Patent Application
20180105891
The present invention relates to a steel sheet with excellent cold workability during forming and a method for manufacturing the sheet.
Automotive parts, knives, and other mechanical parts are manufactured through working processes such as punching, bending, and pressing. In the working processes, improvement of workability is required for a material carbon steel sheet, in order to improve product quality and stability and/or cost reduction.
Generally, a carbon steel sheet is subjected to cold rolling and spheroidizing annealing, so as to produce a soft carbon steel sheet with excellent workability made of ferrite and spheroidized carbide. Many technologies for improving the workability of carbon steel sheets have been proposed so far.
For example, Patent Document 1 discloses a high-carbon steel sheet for precision punching and a method for producing the sheet, wherein the sheet comprises, in terms of % by mass, C: 0.15 to 0.90%, Si: 0.40% or less, Mn: 0.3 to 1.0%, P: 0.03% or less, total Al: 0.1% or less, Ti: 0.01 to 0.05%, B: 0.0005 to 0.0050%, N: 0.01% or less, and Cr: 1.2% or less, has a structure in which carbides having an average carbide grain size of 0.4 to 1.0 μm and a carbide spheroidization ratio of 80% or more are dispersed in a ferrite matrix, and has a notched tensile elongation of 20% or more.
Patent Document 2 discloses a medium- to high-carbon steel sheet with excellent workability and a method for producing the sheet, wherein the sheet comprises C: 0.3 to 1.3 wt %, Si: 1.0 wt % or less, Mn: 0.2 to 1.5 wt %, P: 0.02 wt% or less, and S: 0.02 wt % or less, has a structure in which carbides are dispersed so that the relationship CGB/CIG≤0.8 holds between the carbide number CGB on the ferrite crystal grain boundary and the carbide number CIG in the ferrite crystal grains, and has a cross-sectional hardness of 160 HV or less.
Patent Document 3 discloses a medium- to high-carbon steel sheet with excellent workability, wherein the sheet comprises C: 0.30 to 1.00 wt %, Si: 1.0 wt % or less, Mn: 0.2 to 1.5 wt %, P: 0.02 wt % or less, and S: 0.02 wt % or less, has a structure in which carbides are dispersed in ferrite so that the relationship CGB/CIG≤0.8 holds between the carbide number CGB on the ferrite crystal grain boundary and the carbide number CIG in the ferrite crystal grains, and simultaneously 90% or more of the total carbides are occupied by spheroidized carbides having a long axis/short axis of 2 or less.
Patent Documents 1 to 3 describe that the greater the proportion of carbides in ferrite grains, the more the workability is improved.
In addition, Patent Document 4 discloses a steel sheet having excellent FB workability, mold life, and cold formability after FB processing, wherein the sheet comprises C: 0.1 to 0.5 wt %, Si: 0.5 wt % or less, Mn: 0.2 to 1.5 wt %, P: 0.03 wt % or less, S: 0.02 wt % or less, has a structure based on ferrite and carbide, and the amount Sgb of the carbide present on the ferrite grain boundary is 40% or more, the above Sgb being defined by Sgb={Son/(Son+Sin)}×100 (wherein Son is the total area occupied by the carbides present on the grain boundary among the carbides present per unit area and Sin is the total area occupied by the carbides present on the grain boundary among the carbides present per unit area).
However, in the technology described in Patent Document 1, annealing is performed at a temperature of the AC1 point or higher for softening in order to coarsen ferrite grain size and carbide. But when annealing is performed at a temperature of the AC1 point or higher, rod-like/plate-like carbides may precipitate during annealing. The carbides, even though capable of reducing hardness, deteriorate workability, which is disadvantageous in terms of workability.
The technologies described in Patent Documents 2 and 3 consider that the deterioration of workability is caused by the low carbide spheroidization ratio of carbides precipitated on the grain boundary, but do not take into account the problem of improving the spheroidization ratio of grain boundary carbides. Techniques described in Patent Document 4 only specify the tissue factor, and Patent Document 4 does not discuss the relationship between workability and mechanical properties.
The technology described in Patent Document 5 is an invention made by focusing on the relationship between fine blanking workability and the amount of carbide present in ferrite grains and ferrite grain size. However, Patent Document 5 does not discuss what effect the aggregate structure has on the plastic anisotropy.
Patent Document 6 discloses a hot-rolled steel sheet in which the development of an aggregate structure otherwise developed by rolling is suppressed and a method for manufacturing the sheet. However, Patent Document 6 does not discuss the relationship between the aggregate structure other than the aggregate structure developed by rolling and the cold forgeability.
The technology described in Patent Document 7 is an invention made by considering that the hardness and the total elongation of a high-carbon hot-rolled steel sheet prior to quenching are greatly influenced by the cementite density in the ferrite grains. The hot-rolled steel sheet described in Patent Document 7 is characterized in that it has a microstructure composed of ferrite and cementite, said microstructure having a cementite density of 0.10 strips/μm2 or less in the ferrite grains. However, Patent Document 7 does not discuss what effect the aggregate texture has on the plastic anisotropy.
The technology described in Patent Document 8 is an invention made by considering that the Ceq value is related not only to mechanical properties and weldability but also to the fatigue crack growth rate in steels having a fine structure. Patent Document 8 discloses that by limiting the range of the Ceq value to a range of 0.28% to 0.65%, the fatigue resistance of the steel material is improved and simultaneously weldability is secured. However, Patent Document 8 does not discuss what effect the aggregate texture has on the plastic anisotropy.
[Patent Document 1] Japanese Patent No. 4465057
[Patent Document 2] Japanese Patent No. 4974285
[Patent Document 3] Japanese Patent No. 5197076
[Patent Document 4] Japanese Patent No. 5194454
[Patent Document 5] Japanese Unexamined Patent Publication No. 2007-270331
[Patent Document 6] Japanese Unexamined Patent Publication No. 2009-263718
[Patent Document 7] Japanese Unexamined Patent Publication No. 2015-17294
[Patent Document 8] Japanese Unexamined Patent Publication No. 2004-27355
In view of the current state of the prior art, it is an object of the present invention to address the problem of improving the cold workability of a steel sheet during forming, and to provide a steel sheet that has solved the problem and a method for manufacturing the sheet.
The present inventors have conducted intensive and extensive studies on methods for solving the above-mentioned problems. As a result, the present inventors have found that by controlling the dispersion state of the carbide in the structure of the steel sheet before cold working through the optimization of the manufacturing conditions in the steps from hot rolling to annealing, the carbide can be precipitated on the ferrite boundary and simultaneously the aggregate structure in the hot rolled steel plate can be controlled, thereby leading to enhanced cold workability.
Further, we have found after intensive and extensive research that it is difficult to manufacture a steel sheet that satisfies the above-mentioned conditions merely by devising hot rolling conditions and annealing conditions separately, and that it can be manufactured by optimizing the above conditions in mutual cooperation in an integrated process of the hot rolling and annealing steps.
The present invention has been made based on the above findings, and the gist thereof lies in:
(1) A steel sheet having an excellent cold workability during forming, comprising, in terms of % by mass:
a balance of Fe and inevitable impurities,
wherein (a) a ratio of the number of carbides at a ferrite grain boundary relative to the number of carbides in the ferrite grain is more than 1,
wherein (b) a diameter of the ferrite grain is 5 μm or more and 50 μm or less,
wherein (c) an in-plane anisotropy |Δr| of the r value standardized according to JIS Z 2254 is 0.2 or less,
wherein (d) a Vickers hardness of the steel sheet is 100 HV or more and 150 HV or less, and
wherein (e) a ratio of X-ray diffraction intensity of the {311} <011> orientation at the ½-thickness portion of the steel sheet relative to the X-ray diffraction intensity obtained when a sample with a random orientation distribution of crystal grains in the steel sheet is subjected to X-ray diffraction is 3.0 or less.
(2) The steel sheet with excellent cold workability during forming described in the above (1) further comprising, in terms of % by mass, one or a plurality of:
(3) A method for producing a steel sheet with excellent cold workability during forming according to the above (1) or (2), the method comprising:
subjecting a steel strip having an ingredient composition according to claim 1 or 2 to hot rolling by heating, followed by completing the finish hot rolling at a temperature range of 800° C. or higher and 900° C. or lower;
coiling the hot-rolled steel sheet at a temperature of 400° C. or higher and 550° C. or lower;
pickling the hot-rolled steel sheet, and then subjecting the hot-rolled steel sheet to a two-step type annealing in which the hot-rolled steel sheet is retained in two temperature ranges,
wherein the two-step type annealing comprises
(i) subjecting the hot-rolled steel sheet to a first step annealing performed by retaining said hot-rolled steel at a temperature range of 650° C. or higher and 720° C. or lower for 3 hours or longer and 60 hours or shorter, and then a second step annealing performed by retaining the hot-rolled steel at a temperature range of 725° C. or higher and 790° C. or lower for 3 hours or longer and 50 hours or shorter, and thereafter
(ii) cooling the hot-rolled steel sheet to 650° C. or lower at a cooling rate of 1° C./hour or more and 30° C./hour or less.
(4) The method for producing a steel sheet described in the above (3), wherein the steel sheet has a cross-sectional shrinkage percentage of 40% or more.
According to the present invention, a steel sheet with excellent cold workability during forming can be manufactured and provided.
A steel sheet with excellent cold workability during forming according to the present invention (hereinafter may be referred to as “the inventive steel sheet”) comprises, in terms of % by mass:
a balance of Fe and inevitable impurities,
the above sheet being characterized in that:
(a) the ratio of the number of carbides at a ferrite grain boundary relative to the number of carbides in the ferrite grain exceeds 1,
(b) the ferrite grain diameter is 5 μm or more and 50 μm or less,
(c) the in-plane anisotropy |Δr| of the r value standardized according to JIS Z 2254 is 0.2 or less,
(d) the Vickers hardness is 100 HV or more and 150 HV or less, and
(e) the ratio of X-ray diffraction intensity of the {311} <011>orientation at the ½-thickness portion of the steel sheet relative to the X-ray diffraction intensity obtained when a sample with a random orientation distribution of crystal grains in the steel sheet is subjected to X-ray diffraction is 3.0 or less.
The method (hereinafter may be referred to as “the inventive method”) of the present invention for producing a steel sheet with excellent cold workability during forming is a method for producing the inventive steel sheet,
wherein a hot-rolled steel strip that has been obtained by subjecting a steel strip having an ingredient composition of the inventive steel sheet to hot rolling by heating, followed by completing the finish hot rolling at a temperature range of 800° C. or higher and 900° C. or lower, and by coiling the resulting hot-rolled steel sheet at a temperature of 400° C. or higher and 550° C. or lower is, after pickling, subjected to two-step type annealing in which the sheet is retained in two temperature ranges, whereupon
(i) the hot-rolled steel sheet is subjected to a first step annealing performed by retaining said hot-rolled steel at a temperature range of 650° C. or higher and 720° C. or lower for 3 hours or longer and 60 hours or shorter, and then subjected to a second step annealing performed by retaining the hot-rolled steel at a temperature range of 725° C. or higher and 790° C. or lower for 3 hours or longer and 50 hours or shorter, and thereafter
(ii) the sheet is cooled down to 650° C. or lower at a cooling rate of 1° C./hour or more and 30° C./hour or less.
Hereinafter, the inventive steel sheet and the inventive manufacturing method will be described.
First, the reasons for limiting the ingredient composition of the inventive steel sheet will be described. The percentage relating to the ingredient composition means % by mass.
C is an element that forms carbide in steel, and is effective for strengthening steel and refining ferrite grains. In order to prevent the surface of the steel sheet from being textured by cold working and ensure the aesthetic appearance of surface of cold forged parts, it is necessary to suppress the coarsening of ferrite grain size. However, when its content is less than 0.10%, the volume fraction of the carbide is insufficient and the coarsening of carbides during annealing cannot be suppressed. Therefore, C is set to 0.10% or more, and preferably 0.12% or more.
On the other hand, when it exceeds 0.40%, the volume fraction of the carbide increases, a large amount of cracks serving as fracture starting points are formed when a load is instantaneously applied, and thus the impact resistance property decreases. Therefore, C is set to 0.40% or less, and preferably 0.38% or less.
Si is an element that acts as a deoxidizing agent and also affects the form of the carbide. In order to reduce the number of carbides in the ferrite grain and increase the number of carbides on the ferrite grain boundaries, it is necessary to generate an austenite phase during annealing in the two-step type annealing, and, after transiently dissolving the carbides, to cool gradually to promote the precipitation of carbides at the ferrite grain boundaries.
In the inventive steel sheet, the amount of Si may preferably be as small as possible. However, when it is reduced to less than 0.01%, the manufacturing cost increases. Therefore, Si is set to 0.01% or more.
On the other hand, when it exceeds 0.30%, the ductility of ferrite lowers and breaking may easily occur during cold working, resulting in reduced cold workability. Therefore, Si is set to 0.30% or less, and preferably 0.28% or less.
Mn is an element that controls the figuration of carbides in the two-step type annealing. When its content is less than 0.30%, it is difficult to precipitate carbides at the ferrite grain boundaries in slow cooling after the second-step annealing. Therefore, Mn is set to 0.30% or more, and preferably 0.33% or more.
On the other hand, when it exceeds 1.00%, the hardness of ferrite increases and the cold workability deteriorates. Therefore, Mn is set to 1.00% or less, and preferably 0.96% or less.
P is an element that segregates at the ferrite grain boundaries and suppresses the formation of grain boundary carbides. The amount of P may preferably be as small as possible. However, when P is reduced to less than 0.0001% in the refining process, the refining cost may greatly increase. Therefore, it is set to 0.0001% or more, and preferably 0.0013% or more.
On the other hand, when it exceeds 0.020%, the number percentage of the grain boundary carbides decreases and the cold workability deteriorates. Therefore, P is set to 0.020% or less, and preferably 0.018% or less.
S is an element that forms a non-metallic inclusion such as MnS. Since a non-metallic inclusion serves as the starting point for break generation during cold forging, the amount of S may preferably be as small as possible. However, when it is reduced to less than 0.0001%, the refining cost greatly increases. Therefore, S is set to 0.0001% or more, and preferably 0.0012% or more.
On the other hand, when it exceeds 0.010%, cold workability deteriorates. Therefore, S is set to 0.010% or less, and preferably 0.007% or less.
Al is an element that acts as a deoxidizing agent for steel and stabilizes ferrite. When its content is less than 0.001%, a sufficient addition effect cannot be obtained. Therefore, Al is set to 0.001% or more, and preferably 0.004% or more.
On the other hand, when it exceeds 0.10%, the number percentage of carbides on the grain boundary decreases and the cold workability deteriorates. Therefore, Al is set to 0.10% or less, and preferably 0.08% or less.
In addition to the above elements, the inventive steel sheet may contain one or a plurality of N: 0.0001 to 0.010%, 0: 0.0001 to 0.020%, Cr: 0.001 to 0.50%, Mo: 0.001 to 0.10%, Nb: 0.001 to 0.10%, V: 0.001 to 0.10%, Cu: 0.001 to 0.10%, W: 0.001 to 0.10%, Ta: 0.001 to 0.10%, Ni: 0.001 to 0.10%, Sn: 0.001 to 0.050%, Sb: 0.001 to 0.050%, As: 0.001 to 0.050%, Mg: 0.0001 to 0.050%, Ca: 0.001 to 0.050%, Y: 0.001 to 0.050%, Zr: 0.001 to 0.050%, La: 0.001 to 0.050%, and Ce: 0.001 to 0.050%, in order to improve the properties of the inventive steel sheet.
N is an element that, when present in large amounts, causes the embrittlement of ferrite. The amount of N may preferably be as small as possible. However, when it is reduced to less than 0.0001%, the refining cost greatly increases. Therefore, N should be 0.0001% or more, and preferably 0.0006% or more. On the other hand, when it exceeds 0.010%, ferrite embrittles and the cold forgeability deteriorates. Therefore, N should be 0.010% or less, and preferably 0.007% or less.
O is an element that, when present in large amounts, forms coarse oxides in steel. The amount of O may preferably be as small as possible. However, when it is reduced to less than 0.0001%, the refining cost increases greatly. Therefore, O is set to 0.0001% or more, and preferably 0.0011% or more. On the other hand, when it exceeds 0.020%, coarse oxides are formed in the steel, the oxides serving as the starting point for break generation during cold working. Therefore, 0 is set to 0.020% or less, and preferably 0.017% or less.
Cr is an element which enhances quenchability and contributes to the improvement of strength and which is thickened to carbide and forms stable carbide even in the austenitic phase. When its content is less than 0.001%, the sufficient effect of improving quenchability cannot be obtained. Therefore, Cr is set to 0.001% or more, and preferably 0.007% or more. On the other hand, when it exceeds 0.50%, the carbide becomes stabilized thereby delaying the dissolution of the carbide during quenching, and thus, it is feared that the desired quenching strength may not be achieved. Therefore, Cr is set to 0.50% or less, and preferably 0.45% or less.
Like Mn, Mo is an element effective for controlling the figuration of carbides. When its content is less than 0.001%, a sufficient addition effect cannot be obtained. Therefore, Mo is set to 0.001% or more, and preferably 0.010% or more. On the other hand, when it exceeds 0.10%, the in-plane anisotropy of the r value deteriorates and the cold workability deteriorates. Therefore, Mo is set to 0.10% or less, and preferably 0.08% or less.
Nb is an element which is effective for controlling the figuration of carbides and which refines the structure, thereby contributing to the enhancement of its toughness. When its content is less than 0.001%, a sufficient addition effect cannot be obtained. Therefore, Nb should be 0.001% or more, and preferably 0.004% or more. On the other hand, when it exceeds 0.10%, a large number of fine Nb carbides precipitate, which leads to excessively increased strength. It also causes the reduction in the number ratio of grain boundary carbides, and the deterioration in cold forgeability. Therefore, Nb is set to 0.10 or less, and preferably 0.08% or less.
Like Nb, V is an element which is effective for controlling the figuration of carbides and which refines the structure, thereby contributing to the enhancement of its toughness. When its content is less than 0.001%, a sufficient addition effect cannot be obtained. Therefore, V is set to 0.001% or more, and preferably 0.004% or more. On the other hand, when it exceeds 0.10%, a large number of fine V carbides precipitate, which leads to excessively increased strength, to the reduced number ratio of grain boundary carbides, and to the deteriorated cold forgeability. Therefore, V is set to 0.10 or less, and preferably 0.08% or less.
Cu is an element which segregates at the ferrite crystal grain boundary and forms fine precipitates thereby to contribute to the enhancement of strength. When its content is less than 0.001%, a sufficient effect of enhancing strength cannot be obtained. Therefore, Cu is set to 0.001% or more, and preferably 0.005% or more. On the other hand, when it exceeds 0.10%, red heat embrittlement occurs and the productivity by hot rolling decreases. Therefore, Cu is set to 0.10% or less, and preferably 0.08% or less.
Like Nb and V, W is also an element effective for controlling the figuration of carbides. When its content is less than 0.001%, a sufficient addition effect cannot be obtained. Therefore, W is set to 0.001% or more, and preferably 0.003% or more. On the other hand, when it exceeds 0.10%, a large number of fine W carbides precipitate, which leads to excessively increased strength, to the reduced number ratio of grain boundary carbides, and to the deteriorated cold forgeability. Therefore, W is set to 0.10 or less, and preferably 0.08% or less.
Like Nb, V and W, Ta is also an element effective for controlling the figuration of carbides. When its content is less than 0.001%, a sufficient addition effect cannot be obtained. Therefore, W is set to 0.001% or more, and preferably 0.005% or more. On the other hand, when it exceeds 0.10%, a large number of fine W carbides precipitate, which leads to excessively increased strength, to the reduced number ratio of grain boundary carbides, and to the deteriorated cold forgeability. Therefore, Ta is set to 0.10 or less, and preferably 0.08% or less.
Ni is an element effective for improving the toughness of parts. When its content is less than 0.001%, a sufficient addition effect cannot be obtained. Therefore, Ni is set to 0.001% or more, and preferably 0.003% or more. On the other hand, when it exceeds 0.10%, the number ratio of grain boundary carbides decreases and the cold forgeability deteriorates. Therefore, Ni is set to 0.10% or less, and preferably 0.08% or less.
Sn is an element contaminated from a steel raw material (scrap). It segregates at the grain boundary, leading to the decreased number ratio of grain boundary carbides. Therefore, its content may preferably be as small as possible. However, when it is reduced to less than 0.001%, the refining cost will be greatly increased. Therefore, Sn is set to 0.001% or more, and preferably 0.002% or more. On the other hand, when it exceeds 0.050%, ferrite embrittles and cold forgeability deteriorates. Therefore, Sn is set to 0.050% or less, and preferably 0.040% or less.
Like Sb, Sb is an element contaminated from a steel raw material (scrap). It segregates at the grain boundary, leading to the decreased number ratio of grain boundary carbides. Therefore, its content may preferably be as small as possible. However, when it is reduced to less than 0.001%, the refining cost will be greatly increased. Therefore, Sb is set to 0.001% or more, preferably 0.002% or more. On the other hand, when it exceeds 0.050%, the cold forgeability deteriorates. Therefore, Sb is set to 0.050% or less, and preferably 0.040% or less.
Like Sn and Sb, As is an element contaminated from a steel raw material (scrap). It segregates at the grain boundary, thereby leading to a decrease in the number ratio of grain boundary carbides. Therefore, its content may preferably be as small as possible. However, when it is reduced to less than 0.001%, the refining cost increases greatly. Therefore, As is set to 0.001% or more, and preferably 0.002% or more. On the other hand, when it exceeds 0.050%, the number ratio of the grain boundary carbides decreases and the cold forgeability deteriorates. Therefore, As is set to 0.050% or less, and preferably 0.040% or less.
Mg is an element that can control the figuration of sulfides with the addition of its trace amount. When its content is less than 0.0001%, a sufficient addition effect cannot be obtained. Therefore, Mg is set to 0.0001% or more, and preferably 0.0008% or more. On the other hand, when it exceeds 0.050%, ferrite embrittles and the cold forgeability deteriorates. Therefore, Mg is set to 0.050% or less, and preferably 0.040% or less.
Like Mg, Ca is an element that can control the figuration of sulfides with the addition of its trace amount. When its content is less than 0.001%, a sufficient addition effect cannot be obtained. Therefore, Ca is set to 0.001% or more, and preferably 0.003% or more. On the other hand, when it exceeds 0.050%, coarse Ca oxides are formed, which serve as starting points of break generation during cold forging. Therefore, Ca is set to 0.050% or less, and preferably 0.040% or less.
Like Mg and Ca, Y is an element that can control the figuration of sulfides with the addition of its trace amount. When its content is less than 0.001%, a sufficient addition effect cannot be obtained. Therefore, Y is set to 0.001% or more, and preferably 0.003% or more. On the other hand, when it exceeds 0.050%, coarse Y oxides are formed, which serve as starting points of break generation during cold working. Therefore, Y is set to 0.050% or less, and preferably 0.035% or less.
Like Mg, Ca and Y, Zr is an element that can control the figuration of sulfides with the addition of its trace amount. When its content is less than 0.001%, a sufficient addition effect cannot be obtained. Therefore, Zr is set to 0.001% or more, and preferably 0.004% or more. On the other hand, when it exceeds 0.050%, coarse Zr oxides are formed, which serve as starting points for break generation during cold working. Therefore, Zr is set to 0.050% or less, and preferably 0.045% or less.
La is an element that can control the figuration of sulfides with the addition of its trace amount, but it is also an element that segregates at the grain boundary and causes a decrease in the number ratio of grain boundary carbides. When its content is less than 0.001%, a sufficient effect of controlling figuration cannot be obtained. Therefore, La is set to 0.001% or more, and preferably 0.004% or more. On the other hand, when it exceeds 0.050%, the number ratio of grain boundary carbides decreases and the cold workability deteriorates. Therefore, La is set to 0.050% or less, and preferably 0.045% or less.
Like La, Ce is an element that can control the figuration of sulfides with the addition of its trace amount, but it is also an element that segregates at the grain boundary and causes a decrease in the number ratio of grain boundary carbides. When its content is less than 0.001%, a sufficient effect of controlling figuration cannot be obtained. Therefore, Ce is set to 0.001% or more, and preferably 0.004% or more. On the other hand, when it exceeds 0.050%, the number ratio of grain boundary carbides decreases and the cold forgeability deteriorates. Therefore, Ce is set to 0.050% or less, and preferably 0.045% or less.
The remainder of the ingredient composition of the inventive steel sheet is Fe and unavoidable impurities.
It is a novel finding by the inventors that the inventive steel sheet has excellent cold workability during forming, because, in addition to the above ingredient composition, it was found, as a result of optimum hot rolling and annealing, that
(a) the ratio of the number of carbides at the ferrite grain boundary relative to the number of carbides in the ferrite grain exceeds 1,
(b) the ferrite grain diameter is 5 μm or more and 50 μm or less,
(c) the in-plane anisotropy |Δr| of the r value standardized according to JIS Z 2254 is 0.2 or less,
(d) the Vickers hardness is 100 HV or more and 150 HV or less, and
(e) the ratio of X-ray diffraction intensity of the {311} <011> orientation at the ½-thickness portion of the steel sheet relative to the X-ray diffraction intensity obtained when a sample with a random orientation distribution of crystal grains in the steel sheet is subjected to X-ray diffraction is 3.0 or less.
The above (a) to (e) will be described below.
(a) The ratio of the number of carbides at the ferrite grain boundary relative to the number of carbides in the ferrite grain exceeds 1:
The inventive steel sheet has a structure which is substantially composed of ferrite and carbide, and in which the ratio of the number of carbides at the ferrite grain boundary relative to the number of carbides in the ferrite grain exceeds 1. Carbides are, in addition to cementite (Fe3C) that is a compound of iron and carbon, compounds obtained by replacing Fe in cementite with an element such as Mn and Cr, and alloy carbides (M23C6, M6Co, MC, etc., wherein M is Fe and another additive metal element).
When a steel sheet is formed into a predetermined part shape, a shear band is formed in the macrostructure of the steel sheet, and slip deformation is generated and concentrated in the vicinity of the shear band. The slip deformation involves propagation of dislocations, and regions with high dislocation density are formed in the vicinity of the shear band. As the strain amount applied to the steel sheet increases, the slip deformation is promoted and thereby the dislocation density increases. In cold forging, strong processing exceeding an equivalent strain of 1 is applied.
For this reason, in the conventional steel sheet, generation of voids and/or cracks due to the increased dislocation density could not be prevented, and it was difficult to improve cold forgeability.
In order to solve the above challenging problems, it is effective to suppress the formation of shear bands during forming. From the viewpoint of a microstructure, shear band formation is a phenomenon in which a slip generated in one crystal grain crosses the crystal grain boundary and propagates continuously to an adjacent crystal grain. Therefore, in order to suppress the formation of a shear band, it is necessary to prevent the propagation of slippage beyond crystal the grain boundary.
Carbides in the steel sheet are tenacious particles that hinder slippage. Therefore, the presence of carbides at the ferrite grain boundaries would make it possible, for the first time, to suppress the formation of a shear band and thereby to improve cold forgeability.
Based on the theory and principle, it is considered that cold forgeability is strongly influenced by the coverage rate of carbides at the ferrite grain boundaries. Therefore, it becomes necessary to measure the coverage rate with high accuracy.
In order to measure the coverage rate of carbides at the ferrite grain boundaries in a three-dimensional space, serial sectioning SEM observation or repeated three-dimensional EBSP observation become essential in which sample cutting by FIB and observation are repeated in the scanning electron microscope. However, these methods take a huge amount of measurement time and the accumulation of technical know-how becomes indispensable. We clarified this fact and concluded that common analytical methods are not suitable.
Therefore, as a result of searching a simple and highly accurate evaluation index, the present inventors have found that cold forgeability can be evaluated by using, as an index, the ratio of the number of carbides at the ferrite grain boundary relative to the number of carbides in the ferrite grain, and that cold forgeability can be remarkably improved when the ratio of the number of carbides at the ferrite grain boundary relative to the number of carbides in the ferrite grain is more than 1.
Any of buckling, folding and convolution of a steel sheet that occurs during cold working is caused by the localization of strain accompanying the formation of a shear band. Therefore, by allowing the carbide to exist at the ferrite grain boundaries, the formation of the shear band and the localization of strain can be alleviated and the generation of buckling, folding and convolution can be suppressed.
When the spheroidization percentage of carbides on the crystal grain boundary is less than 80%, strains are concentrated locally on the rod-shaped or plate-shaped carbides, and voids and/or cracks are likely to occur. Therefore, the carbide spheroidization ratio on the crystal grain boundary may preferably be 80% or more, and more preferably 90% or more.
When the average particle diameter of the carbide in the ferrite grain and the carbide at the ferrite grain boundaries is less than 0.1 μm, the hardness of the steel sheet remarkably increases and the workability deteriorates. Therefore, the average particle diameter of the carbide may preferably be 0.1 μm or more, and more preferably 0.17 μm or more. On the other hand, when the average particle diameter of the carbide exceeds 2.0 μm, fissures occur with the coarse carbide serving as a starting point during cold working, and thus the cold workability deteriorates. Therefore, the average particle diameter of the carbide may preferably be 2.0 μm or less, and more preferably 1.95 μm or less.
Subsequently, the method of observing and measuring the structure will be described.
Observation of the carbide is carried out by a scanning electron microscope. Prior to observation, samples for structure observation are polished by wet polishing with emery paper and polishing with diamond abrasive grains having an average particle size of 1 μm. After polishing the observation surface to a mirror finish, the structure is etched with a 3% nitric acid-alcohol solution.
Among the magnification for observation, within 3000 times, a magnification capable of discriminating between ferrite and carbide is selected. At the selected magnification, eight images with a viewing field of 30 μm×40 μm are randomly photographed at the ¼ plate layer thickness.
With respect to the tissue image obtained, the area of each carbide contained in the region is measured in detail by an image analysis software represented by Mitsuya Shoji Co. Ltd. (Win ROOF). A circle equivalent diameter (=2×√(area/3.14)) is obtained from the area of each carbide, and the average value is taken as the carbide particle diameter.
Further, the spheroidization ratio of the carbide was determined by approximating the carbide to an ellipse having an equal area and equal moment of inertia, and then by calculating the proportion of the carbides in which the ratio of the maximum length to the maximum length in the perpendicular direction is less than 3.
In order to suppress the effect of measurement error due to noise, carbides having an area of 0.01 μm2 or more among the carbides in grains and grain boundaries were counted and the carbides having an area of 0.01 μm2 or less were excluded from evaluation.
The number of carbides present on the ferrite grain boundary was counted, and from the total number of carbides the number of carbides in the ferrite grain was determined by subtracting the number of carbides on the ferrite grain boundary. Based on the measured number, the ratio of the number of carbides on the grain boundary relative to the number of carbides in the ferrite grain was determined.
(b) The ferrite grain diameter is 5 μm or more and 50 μm or less:
In the structure after annealing the cold rolled steel sheet, the cold workability can be improved by setting the ferrite grain diameter to 5 μm or more. When the ferrite grain size is less than 5 μm, the hardness increases and fissures and cracks tend to generate easily during cold working. Therefore, the ferrite grain size is set to 5 μm or more, and preferably 7 μm or more.
On the other hand, when it exceeds 50 μm, the number of carbides on the crystal grain boundary that suppress slippage propagation decreases and the cold workability deteriorates. Therefore, that the ferrite grain size is set to 50 μm or less, and preferably 37 μm or less.
The ferrite grain diameter is measured in the above-described polishing method, wherein the observation surface of the sample is polished to a mirror surface, followed by etching with a 3% nitric acid-alcohol solution. The structure of the observation surface is then examined with an optical microscope or a scanning electron microscope, and a line segment method is then applied to the image photographed to determine the ferrite grain diameter.
(c) The in-plane anisotropy |Δr| of the r value standardized according to JIS Z 2254 is 0.2 or less:
The in-plane anisotropy |Δr| of the plastic strain ratio (r value) of the steel sheet is measured in a method in accordance with JIS Z 2254. The r value (0° direction: r0, 45° direction: r45, 90° direction: r90) measured by taking test strips from each direction of 0° direction, 45° direction and 90° direction with respect to the rolling direction was used to calculate the following equation.
|Δr|=(r0−2r45+r90)/2
By setting the in-plane anisotropy |Δr| of the plastic strain ratio (r value) of the steel sheet to 0.2 or less, the cold workability can be improved. When |Δr| exceeds 0.2, the thickness of parts and the height of the earing become uneven during drawing. Therefore, the in-plane anisotropy |Δr| is set to 0.2 or less.
(d) The Vickers hardness is 100 HV or more and 150 HV or less:
By setting the Vickers hardness of the steel sheet to 100 HV or more and 150 HV or less, the cold workability can be improved. When the Vickers hardness is less than 100 HV, buckling can easily occur during cold working. Therefore, the Vickers hardness is set to 100 HV or more, and preferably 110 HV or more.
On the other hand, when the Vickers hardness exceeds 150 HV, the ductility decreases and the internal breaking tends to occur easily during cold forging. Therefore, the Vickers hardness is set to 150 HV or less, and preferably 146 HV or less.
(e) The ratio of X-ray diffraction intensity of the {311} <011> orientation at the ½-thickness portion of the steel sheet relative to the X-ray diffraction intensity obtained when a sample with a random orientation distribution of crystal grains in the steel sheet is subjected to X-ray diffraction is 3.0 or less:
In cold forging, in addition to controlling the figuration of carbides, the draw formability during cold forging must be secured. In order to improve the draw formability during cold forging, plastic anisotropy such as in-plane anisotropy |Δr| must be improved. For that purpose, the aggregate structure of a hot-rolled steel sheet must be controlled. For evaluation of the aggregate structure, analysis by X-ray diffraction on a plane parallel to the plate surface at the ½ thickness portion of the hot-rolled steel plate is used.
One surface of a hot-rolled steel plate is ground to a ½ plate thickness surface in parallel to the surface to expose a ½ plate thickness surface, followed by the analysis of the ½ plate thickness surface by X-ray diffraction. As the X-ray diffraction, X-ray diffraction by Mo bulb may be used. Diffraction intensities of diffraction orientations {110}, {220}, {211} and {310} by reflection are obtained, and based thereon, the orientation distribution function (ODF) is created.
The X-ray diffraction intensity ratio is determined by using the diffraction intensity data of the ½ plate thickness surface obtained from the ODF and the diffraction intensity data of random orientation of the hot-rolled steel sheet. Specifically, as a standard sample in which the metallic structure has no accumulation in a specific direction, a sample obtained by sintering powder iron of a hot-rolled steel sheet to be measured or the powder before sintering is used to determine the diffraction intensity under the same conditions as when the diffraction intensity data of the ½ plate thickness surface was obtained. The part to be collected as the standard sample is not particularly limited and may be any part of the hot-rolled steel sheet. The X-ray diffraction intensity ratio in a specific orientation is a numerical value obtained by dividing the diffraction intensity in the specific direction of the ½ plate thickness surface obtained from the ODF by the diffraction intensity of the standard sample.
When the X-ray diffraction intensity ratio of the {311} <011> orientation obtained by the above-described ODF analysis is set to I1, it is necessary that this I1 is 3.0 or less, and preferably 2.5 or less for the random aggregate structure during hot rolling. When a random aggregate structure having I1 of 3.0 or less can be obtained, the plasticity anisotropy is reduced and the cold formability is improved.
Next, the inventive manufacturing method will be described.
The manufacturing method according to the present invention is characterized in that the hot rolling and the annealing are consistently managed to control the structure. After continuously casting a steel strip having a predetermined ingredient composition, the steel strip is subjected to hot rolling by heating to complete finish hot rolling at a temperature range of 800° C. or higher to 900° C. or lower, coiled at 400° C. or higher and 550° C. or lower to obtain a hot-rolled steel sheet. The hot-rolled steel sheet is, after pickling, subjected to a two-step type annealing in which the hot-rolled steel sheet is maintained in two temperature ranges, whereupon
(i) the hot-rolled steel sheet is subjected to a first step annealing performed by retaining said hot-rolled steel at a temperature range of 650° C. or higher and 720° C. or lower for 3 hours or longer and 60 hours or shorter, and then subjected to a second step annealing performed by retaining the hot-rolled steel at a temperature range of 725° C. or higher and 790° C. or lower for 3 hours or longer and 50 hours or shorter, and thereafter
(ii) the hot-rolled steel sheet is cooled to 650° C. or lower at a cooling rate of 1° C./hour or more and 30° C./hour or less,
and thus a steel sheet excellent in cold workability during forming can be produced.
By the hot rolling and annealing mentioned above, a structure composed of fine pearlite and bainite can be formed as the structure of the steel sheet.
The processing conditions will be described below.
Heating temperature of a steel strip: 1000° C. or higher and 1250° C. or lower
The heating temperature of the steel strip subjected to hot rolling may preferably be 1000° C. or higher and 1250° C. or lower, and the heating time may preferably be 0.5 hour or longer and 3 hours or shorter.
When the heating temperature is lower than 1000° C. or the heating time is shorter than 0.5 hour, the microsegregation and/or macrosegregation formed by casting are not eliminated, and regions in which Si, Mn, etc., are locally concentrated inside the steel material may remain, and thus the impact resistance property of the steel material is lowered. Therefore, the heating temperature may preferably be 1000° C. or higher, and preferably 0.5 hour or longer.
On the other hand, when the heating temperature exceeds 1250° C. or the heating time exceeds 3 hours, decarburization from the surface layer of the steel strip becomes conspicuous, and austenite grains in the surface layer grow abnormally during heating before carburizing and quenching, and the impact resistance property of the steel strip is deteriorated. Thus the heating temperature may preferably be 1250° C. or lower, and the heating time may preferably be 3 hours or shorter.
Finish hot rolling temperature: 800° C. or higher and 900° C. or lower
Finish hot rolling is completed at 800° C. or higher and 900° C. or lower. When the finish hot rolling temperature is lower than 800° C., the deformation resistance of the steel strip increases, the rolling load increases markedly, the wear amount of the roll increases, and the productivity decreases. Therefore, the finish hot rolling temperature is set to 800° C. or higher, and preferably 820° C. or higher.
On the other hand, when the finish hot rolling temperature exceeds 900° C., thick scales are generated during plate passing on the ROT (Run Out Table), scratches are generated on the surface of the steel sheet due to the scale, and cracks are generated starting from scratches when an impact load is applied after cold forging and carburizing and annealing, leading to reduced impact resistance property of the steel sheet. Therefore, the finish hot rolling temperature is set to 900° C. or lower, and preferably 880° C. or lower.
Cooling rate on ROT: 10° C./sec or more and 100° C./sec or less
The cooling rate at the time of cooling the hot-rolled steel sheet on the ROT after finish hot rolling may preferably be 10° C./sec or more and 100° C./sec or less. When the cooling rate is less than 10° C./sec, thick scales are generated during cooling and the occurrence of scratches on the surface of the steel sheet due to the scales cannot be suppressed. Therefore, the cooling rate is set to 10° C./sec or more, and more preferably 20° C./sec or more.
On the other hand, when the cooling rate exceeds 100° C./sec, the steel sheet is cooled at a cooling rate exceeding 100° C./sec from the surface layer to the inside of the steel sheet, the outermost layer part of the steel sheet is excessively cooled, and a low-temperature transformed structure such as bainite or martensite is formed.
At the time of discharging the hot-rolled coil cooled from 100° C. to room temperature after coiling, microcracks are generated in the low-temperature transformed structure. It is difficult to remove the microcracks in the subsequent pickling step and cold rolling step, and fissures progress from the microcracks as a starting point during cold working, leading to reduced cold workability. Therefore, the cooling rate may preferably be 100° C./sec or less.
Note that the above cooling rate refers to the cooling capacity from the cooling facility at each water injection zone from the point at which the hot-rolled steel sheet after the finish hot rolling is cooled at the water injection zone after passing through the water-free zone to a point at which it is cooled to the coiling target temperature on the ROT, and does not refer to the average cooling rate from the water injection starting point to the temperature at which it is coiled by the coiling device.
Coiling temperature: 400° C. or higher and 550° C. or lower
The coiling temperature is set to 400° C. or higher and 550° C. or lower. When the coiling temperature is lower than 400° C., the austenite which was not transformed before coiling is transformed into hard martensite, cracks are generated in the surface layer of the steel sheet during discharge of the hot-rolled coil, leading to reduced workability. Therefore, the coiling temperature is set to 400° C. or higher, and preferably 430° C. or higher.
On the other hand, when the coiling temperature exceeds 550° C., pearlite having a large lamellar spacing is generated and thick needle-shaped carbides having high thermal stability are formed, and even after the two-step type annealing, needle-shaped carbides remain. Since fissures are generated during cold working with these needle-shaped carbides as a starting point, the coiling temperature is set to 550° C. or lower, and preferably 520° C. or lower.
The hot-rolled coil manufactured under the above conditions is annealed, after pickling, in a two-step type annealing which retains the coil in two temperature ranges. The first-step annealing and the second-step annealing may be either box annealing or continuous annealing. By controlling the stability of carbides by the two-step type annealing, the formation of carbides on the ferrite grain boundary and the spheroidization ratio of carbides on the ferrite grain boundary can be enhanced.
The two-step type annealing will be described below.
The first step annealing is carried out in a temperature range of the Aci point or lower to coarsen carbides and enrich alloy elements to increase the thermal stability of carbides. Thereafter, the temperature is raised to a range from AC1 point or higher to A3 point or lower to generate austenite in the structure.
Thereafter, by gradual cooling, the austenite is transformed into ferrite and the carbon concentration in the austenite is increased. By proceeding slow cooling, carbon atoms are adsorbed to the carbides remaining in the austenite, and thus the carbide and austenite come to cover the grain boundary of the ferrite. Finally it becomes possible to form a structure in which many spheroidized carbides are present in the grain boundary of the ferrite.
When the residual carbides are small in quantity while maintaining the temperature range of AC1 point or higher to A3 point or lower, pearlite, rod-shaped carbides and plate-like carbides are produced during cooling. When these pearlite, rod-shaped carbides and plate-like carbides are produced, the workability of the steel sheet is remarkably deteriorated. Therefore, increasing the number of residual carbides in the temperature range from AC1 point or higher to A3 point or lower is an important factor to enhance the workability of the steel sheet.
By using a steel sheet structure obtained under the above hot rolling condition, the thermal stability of carbides at a temperature of AC1 point or lower can be secured. Therefore, an increase in the number of residual carbides in the temperature range from AC1 point or higher to A3 point or lower can be targeted.
Hereinafter, an annealing condition for the two-step type annealing will be described.
First step annealing
Temperature range: 650° C. or higher and 720° C. or lower
Retention time: 3 hours or longer and 60 hours or shorter
In the first step annealing, the annealing temperature is set to 650° C. or higher and 720° C. or lower. When the annealing temperature of the first step is lower than 650° C., the stability of the carbide becomes insufficient and it becomes difficult to allow the carbide to remain in the austenite in the second step annealing. Therefore, the temperature of the first step annealing is set to 650° C. or higher, and preferably 670° C. or higher.
On the other hand, when the temperature of the first step annealing exceeds 720° C., austenite is generated before enhancing the stability of the carbide, which makes it difficult to control the required change in the structure. Therefore, the first step annealing temperature is set to 720° C. or lower, and preferably 700° C. or lower.
The retention time at the first step is 3 hours or longer and 60 hours or shorter. When the retention time is lower than 3 hours, the stability of the carbide is insufficient and it becomes difficult to allow the carbide to remain at the second step annealing. Therefore, the retention time of the first step is set to 3 hours or longer. On the other hand, when the retention time of the first step exceeds 60 hours, improvement of the stability of the carbide cannot be expected and furthermore the productivity is lowered. Therefore, the retention time of the first step is set to 60 hours or shorter, and preferably 55 hours or shorter.
The annealing atmosphere is not limited to a specific atmosphere. For example, it may be either a nitrogen atmosphere having a nitrogen content of 95% or more, a hydrogen atmosphere having a hydrogen content of 95% or more, or an atmospheric atmosphere.
Second step annealing
Temperature range: 725° C. or higher and 790° C. or lower
Retention time: 3 hours or longer and 50 hours or shorter
In the second step annealing, the annealing temperature is set to 725° C. or higher and 790° C. or lower. When the second-step annealing temperature is lower than 725° C., the amount of austenite produced is small and the number ratio of carbides on the ferrite grain boundary is lowered. Therefore, the second-step annealing temperature is set to 725° C. or higher, and preferably 745° C. or higher.
On the other hand, when the second-step annealing temperature exceeds 790° C., it becomes difficult to allow the carbide to remain in the austenite and to control the required structure change. Therefore, the second-step annealing temperature is set to 790° C. or lower, and preferably 770° C. or lower.
The retention time of the second step is set to 3 hours or longer and 50 hours or shorter. When the retention time of the second step is less than 3 hours, the amount of austenite produced is small, dissolution of the carbide in the ferrite grains is insufficient, and it becomes difficult to increase the number ratio of carbides on the ferrite grain boundary. Therefore, the retention time of the second step is set to 3 hours or longer, and preferably 5 hours or longer.
On the other hand, when the retention time of the second step exceeds 50 hours, it becomes difficult to allow the carbide to remain in the austenite. Therefore, the retention time of the second step is set to 50 hours or shorter, and preferably is 46 hours or shorter.
The annealing atmosphere is not limited to a specific atmosphere. For example, it may be either a nitrogen atmosphere having a nitrogen content of 95% or more, a hydrogen atmosphere having a hydrogen content of 95% or more, or an atmospheric atmosphere.
After completion of the two-step type annealing, the hot-rolled steel sheet is cooled, whereupon it is cooled to 650° C. at a cooling rate of 1° C./hour or more to 30° C./hour or less.
Cooling rate to a temperature of 650° C. or lower: 1° C./hour or more and 30° C./hour or less
Since the temperature range for controlling the structure change by slow cooling is sufficient up to 650° C., it is only necessary to control the cooling rate in the temperature range up to 650° C. After reaching a temperature of 650° C. or lower, it may be cooled to room temperature within the above range without controlling the cooling rate.
It may be preferable that the cooling rate is slow in order to gradually cool the austenite produced in the second step annealing to transform into ferrite and allow carbon to be adsorbed to the carbides remaining in the austenite. However, when the cooling rate is less than 1° C./hour, the time required for cooling increases and the productivity decreases. Therefore, the cooling rate is 1° C./hour or more, and preferably 5° C./hour.
On the other hand, when the cooling rate exceeds 30° C./hour, austenite transforms to pearlite, the hardness of the steel sheet increases, the cold forgeability deteriorates, and the impact resistance property of the steel sheet after carburizing quenching and tempering decreases. Therefore, the cooling rate is set to 30° C./hour or less, and preferably 26° C./hour or less.
Further, according to the inventive production method, a steel sheet with excellent cold workability during forming can be produced in which the ingredient composition is, in terms of % by mass, comprising: C: 0.10 to 0.40%, Si: 0.01 to 0.30%, Mn: 0.30 to 1.00% , P: 0.0001 to 0.020%, S: 0.0001 to 0.010%, and Al: 0.001 to 0.10%, the balance being Fe and unavoidable impurities, the metal structure is substantially composed of ferrite and spheroidized carbides, and (a) the ratio of the number of carbides at the ferrite grain boundary to the number of carbides in the ferrite grain exceeds 1, (b) the ferrite grain size is 5μm or more and 50 μm or less, (c) the in-plane anisotropy |Δr| of the r value standardized according to JIS Z 2254 is 0.2 or less, (d) the Vickers hardness is 100 HV or more and 150 HV or less, the cross-sectional shrinkage percentage is 40% or more, and the ratio of X-ray diffraction intensity of the {131} <011> orientation at the ½-thickness portion of the steel sheet relative to the X-ray diffraction intensity obtained when a sample having the random orientation distribution of crystal grains in the steel sheet is subjected to X-ray diffraction is 3.0 or less.
The cross-sectional shrinkage percentage is defined by the following formula (1). A large value of this ratio means that the local deformability is high, and as the value of the formula (1) increases, the workability of the steel sheet increases.
Sectional shrinkage percentage (%)=100-(cross-sectional area at tensile fracture/initial cross-sectional area)×100 Equation (1)
As described above, the present invention is characterized in that by rolling control and heat treatment after rolling, a structure in which carbides (that is, cementite) are uniformly dispersed is formed, so that the crystal anisotropy can be eliminated. Therefore, in the present invention, the random intensity ratio of the {311} <011> orientation at the ½ plate thickness portion of the steel sheet can be made 3.0 or less.
Next, examples will be described, but the level of examples is an example of conditions adopted for confirming the feasibility and effectiveness of the present invention, and the present invention is not limited to this one condition example. The present invention can adopt various conditions as long as the object of the present invention is achieved without departing from the gist of the present invention.
In order to investigate the effect of hot rolling conditions, a continuous cast strip (steel ingot) having the ingredient composition shown in Table 1 was subjected to hot rolling under the conditions shown in Table 2 to produce a hot-rolled coil having a thickness of 3.0 mm. Incidentally, the steel type described as “Developed steel” in the column of “Remarks” in Table 1 has a composition included in the composition range of the steel sheet according to the present invention. Also, the steel type described as “Comparative steel” in the column of “Remarks” in Table 1 has a composition outside the composition range of the steel sheet according to the present invention. In addition, the ingredients that do not satisfy the composition conditions of the steel sheet according to the present invention are underlined.
A sample for characterization was prepared as follows: a hot-rolled coil, after pickling, was placed in a box-type annealing furnace, the atmosphere was controlled to 95% hydrogen-5% nitrogen, the coil was heated from room temperature to 705° C. and was retained for 36 hours to make the temperature distribution uniform in the hot-rolled coil. The coil was then heated to 760° C. and retained at 760° C. for 10 hours, and was then cooled to 650° C. at a cooling rate of 10° C./hour, then furnace-cooled to room temperature to prepare the sample for characterization. The structure of the sample was measured by the method described above.
A
0.618
I
1.180
0.480
K
0.0360
M
1.21
N
1.27
O
0.51
A
Comparative
steel
I
Comparative
steel
K
Comparative
steel
M
Comparative
steel
N
Comparative
steel
O
Comparative
steel
A-1
5.5
0.44
Comparative
steel
D-1
742
5.6
0.48
Comparative
steel
E-1
617
29.7
Comparative
steel
F-1
937
Comparative
steel
I-1
5.0
0.39
Comparative
steel
K-1
31.4
Comparative
steel
M-1
172
34.9
4.6
0.37
Comparative
steel
N-1
152
38.5
Comparative
steel
O-1
38.4
Comparative
steel
R-1
348
Comparative
steel
T-1
671
38.7
Comparative
steel
AC-1
381
Comparative
steel
The cold workability was evaluated using the notched tensile test and the in-plane anisotropy of the r value. In the notched tensile test, a notched tensile test strip was taken from an as-annealed material with a thickness of 3 mm, and a tensile test was performed in the rolling direction to determine the cross-sectional shrinkage percentage, and the local deformability was evaluated. When the cross-sectional shrinkage percentage is 40% or more, it was rated as superior.
Further, the in-plane anisotropy of the r value was rated as superior when the in-plane anisotropy |Δr| of the r value standardized according to JIS Z 2254 of an as-annealed material with a thickness of 3 mm was 0.2 or less.
In order to determine the X-ray diffraction intensity ratio (I1) of {311} <011>, X-ray diffraction with an Mo tube was performed from the center of the plate thickness of each sample followed by an ODF analysis. Based on the results obtained by the ODF analysis, the I1 was determined.
Table 2 shows, for each of the samples prepared, the results of the carbide diameter, the ferrite grain diameter, the Vickers hardness, the ratio of the number of carbides at the ferrite grain boundary relative to the number of carbides in the ferrite grain, the cross-sectional shrinkage percentage, the X-ray diffraction intensity ratio of {311} <011> and in-plane anisotropy. Among the samples in Table 2, those indicated as “Inventive steel” in the Remarks column satisfy the requirements of the steel sheet according to the present invention, and those indicated as “Comparative steel” in the Remarks column do not satisfy the requirements of the steel sheet according to the present invention. In Table 2, the measurement results that do not satisfy the requirements of the steel sheet according to the present invention and the manufacturing conditions that do not satisfy the requirements of the steel sheet manufacturing method according to the present invention are underlined.
As shown in Table 2, in any of the inventive steels B-1, C-1, G-1, H-1, J-1, L-1, P-1, Q-1, S-1, U-1, W-1, X-1, Y-1, Z-1, AA-1, AB-1 and AD-1, the ratio of the number of carbides at the ferrite grain boundary relative to the number of carbides in the ferrite grain exceeds 1, and the Vickers hardness is 150 HV or less. In addition, in any of the inventive steels, the cross-sectional shrinkage percentage exceeds 40% and the in-plane anisotropy |Δr| of the r value is 0.2 or less. Thus, they have excellent cold workability. Furthermore, since it was confirmed that scale scratches were not generated on the steel sheet surface in any of the inventive steels, these steels can be suitably used for cold working.
On the other hand, in the Comparative steel A-1, since the Al content is high and the A3 point decreased, recrystallization during finish hot rolling was inhibited and |Δr| deteriorated. Thus, the cold workability is low. In the Comparative steel I-1, the contents of Mo and Cr are high, recrystallization during finish hot rolling was inhibited, and |Δr| deteriorated. In the comparative steels K-1 and N-1, the content of S or Mn is high, coarse MnS was formed in the steel, and the cold workability is low. In the Comparative steel M-1, the content of Si was high and hardness increased, and thus cold workability is low. Also, in the Comparative steel M-1, since the A3 point rose, recrystallization during finish hot rolling was hindered and |Δr| deteriorated.
In the Comparative steel O-1, C is high, the volume fraction of carbides increased, a large amount of cracks as the starting point of fractures were generated, and the cross-sectional shrinkage percentage was low. Thus, the cold workability is low. In the Comparative steel D-1, the finish temperature of hot rolling was low and the productivity decreased. In the Comparative steel F-1, the finish temperature of hot rolling was high, and scale scratches were generated on the surface of the steel sheet.
In the Comparative steels R-1 and AC-1, the coiling temperature of hot rolling was low, the low-temperature transformation structure such as bainite and martensite increased resulting in brittled steel, and breaks frequently occurred when the hot-rolled coil was discharged resulting in a decrease in productivity. In the Comparative steels E-1 and T-1, the coiling temperature of hot rolling was high, thick pearlite with lamellar spacing and needle-shaped coarse carbides with high thermal stability were produced in the hot rolled structure. Since these carbides remained in the steel sheet even after the two-step type annealing, the cross-sectional shrinkage percentage was low and thus the cold workability is low.
Subsequently, in order to investigate the effect of annealing conditions, steel strips (slabs) having the ingredient composition shown in Table 1 were heated at 1240° C. for 1.8 hours and then subjected to hot rolling. After completing finish hot rolling at 890° C., they were cooled to 520° C. at a cooling rate of 45° C./sec on ROT and coiled at 510° C. to produce a hot-rolled coil with a thickness of 3.0 mm. And under the conditions shown in Table 3, a hot-rolled sheet-annealed sample with a thickness of 3.0 mm was prepared.
For each of the samples prepared, the carbide diameter, the ferrite grain diameter, the Vickers hardness, the ratio of the number of carbides at the ferrite grain boundary relative to the number of carbides in the ferrite grain, the cross-sectional shrinkage percentage, the X-ray diffraction intensity ratio of {311} <011> and the in-plane anisotropy were determined in the same manner as the inventive steels and the comparative steels in Table 2. The results are shown in Table 3.
A-2
E-2
34
H-2
I-2
J-2
811
K-2
M-2
N-2
O-2
Q-2
T-2
741
Y-2
54
Z-2
706
AB-2
68
AC-2
637
A-2
4.6
0.38
Comp. steel
E-2
161
38.6
Comp. steel
H-2
0.4
Comp. steel
I-2
3.8
0.28
Comp. steel
J-2
151
32.6
Comp. steel
K-2
20.7
Comp. steel
M-2
183
38.6
Comp. steel
N-2
153
39.6
Comp. steel
O-2
37.4
Comp. steel
Q-2
37.6
Comp. steel
T-2
34.7
Comp. steel
Y-2
29.2
Comp. steel
Z-2
0.5
Comp. steel
AB-2
Comp. steel
AC-2
31.2
Comp. steel
As shown in Table 3, in any of the inventive steels B-2, C-2, D-2, F-2, G-2, L-2, P-2, R-2, S-2, U-2, W-2, X-2, AA-2 and AD-2, the ratio of the number of carbides at the ferrite grain boundary to the number of carbides in the ferrite grain exceeds 1, and the Vickers hardness is 150 HV or less. In addition, in any of the inventive steels, the cross-sectional shrinkage percentage exceeds 40% and the in-plane anisotropy |Δr| of the r value is 0.2 or less. Thus, they have excellent cold workability.
On the other hand, in the Comparative steel A-2, since the Al content is high and the A3 point decreased, recrystallization during finish hot rolling was inhibited and |Δr| deteriorated. Thus, the cold workability is low. In the Comparative steel 1-2, the contents of Mo and Cr are high, recrystallization during finish hot rolling was inhibited, and |Δr| deteriorated. In the comparative steels K-2 and N-2, the content of S or Mn is high, coarse MnS was formed in the steel. Thus, the cold workability deteriorated. In the Comparative steel M-2, the content of Si was high and hardness increased. Thus, the cold workability is low. Also, in the Comparative steel M-2, since the A3 point decreased, recrystallization during finish hot rolling was hindered and |Δr| deteriorated.
In the Comparative steel O-2, C is high, the volume fraction of carbides increased, a large amount of cracks as the starting point of fracture were generated, and the cross-sectional shrinkage percentage was low. Thus, the cold workability is low.
In the Comparative steel AC-2, since the annealing temperature in the first-step annealing during the two-step type box annealing is low, the treatment of carbide coarsening at the Ac1 temperature or lower is insufficient, and the thermal stability of carbides is insufficient, thus the carbides remaining at the second step of annealing decreases, the pearlite transformation cannot be suppressed in the structure after the slow cooling, and the cross-sectional shrinkage percentage is low. Thus, the cold workability is low.
In the Comparative steel T-2, since the annealing temperature in the first step annealing during the two-step type box annealing is high, austenite is generated during annealing and the stability of carbide cannot be increased, so that carbides remaining during the second step annealing decrease, and pearlite transformation cannot be suppressed in the structure after the slow cooling, and the cross-sectional shrinkage percentage is low. Thus, the cold forgeability is low.
In the Comparative steel Q-2, since the retention time in the first step annealing during annealing of the two-step type is short, the treatment of the carbide coarsening at the Ac1 temperature or lower is insufficient, and the thermal stability of the carbide is insufficient, and thus the carbide remaining at the second step of annealing decreases and the pearlite transformation cannot be suppressed in the structure after the slow cooling, and the cross-sectional shrinkage percentage is low. And thus the cold workability is low. In the comparative steel AB-2, the retention time during the first stage box annealing of the two-step type is long and the productivity is low.
In the Comparative steel Z-2, since the annealing temperature during the second-step annealing during the two-step type box annealing is low, and the amount of austenite produced is small, so that the proportion of the number of carbides in the grain boundary cannot be increased. Thus, the cold workability is low. In the Comparative steel J-2, the annealing temperature during the second-step annealing during the two-step type annealing is high, the amount of the carbide remaining is decreased due to the promoted dissolution of carbides, and pearlite transformation cannot be suppressed in the structure after the slow cooling, the Vickers hardness is too high, and the cross-sectional shrinkage percentage is low. Thus, the cold forgeability is low.
In the Comparative steel H-2, since the annealing temperature during the second-step annealing during the two-step type annealing is low, and the amount of austenite produced is small, so that the proportion of the number of carbides in the grain boundary cannot be increased. Thus, the cold workability is low. In the Comparative steel Y-2, the retention time during the second-step annealing during the two-step type annealing is long, the amount of carbides remaining is decreased due to the promoted dissolution of carbides, and pearlite transformation cannot be suppressed in the structure after the slow cooling, and the cross-sectional shrinkage percentage is low. Thus, the cold forgeability is low. In the Comparative steel E-2, the cooling rate from the second-step annealing during the two-step type annealing to 650° C. is fast, pearlite transformation occurred during cooling, the Vickers hardness is too high, and the cross-sectional shrinkage percentage is low. Thus, the cold workability is low.
In any of the comparative steels A-1, D-1, I-1, M-1, A-2 and I-2, the X-ray diffraction intensity ratio of {311} <011> is greater than 3.0. In these comparative steels, the in-plane anisotropy |Δr| exceeds 0.2, and thus the cold workability is low. As described above, by performing analysis by X-ray diffraction on a plane parallel to the plate surface at the ½ plate thickness portion of the hot-rolled steel sheet, the degree of plastic anisotropy such as the in-plane anisotropy |Δr| or the quality of cold workability of the hot-rolled steel sheet to be cold worked can be determined before cold working.
As described above, according to the present invention, a steel sheet with excellent cold workability during forming can be manufactured and provided. The steel sheet of the present invention is a steel sheet suitable as a material for automotive parts, blades, and other mechanical parts manufactured through processing steps such as punching, bending, pressing, etc. Therefore, the present invention has excellent industrial applicability.
| Number | Date | Country | Kind |
|---|---|---|---|
| 2015-081101 | Apr 2015 | JP | national |
| Filing Document | Filing Date | Country | Kind |
|---|---|---|---|
| PCT/JP2016/061608 | 4/8/2016 | WO | 00 |
