structural member, hole expandability is required along with elongation for finishing it into a complicated shape. If the hole expandability is too small, the material may crack at cold forming. On the other hand, the higher the hole expandability, the better, but if trying to excessively increase the hole expandability, a large amount of retained austenite becomes required in the microstructure and due to this, the yield strength at the time of bending deformation may fall. The hole expansion rate λ of steel sheet is not particularly limited, but may be 20% or more and may be 25% or more. Further, it may be 90% or less and may be 80% or less.
When cold forming a steel sheet material to produce a structural member, bendability is required for finishing it into a complicated shape. If the VDA bending angle is too small, the material may crack at cold forming. The higher the bendability, the better. The VDA bending angle of steel sheet is not particularly limited, but may be 450 or more and may be 500 or more.
The sheet thickness is a factor affecting the rigidity of a steel member after being formed. The greater the sheet thickness, the higher the rigidity of the member. If the sheet thickness is too small, a drop in the rigidity may be invited and, due to the effect of unavoidable nonferrous inclusions present inside the steel sheet, the press formability may fall. On the other hand, if the sheet thickness is too large, the press forming load may increase and wear of the dies and a drop in the productivity may be invited. The thickness of the steel sheet is not particularly limited, but may be 0.2 mm or more and may be 6.0 mm or less.
Next, the methods of examination and measurement of the microstructure prescribed above and the methods of measurement and evaluation of the properties prescribed above will be explained.
The microstructure is examined by a scan type electron microscope. Before examination, a sample for examination of the microstructure is polished by wet polishing by emery paper and polishing by diamond abrasives having a 1 μm average particle size to finish the examined surface to a mirror surface, then the microstructure is etched by a 3% nitric alcohol solution. The power of the examination is made 3000×. Ten 30 μm×40 μm fields at a thickness ¼ position from the surface side of the steel sheet are randomly captured. The ratios of the structures are found by the point count method. In the obtained structural images, a total of 100 lattice points arranged at vertical 3 μm and horizontal 4 μm intervals are determined, the structures present under the lattice points are judged, and the ratios of structures contained in the steel sheet are found from the average value of 10 images. Ferrite forms lump shaped crystal grains and does not contain long axis 100 nm or more Fe-based carbides inside. Bainite forms collections of lath shaped crystal grains and does not contain long axis 20 nm or more Fe-based carbides inside or contains long axis 20 nm or more Fe-based carbides inside wherein the carbides fall under single variants, that is, groups of Fe-based carbides stretched in one direction. Here, the “groups of Fe-based carbides stretched in one direction” mean groups of Fe-based carbides with differences in the stretched direction of within 5°. In bainite, bainite surrounded by grain boundaries of difference of direction of 15° or more is counted as one bainite grain. Pearlite forms structures containing cementite precipitated in lines and regions captured by bright contrast in the secondary electron image are deemed as “pearlite” whereby the area ratio is calculated.
The fresh martensite and tempered martensite are examined by scan type and transmission type electron microscopes. Structures containing Fe-based carbides inside (Fe-based carbides of 1/μm2 or more) are identified as tempered martensite, and structures containing almost no Fe-based carbides (Fe-based carbides of less than 1/μm2) are identified as fresh martensite. Fe-based carbides having various crystal structures have been reported. Any of these Fe-based carbides may be contained. Depending on the heat treatment conditions, sometimes there will be several types of Fe-based carbides. In the present application, the area ratio A1 of the total of ferrite, pearlite, and bainite is measured by the method explained above, the area ratio A2 of the retained austenite is measured by the method explained below, and the balance obtained by subtracting the total value of the area ratios A1 and A2 from 100% is deemed the area ratio of the total of the fresh martensite and tempered martensite.
The area ratio of the retained austenite is determined in the following way by X-ray analysis. First, the part from the surface of the steel sheet down to ¼ of the thickness of the steel sheet is removed by mechanical polishing and chemical polishing. The chemically polished surface is measured using MoKα rays as the characteristic X-rays. Further, the following formula is used to calculate the area ratio of retained austenite at the center part of sheet thickness from the ratio of integral strength of the diffraction peaks of the (200) and (211) planes of the body centered cubic (bcc) phase and (200), (220), and (311) planes of the face centered cubic (fcc) phase.
The sample used for X-ray diffraction may be prepared by reducing the thickness of the steel sheet down to a predetermined sheet thickness by mechanical polishing etc., then removing the strain by chemical polishing, electrolytic polishing, etc. and simultaneously preparing the sample for measurement in accordance with the above-mentioned methods so that a suitable surface becomes the measurement surface in the range of sheet thickness of ⅛ to ⅜. While natural, the above-mentioned X-ray strength is not limited to just near the sheet thickness ¼. By satisfaction of as much of the thickness as possible, the material anisotropy is reduced much more. However, by measurement within a range of ⅛ to ⅜ from the surface of the steel sheet, it is possible to represent the material properties of the steel sheet as a whole. Therefore, ⅛ to ⅜ of the sheet thickness is made the measurement range.
Further, the area ratio of acicular retained austenite in the retained austenite can be measured by, for example, EBSD.
The number densities of precipitates in the tempered martensite at the first surface of the front side and second surface of the back side of the steel sheet are measured as follows: First, from the front surface or back surface of the steel sheet (meaning front surface or back surface of the base material steel sheet. For example, in the case of a surface treated steel sheet having a plating or other surface treated layer, meaning the front surface or back surface of the base material steel sheet after removal of the surface treated layer), material for examination is taken at a depth position of ⅛ of the thickness in the sheet thickness direction. This is prepared into a test piece for examination of a thin film or extraction replica. The test piece is examined by a transmission type electron microscope at a 10,000× power to acquired captured images of a minimum of 30 fields, these examined images are measured for the number density of precipitates per unit area, and the value found by finding the arithmetic average of the number densities of 30 fields is made the number density of precipitates of the first surface of the front side or second surface of the back side. Note that, the field examined by a transmission type electron microscope at 10,000× power is a rectangular region of sides of 600 nm or so. The area of the 30 fields used for measurement of the number densities of precipitates becomes about 10.8 μm2.
(Method of Measurement of Yield Strength YP, Tensile Strength TS, and Total Elongation t-El)
A tensile test for measuring the yield strength, tensile strength, and total elongation is performed based on JIS Z 2241: 2011 by taking a JIS No. 5 test piece from a direction where the long direction of the test piece becomes parallel with a direction perpendicular to the rolling of the steel strip.
The hole expandability is evaluated by punching a diameter 10 mm circular hole under conditions of a clearance of 12.5%, arranging the sheet so that that the burr becomes the die side, inserting a 600 conical punch, and finding the hole expansion rate λ (%). Hole expansion tests are conducted five times and the average value used as the hole expansion rate.
The bendability was evaluated by performing a test based on the provisions of the standard 238-100 of the Verband der Automobilindustrie (VDA) 2 using a width 60 mm test piece and measuring the maximum bending angle α as the VDA bending angle. Further, the bending strength is evaluated by the value dividing the load at the bending angle 5° by the sheet thickness.
The method of production of steel sheet according to the present embodiment uses a material having the above-mentioned chemical composition for integrated management of hot rolling, cold rolling, and annealing. Specifically, the method of production of steel sheet according to the present embodiment includes the step of hot rolling and coiling a steel slab having a chemical composition the same as the chemical composition explained above relating to the steel sheet and pickling the obtained hot rolled steel sheet, cold rolling, annealing, then again annealing it. More specifically, the method of production of steel sheet according to the present embodiment comprises
In the present embodiment, a steel slab obtained by the continuous casting method and other known methods is hot rolled to obtain hot rolled steel sheet. Here, the finish rolling temperature of the hot rolling is a factor affecting control of the texture of old austenite grain size. From the viewpoint of the growth of the rolling texture of austenite and inviting anisotropy of the steel material properties, the finish rolling temperature is preferably 650° C. or more. Further, from the aim of suppressing unevenness in the texture due to abnormal grain growth of austenite, the finish rolling temperature is preferably 950° C. or less.
The temperature when coiling the hot rolled steel sheet (coiling temperature of hot rolled coil) controls the state of formation of oxide scale at the hot rolled steel sheet and is a factor affecting the strength of hot rolled steel sheet. The thinner the thickness of the scale formed on the hot rolled steel sheet surface, the better. From this, the lower the coiling temperature, the better. Further, if extremely lowering the coiling temperature, special facilities become required. Further, if the coiling temperature is too high, as explained above, the oxide scale formed on the surface of the hot rolled steel sheet becomes remarkably thick. From the above viewpoint, the temperature when coiling the hot rolled steel sheet may be 700° C. or less and may be 680° C. or less. Further, it may be 0° C. or more and may be 20° C. or more.
The pickling of the hot rolled steel sheet has as its object to remove scale etc. The pickling may be performed under known pickling conditions.
In cold rolling, if the total of the rolling reduction becomes too large, the ductility of the base material steel sheet is lost and the risk of fracture of the base material steel sheet at cold rolling becomes higher. On this point, the total of the rolling reduction in the cold rolling is preferably 85% or less. On the other hand, to make the recrystallization at the annealing step sufficiently proceed, the total of the rolling reduction is preferably made 20% or more, more preferably 30% or more. For the purpose of reducing the cold rolling load before cold rolling, the sheet may also be annealed at 700° C. or less.
In the first annealing (Q-annealing), the cold rolled steel sheet as a base material steel sheet is heated to the Ac3 point or more and 1000° C. or less (that is, the austenite single phase region and 1000° C. or less). The reason for making the peak temperature of heating the Ac3 point or more is to heat the base material steel sheet to the austenite single phase region and thereby obtain an area ratio of 90% or more of martensite structures by the subsequent quenching and promote the precipitation of c carbides by aging. Due to this, at the holding at a low temperature, a mainly martensite microstructure cannot be obtained and the bending strength remarkably falls. On the other hand, if heating to over 1000° C., the surface layer of the steel sheet is decarburized and the strength falls sometimes, so the bending strength falls.
In the first annealing (Q-annealing), the sheet is preferably held at the Ac3 point or more and 1000° C. or less of heating temperature for 5 seconds or more. If the holding time is too short, austenite transformation of the base material steel sheet does not sufficiently proceed and in addition Mn and other substitution type elements stabilizing austenite become insufficiently concentrated in the austenite, so the retained austenite may become unstable and the drop in ductility of the steel sheet may become remarkable. From these viewpoints, the holding time is more preferably 10 seconds or more, more preferably 20 seconds or more.
In the first annealing (Q-annealing), to add a decarburized layer to the surface layer of steel sheet to improve the bendability, it is also possible to control the oxygen potential at one or both of the heated strip and soaked strip at the time of annealing. Specifically, the annealing is preferably performed in an atmosphere including 0.1 to 30 vol % of hydrogen and dew point −40 to 20° C. H2O and having a balance of nitrogen and impurities. More preferably, the atmosphere includes 0.5 to 20 vol % of hydrogen and dew point −30 to 15° C. H2O, more preferably 1 to 10 vol % of hydrogen and dew point −20 to 10° C. H2O.
In the first annealing (Q-annealing), at the time of cooling after soaking, the sheet is preferably cooled from 750° C. to 550° C. by an average cooling rate of 100° C./s or less. The lower limit value of the average cooling rate is not particularly prescribed so long as an area ratio 90% or more of martensite structures are obtained, but, for example, may be 3° C./s. The reason why the lower limit value of the average cooling rate is made 3° C./s is to keep ferrite transformation from occurring at the base material steel sheet and the area ratio of martensite becoming less than 90% at the steel microstructure after Q-annealing. More preferably, the rate is 10° C./s or more, still more preferably 15° C./s or more, still more preferably 20° C./s or more. On the other hand, if the cooling rate from 750° C. to 550° C. is too fast, low temperature transformed structures form at the steel sheet surface layer as well and become causes of variation of hardness. On this point, the average cooling rate is preferably 100° C./s or less, more particularly 80° C./s or less, still more preferably 50° C./s or less. Further, at 750° C. or more, ferrite transformation becomes remarkably hard to occur, so the cooling rate is not limited. Further, at 550° C. or less in temperature, low temperature transformed structures are obtained, so the cooling rate is not limited.
Further, after the above cooling, the sheet may be further cooled to 25° C. to 550° C. in temperature, then made to dwell there at 150° C. to 550° C. in temperature region. The lower limit of the cooling stop temperature was made 25° C. not only because excessive cooling requires massive capital investment, but also because the effect becomes saturated. The dwell time is not particularly limited, but, for example, may be 30 seconds to 500 seconds.
Applying to steel sheet, controlled to a microstructure of mostly martensite by a first annealing, a bending deformation of a bending radius R of 2.0 m or less and in that state as is holding the sheet at 0 to 40° C. for 20 hr or more is an important factor in raising the bending strength of the steel sheet. The carbon atoms dissolved in the martensite during this treatment form clusters or transition carbides and become nuclei for precipitation of carbides at the time of raising the temperature in the following second annealing. To make the carbides finely disperse and increase the bending strength, it is important that the clusters or transition carbides forming nuclei for precipitation of carbides be present finely and at a high density. To promote the formation of clusters or transition carbides, use of tensile strain is extremely effective. This effect is easily obtained at a bending deformation of a bending radius R of 2.0 m or less. The bending radius R may be 1.8 m or less, may be 1.5 m or less, and may be 1.3 m or less. On the other hand, if the bending radius R becomes more than 2.0 m, this effect becomes hard to obtain. For example, it is possible to apply the above bending deformation to steel sheet by coiling up the steel sheet (steel strip) after the first annealing to obtain a coil.
Further, if the holding temperature is less than 0° C., clustering of carbon atoms and formation of transition carbides are suppressed. If the holding temperature is more than 40° C., the transition carbides are formed coarsely (number of nuclei is reduced), so in the second annealing, it becomes hard to obtain fine carbides and the bending strength may fall. If the holding temperatures at the aging treatment 1 and the later explained aging treatment 2 are within 0 to 40° C. in range, the difference in the numbers of precipitates between the aging treatment 1 and the aging treatment 2 becomes smaller and the difference between the number density of precipitates at the first surface at the front side of the steel sheet and the number density of precipitates at the second surface of the back side of the steel sheet becomes within 10%. The holding temperature may be 5° C. or more, may be 10° C. or more, may be 35° C. or less, or may be 30° C. or less.
Furthermore, if the holding time is less than 20 hr, the number of nuclei for formation is not stable. Further, a sufficient amount of nuclei is not formed. Sometimes it is difficult to keep the difference between the number density of precipitates at the first surface at the front side of the steel sheet and the number density of precipitates at a second surface of the back side of the steel sheet to within 10%. The longer the holding time, the better. It may be 30 hr or more, may be 40 hr or more, and may be 50 hr or more. Note that, if the holding time is more than 300 hr, the effect of clustering of carbon atoms or formation of transition carbides becomes saturated. If holding the sheet for more than this, it becomes hard for any remarkable change in the form (size) of precipitates to occur, so the holding time may also be 300 hr or less. Further, if the holding time becomes long, the precipitates become larger, but there is no large change in the number of precipitates. That is, if the holding time at the aging treatment 1 and later explained aging treatment 2 is 20 hr or more, the number of nuclei of formation stabilizes and the difference between the number density of precipitates at the first surface at the front side of the steel sheet and the number density of precipitates at the second surface at the back side of the steel sheet becomes within 10%.
If applying bending deformation to steel sheet to promote aging, clustering of carbon atoms and precipitation of transition carbides remarkably occur in the regions receiving tensile deformation, so if the coiling and uncoiling process is performed only one time by the above aging treatment 1, precipitates finely disperse at only one surface of the steel sheet. For this reason, in the method of production according to the present embodiment, after the aging treatment 1 for making precipitates precipitate and disperse at one side of the front side and back side of steel sheet, aging treatment 2 is performed for making precipitates precipitate and disperse at the other side of the front side and back side of the steel sheet. For example, after the first annealing, the sheet may be coiled up into a shape so that the front side of the sheet becomes the outside and the back side becomes the inside and tensile deformation given at the front side of the sheet so that the bending radius R becomes 2.0 m or less as aging treatment 1, then the coil may be uncoiled and the sheet again coiled up into a shape so that the back side of the sheet becomes the outside and the front side becomes the inside and tensile deformation given at the back side of the sheet so that the bending radius R becomes 2.0 m or less as aging treatment 2. In the aging treatment 2 as well, in the same way as the holding conditions in the aging treatment 1, by performing bending deformation of a bending radius R of 2.0 m or less, the difference between the number density of precipitates at the first surface at the front side of the steel sheet and the number density of precipitates at the second surface of the back side of the steel sheet can be kept within 10%. The bending radius R may be 1.8 m or less, may be 1.5 m or less, and may be 1.3 m or less. On the other hand, if the bending radius R is more than 2 m, it is difficult to keep the difference of the number densities of the precipitate to within 10%.
Further, in the aging treatment 2 as well, in the same way as the aging treatment 1, if the holding temperature is less than 0° C., clustering of carbon atoms and formation of transition carbides are suppressed. If the holding temperature is more than 40° C., the transition carbides are formed coarsely (number of nuclei is reduced), so in the second annealing, it becomes hard to obtain fine carbides and sometimes the bending strength falls. As explained above, if the holding temperatures at the aging treatment 1 and the later explained aging treatment 2 are within 0 to 40° C. in range, the difference between the number density of precipitates at the first surface at the front side of the steel sheet and the number density of precipitates at the second surface of the back side of the steel sheet becomes within 10%. The holding temperature may be 5° C. or more, may be 10° C. or more, may be 35° C. or less, or may be 30° C. or less.
Furthermore, in the aging treatment 2 as well, like the aging treatment 1, if the holding time is less than 20 hr, the number of nuclei for formation is not stable whereby it may be difficult to keep the difference between the number density of precipitates at the first surface at the front side of the steel sheet and the number density of precipitates at the second surface of the back side of the steel sheet to within 10%. The longer the holding time, the better. It may be 30 hr or more, may be 40 hr or more, and may be 50 hr or more. Note that, in the same way as the aging treatment 1, if the holding time is more than 300 hr, the effect of clustering of carbon atoms or formation of transition carbides becomes saturated. If holding the sheet for more than this, it becomes hard for any remarkable change in the form (size) of precipitates to occur, so the holding time may also be 300 hr or less.
In the second annealing (IA annealing), the holding temperature is a temperature at which a dual phase region of ferrite and austenite is formed. For example, it is preferably 720° C. or more and 860° C. or less. If the annealing temperature is less than 720° C., austenite is not sufficiently formed. In this case, the martensite obtained in the first annealing (Q-annealing) is tempered and precipitation of carbides is invited, so the predetermined area ratio of retained austenite may no longer be satisfied. Further, the area ratio of austenite at the maximum heating temperature (annealing temperature) also decreases, so the carbon required for obtaining retained austenite may no longer be concentrated at the austenite and 10.0% or more of retained austenite may no longer be secured. On the other hand, if the annealing temperature is more than 860° C., austenite is excessively formed and tempered martensite including precipitates is decreased, so the bending strength may fall. In addition, 10.0% or more of retained austenite may no longer be secured. For this reason, the upper limit of the holding temperature in the second annealing is preferably 860° C. The annealing may be performed in the air atmosphere and may be performed in an atmosphere controlled in hydrogen concentration and dew point for the purpose of improving the adhesion of the plating.
In the second annealing (IA annealing), the sheet is preferably held at a 720° C. or more and 860° C. or less heating temperature for 5 seconds or more. If the holding time is too short, austenite transformation of the base material steel sheet does not sufficiently proceed and in addition Mn and other substitution type elements stabilizing austenite become insufficiently concentrated in austenite, so the retained austenite becomes unstable and the drop in ductility of the steel sheet may become remarkable. From these viewpoints, the holding time is more preferably 10 seconds or more, more preferably 20 seconds or more.
In the second annealing (IA annealing) as well, like in the first annealing (Q-annealing), to add a decarburized layer to the surface layer of the steel sheet to improve the bendability, it is also possible to control the oxygen potential at one or both of the heated strip and soaked strip at the time of annealing. Specifically, the annealing is preferably performed in an atmosphere including 0.1 to 30 vol % of hydrogen and dew point −40 to 20° C. H2O and having a balance of nitrogen and impurities. More preferably, the atmosphere includes 0.5 to 20 vol % of hydrogen and dew point −30 to 15° C. H2O, more preferably 1 to 10 vol % of hydrogen and dew point −20 to 10° C. H2O.
In the second annealing (IA annealing), at the time of cooling after soaking, the sheet is preferably cooled from 750° C. to 550° C. by an average cooling rate of 100° C./s or less. The lower limit value of the average cooling rate is not particularly prescribed, but, for example, may be 2.5° C./s. The reason for making the lower limit value of the average cooling rate 2.5° C./s is to suppress ferrite transformation from the acicular austenite at which alloy elements concentrate in the base material steel sheet and softening of the base material steel sheet. If the average cooling rate is too slow, the strength easily falls. More preferably, the rate is 5° C./s or more, still more preferably 10° C./s or more, even more preferably 20° C./s or more. On the other hand, if the cooling rate from 750° C. to 550° C. is too fast, low temperature transformed structures are formed on the surface layer of the steel sheet and become causes of variation in hardness. On this point, the average cooling rate is preferably 100° C./s or less, more preferably, it is 80° C./s or less, still more preferably 50° C./s or less. Further, at 750° C. or more, ferrite transformation becomes remarkably hard to occur, so the cooling rate is not limited. Further, at 550° C. or less in temperature, low temperature transformed structures are obtained, so the cooling rate is not limited.
Further, after the above cooling, the sheet may be further cooled to 25° C. to 550° C., then reheated to 150° C. to 550° C. and made to dwell there. If cooling at the above temperature range, martensite is formed from the nontransformed austenite during the cooling. After that, by reheating, carbon concentrates from the martensite to the nontransformed austenite and the strength-ductility balance of the steel sheet is improved. The lower limit of the cooling stop temperature is made 25° C. because not only does excessive cooling require massive capital investment, but also the effect becomes saturated. The dwell time is not particularly limited, but, for example, may be 30 seconds to 500 seconds.
In the method of production of steel sheet according to the present embodiment, it is preferable to obtain retained austenite in acicular form by the second annealing (IA annealing). For example, while the area ratio of retained austenite obtained by the IA annealing becomes 10 to 50%, the temperature is controlled so as to make the change of temperature per second of the steel sheet within ±3° C. or more in the holding process of the IA annealing and thereby make the alloy elements segregate at the interface of the ferrite and austenite while holding in the dual phase region and make the ease of mobility of the interfaces decline, whereby acicular retained austenite is obtained at room temperature.
Furthermore, after reheating and before immersion in the plating bath, the steel sheet may be made to dwell at a 350 to 550° C. temperature region. The dwell operation at this temperature region not only contributes to tempering of the martensite, but also eliminates uneven temperature in the width direction of the sheet and improves the appearance after plating. Note that, if the cooling stop temperature is 350° C. to 550° C., a dwell operation may be performed without reheating.
The time for dwelling is preferably 30 seconds or more and 300 seconds or less for obtaining this effect.
In the series of annealing steps, after the cold rolled steel sheet or plated steel sheet comprised of cold rolled steel sheet given plating treatment is cooled down to room temperature or in the middle of cooling down to room temperature (however, Ms or less), reheating may be started and the sheet held at a temperature region of 150° C. or more and 400° C. or less for 2 seconds or more. According to this step, it is possible to temper the martensite formed during the cooling after the reheating to obtain tempered martensite and thereby improve the hydrogen embrittlement resistance. If performing the tempering step, if the holding temperature is too low or if the holding time is too short, the martensite is not sufficiently tempered and there is almost no change in the microstructure and mechanical properties. On the other hand, if the holding temperature is too high, the dislocation density in the tempered martensite ends up falling and a drop in tensile strength is invited. For this reason, if performing tempering, the sheet is preferably held at 150° C. or more and 400° C. or less for 2 seconds or more. The tempering may be performed in a continuous annealing facility and also may be performed at a separate facility off-line after continuous annealing At this time, the tempering time differs depending on the tempering temperature. That is, the lower the temperature, the longer the time. The higher the temperature, the shorter the time.
The steel sheet may, as necessary, be heated or cooled to (hot dip coating bath temperature-40°) C to (hot dip coating bath temperature+50°) C and hot dip galvanized. By the hot dip galvanization step, the surface of the steel sheet is formed with a hot dip galvanized layer. In this case, the corrosion resistance of the cold rolled steel sheet is improved, so this is preferable. In the present embodiment, the type of the plated layer is not limited to a hot dip galvanized layer. Various types of coated layers may be employed. Further, the timing of plating the steel sheet is not particularly prescribed. For example, in the method of production according to the present embodiment, in the IA annealing, the cold rolled steel sheet is held in the dual phase region of ferrite and austenite, then is cooled down to room temperature. In that process, the front and back surfaces of the steel sheet may be formed with coated layers comprised of zinc, aluminum, magnesium, or their alloys. Alternatively, the front and back surfaces of the steel sheet after annealing may be formed with coated layers after annealing.
The temperature of the steel sheet at the time of immersion in a hot dip galvanization bath is preferably a temperature range from a temperature 40° C. lower than the hot dip galvanization bath temperature (hot dip galvanization bath temperature−40° C.) to a temperature 50° C. higher than the hot dip galvanization bath temperature (hot dip galvanization bath temperature+50° C.). If the temperature is lower than the hot dip galvanization bath temperature-40° C., the heat removal at the time of immersion in the plating bath is large, part of the molten zinc ends up solidifying, and the plating appearance is degraded. If the sheet temperature before immersion is lower than the hot dip galvanization bath temperature−40° C., any method may be used for further heating before immersion in the plating bath, the sheet temperature may be controlled to the hot dip galvanization bath temperature−40° C. or more, and then the sheet may be immersed in the plating bath. Further, if the steel sheet temperature at the time of immersion in the plating bath is more than the hot dip galvanization bath temperature+50° C., sometimes problems in operation accompanying the rise in temperature of the plating bath are induced.
The plating bath is preferably mainly comprised of Zn and has an effective amount of Al (value of total amount of Al in the plating bath minus total amount of Fe) of 0.050 to 0.250 mass %. If the effective amount of Al in the plating bath is too small, Fe excessively penetrates the plating layer and the plating adhesion is liable to fall. On the other hand, if the effective amount of Al in the plating bath is too large, Al-based oxides obstructing movement of Fe atoms and Zn atoms are formed at the boundary of the steel sheet and plated layer and the plating adhesion is liable to fall. The effective amount of Al in the plating bath is more preferably 0.065 mass % or more and more preferably 0.180 mass % or less.
(Temperature of Steel Sheet after Immersion in Plating Bath)
If alloying the hot dip galvanized layer, the steel sheet formed with the hot dip galvanized layer is heated to 450 to 600° C. in temperature range. If the alloying temperature is too low, alloying is liable to not sufficiently proceed. On the other hand, if the alloying temperature is too high, the alloying proceeds to much. Due to the formation of the F phase, the concentration of Fe in the plated layer exceeds 15% and the corrosion resistance is liable to be degraded. The alloying temperature is more preferably 470° C. or more and is more preferably 550° C. or less. The alloying temperature has to be changed in accordance with the chemical composition of the steel sheet, so should be set while checking the concentration of Fe in the plated layer.
To further improve the plating adhesion, before the annealing on the hot dip galvanization line etc., the steel sheet may also be given a plating comprised of one or more of Ni, Cu, Co, and Fe.
To improve the coatability and weldability, the plated surface of hot dip galvanized steel sheet comprised of steel sheet and a hot dip galvanized layer formed on the surface of the steel sheet or the plated surface of hot dip galvannealed steel sheet comprised of steel sheet and a hot dip galvannealed layer formed on the surface of the steel sheet may be given a top layer plating or given various treatment, for example, chromate treatment, phosphate treatment, treatment for improving the lubrication ability, treatment for improving the weldability, etc.
Furthermore, for the purpose of correcting the shape of the steel sheet and raising the ductility due to introduction of movable dislocations, skin pass rolling may also be performed. The rolling reduction of the skin pass rolling after heat treatment is preferably 0.1 to 1.5%. If less than 0.10%, the effect is small and control is difficult, so this becomes the lower limit. If more than 1.5%, the productivity remarkably falls, so this is made the upper limit. The skin pass may be performed in-line or may be performed off-line.
Note that, the difference of hardnesses of the front and back of steel sheet has no substantial relationship with the difference of the number densities of precipitates of steel sheet. That is, even if making the difference of hardnesses of the front and back of steel sheet small, it is not possible to make the difference of the number densities of the precipitates at the front and back of steel sheet smaller, the bending strength of steel sheet is also not necessarily improved, and the difference of the bending strengths of the front and back of the steel sheet also cannot be made smaller. By making the difference of the number densities of the precipitates at the front and back of steel sheet smaller like in the steel sheet according to the present embodiment, it is possible to improve the bending strength of steel sheet and possible to reduce the difference in bending strength at the front and back of steel sheet. To make the difference of the number densities of the precipitates at the front and back of steel sheet smaller, it is effective, like in the method of production of steel sheet according to the present embodiment, to perform aging treatments 1 and 2 between the first annealing of the cold rolled steel sheet (Q-annealing) and second annealing (IA annealing). In the past, reduction of the difference of the number densities of the precipitates at the front and back of steel sheet has not been studied. Performing Q-annealing, aging treatments 1 and 2, and IA annealing like in the method of production according to the present embodiment has not been envisioned at all.
Below, examples according to the present invention will be shown. The present invention is not limited to these illustrations of conditions. The present invention may employ various conditions so long as not deviating from its gist and achieving its object.
Steels having various chemical compositions were smelted to produce steel slabs. These steel slabs were inserted to a furnace heated to 1220° C. and held there for 60 minutes as soaking treatment, then were taken out into the air and hot rolled to obtain thickness 2.8 mm steel sheets. In the hot rolling, the end temperature of the finish rolling was 910° C. The sheets were cooled down to 550° C. and coiled. Next, the hot rolled steel sheets were pickled to remove the oxide scale and were cold rolled by a rolling reduction of 45.0% to finish them to thicknesses of 1.54 mm. Furthermore, the cold rolled steel sheets were Q-annealed, specifically were increased in temperature to 930° C. and held in that temperature range for 90 seconds. Next, the cold rolled steel sheets were cooled and made to dwell at 280° C. and were coiled up to maximum radius 1.4 m coils. The area ratios of martensite in the coiled steel sheets was, in each of the steel compositions, 90% or more. The coils were held at the 6° C. to 22° C. temperature range for 38 hours as aging treatment 1, then the coil was paid out and coiled to again prepare maximum radius 1.4 m coils. The steel sheets were bent in opposite directions to the aging treatment 1 while again being held at the 6° C. to 22° C. temperature range for 38 hours as aging treatment 2. Next, the aging treated steel sheets subjected to the two aging treatments were provided to IA annealing, specifically were increased to 785° C. and held at that temperature for 130 seconds. Next, the aging treated steel sheets were cooled to 270° C., were reheated to 390° C. and made to dwell at that temperature for 140 seconds, were cooled down to room temperature, then were skin pass rolled. The chemical compositions obtained by analysis of samples taken from the obtained steel sheets were as shown in Tables 1 to 3. Note that, in Tables 1 to 3, “−” mean the detection limit values or less. The balances aside from the constituents shown in Tables 1 to 3 are comprised of Fe and impurities. Further, Table 4 shows the results of evaluation of the properties of the steel sheets worked and heat treated.
Further, in Table 4, the methods of measurement of the “area ratio” at the different structures and phases of the steel microstructure, “yield strength YS”, “tensile strength TS”, “total elongation t-El”, “hole expansion rate λ”, and “difference of number densities of precipitates in tempered martensite” were as explained above. Regarding the “yield strength YS”, samples of 600 MPa or more were judged as “passing”. Further, regarding the “bending strength (bending resistance)”, the two of (1) the value itself of the bending strength of the steel sheet and (2) the difference of the bending strengths of the front and back surfaces of the steel sheet were used as indicators for evaluation. Among these, (1) the value of the “bending strength” was judged based on the load at the time of the above VDA bending. Specifically, bending tests were conducted from both of the front and back surfaces of the steel sheet. Cases where the loads at the time of application of a 5° bending angle at both surfaces were 1400N or more per 1 mm thickness were judged as “A”, cases where they were 900N or more and less than 1400N were judged as “B”, and cases where they were less than 900N were judged as “C”. A or B was judged as “passing”. Furthermore, regarding the (2) “differences of the bending strengths of the front and back surfaces of the steel sheets” as well, this was judged based on the load at the time of VDA bending. Specifically, samples with a difference of within 3% of the loads of the front and back surfaces at the time of applying predetermined bending to the front and back surfaces of the steel sheet were judged as “A”, ones of more than 3% to 8% or less were judged as “B”, and ones of more than 8% were judged as “C”. A or B was judged as “passing”.
From the results shown in Tables 1 to 4, the following will be understood.
In BA-1, the C content was too small at the steel sheet, so the yield strength YS and the tensile strength TS of the steel fell and sufficient bending resistance (bending strength) could not be secured.
In BB-1, the C content was too large at the steel sheet, so the elongation fell and brittle fracture of the steel was invited resulting in bending resistance falling.
In BC-1, the Si content was too small at the steel sheet, so it is believed the stability of the retained austenite structures with respect to working fell and precipitation of carbides in the martensite structures at the time of aging could not be suppressed. As a result, the area ratio of the total of the ferrite, pearlite, and bainite increased and the yield strength YS and tensile strength TS of the steel sheet fell. Further, the area ratio of the retained austenite fell and when bending deformation was applied to the steel sheet, the effect of raising the strength by work induced transformation from retained austenite to martensite could not be obtained and sufficient bending resistance could not be secured.
In BD-1, the Si content was too large at the steel sheet, so it is believed that at the aging treatment, formation of c carbides was suppressed. As a result, sufficient bending resistance could not be secured.
In BE-1, the Mn content was too small at the steel sheet, so ferrite transformation easily occurred in the cooling process of the Q-annealing, the ratio of martensite structures after Q-annealing fell, and the area ratio of retained austenite in the finally obtained steel sheet fell. As a result, the yield strength YS and tensile strength TS of the steel sheet fell. Further, when bending deformation was applied to the steel sheet, the effect of raising the strength by work induced transformation from retained austenite to martensite could not be obtained and sufficient bending resistance could not be secured.
In BF-1, the Mn content was too large at the steel sheet, so the concentrated Mn layer formed due to microsegregation and center segregation remarkably appeared in the steel sheet and, due to the difference in solidification speeds at the front and back surfaces of the slab, differences appeared in the state of distribution of the concentrated Mn layer at the front and back surfaces of the steel sheet, so differences also appeared in the number density of precipitates in the tempered martensite due to differences in formation of segregated Mn bands and differences presumably appeared in the bending strength of the front and back surfaces. As a result, sufficient bending resistance could not be secured.
In BG-1, the P content was too large at the steel sheet, so presumably brittle fracture of the steel sheet was invited and a drop in the bending strength at the time of bending deformation was promoted. As a result, sufficient bending resistance could not be secured.
In BH-1, the S content was too large at the steel sheet, so presumably nonmetallic inclusions were formed and a drop in ductility of the steel sheet was invited and formation of voids starting from the nonmetallic inclusions was invited at the time of bending deformation. As a result, sufficient bending resistance could not be secured.
In BI-1, the Al content was too large at the steel sheet, so presumably ferrite transformation and bainite transformation were excessively promoted at the cooling process in the annealing. As a result, the yield strength YS and tensile strength TS of the steel sheet fell and sufficient bending resistance could not be secured.
In BJ-1, the N content was too large at the steel sheet, so presumably it bonded with A1 to form a large amount of AlN and that these nitrides became starting points for formation of voids at the time of bending deformation. As a result, sufficient bending resistance could not be secured.
In BK-1, the Ti content was too large at the steel sheet, so presumably precipitation of coarse carbides became greater and that these carbides became starting points for formation of voids at the time of bending deformation. As a result, sufficient bending resistance could not be secured.
In BL-1, the Co content was too large at the steel sheet, so presumably fine Co carbides precipitated in large numbers and that these carbides became starting points for formation of voids at the time of bending deformation. As a result, sufficient bending resistance could not be secured.
In BM-1, the Ni content was too large at the steel sheet, so presumably the peelability of oxide scale at the time of hot rolling was affected and formation of defects at the surfaces of the steel sheet was promoted. As a result, sufficient bending resistance could not be secured.
In BN-1, the Mo content was too large at the steel sheet, so presumably fine Mo carbides precipitated in large numbers and these carbides became starting points for formation of voids at the time of bending deformation. As a result, sufficient bending resistance could not be secured.
In BO-1, the Cr content was too large at the steel sheet, so formation of retained austenite was promoted. Due to the presence of excessive retained austenite, sufficient bending resistance could not be secured.
In BP-1, the O content was too large at the steel sheet, so presumably a large amount of oxides were formed as inclusions, the oxides at the punched edges or cut surfaces caused notches or coarse dimples to be formed at the end faces, at the time of bending deformation, stress concentration was invited and these became starting points of crack formation. As a result, sufficient bending resistance could not be secured.
In BQ-1, the B content was too large at the steel sheet, so presumably formation of coarse B inclusions in the steel was invited and these inclusions became starting points of formation of voids. As a result, sufficient bending resistance could not be secured.
In BR-1, the Nb content was too large at the steel sheet, so presumably fine, hard Nb carbides were precipitated in large numbers and these carbides became starting points for formation of voids at the time of bending deformation. As a result, sufficient bending resistance could not be secured.
In BS-1, the V content was too large at the steel sheet, so presumably a greater amount of carbonitrides was precipitated and these carbonitrides became starting points for formation of voids at the time of bending deformation. As a result, sufficient bending resistance could not be secured.
In BT-1, the Cu content was too large at the steel sheet, so presumably the strength of the steel sheet rose and the ductility fell. As a result, the bending strength at the time of bending deformation fell and sufficient bending resistance could not be secured.
In BU-1, the W content was too large at the steel sheet, so presumably coarse carbides became starting points for easier formation of voids. As a result, sufficient bending resistance could not be secured.
In BV-1, the Ta content was too large at the steel sheet, so presumably fine Ta carbides were precipitated in large numbers and these carbides became starting points for easier formation of voids. As a result, sufficient bending resistance could not be secured.
In BW-1, the Sn content was too large at the steel sheet, so presumably the steel sheet became brittle and thereby the bending strength at the time of bending deformation fell. Sufficient bending resistance could not be secured.
In BX-1, the Sb content was too large at the steel sheet, so presumably the steel sheet became brittle and thereby the bending strength at the time of bending deformation fell. Sufficient bending resistance could not be secured.
In BY-1, the As content was too large at the steel sheet, so presumably the steel sheet became brittle and thereby the bending strength at the time of bending deformation fell. The sufficient bending resistance could not be secured.
In BZ-1, the Mg content was too large at the steel sheet, so presumably coarse inclusions were formed and these inclusions became starting points of formation of voids. As a result, sufficient bending resistance could not be secured.
In CA-1, the Ca content was too large at the steel sheet, so presumably the steel sheet became brittle and thereby the bending strength at the time of bending deformation fell. Sufficient bending resistance could not be secured.
In CB-1, the Y content was too large at the steel sheet, so presumably coarse Y inclusions were formed and these inclusions became starting points of formation of voids. As a result, sufficient bending resistance could not be secured.
In CC-1, the Zr content was too large at the steel sheet, so presumably coarse Zr inclusions were formed and these inclusions became starting points of formation of voids. As a result, sufficient bending resistance could not be secured.
In CD-1, the La content was too large at the steel sheet, so presumably La inclusions were formed and these inclusions became starting points of formation of voids. As a result, sufficient bending resistance could not be secured.
In CE-1, the Ce content was too large at the steel sheet, so presumably Ce inclusions were formed and these inclusions became starting points of formation of voids. As a result, sufficient bending resistance could not be secured.
As opposed to this, in each of A-1 to AZ-1, steel sheet having a predetermined chemical composition was produced under predetermined conditions whereby a predetermined metallographic structure was obtained at the steel sheet and the steel sheet became excellent in mechanical properties and bending resistance. Further, in the steel sheets obtained in Examples 1, ones with tempered martensite present had number densities of precipitates in a range of 1/μm2 or more and 300/μm2 or less.
Furthermore, to investigate the effects of the production conditions, steel types A to AZ for which excellent properties were observed in Table 1 were hot rolled at the finish temperature described in Tables 5 to 7 to prepare thickness 2.8 mm hot rolled steel sheets. These hot rolled steel sheets were coiled, pickled, and cold rolled to prepare cold rolled steel sheets. These cold rolled steel sheets were annealed and aged. Further, any plating treatments could be performed to obtain steel samples for evaluation of properties. Here, the plated steel sheets were immersed in a hot dip galvanization bath, then held at the temperature shown in Tables 5 to 7 to prepare hot dip galvannealed steel sheets comprised of steel sheets on the surface of which an alloy plated layer of iron and zinc was formed. Further, in some of the steel sheets, in the process of cooling the steel sheets after being held at the respective dwell temperatures down to room temperature in the annealing of the cold rolled steel sheets, the steel sheets cooled down once to 150° C. were reheated to predetermined temperatures and held there for 2 seconds or more for tempering. The obtained results are shown in Tables 5 to 7. Note that, the methods of evaluation of the properties were similar to the case of Examples 1.
From the results shown in Tables 5 to 7, the following will be understood.
In each of A-2 and X-2, the aging time at the aging treatment 1 was too short, so the difference in amounts of precipitation of the precipitates between the aging treatment 1 and the aging treatment 2 became larger. As a result, the difference of the number densities of precipitates in the tempered martensite at the front and back of the steel sheet became greater and the bending resistance of the steel sheet fell.
In each of C-2 and S-3, the annealing holding temperature at the Q-annealing was too high, so presumably the surface layers of the steel sheet became decarburized. As a result, the strength of the steel sheet fell and sufficient bending resistance could not be secured.
In each of I-2 and W-2, the annealing holding temperature at the IA annealing was too high (the annealing holding temperature was a temperature outside the range of the dual phase region of ferrite and austenite), so austenite excessively formed and tempered martensite including precipitates fell, resulting in a drop in the bending strength. Further, it was not possible to secure 10.0% or more of retained austenite and the elongation fell.
In each of L-2 and Z-3, the annealing holding temperature at the IA annealing was too low (the annealing holding temperature was a temperature outside the range of the dual phase region of ferrite and austenite), so austenite could not be sufficiently formed and the martensite obtained by Q-annealing was tempered, so 10.0% or more of retained austenite could not be secured. As a result, when bending deformation was applied to the steel sheet, the effect of raising the strength by work induced transformation from retained austenite to martensite could not be obtained and sufficient bending resistance could not be secured.
In each of E-3 and AX-2, the annealing holding temperature at the Q-annealing was low and the base material steel sheet could not be heated to the austenite single phase region, so the martensite area ratio after Q-annealing fell. As a result, in the aging treatment, a sufficient amount of c carbides could not be made to precipitate and sufficient bending resistance could not be secured. Further, when the area ratio of the retained austenite fell and bending deformation was applied to the steel sheet, the effect of raising the strength by work induced transformation from retained austenite to martensite could not be obtained. On this point as well, sufficient bending resistance could not be secured.
In each of G-3 and U-2, the bending radius at the aging treatment 1 was too large, so presumably sufficient tensile strain was not generated for promoting the clustering or formation of transition carbides. As a result, a sufficient amount of E carbides could not be made to precipitate at the aging treatment 1, the difference of the number densities of precipitates in the tempered martensite at the front and back of the steel sheet became large, and sufficient bending resistance could not be secured.
In M-3, the bending radius at the aging treatment 2 was too large, so presumably sufficient tensile strain was not generated for promoting the clustering or transition carbides. As a result, a sufficient amount of c carbides could not be made to precipitate at the aging treatment 2, the difference of the number densities of precipitates in the tempered martensite at the front and back of the steel sheet became large, and sufficient bending resistance could not be secured.
In each of N-3 and AF-2, the aging time at the aging treatment 2 was too short, so the difference in the amounts of precipitation of the precipitates between the aging treatment 1 and the aging treatment 2 became larger and, as a result, the difference in the number densities of precipitates in the tempered martensite at the front and back of the steel sheet became larger and the bending resistance of the steel sheet fell.
In each of K-4 and AW-4, at least one of the aging treatments 1 and 2 was omitted, so the number density of precipitates in the tempered martensite at least at one of the front and back of the steel sheet was not controlled and the difference in the number densities of precipitates in the tempered martensite at the front and back of the steel sheet became larger. As a result, the bending resistance of the steel sheet fell.
As opposed to this, in the examples other than the above, by producing the steel sheet having a predetermined chemical composition under predetermined conditions, a predetermined microstructure was obtained at the steel sheet and the shapeability and bending resistance of the steel sheet were excellent. Further, in the steel sheets obtained at Examples 2, aging treatment was performed and, in those in which tempered martensite was present, the number density of precipitates was within the range of 1/μm2 or more and 300/μm2 or less.
From the above results, the steel sheet satisfying the following requirements (I) to (IV) can be said to be excellent in strength, elongation, and other mechanical properties and excellent in bending strength.
Further, the steel sheet satisfying the above requirements (I) to (IV) can be produced by the following method:
A method of production of steel sheet comprising
| Number | Date | Country | Kind |
|---|---|---|---|
| 2021-141503 | Aug 2021 | JP | national |
| Filing Document | Filing Date | Country | Kind |
|---|---|---|---|
| PCT/JP2022/016848 | 3/31/2022 | WO |
