Patent Application
20170306434
The present invention relates to a high-carbon steel sheet with improved formability and a method of manufacturing the same.
A high-carbon steel sheet is used for various steel products, which are a driving system component for automobile such as a chain, a gear and a clutch, a saw, a knife, and others. When the steel products are manufactured, forming and heat treatments of a high-carbon steel sheet are performed. As the forming, punching, tensile forming, compressing, shearing, and so on are performed, and as the heat treatment, quenching, tempering, carburizing, nitriding, soft-nitriding, and so on are performed. A strength of a high-carbon steel sheet is higher than that of a mild steel sheet, and therefore a metal mold used for forming of a high-carbon steel sheet is more easily worn than a metal mold used for forming of a mild steel sheet. Further, a high-carbon steel sheet cracks more easily than a mild steel sheet during forming.
For suppressing the wearing of a metal mold, improving lubricity on a surface of a high-carbon steel sheet is effective, and for suppressing the cracking during forming, softening of a high-carbon steel sheet is effective. Thus, some techniques have been proposed aiming at an improvement in lubricity and softening (Patent Literatures 1 to 5).
However, these prior techniques cause a significant increase in cost, and therefore are not preferred.
Although a carbon steel sheet aiming at an improvement in punchability has been described in Patent Literature 6 and a high-carbon steel sheet aiming at an improvement in formability has been described in Patent Literature 7, it is not possible for them to obtain sufficient formability.
Patent Literature 1: Japanese Laid-open Patent Publication No. 2010-174252
Patent literature 2: Japanese Laid-open Patent Publication No, 2009-215612
Patent Literature 3: Japanese Laid-open Patent Publication No. 2011-168842
Patent Literature 4: Japanese Laid-open Patent Publication No. 2010-255066
Patent Literature 5: Japanese Laid-open Patent Publication No, 2000-34542
Patent Literature 6: Japanese Laid-open Patent Publication No. 2000-265240
Patent Literature 7: Japanese Laid-open Patent Publication No. 10-147816
An object of the present invention is to provide a high-carbon steel sheet capable of obtaining excellent formability while avoiding a significant increase in cost, and a method of manufacturing the same.
The present inventors conducted earnest studies repeatedly to solve the above-described problem, and consequently found out that it is important that a high-carbon steel sheet contains a specific amount of B, that a coefficient of micro-friction of ferrite on a surface is a specific one, and that form of cementite is a specific one. Further, it was also found out that, in order to manufacture such a high-carbon steel sheet, it is important to perform hot-rolling and annealing under specific conditions while assuming hot-rolling and annealing as what is called a consecutive process. Then, the inventors of the present application devised the following various aspects of the invention based on these findings.
(1)
A high-carbon steel sheet, including:
a chemical composition represented by, in mass %:
C: 0.30% to 0.70%,
Si: 0.07% to 1.00%,
Mn: 0.20% to 3.00%,
Ti: 0.010% to 0.500%,
Cr: 0.01% to 1.50%,
B: 0.0004% to 0.0035%,
F: 0.025% or less,
Al: 0.100% or less,
S: 0.0100% or less,
N: 0.010% or less,
Cu: 0.500% or less,
Nb: 0.000% to 0.500%,
Mo: 0.000% to 0.500%,
V: 0.000% to 0.500%,
W: 0.000% to 0.500%,
Ta: 0.000% to 0.500%,
Ni: 0.000% to 0.500%,
Mg: 0.000% to 0.500%,
Ca: 0.000% to 0.500%,
Y: 0.000% to 0.500%,
Zr: 0.000% to 0.500%,
La: 0.000% to 0.500%,
Ce: 0.000% to 0.500%, and
balance: Fe and impurities; and
a structure represented by:
a spheroidized ratio of cementite: 80% or more; and
an average diameter of cementite: 0.3 μm to 2.2 μm, wherein
a coefficient of micro-friction of ferrite on a surface of the steel sheet is less than 0.5.
(2)
The high-carbon steel sheet according to (1), wherein
in the chemical composition,
Nb: 0.001% to 0.500%,
Mo: 0.001% to 0.500%,
V: 0.001% to 0.500%,
W: 0.001% to 0.500%,
Ta: 0.001% to 0.500%,
Ni: 0.001% to 0.500%,
Mg: 0.001% to 0.500%,
Ca: 0.001% to 0.500%,
Y: 0.001% to 0.500%,
Zr: 0.001% to 0.500%,
La: 0.001% to 0.500%, or
Ce: 0.001% to 0.500%, or
any combination thereof is satisfied.
(3)
A method of manufacturing a high-carbon steel sheet, including:
hot-rolling of a slab so as to obtain a hot-rolled steel sheet;
pickling of the hot-rolled steel sheet; and
annealing of the hot-rolled steel sheet after the pickling,
the slab including a chemical composition represented by, in mass %:
C: 0.30% to 0.70%,
Si: 0.07% to 1.00%,
Mn: 0.20% to 3.00%,
Ti: 0.010% to 0.500%,
Cr: 0.01% to 1.50%,
B: 0.0004% to 0.0035%,
P: 0.025% or less,
Al: 0.100% or less,
S: 0.0100% or less,
N: 0.010% or less,
Cu: 0.500% or less,
Nb: 0.000% to 0.500%,
Mo: 0.000% to 0.500%,
V: 0.000% to 0.500%,
W: 0.000% to 0.500%,
Ta: 0.000% to 0.500%,
Ni: 0.000% to 0.500%,
Mg: 0.000% to 0.500%,
Ca: 0.000% to 0.500%,
Y: 0.000% to 0.500%,
Zr: 0.000% to 0.500%,
La: 0.000% to 0.500%,
Ce: 0.000% to 0.500%, and
balance: Fe and impurities, wherein
in the hot-rolling,
the slab is heated at a temperature of 1000° C. or more and less than 1150° C.,
a finish rolling temperature is 830° C. or more and 950° C. or less, and
a coiling temperature is 450° C. or more and 700° C. or less, and
the annealing comprises:
retaining the hot-rolled steel sheet at a temperature of 730° C. or more and 770° C. or less for 3 hours or more and 60 hours or less; and
then cooling the hot-rolled steel sheet down to 650° C. at a cooling rate of 1° C./hr or more and 60° C./hr or less.
(4)
The method of manufacturing the high-carbon steel sheet according to (3), wherein
in the chemical composition,
Nb: 0.001% to 0.500%,
Mo: 0.001% to 0.500%,
V: 0.001% to 0.500%,
W: 0.001% to 0.500%,
Ta: 0.001% to 0.500%,
Ni: 0.001% to 0.500%,
Mg: 0.001% to 0.500%,
Ca: 0.001% to 0.500%,
Y: 0.001% to 0.500%,
Zr: 0.001% to 0.500%,
La: 0.001% to 0.500%, or
Ce: 0.001% to 0.500%, or
any combination thereof is satisfied.
According to the present invention, a B content, a coefficient of micro-friction of ferrite on a surface and others are appropriate, thereby making it possible to obtain excellent formability while avoiding a significant increase in cost.
Hereinafter, there will be explained an embodiment of the present invention.
First, chemical compositions of a high-carbon steel sheet according to the embodiment of the present invention and a slab (steel ingot) used for manufacturing the same will be explained. Although details will be described later, the high-carbon steel sheet according to the embodiment of the present invention is manufactured by going through hot-rolling of the slab, annealing, and the like. Accordingly, the chemical compositions of the high-carbon steel sheet and the slab are appropriate for the above-stated processes in addition to properties of the high-carbon steel sheet. In the following description, “%” being a unit of content of each element contained in the high-carbon steel sheet and the slab used for manufacturing the same means “mass %” unless otherwise mentioned. The high-carbon steel sheet according to the embodiment and the slab used for manufacturing the same include a chemical composition represented by C: 0.30% to 0.70%, Si: 0.07% to 1.00%, Mn: 0.20% to 3.00%, Ti: 0.010% to 0.500%, Cr: 0.01% to 1.50%, B: 0.0004% to 0.0035%, P: 0.025% or less, Al: 0.100% or less, S: 0.0100% or less, N: 0.010% or less, Cu: 0.500% or less, Nb: 0.000% to 0.500%, Mo: 0.000% to 0.500%, V: 0.000% to 0.500%, W: 0.000% to 0.500%, Ta: 0.000% to 0.500%, Ni: 0.000% to 0.500%, Mg: 0.000% to 0.500%, Ca: 0.000% to 0.500%, Y: 0.000% to 0.500%, Zr: 0.000% to 0.500%, La: 0.000% to 0.500%, Ce: 0.000% to 0.500%, and balance: Fe and impurities. As the impurities, ones contained in raw materials such as ore and scrap, and ones contained during a manufacturing process are exemplified. For example, when scrap is used as a raw material, Sn, Sb or As or any combination thereof may mix in by 0.003% or more. If the content is 0.03% or less, none of them hinder the effect of the embodiment, and thus may be tolerated as impurities. Further, O may be tolerated as an impurity up to 0.0025%. O forms oxide, and when oxides aggregate and become coarse, sufficient formability is not obtained. Therefore, the O content is the lower the better. However, it is technically difficult to decrease the O content to less than 0.0001%.
(C: 0.30% to 0.70%)
C bonds to Fe to form cementite having a small friction coefficient, and thus is an important element when securing macro-lubricity of the high-carbon steel sheet. When the C content is less than 0.30%, the amount of cementite is insufficient, resulting in that sufficient lubricity cannot be obtained and adhesion to a metal mold occurs during forming. Thus, the C content is 0.30% or more, and preferably 0.35% or more. When the C content is greater than 0.70%, the amount of cementite is excessive, resulting in that a crack originating from the cementite occurs easily during forming. Thus, the C content is 0.70% or less, and preferably 0.65% or less.
(Si: 0.07% to 1.00%)
Si operates as a deoxidizes, and is effective for suppressing excessive coarsening of cementite during annealing. When the Si content is less than 0.07%, the effect by the above-described operation cannot be obtained sufficiently. Thus, the Si content is 0.07% or more, and preferably 0.10% or more. When the Si content is greater than 1.00%, the ductility of ferrite is low and a crack originating from transgranular fracture of ferrite occurs easily during forming. Thus, the Si content is 1.00% or less, and preferably 0.80% or less.
(Mn: 0.20% to 3.00%)
Mn is important for controlling pearlite transformation. When the Mn content is less than 0.20%, the effect by the above-described operation cannot be obtained sufficiently. That is, when the Mn content is less than 0.20%, pearlite transformation occurs in cooling after dual-phase annealing and a spheroidized ratio of cementite becomes insufficient. Thus, the Mn content is 0.20% or more, and preferably 0.25% or more. When the Mn content is greater than 3.00%, the ductility of ferrite is low and a crack originating from transgranular fracture of ferrite occurs easily during forming. Thus, the Mn content is 3.00% or less, and preferably 2.00% or less.
(Ti: 0.010% to 0.500%)
Ti forms a nitride in molten steel, and effective for preventing formation of BN. When the Ti content is less than 0.010%, the effect by the above-described operation cannot be obtained sufficiently. Thus, the Ti content is 0.010% or more, and preferably 0.040% or more. When the Ti content is greater than 0.500%, a crack originating from a coarse oxide of Ti occurs easily during forming. This is because during continuous casting, coarse oxides of Ti are formed to get involved inside the slab. Thus, the Ti content is 0.500% or less, and preferably 0.450% or less.
(Cr: 0.01% to 1.50%)
Cr has a high affinity with N, effective for suppressing formation of BN, and effective also for controlling pearlite transformation. When the Cr content is less than 0.01%, the effect by the above-described operation cannot be obtained sufficiently. Thus, the Cr content is 0.01% or more, and preferably 0.05% or more. When the Cr content is greater than 1.50%, spheroidizing of cementite during annealing is hindered and coarsening of cementite is suppressed drastically. Thus, the Cr content is 1.50% or less, and preferably 0.90% or less.
(B: 0.0004% to 0.0035%)
B lowers the coefficient of micro-friction of ferrite on the surface of the high-carbon steel sheet. B segregates to and concentrates at an interface between ferrite and cementite during later-described annealing and suppresses peeling at the interface during forming, and B is also effective for preventing a crack. When the B content is less than 0.0004%, the effect by the above-described operation cannot be obtained sufficiently. Thus, the B content is 0.0004% or more, and preferably 0.0008% or more. When the B content is greater than 0.0035%, a crack originating from boride such as carbide of Fe and B occurs easily during forming. Thus, the B content is 0.0035% or less, and preferably 0.0030% or less.
(P: 0.025% or less)
P is not an essential element and is contained as an impurity in the steel sheet, for example. P strongly segregates to the interface between ferrite and cementite, and thereby the segregation of B to the interface is hindered and peeling at the interface is caused. Therefore, the P content is the smaller the better. When the P content is greater than 0.025%, adverse effects are particularly prominent. Thus, the P content is 0.025% or less. Decreasing the P content takes refining cost, and it requires a considerable refining cost to decrease the P content to less than 0.0001%. Thus, the P content may be 0.0001% or more.
(Al: 0.100% or less)
Al operates as a deoxidizer in steelmaking and is effective for fixing N, but is not an essential element of the high-carbon steel sheet and is contained as an impurity in the steel sheet, for example. When the Al content is greater than 0.100%, the ductility of ferrite is low and a crack originating from transgranular fracture of ferrite occurs easily during forming, and strength is excessive to cause an increase in forming load. Thus, the Al content is set to 0.100% or less. When the Al content of the high-carbon steel sheet is less than 0.001%, fixation of N sometimes may be insufficient. Thus, the Al content may be 0.001% or more.
(S: 0.0100% or less)
S is not an essential element and is contained as an impurity in the steel sheet, for example. S forms coarse non-metal inclusions such as MnS to impair formability. Therefore, the S content is the smaller the better. When the S content is greater than 0.0100%, adverse effects are particularly prominent. Thus, the S content is 0.0100% or less. Decreasing the S content takes refining cost, and it requires a considerable refining cost to decrease the S content to less than 0.0001%. Thus, the S content may be 0.0001% or more.
(N: 0.010% or less)
N is not an essential element and is contained as an impurity in the steel sheet, for example. N lowers an amount of solid-solution B due to formation of BN so as to cause adhesion to the metal mold, cracking during forming, and the like. Therefore, the N content is the smaller the better. When the N content is greater than 0.010%, adverse effects are particularly prominent. Thus, the N content is set to 0.010% or less. Decreasing the N content takes refining cost, and it requires a considerable refining cost to decrease the N content to less than 0.001%. Thus, the N content may be 0.001% or more.
(Cu: 0.000% to 0.500%)
Cu is not an essential element and is mixed from scrap or the like to be contained as an impurity in the steel sheet, for example. Cu causes an increase in strength and brittleness in hot working. Therefore, the Cu content is the smaller the better. When the Cu content is greater than 0.500%, adverse effects are particularly prominent. Thus, the Cu content is 0.500% or less. Decreasing the Cu content takes refining cost, and it requires a considerable refining cost to decrease the Cu content to less than 0.001%. Thus, the Cu content may be 0.001% or more.
Nb, Mo, V, W, Ta, Ni, Mg, Ca, Y, Zr, La, and Ce are not essential elements, and are optional elements that may be appropriately contained in the high-carbon steel sheet and the slab up to a specific amount.
(Nb: 0.000% to 0.500%)
Nb forms a nitride and is effective for suppressing formation of BN. Thus, Nb may be contained. However, when the Nb content is greater than 0.500%, the ductility of ferrite is low to make it impossible to obtain sufficient formability. Thus, the Nb content is 0.500% or less. In order to securely obtain the effect by the above-described operation, the Nb content is preferably 0.001% or more.
(Mo: 0.000% to 0.500%)
Mo is effective for improving hardenability. Thus, Mo may be contained. However, when the Mo content is greater than 0.500%, the ductility of ferrite is low to make it impossible to obtain sufficient formability. Thus, the Mo content is 0.500% or less. In order to securely obtain the effect by the above-described operation, the Mo content is preferably 0.001% or more.
(V: 0.000% to 0.500%)
V forms a nitride and is effective for suppressing formation of BN similarly to Nb. Thus, V may be contained. However, when the V content is greater than 0.500%, the ductility of ferrite is low to make it impossible to obtain sufficient formability. Thus, the V content is 0.500% or less. In order to securely obtain the effect by the above-described operation, the V content is preferably 0.001% or more.
(W: 0.000% to 0.500%)
W is effective for improving hardenability similarly to Mo. Thus, W may be contained. However, when the N content is greater than 0.500%, the ductility of ferrite is low to make it impossible to obtain sufficient formability. Thus, the W content is 0.500% or less. In order to securely obtain the effect by the above-described operation, the W content is preferably 0.001% or more.
(Ta: 0.000% to 0.500%)
Ta forms a nitride and is effective for suppressing formation of BN similarly to Nb and V. Thus, Ta may be contained. However, when the Ta content is greater than 0.500%, the ductility of ferrite is low to make it impossible to obtain sufficient formability. Thus, the Ta content is 0.500% or less. In order to securely obtain the effect by the above-described operation, the Ta content is preferably 0.001% or more.
(Ni: 0.000% to 0.500%)
Ni is effective for improving toughness and improving hardenability. Thus, Ni may be contained. However, when the Ni content is greater than 0.500%, the coefficient of micro-friction of ferrite is high to cause adhesion to the metal mold easily. Thus, the Ni content is 0.500% or less. In order to securely obtain the effect by the above-described operation, the Ni content is preferably 0.001% or more.
(Mg: 0.000% to 0.500%)
Mg is effective for controlling the form of sulfide. Thus, Mg may be contained. However, Mg forms oxide easily, and when the Mq content is greater than 0.500%, sufficient formability cannot be obtained due to a crack originating from the oxide. Thus, the Mg content is 0.500% or less. In order to securely obtain the effect by the above-described operation, the Mg content is preferably 0.001% or more.
(Ca: 0.000% to 0.500%)
Ca is effective for controlling the form of sulfide similarly to Mg. Thus, Ca may be contained. However, Ca forms oxide easily, and when the Ca content is greater than 0.500%, sufficient formability cannot be obtained due to a crack originating from the oxide. Thus, the Ca content is 0.500% or less. In order to securely obtain the effect by the above-described operation, the Ca content is preferably 0.001% or more.
(Y: 0.000% to 0.500%)
Y is effective for controlling the form of sulfide similarly to Mg and Ca. Thus, Y may be contained. However, Y forms oxide easily, and when the Y content is greater than 0.500%, sufficient formability cannot be obtained due to a crack originating from the oxide. Thus, the Y content is 0.500% or less. In order to securely obtain the effect by the above-described operation, the Y content is preferably 0.001% or more.
(Zr: 0.000% to 0.500%)
Zr is effective for controlling the form of sulfide similarly to Mg, Ca, and Y. Thus, Zr may be contained. However, Zn forms oxide easily, and when the Zr content is greater than 0.500%, sufficient formability cannot be obtained due to a crack originating from the oxide. Thus, the Zr content is 0.500% or less. In order to securely obtain the effect by the above-described operation, the Zr content is preferably 0.001% or more.
(La: 0.000% to 0.500%)
La is effective for controlling the form of sulfide similarly to Mg, Ca, Y, and Zr. Thus, La may be contained. However, La forms oxide easily, and when the La content is greater than 0.500%, sufficient formability cannot be obtained due to a crack originating from the oxide. Thus, the La content is 0.500% or less. In order to securely obtain the effect by the above-described operation, the La content is preferably 0.001% or more.
(Ce: 0.000% to 0.500%)
Ce is effective for controlling the form of sulfide similarly to Mg, Ca, Y, Zr, and La. Thus, Ce may be contained. However, Ce forms oxide easily, and when the Ce content is greater than 0.500%, sufficient formability cannot be obtained due to a crack originating from the oxide. Thus, the Ce content is 0.500% or less. In order to securely obtain the effect by the above-described operation, the Ce content is preferably 0.001% or more.
Thus, Nb, Mo, V, W, Ta, Ni, Mg, Ca, Y, Zr, La and Ce are optional elements, and it is preferred that “Nb: 0.001% to 0.500%,” “Mo: 0.001% to 0.500%,” “V: 0.001% to 0.500%,” “W: 0.001% to 0.500%,” “Ta: 0.001% to 0.500%,” “Ni: 0.001% to 0.500%,” “Mg: 0.001% to 0.500%,” “Ca: 0.001% to 0.500%,” “Y: 0.001% to 0.500%,” “Zr: 0.001% to 0.500%,” “La: 0.001% to 0.500%,” or “Ce: 0.001% to 0.500%,” or any combination thereof be satisfied.
Next, the coefficient of micro-friction of ferrite on the surface of the high-carbon steel sheet according to the embodiment is explained. The coefficient of micro-friction of ferrite on the surface of the high-carbon steel sheet according to the embodiment is less than 0.5.
(Coefficient of Micro-Friction of Ferrite on the Surface: Less than 0.5)
The coefficient of micro-friction of ferrite on the surface closely relates to adhesion of the high-carbon steel sheet to the metal mold during forming. When the coefficient of micro-friction of ferrite is 0.5 or more, micro-adhesion occurs between the high-carbon steel sheet and the metal mold during forming using the metal mold. As a result, when forming such as punching is performed with several thousands to several tens of thousands of shots by using the metal mold, adhesive matters are accumulated on the metal mold during the forming, and a flaw occurs on either the metal mold or the high-carbon steel sheet or on the both and formability deteriorates. Thus, the coefficient of micro-friction of ferrite is less than 0.5. From the viewpoint of formability, the coefficient of micro-friction is the lower the better. The coefficient of micro-friction often tends to be 0.35 or more, though it depends on a method of manufacturing the high-carbon steel sheet and others.
A coefficient of micro-friction may be measured using a nanoindenter. That is, a kinetic friction force F to occur when a diamond indenter loads a normal load P of 10 μN onto a surface of a high-carbon steel sheet and is moved horizontally is obtained. A moving speed then is 1 μm/second, for example. A coefficient of micro-friction μ (kinetic friction coefficient) is calculated by Expression (1) below. “TI-900 TriboIndenter” made by Omicron, Inc. may be used as a nanoindenter, for example.
F=μP Expression(1)
Next, a structure of the high-carbon steel sheet according to the embodiment is explained. The high-carbon steel sheet according to the embodiment includes a structure represented by a spheroidized ratio of cementite: 80% or more and an average diameter of cementite: 0.3 μm to 2.2 μm.
(Spheroidized Ratio of Cementite: 80% or More)
Stress concentration sometimes originates from cementite during forming, and stress is likely to concentrate locally in acicular cementite particularly. When the spheroidized ratio of cementite is less than 80%, acicular cementite, in which stress is likely to concentrate, is contained in large amounts, and thus stress concentration occurs easily and peeling occurs at an interface between ferrite and cementite, resulting in that sufficient formability cannot be obtained. Thus, the spheroidized ratio of cementite is 80% or more, and preferably 85% or more. From the viewpoint of formability, the spheroidized ratio of cementite is preferred to be as higher as possible, and may be 100%. However, when the spheroidized ratio of cementite is attempted to become 100%, productivity could decrease, and the spheroidized ratio of cementite is preferably 80% or more and less than 100% from the viewpoint of productivity.
(Average Diameter of Cementite: 0.3 μm to 2.2 μm)
The average diameter of cementite closely relates to the degree of the stress concentration to cementite. When the average diameter of cementite is less than 0.3 μm, an Orowan loop is formed by dislocation occurred during forming with respect to cementite, and thereby a dislocation density in the vicinity of cementite increases and voids occur. Thus, the average diameter of cementite is 0.3 μm or more, and preferably 0.5 μm or more. When the average diameter of cementite is greater than 2.2 μm, dislocations occurred during forming are accumulated in large amounts, local stress concentration is generated and a crack occurs. Thus, the average diameter of cementite is 2.2 μm or less, and preferably 2.0 μm or less.
The spheroidized ratio and the average diameter of cementite may be measured by structure observation using a scanning electron microscope. In preparing of a sample for structure observation, an observation surface is mirror finished by wet polishing with an emery paper and polishing with diamond abrasive grains having a size of 1 μm, then the observation surface is etched with an etching solution of 3 vol % of nitric acid and 97 vol % of alcohol. An observation magnification is between 3000 times to 10000 times, for example, 10000 times, 16 visual fields where 500 or more grains of cementite exist on the observation surface are selected, and structure images of them are taken. Then, an area of each cementite in the structure image is measured by using image processing software. “Win ROOF” made by MITANI Corporation may be used as an image processing software, for example. Any cementite grain having an area of 0.01 μm2 or less is excluded from the target of evaluation in order to suppress an influence of measurement error by noise in the measuring. Then, the average area of cementite as an evaluation target is obtained, and the diameter of a circle with which this average area can be obtained is obtained, thereby taking this diameter as the average diameter of cementite. The average area of cementite is a value obtained by dividing the total area of cementite as the evaluation target by the number of grains of cementite in question. Further, any cementite having a ratio of major axis length to minor axis length of 3 or more is assumed as an acicular cementite, any cementite having the ratio of less than 3 is assumed as a spherical cementite grain, and a value obtained by dividing the number of spherical cementite by the number of all cementite is taken as the spheroidized ratio of cementite.
Next, a method of manufacturing the high-carbon steel sheet according to the embodiment is explained. The manufacturing method includes hot-rolling of a slab including the above chemical composition so as to obtain a hot-rolled steel sheet, pickling of the hot-rolled steel sheet, and thereafter annealing of the hot-rolled steel sheet. In the hot-rolling, the slab is heated at a temperature of 1000° C. or more and less than 1150° C., a finish rolling temperature is 830° C. or more and 950° C. or less, and a coiling temperature is 450° C. or more and 700° C. or less. In the annealing, the hot-rolled steel sheet is retained at a temperature of 730° C. or more and 770° C. or less for 3 hours or more and 60 hours or less, and then, the hot-rolled steel sheet is cooled down to 650° C. at a cooling rate of 1° C./hr or more and 60° C./hr or less. An atmosphere of the annealing may be one containing hydrogen by 75 vol % or more at a temperature higher than 400° C., for example, but is not limited to that.
Here, an outline of changes in the steel sheet from the hot-rolling to the cooling is explained.
In an example illustrated in
At a time tA after completion of the slab heating S11, B atoms 13 segregate to an interface between austenite 12 and austenite 12, as illustrated in
According to the manufacturing method, the surface 15 of the steel sheet is covered with many of the B atoms 13 bonded to each other by the covalent bonding 14, and thereby the coefficient of micro-friction of ferrite on the surface 15 can be less than 0.5.
(Slab Heating Temperature: 1000° C. or More and Less than 1150° C.)
When the slab heating temperature is higher than 1150° C., oxygen easily diffuses into the inside of the slab from the surface of the slab to bond to B in the slab. That is, as illustrated in
(Finish Rolling Temperature: 830° C. or More and 950° C. or Less)
When the finish rolling temperature is higher than 950° C., coarse scales are generated until completion of coiling on a run out table (ROT), for example. The coarse scales can be removed by pickling, but traces of large irregularities are left, resulting in that adhesion to the metal mold sometimes occurs during forming due to the traces. Further, when coarse scales are generated, irregular flaw is caused on the surface of the steel sheet in the coiling, resulting in that due to the flaw, adhesion to the metal mold sometimes occurs during forming. Thus, the finish rolling temperature is 950° C. or less, and preferably 940° C. or less. When the finish rolling temperature is lower than 830° C., adhesiveness of scales generated until completion of coiling to the steel sheet is extremely high, thus making it difficult to remove the scales by pickling. The scales may be removed by performing strong pickling, but the strong pickling makes the surface of the steel sheet rough, resulting in that adhesion to the metal mold sometimes occurs during forming. Further, when the finish rolling temperature is lower than 830° C., recrystallization of austenite is not completed by the coiling, so that anisotropy of the hot-rolled steel sheet increases. The anisotropy of the hot-rolled steel sheet is carried over even after annealing, and thus sufficient formability cannot be obtained. Thus, the finish rolling temperature is 830° C. or more, and preferably 840° C. or more.
(Coiling Temperature: 450° C. or More and 700° C. or Less)
When the coiling temperature is higher than 700° C., coarse lamellar pearlite is formed in the hot-rolled steel sheet to hinder spheroidizing of cementite during annealing, resulting in that the spheroidized ratio of 80% or more cannot be obtained. Thus, the coiling temperature is 700° C. or less. Further, when the coiling temperature is higher than 570° C., coarse scales are generated until completion of coiling. Therefore, adhesion to the metal mold sometimes occurs during forming for a reason similar to the case where the finish rolling temperature is higher than 950° C. Thus, the coiling temperature is preferably 570° C. or less, and further preferably 550° C. or less. When the coiling temperature is lower than 450° C., adhesiveness of scales generated until completion of coiling to the steel sheet is extremely high, thus making it difficult to remove the scales by pickling. The scales may be removed by performing strong pickling, but the strong pickling makes the surface of the steel sheet rough, resulting in that adhesion to the metal mold sometimes occurs during forming. Further, when the coiling temperature is lower than 450° C., the hot-rolled steel sheet becomes brittle and the hot-rolled steel sheet may crack when a coil is uncoiled in pickling, resulting in that a sufficient yield cannot be obtained. Thus, the coiling temperature is 450° C. or more, and preferably 460° C. or more.
A rough-rolled bar may be heated near an inlet of a finishing mill in order to ensure qualities in a longitudinal direction and a width direction of a hot-rolled coil obtained by coiling (to reduce variation of quality or the like). An apparatus to be used for the heating and a method of the heating are not limited in particular, but heating by high-frequency induction heating is desirably performed. A preferred temperature range of the heated rough-rolled bar is between 850° C. and 1100° C. Temperatures less than 850° C. are close to a transformation temperature from austenite to ferrite, and therefore, when the temperature of the heated rough bar is lower than 850° C., heat generation and heat absorption due to transformation and reverse transformation sometimes occur, resulting in that temperature controlling is unstable and it is difficult to uniformize a temperatures in the longitudinal direction and the width direction of the hot-rolled coil. Therefore, if the rough-rolled bar is heated, the heating temperature is preferably 850° C. or more. Increasing the temperature of the rough-rolled bar to temperature higher than 1100° C. takes excessive time, and the productivity decreases. Therefore, if rough-rolled bar is heated, the heating temperature is preferably 1100° C. or less.
(Annealing Retention Temperature: 730° C. or More and 770° C. or Less)
When the annealing retention temperature is lower than 730° C., the austenite 12 is not formed sufficiently, and as illustrated in
(Annealing Retention Time: 3 Hours or More and 60 Hours or Less)
When the annealing retention time is less than 3 hours, as illustrated in
(Cooling Rate Down to 650° C.: 1° C./Hr or More and 60° C./Hr or Less)
When the cooling rate down to 650° C. is less than 1° C./hr, as illustrated in
According to the embodiment, excellent lubricity can be obtained, and therefore it is possible to suppress adhesion of the high-carbon steel sheet to the metal mold and suppress wearing of the metal mold. Further, according to the embodiment, it is also possible to suppress cracking during forming.
It should be noted that all of the above-described embodiments merely illustrate concrete examples of implementing the present invention, and the technical scope of the present invention is not to be construed in a restrictive manner by these embodiments. That is, the present invention may be implemented in various forms without departing from the technical spirit or main features thereof.
Next, examples of the present invention will be described. Conditions in the examples are condition examples employed for confirming feasibility and effect of the present invention, and the present invention is not limited to these condition examples. The present invention can employ various conditions as long as the object of the present invention is achieved without departing from the spirit of the invention.
In a first experiment, hot-rolling of a slab (Steel type A to Y, BK) including a chemical composition listed in Table 1 was performed, thereby obtaining a hot-rolled steel sheet having a thickness of 4 mm. In the hot-rolling, the slab heating temperature was 1130° C., the time thereof was 1 hour, the finish rolling temperature was 850° C., and the coiling temperature was 520° C. Then, cooling was performed down to a temperature of less than 60° C., and pickling using sulfuric acid was performed. Thereafter, annealing of the hot-rolled steel sheet was performed to then obtain a hot-rolled annealed steel sheet. In the annealing, the hot-rolled steel sheet was retained for 15 hours at 750° C., and then was cooled down to 650° C. at a cooling rate of 30° C./hr. Subsequently, cooling was performed down to a temperature of less than 60° C. In this manner, various high-carbon steel sheets were manufactured. Blank fields in Table 1 indicate that the content of the element is less than a detection limit, and the balance is Fe and impurities. For example, the Cr content of Steel type BK may be regarded as 0.00%. An underline in Table 1 indicates that the numeric value is out of the range of the present invention.
0.29
0.011
0.110
0.0036
0.18
0.026
1.05
0.0001
0.0002
0.06
0.0122
3.10
1.55
0.78
0.006
0.519
Then, the coefficient of micro-friction of ferrite, and the spheroidized ratio and the average diameter of cementite of each of the high-carbon steel sheets were measured. A friction coefficient of cementite was also measured in measuring the coefficient of micro-friction of ferrite. Results of them are listed in Table 2. An underline in Table 2 indicates that the item is out of the range of the present invention.
Further, evaluation of adhesion suppressive performance and evaluation of crack sensitivity of each of the high-carbon steel sheets were performed as formability evaluation. In the evaluation of adhesion suppressive performance, a draw bead test was performed. That is, an indentation bead with a tip having a 20 mm radius R was pressed against the high-carbon steel sheet with a load of 10 kN and was pulled out. Then, presence or absence of adhesive matter on the tip of the indentation bead was observed, and one with presence of adhesive matter was evaluated as X and one with no presence was evaluated as ◯. The presence of adhesive matter in this test indicates that in press forming with several thousands to several tens of thousands of shots, an adhesive matter occurs early on the metal mold to deteriorate formability. In the evaluation of crack sensitivity, a compression test was performed. That is, a cylindrical test piece having a 10 mm diameter and a 4 mm height was cut out from the high-carbon steel sheet so that a height direction of the test piece was parallel to a sheet thickness direction, and the test piece was compressed to 1 mm in height. Then, an appearance observation and a sectional structure observation were performed, and then one in which cracking appeared in the appearance during compression or after compression and one in which a crack of 1 mm or more was present in the sectional structure observation were evaluated as X, and one other than the above was evaluated as ◯. Results of them are also listed in Table 2.
J
K
0.72
L
M
N
O
P
Q
0.79
R
0.79
S
2.56
T
U
V
65.8
0.26
W
X
0.62
Y
BK
0.69
As listed in Table 2, Sample No. 1 to Sample No. 9 were each within the range of the present invention, thus being able to obtain good adhesion suppressive performance and crack sensitivity.
On the other hand, in Sample No. 10, the C content of Steel type J was too low, and thus the amount of cementite was insufficient, sufficient lubricity was not able to be obtained, and adhesion to the metal mold occurred during forming. In Sample No. 11, the N content of Steel type K was too high, and thus BN precipitated, the amount of solid-solution B was insufficient, the coefficient of micro-friction of ferrite was low, and adhesion and cracking during the compression test occurred. In Sample No. 12, the Al content of Steel type L was too high, and thus the ductility of ferrite was low and a crack originating from transgranular fracture of ferrite occurred during the compression test. In Sample No. 13, the B content of Steel type M was too high, and thus boride was formed and a crack originating from the boride is occurred during the compression test. In Sample No. 14, the Mn content of Steel type N was too low, and thus pearlite transformation occurred during cooling in the annealing, the spheroidized ratio of cementite was low, and a crack originating from acicular cementite occurred during the compression test. In Sample No. 15, the P content of Steel type O was too high, and thus segregation of B to the interface between ferrite and cementite was hindered and cracking occurred during the compression test. In Sample No. 16, the Si content of Steel type P was too high, and thus the ductility of ferrite was low and a crack originating from transgranular fracture of ferrite occurred during the compression test. In Sample No. 77 and Sample No. 18, each B content of Steel type Q and Steel type R was too low, and thus the coefficient of micro-friction of ferrite was low and adhesion and cracking during the compression test occurred. In Sample No. 19, the Si content of Steel type S was too low, and thus cementite became coarse excessively during annealing and a crack originating from the coarse cementite occurred during the compression test. In Sample No. 20, the S content of Steel type T was too high, and thus coarse sulfides being non-metal inclusions were formed and a crack originating from the coarse sulfide occurred during the compression test. In Sample No. 21, the Mn content of Steel type U was too high, and thus the ductility of ferrite was low and a crack originating from transgranular fracture of ferrite occurred during the compression test. In Sample No. 22, the Cr content of Steel type V was too high, and thus spheroidizing of cementite during annealing was hindered, coarsening of cementite was suppressed, and a crack originating from micro acicular cementite occurred during the compression test. In Sample No. 23, the C content of Steel type W was too high, and thus the amount of cementite was excessive and a crack originating from the cementite occurred during the compression test. In Sample No. 24, the Ti content of Steel type X was too low, and thus BN precipitated, the amount of solid-solution B was insufficient, the coefficient of micro-friction of ferrite was low, and adhesion and cracking during the compression test occurred. In Sample No. 25, the Ti content of Steel type Y was too high, and thus coarse oxides of Ti were formed and a crack originating from the coarse oxide of Ti occurred during the compression test. In Sample No. 26, the Cr content of Steel type BK was too low, and thus BN precipitated, the amount of solid-solution B was insufficient, the coefficient of micro-friction of ferrite was low, and adhesion to the metal mold occurred during forming.
In a second experiment, hot-rolling of a slab (Steel type Z to BJ) including a chemical composition listed in Table 3 was performed, thereby obtaining a hot-rolled steel sheet having a thickness of 4 mm. In the hot-rolling, the slab heating temperature was 1130° C., the time thereof was 1 hour, the finish rolling temperature was 850° C., and the coiling temperature was 520° C. Then, cooling was performed down to a temperature of less than 60° C., and pickling using sulfuric acid was performed. Thereafter, annealing of the hot-rolled steel sheet was performed to then obtain a hot-rolled annealed steel sheet. In the annealing, the hot-rolled steel sheet was retained for 15 hours at 750° C., and then was cooled down to 650° C. at a cooling rate of 30° C./hr. Subsequently, cooling was performed down to a temperature of less than 60° C. In this manner, various high-carbon steel sheets were manufactured. Blank fields in Table 3 indicate that the content of the element is less than a detection limit, and the balance is Fe and impurities. An underline in Table 3 indicates that the numeric value is out of the range of the present invention.
0.22
0.506
0.510
0.0003
0.511
0.17
0.516
0.0038
0.523
0.532
0.503
1.70
0.14
0.505
0.507
0.0104
0.512
0.009
0.509
Then, in the same manner as in the first experiment, the coefficient of micro-friction of ferrite, and the spheroidized ratio and the average diameter of cementite of each of the high-carbon steel sheets were measured, and further, the evaluation of adhesion suppressive performance and the evaluation of crack sensitivity were performed. Results of them are listed in Table 4. An underline in Table 4 indicates that the item is out of the range of the present invention.
AM
AN
AO
AP
AQ
0.58
AR
AS
79.0
AT
AU
AV
0.73
AW
AX
AY
64.0
0.24
AZ
67.4
BA
BB
BC
BD
BE
0.69
BF
2.32
BG
BH
BI
BJ
As listed in Table 4, Samples No. 31 to No. 43 were each within the range of the present invention, thus being able to obtain good adhesion suppressive performance and crack sensitivity.
On the other hand, in Sample No. 44, the C content of Steel type AM was too low, and thus the amount of cementite was insufficient, sufficient lubricity was not able to be obtained, and adhesion to the metal mold occurred during forming. In Sample No. 45, the Cu content of Steel type AN was too high, and thus a flaw occurred during hot-rolling and adhesion originating from the flaw occurred. In Sample No. 46, the Ca content of Steel type AO was too high, and thus coarse oxides of Ca were formed and a crack originating from the coarse oxide of Ca occurred during the compression test. In Sample No. 47, the Mo content of Steel type AP was too high, and thus the ductility of ferrite was low and a crack originating from transgranular fracture of ferrite occurred during the compression test. In Sample No. 48, the B content of Steel type AQ was too low, and thus the coefficient of micro-friction of ferrite was low and adhesion and cracking during the compression test occurred. In Sample No. 49, the Nb content of Steel type AR was too high, and thus the ductility of ferrite was low and a crack originating from transgranular fracture of ferrite occurred during the compression test. In Sample No. 50, the Mn content of Steel type AS was too low, and thus pearlite transformation occurred during cooling in the annealing, the spheroidized ratio of cementite was low, and a crack originating from acicular cementite occurred during the compression test. In Sample No. 51, the Ce content of Steel type AT was too high, and thus coarse oxides of Ce were formed and a crack originating from the coarse oxide of Ce occurred during the compression test. In Sample No. 52, the B content of Steel type AU was too high, and thus boride was formed and a crack originating from the boride occurred during the compression test. In Sample No. 53, the Ni content of Steel type AV was too high, and thus the coefficient of micro-friction of ferrite was high and adhesion occurred. In Sample No. 54, the V content of Steel type AW was too high, and thus the ductility of ferrite was low and a crack originating from transgranular fracture of ferrite occurred during the compression test. In Sample No. 55, the Zr content of Steel type AX was too high, and thus coarse oxides of Zr were formed and a crack originating from the coarse oxide of Zr occurred during the compression test. In Sample No. 56, the Cr content of Steel type AY was too high, and thus spheroidizing of cementite during annealing was hindered, coarsening of cementite was suppressed, and a crack originating from micro acicular cementite occurred during the compression test. In Sample No. 57, the Mn content of Steel type AZ was too low, and thus pearlite transformation occurred during cooling in the annealing, the spheroidized ratio of cementite was low, and a crack originating from acicular cementite occurred during the compression test. In Sample No. 58, the Y content of Steel type BA was too high, and thus coarse oxides of Y were formed and a crack originating from the coarse oxide of Y occurred during the compression test. In Sample No. 59, the La content of Steel type BB was too high, and thus coarse oxides of La were formed and a crack originating from the coarse oxide of La occurred during the compression test. In Sample No. 60, the S content of Steel type BC was too high, and thus coarse sulfides being non-metal inclusions were formed and a crack originating from the coarse sulfide occurred during the compression test. In Sample No. 61, the W content of Steel type BD was too high, and thus the ductility of ferrite was low and a crack originating from transgranular fracture of ferrite occurred during the compression test. In Sample No. 62, the Ti content of Steel type BE was too low, and thus EN precipitated, the amount of solid-solution B was insufficient, the coefficient of micro-friction of ferrite was low, and adhesion and cracking during the compression test occurred. In Sample No. 63, the Si content of Steel type BE was too low, and thus cementite became coarse excessively and a crack originating from the coarse cementite occurred during the compression test. In Sample No. 64, the P content of Steel type BG was too high, and thus segregation of B to the interface between ferrite and cementite was hindered and cracking occurred during the compression test. In Sample No. 65, the Ta content of Steel type BH was too high, and thus the ductility of ferrite was low and a crack originating from transgranular fracture of ferrite occurred during the compression test. In Sample No. 66, the Mg content of Steel type BI was too high, and thus coarse oxides of Mg were formed and a crack originating from the coarse oxide of Mg occurred during the compression test. In Sample No. 67, the C content of Steel type BJ was too high, and thus the amount of cementite was excessive and a crack originating from the cementite occurred during the compression test.
In a third experiment, hot-rolling and annealing were performed under various conditions on the steel types that were within the range of the present invention (Steel types A to I and Steel types Z to AL) out of the steel types used in the first experiment and the steel types used in the second experiment so as to manufacture high-carbon steel sheets. Conditions of them are listed in Table 5 to Table 7. An underline in Table 5 to Table 7 indicates that the numeric value is out of the range of the present invention.
786
706
714
68.1
776
435
721
983
1175
62.3
1158
823
718
801
444
62.6
962
68.3
Then, in the same manner as in the first experiment, the coefficient of micro-friction of ferrite, and the spheroidized ratio and the average diameter of cementite of each of the high-carbon steel sheets were measured, and further, the evaluation of adhesion suppressive performance and the evaluation of crack sensitivity were performed. Results of them are listed in Table 8 to Table 10. An underline in Table 8 to Table 10 indicates that the item is out of the range of the present invention.
0.72
2.25
0.64
78.3
0.63
0.29
0.59
0.71
75.4
0.61
62.3
0.73
2.36
0.64
0.72
0.61
0.71
0.61
64.1
0.69
0.59
0.63
0.68
2.45
0.66
0.59
0.66
2.55
0.60
0.22
0.54
2.38
0.62
0.72
0.58
2.46
As listed in Table 8, Samples No. 72, No. 74, No. 77 to No. 80, No. 82, No. 83, No. 85, and No. 88 to No. 92 were each within the range of the present invention, thus being able to obtain good adhesion suppressive performance and crack sensitivity. As listed in Table 9, Samples No. 103, No. 105, No. 106, No. 108 to No. 111, No. 114 to No. 117, and No. 120 to No. 122 were each also within the range of the present invention, thus being able to obtain good adhesion suppressive performance and crack sensitivity. As listed in Table 10, Samples No. 131, No. 133, No. 134, No. 136, No. 139, No. 141 to No. 143, No. 145, No. 147, No. 148, No. 151, and No. 152 were each also within the range of the present invention, thus being able to obtain good adhesion suppressive performance and crack sensitivity.
On the other hand, in Sample No. 71, the annealing retention temperature was too high, and thus volume expansion was large, a hot-rolled coil was uncoiled to cause abrasions, and a tightening mark caused by a tightening band also occurred. Further, the thickness of the film of crystals of B greatly varied and the coefficient of micro-friction of ferrite was large. Therefore, adhesion occurred. Further, cementite became coarse excessively and a crack originating from the coarse cementite occurred during the compression test. In Sample No. 73, the coiling temperature was too high, and thus coarse lamellar pearlite was formed in the hot-rolled steel sheet, spheroidizing of cementite during annealing was hindered, and the spheroidized ratio of cementite was low. Further, large irregularities were formed with removal of scales and the coefficient of micro-friction of ferrite was large. Therefore, adhesion and cracking during the compression test occurred. In Sample No. 75, the annealing retention time was too short, and thus the coefficient of micro-friction of ferrite was large and the average diameter of cementite was small. Therefore, adhesion and cracking during the compression test occurred. In Sample No. 76, the slab heating temperature was too low, and thus segregations of B, Mn, and others were not eliminated and the coefficient of micro-friction of ferrite was large. Therefore, adhesion and cracking during the compression test occurred. In Sample No. 81, the coiling temperature was too high, and thus adhesion and cracking during the compression test occurred similarly to Sample No. 73. In Sample No. 84, the cooling rate was too high, and thus pearlite transformation occurred during cooling and a crack originating from acicular cementite occurred during the compression test. Further, a good film of crystals of B was not formed on the surface of the high-carbon steel sheet, the coefficient of micro-friction of ferrite was high, and adhesion occurred. In Sample No. 86, the annealing retention temperature was too high, and thus adhesion and cracking during the compression test occurred, similarly to Sample No. 81. In Sample No. 87, the coiling temperature was too low, and thus as a result of removal of scales, the surface of the steel sheet became rough and adhesion occurred.
In Sample No. 101, the annealing retention temperature was too low, and thus the segregation of B to the interface between ferrite and austenite was suppressed, the coefficient of micro-friction of ferrite was large, and adhesion occurred. Further, the segregation of B to the interface between ferrite and cementite was also suppressed and cracking occurred during the compression test. In Sample No. 102, the finish rolling temperature was too high, and thus large irregularities were formed with removal of scales and the coefficient of micro-friction of ferrite was large. Therefore, adhesion occurred. In Sample No. 104, the slab heating temperature was too high, and thus B atoms were oxidized during slab heating and the coefficient of micro-friction of ferrite was large. Therefore, adhesion occurred. In Sample No. 107, the cooling rate was too high, and thus pearlite transformation occurred during cooling and a crack originating from acicular cementite occurred during the compression test. Further, a good film of crystals of B was not formed on the surface of the high-carbon steel sheet, the coefficient of micro-friction of ferrite was high, and adhesion occurred. In Sample No. 112, the slab heating temperature was too high, and thus adhesion occurred, similarly to Sample No. 104. In Sample No. 113, the finish rolling temperature was too low, and thus anisotropy of the structure was strong and a crack originating from a nonuniform structure occurred during the compression test. Further, as a result of removal of scales, the surface of the steel sheet became rough and adhesion occurred. In Sample No. 118, the annealing retention temperature was too low, and thus adhesion and cracking during the compression test occurred, similarly to Sample No. 101.
In Sample No. 132, the cooling rate was too low, and thus the thickness of the film of crystals of B greatly varied and the coefficient of micro-friction of ferrite was large. Therefore, adhesion occurred. Further, cementite became coarse excessively and a crack originating from the coarse cementite occurred during the compression test. In Sample No. 135, the finish rolling temperature was too low, and thus anisotropy of the structure was strong and a crack originating from a nonuniform structure occurred during the compression test. Further, as a result of removal of scales, the surface of the steel sheet became rough and adhesion occurred. In Sample No. 137, the coiling temperature was too low, and thus as a result of removal of scales, the surface of the steel sheet became rough and adhesion occurred. In Sample No. 138, the annealing retention time was too long, and thus volume expansion was large, a hot-rolled coil was uncoiled to cause abrasions, and a tightening mark caused by a tightening band also occurred. Further, the thickness of the film of crystals of B greatly varied and the coefficient of micro-friction of ferrite was large. Therefore, adhesion occurred. Further, cementite became coarse excessively and a crack originating from the coarse cementite occurred during the compression test. In Sample No. 140, the annealing retention time was too short, and thus the coefficient of micro-friction of ferrite was large and the average diameter of cementite was small. Therefore, adhesion and cracking during the compression test occurred. In Sample No. 144, the cooling rate was too low, and thus adhesion and cracking during the compression test occurred, similarly to Sample No. 132. In Sample No. 146, the finish rolling temperature was too high, and thus, large irregularities were formed with removal of scales and the coefficient of micro-friction of ferrite was large. Therefore, adhesion occurred. In Sample No. 149, the slab heating temperature was too low, and thus segregations of B, Mn, and others were not eliminated and the coefficient of micro-friction of ferrite was large. Therefore, adhesion and cracking during the compression test occurred. In Sample No. 150, the annealing retention time was too long, and thus adhesion and cracking during the compression test occurred, similarly to Sample No. 138.
The present invention may be utilized in, for example, manufacturing industries and application industries of high-carbon steel sheets used for various steel products, such as a driving system component for automobile, a saw, a knife, and others.
| Filing Document | Filing Date | Country | Kind |
|---|---|---|---|
| PCT/JP2014/077544 | 10/16/2014 | WO | 00 |
