KR101600723B1 - Medium carbon steel sheet, quenched member, and method for manufacturing medium carbon steel sheet and quenched member - Google Patents

Abstract

Wherein the carbon steel sheet contains, by mass%, 0.10 to 0.80% of C, 0.01 to 0.3% of Si, 0.3 to 2.0% of Mn, 0.001 to 0.10% of Al and 0.001 to 0.01% of N, The average diameter of the carbide is 0.4 占 퐉 or less and the ratio of the number of carbides having a size of 1.5 times or more of the average diameter of the carbides is 30% or less with respect to the total number of carbides, the carbide has a spheronization rate of 90% And a tensile strength (TS) of 550 MPa or less.

Description

Technical Field [0001] The present invention relates to a carbon steel sheet, a medium carbon steel sheet, a quenching member, and a method of manufacturing the same. BACKGROUND ART [0002]

The present invention relates to a medium carbon steel sheet excellent in cold workability, particularly, in cold forging and a method for producing the same. The present invention also relates to a member (quenching member) obtained by molding and then quenching the carbon steel sheet, and a manufacturing method thereof.

The present application claims priority based on Japanese Patent Application No. 2011-197044 filed on September 9, 2011, the contents of which are incorporated herein by reference.

Since members such as a chain, a gear, a clutch, a saw, and a blade require abrasion resistance and fatigue resistance, it is necessary to increase the strength (particularly, the strength of the surface) of the member by quenching.

Conventionally, the member has been manufactured by forming a steel (steel sheet or steel ingot) containing 0.2% by mass or less of C by hot forging and increasing the strength of the member surface by carburizing and high frequency quenching. In order to ensure workability in hot forging, the amount of C is reduced to 0.2 mass%. In order to compensate for the lack of quenching due to low C, the carbon concentration of the member surface is increased by carburizing have.

BACKGROUND ART [0002] In recent years, there has been a demand for a manufacturing method capable of manufacturing a member having sufficient strength under a condition of a lower number of process water in a shorter period of time at a lower temperature, for the purpose of CO ₂ reduction and cost reduction, .

In order to cope with such a demand, for example, hot forming such as hot forging may be changed to cold forming such as cold forging or carburization may be omitted by increasing the carbon concentration (quenching) in the steel sheet before molding. In this case, a medium carbon steel sheet is required in which sufficient cold workability can be obtained even if the C amount of the steel sheet is increased. Particularly, because recent processing technology has been developed and a molding method having a higher degree of processing than that of the prior art can be adopted for a steel sheet, a steel is required to have a soft and easy to deform (low deformation resistance) (High deformability) is also required. Examples of the molding method having a high degree of processing include a pressing method for increasing the machining accuracy and shortening the machining time by simultaneously applying the compressive load from multiple directions, and the like.

The inventors of the present invention have found that even when a machining area having a considerable deformation of more than 1 is generated during cold working, it can be applied to a forming method having a high degree of machining if a steel plate is free from cracking.

However, in the prior art, it has been difficult to use a medium carbon steel sheet with a high C concentration for the above-described high-process molding method (a forming method in which a machining area with equivalent deformation exceeding 1 occurs).

For example, Patent Documents 1 to 8 disclose a medium carbon steel sheet for obtaining a molded article by processing.

Among them, Patent Document 1, for example, discloses that carbide is dispersed in ferrite so that the carbide spheroidization ratio is 90% or more including 0.1-0.8 mass% of C and 0.01 mass% or less of S: A carbide particle diameter of 0.4 to 1.0 占 퐉, and a ferrite grain diameter of 20 占 퐉 or more as the need arises. Further, for example, Patent Document 2 discloses a method for producing a medium- and high-carbon steel sheet excellent in stretch flangeability. In the steel sheets of Patent Documents 1 and 2, since the carbide is coarsened by annealing to control the average carbide particle diameter to 0.4 to 1.0 탆, the yield ratio is large and cracks starting from coarse carbides tend to occur, It is difficult to apply these steel sheets to a molding method having a high degree of processing.

Patent Document 3 discloses a steel sheet for punching parts having excellent fatigue characteristics. In this patent document 3, a steel sheet cold-rolled at a reduction ratio of 50% or more is annealed at a temperature of Ac1 DEG C or lower in order to control the carbide to 0.3 mu m or less. However, in the method of manufacturing a steel sheet disclosed in Patent Document 3, since the microstructure becomes finer due to a high reduction rate, the hardness of the steel sheet is increased to 200 to 400 HV (yield strength 600 to 1400 MPa) Low deformation resistance) can not be obtained.

Likewise, Patent Documents 4 to 8 disclose a medium carbon steel sheet whose carbide shape is controlled. However, in Patent Documents 4 to 6, since the carbide is precipitated from the low-temperature phase, the particle diameter distribution of the carbide tends to be widened, and cracks starting from the coarse carbide are likely to occur. In Patent Document 7, since no annealing is performed before spheroidizing annealing, coarse carbides that are not substantially spheroidized tend to be generated, and this carbide tends to become a starting point of cracking. In Patent Document 8, since the spheroidizing annealing and the machining are performed at the same time, the particle diameter distribution of the carbide tends to be widened, and cracks starting from the coarse carbide are likely to occur.

In this way, a medium carbon steel sheet applicable to a molding method having a high degree of processing is not found.

Further, even when the hardness of the member surface is increased by quenching, members such as gears are required to have excellent shape accuracy by uniform quenching in the members.

However, in the steel sheet manufacturing methods of Patent Documents 1 and 2, since the temperature range of Ac1 to Ac1 + 100 占 폚 is included at the time of annealing, not only the carbide is coarsened, but also the portion that was austenite and the ferrite A difference occurs in the size of the carbide. According to the difference in the size of the carbide, austenite becomes a seamless texture at the time of quenching in the production of the member, and the quenching of the steel sheet and the shape precision of the member after quenching are reduced.

As described above, there is also a problem that the shape precision of the member after quenching is small in the prior art.

Japanese Patent Application Laid-Open No. 11-80884 Japanese Patent Application Laid-Open No. 11-269552 Japanese Patent Application Laid-Open No. 2001-59128 Japanese Patent Application Laid-Open No. 2003-89846 Japanese Patent Laid-Open No. 9-268344 Japanese Patent Application Laid-Open No. 2004-137527 Japanese Patent Laid-Open No. 2001-329333 Japanese Patent Application Laid-Open No. 2001-355047

SUMMARY OF THE INVENTION In view of the above, it is an object of the present invention to provide a medium carbon steel plate excellent in cold workability and quenching stability, and members having excellent shape accuracy, and a method of manufacturing the same.

The inventors of the present invention have conducted intensive studies on a method for solving the above problems. As a result, it is effective to take measures against fine cracks by ensuring the uniformity of the deformation propagation to improve the cold workability. It is effective to set the mean ferrite grain size to 10 탆 or more, the average diameter of the carbide to 0.4 탆 or less, Or more in terms of temperature and humidity. In addition, in the steel sheet with improved workability, it is also known that the average diameter of carbide is extremely small, and the ratio of coarse carbide particles is also reduced. It is also known that the quenching can be stabilized even under any quenching condition .

The steel sheet satisfying the above conditions is difficult to manufacture even when a single manufacturing condition such as a simple rolling condition or an annealing condition is devised and it is difficult to optimize a plurality of conditions by a so-called continuous process such as a process from hot rolling through cold rolling to annealing The present invention has been completed based on this discovery.

The gist of the present invention is as follows.

(1) A medium carbon steel sheet according to one aspect of the present invention is a steel sheet comprising, by mass%, 0.10 to 0.80% of C, 0.01 to 0.3% of Si, 0.3 to 2.0% of Mn, 0.001 to 0.10% 0.01% or less of S, 0.0025% or less of O, 1.5% or less of Cr, 0.01% or less of B, 0.5% or less of Mo, and 0.01% or less of Nb and 0.12 to 0.5% 0.5% or less of Ta, 0.5% or less of Ni, 0.003% or less of Mg, 0.003% or less of Ca, 0.003% or less of Ca, 0.5% or less of V, 0.5% or less of Ti, 0.03% or less of Zr, 0.03% or less of La, 0.03% or less of La, 0.03% or less of Ce, 0.03% or less of Sn, 0.03% or less of Sb and 0.03% or less of As and the balance of Fe and inevitable impurities Wherein a ratio of the number of carbides having a size of at least 1.5 times an average diameter of the carbides is 30% or less with respect to a total number of the carbides, and a percentage of spheroidizing of the carbides is at least 90% And an average ferrite grain size of 10 mu m , And is less than the tensile strength (TS) it is 550㎫.

(2) In the medium carbon steel sheet according to (1), the yield ratio (YR) may be 60% or less.

(3) In the medium carbon steel sheet according to (1) or (2), the thickness may be 1 to 12.5 mm.

(4) In the method for producing a medium carbon steel sheet according to one aspect of the present invention, the steel sheet contains 0.10 to 0.80% of C, 0.01 to 0.3% of Si, 0.3 to 2.0% of Mn, 0.001 to 0.10% 0.001 to 0.01% of N, 0.12 to 0.5% of Nb, 0.03% or less of P, 0.01% or less of S, 0.0025% or less of O, Cu: not more than 0.5%, W: not more than 0.5%, Ta: not more than 0.5%, Ni: not more than 0.5%, Mg: not more than 0.003%, Ca: not more than 0.003% 0.03% or less of Y, 0.03% or less of Zr, 0.03% or less of La, 0.03% or less of Ce, 0.03% or less of Sn, 0.03% or less of Sb and 0.03% or less of As, Casting a steel having a chemical composition comprising inevitable impurities; Hot-rolled; Air cooling for 2 to 10 seconds immediately after the end of the hot rolling; Cooling from a temperature at the end of the air cooling to a temperature range of 480 to 600 캜 at an average cooling rate of 10 to 80 캜 / s; Winding at a temperature range of 400 to 580 占 폚 and at a temperature lower than the termination temperature of said cooling; Cold rolling at a cold rolling rate of 5% or more and less than 30%; Annealing is performed for 5 to 40 hours at a temperature range of 650 to 720 占 폚.

(5) In the method for producing a medium carbon steel sheet according to (4), the average lamellar thickness of cementite in the pearlite contained in the steel after the winding is 0.02 to 0.5 탆.

(6) A quenching member according to one aspect of the present invention is a quenching member obtained from the medium carbon steel sheet according to the above (1) or (2), and contains 0.10 to 0.80% of C, 0.01 to 0.3% 0.003 to 2.0% of Mn, 0.001 to 0.10% of Al, 0.001 to 0.01% of N and 0.12 to 0.5% of Nb, P: not more than 0.03%, S: not more than 0.01% 0.5% or less of Cu, 0.5% or less of W, 0.5% or less of W, 0.5% or less of Ta, 0.5% or less of Cr, 1.5% or less of B, 0.03% or less; La: 0.03% or less; Ce: 0.03% or less; Sn: 0.03% or less; Sb: 0.03% or less; Or less and 0.03% or less of As, the remaining amount of Fe and inevitable impurities, and the ratio of area occupied by old austenite grains having a grain size of not more than 0.5 times or not less than twice the average grain size of old austenite grains is 30 % Or less.

(7) In the quenching member according to (6), the area ratio of martensite may be 95% or more.

(8) In the method of manufacturing a quenching member according to one aspect of the present invention, the medium carbon steel sheet according to any one of (1) to (3) Heating the member to a temperature higher than the Ac3 transformation point; The member is cooled.

According to each of the above-mentioned aspects, it is possible to provide a medium carbon steel sheet excellent in cold workability and quenching stability against strict cold working, a method for manufacturing the same, and a quenching member excellent in shape accuracy and a method for manufacturing the same.

Here, the quenching stability refers to uniformity of microstructure after quenching, suppression of heat treatment deformation, and uniformity of residual stress in the steel material when quenching members are obtained from the steel sheet. The cold workability (hereinafter, referred to as " cold workability ") for strict cold working refers to the workability in the case where a machining area having a considerable deformation of more than 1 occurs during cold working. Further, the processing (cold working) includes bending, thickness increase, drawing, and the like.

BRIEF DESCRIPTION OF THE DRAWINGS Fig. 1 is a diagram showing the relationship between cold workability and the ratio of the average carbide diameter and the number of coarse carbides. Fig.
2A is a perspective view showing a shape of a test material before cold working for evaluating cold workability;
Fig. 2B is a perspective view showing the shape of the test material after cold working for evaluating the cold workability. Fig.
3 is a view showing the relationship between the ratio of the number of coarse carbides and the percent area percent of ideal austenite.
4 is a graph showing the relationship between the cold workability, the spheroidization ratio of carbide, and the average ferrite grain size.
5 is a diagram showing the relationship between the blend index and the deformation amount.
6A is a side view showing an example of a blank before quenching in the high frequency quenching test.
6B is a perspective view showing an example of a blank before quenching in the high frequency quenching test.
6C is a side view showing an example of a blank after quenching in the high frequency quenching test.
6D is a perspective view showing an example of a specimen after quenching in the high frequency quenching test.
7 is a view showing the relationship between the air cooling time and the number ratio of coarse carbides.
8 is a view showing the relationship between the cold rolling ratio and the Vickers hardness.
9 is a flowchart showing an example of a method for producing a medium carbon steel sheet according to an embodiment of the present invention.
10 is a flowchart showing an example of a method of manufacturing a quenching member according to an embodiment of the present invention.
11 is a schematic view for explaining each variable of the equivalent variation in the equation (1).

Hereinafter, the present invention will be described in detail.

Furthermore, each of the following embodiments was obtained by repeating irradiation of a steel sheet and its manufacturing method based on a component system specified in JIS G4051 (Carbon steel for machine structure), JIS G4401 (Carbon tool steel) or JIS G4802 (Cold rolled steel for spring). In this investigation, the present inventors conducted a cup forming test as a cold working (cold forging) test in which a machining area in which a considerable deformation exceeded 1 at the time of cold working was generated, and the obtained steel sheet was evaluated. As shown in Fig. 11, the equivalent strain (epsilon _e ) is obtained by assuming that the three sides of the rectangular parallelepiped coincide with the x, y and z axes of the Cartesian coordinate system, respectively, (in FIG. 11, expansion) _x, (in Fig. 11, shrinkage) ε _y, (in Fig. 11, shrinkage) ε _z to use a] is expressed by the equation (1).

Further, each elongation strain is an increase ratio of the dimension when the dimension before machining is taken as a reference (that is, 1), which is obtained from the dimensional change of each axial direction material. For example, in a bending test at a bending angle of 90, the equivalent strain is 1 or less, and the equivalent strain to be the object is changed by the test method.

First, a medium carbon steel sheet according to one embodiment of the present invention will be described below.

The reason for limiting the chemical composition (chemical composition) will be described. Here, "% " with respect to the component means% by mass.

(C: 0.10 to 0.80%)

C is an important element in ensuring strength after quenching of a member (a steel sheet after molding). When the amount of C is 0.10% or more, required strength is secured. When the C content is less than 0.10%, ferrite transformation is promoted during hot rolling and coiling, and it becomes difficult to uniformly disperse the cementite particles in the steel material. Therefore, the lower limit of the amount of C is set to 0.10%. On the other hand, if the amount of C exceeds 0.80%, the lamellar thickness of the cementite in the pearlite of the hot-rolled sheet before annealing exceeds 0.5 탆. Therefore, in this case, the cementite becomes difficult to spheroidize, and the proportion of coarse cementite increases to lower the quenching stability. Therefore, the upper limit of the amount of C is set to 0.80%. When the strength or the quenching is further increased, the lower limit of the amount of C is preferably 0.15% or 0.20%, more preferably 0.24%, 0.28% or 0.32%. When the quenching stability is further improved, the upper limit of the amount of C is preferably 0.70% or 0.65%, more preferably 0.60% or 0.55%.

(Si: 0.01 to 0.3%)

Si acts as a deoxidizing agent and is an element effective for improving hardness and increasing strength. When the amount of Si is less than 0.01%, the effect of such addition can not be obtained, so the lower limit of the amount of Si is set to 0.01%. On the other hand, if the amount of Si exceeds 0.3%, the deformability of the ferrite decreases, and cracks tend to occur in the grain during processing, so that the cold workability decreases. Therefore, the upper limit of the amount of Si is set to 0.3%. When the strength or the quenching is further increased, the lower limit of the amount of Si is preferably 0.03% or 0.05%, more preferably 0.08% or 0.10%. When the cold workability is further enhanced, the upper limit of the amount of Si is preferably 0.28%, more preferably 0.25%.

(Mn: 0.3 to 2.0%)

Mn is an important element for controlling the thermal stability of cementite. When the amount of Mn is less than 0.3%, such addition effect can not be obtained, so the lower limit of the amount of Mn is set to 0.3%. On the other hand, if the amount of Mn exceeds 2.0%, the amount of MnS increases, and cracks tend to occur during cold working. In addition, in this case, cementite is liable to remain at the time of quenching, and the austenite grain size is increased. Therefore, the upper limit of the amount of Mn is set to 2.0%. When the thermal stability of the cementite is further increased and the grain size control of the cementite is more stably performed, the lower limit of the Mn content is preferably 0.4% or 0.5%, more preferably 0.6% or 0.7%. When the thermal stability of cementite is suppressed and the quenching stability is further enhanced, the upper limit of the amount of Mn is preferably 1.7% or 1.6%, more preferably 1.5% or 1.4%.

(Al: 0.001 to 0.10%)

Al is an element which acts as a deoxidizing agent and which is effective for fixing N and has a large solid solubility. When the amount of Al is less than 0.001%, the effect of this addition is not sufficiently obtained, so the lower limit of the amount of Al is set to 0.001%. On the other hand, if the Al content exceeds 0.10%, the addition effect saturates, and the deformability of the ferrite decreases, and cracking tends to occur in the grain during processing, so that the cold workability decreases. Therefore, the upper limit of the Al amount is set to 0.10%. Further, in order to fix as much N as possible, the lower limit of the amount of Al may be 0.003%, 0.005%, or 0.010%. The upper limit of the amount of Al may be limited to 0.09%, 0.08%, or 0.07%.

(N: 0.001 to 0.01%)

N is an element that forms nitride. If an excessive amount of N is contained in the steel, not only the cold workability is lowered but also the grain growth of the austenite during quenching heating is suppressed to increase the degree of agglomeration, so the upper limit of the N amount is set to 0.01%. The smaller the amount of N is, the more preferable. However, if the amount of N is reduced to less than 0.001%, the refining cost increases, so the lower limit of the amount of N is set to 0.001%. If necessary, the upper limit of the amount of N may be limited to 0.009%, 0.008%, or 0.007%.

The above chemical element is a basic element (basic element) of steel in this embodiment, and the basic element is controlled (contained or limited) so that the chemical composition including the remaining amount of iron and inevitable impurities is equivalent to that of the present embodiment The basic composition. However, in the present embodiment, the following chemical elements (optional elements) may be further contained in the steel in addition to the basic components (instead of a part of Fe in the remainder portion). In addition, even if these selective elements are mixed in the steel inevitably (for example, the amount of each selected element is lower than the preferable lower limit), the effect of the present embodiment is not impaired.

In other words, the carbon steel sheet according to the present embodiment contains P, S, O, Cr, B, Nb, Mo, V, Ti, Cu, W, Ta, Ni, Mg, Ca, Y, Zr, La, Ce, Sn, Sb and As. Further, since these elements do not necessarily have to be added to the steel, the lower limit of these 22 elements is 0% and is not limited. As a result, only the upper limit of these 22 elements is limited.

(P: 0 to 0.03%)

P is an element that acts on the increase of the strength. If an excessive amount of P is contained in the steel, not only the tensile strength TS is increased but also the toughness is lowered to deteriorate the cold workability. Therefore, the upper limit of the P amount is set to 0.03%. The upper limit of the P amount may be limited to 0.025%, 0.02%, or 0.015% for improvement of toughness and cold workability. However, if the P is reduced to less than 0.001%, the refining cost greatly increases, so that the lower limit of the P amount may be set to 0.001%.

(S: 0 to 0.01%)

S forms nonmetallic inclusions such as MnS and deteriorates the cold workability. Also, S increases the agglomeration of austenite grains by pinning the growth of austenite grains during heating. Therefore, the upper limit of the amount of S is set to 0.01%. The upper limit of the amount of S may be limited to 0.008%, 0.007%, or 0.005% in order to further improve the cold workability. However, if the S is reduced to less than 0.0001%, the refining cost greatly increases, so that the lower limit of the S amount may be set to 0.0001%.

(O: 0 to 0.0025%)

When the oxides coagulate and coarsen, the cold workability decreases, so the O amount (oxygen amount) is 0.0025% or less. This amount is small enough to be judged as an inevitable impurity. The amount of O is preferably as small as possible, and may be limited to 0.002% or less. However, it is technically difficult to reduce the amount of O to less than 0.0001%, so the amount of O may be 0.0001% or more.

At least one selected from the group consisting of Cr, B, Nb, Mo, V, Ti, Cu, W and Ta may be added to the steel in order to strengthen the mechanical properties of the medium carbon steel sheet.

(Cr: 0 to 1.5%)

Cr is an element effective for controlling the increase in the strength of the steel sheet and the thermal stability of the cementite. When Cr is added to steel, when the amount of Cr is less than 0.010%, the effect of the addition is small, so that the lower limit of the amount of Cr may be set to 0.010%. However, when the amount of Cr exceeds 1.5%, the yield ratio (YR) increases due to the growth or inhibition of dissolution of cementite, or the austenite structure during heating becomes coarse, so that the upper limit of the amount of Cr is set to 1.5% do. For lowering the alloy cost, the upper limit of the amount of Cr may be set at 1.2%, 1.0%, 0.8%, 0.6%, or 0.4%.

(B: 0 to 0.01%)

B is an element effective for increasing the quenching property by the addition of a trace amount. When B is added to the steel, the addition effect is not obtained when the B content is less than 0.001%, so the lower limit of the B content may be 0.001%. However, when the amount of B exceeds 0.01%, segregation is promoted during continuous casting, and coarse carbides are produced, and scratches are likely to occur in the slab. Therefore, the upper limit of the amount of B is set to 0.01%. To prevent scratches, the upper limit of the amount of B may be limited to 0.008%, 0.006%, 0.004%, or 0.002%.

(Nb: 0 to 0.5%)

Nb is an element effective for forming carbonitride and preventing remarkable coarsening of austenite grains. When Nb is added to steel, if the amount of Nb is less than 0.01%, the effect of the addition is not sufficiently manifested, so that the lower limit of the amount of Nb may be set to 0.01%. However, when the amount of Nb exceeds 0.5%, in addition to raising the yield ratio (YR), the ratio of fine austenite grains is excessively increased, so the upper limit of the amount of Nb is 0.5%. When the ratio of fine austenite is further reduced and the quenching stability is further enhanced, the upper limit of the amount of Nb is preferably 0.3%, 0.2% or 0.15%.

(Mo: 0 to 0.5%)

Mo is an element effective for forming carbide and preventing remarkable coarsening of austenite grains. When Mo is added to the steel, the effect of addition is not exhibited when the amount of Mo is less than 0.01%. Therefore, the lower limit of the amount of Mo may be set to 0.01%. However, when the amount of Mo exceeds 0.5%, the upper limit of the amount of Mo is set to 0.5% in order to excessively increase the ratio of fine austenite particles in addition to raising the yield ratio (YR). If necessary, the upper limit of the amount of Mo may be limited to 0.4%, 0.3%, 0.2%, or 0.1%.

(V: 0 to 0.5%)

V, like Nb, is an element effective for forming carbonitride and preventing remarkable coarsening of austenite grains. When V is added to the steel, when the V content is less than 0.01%, the effect of the addition is not fully manifested, so that the lower limit of the V content may be set to 0.01%. However, when the amount of V exceeds 0.5%, in addition to raising the yield ratio (YR), the ratio of fine austenite particles is excessively increased, so the upper limit of the amount of V is set to 0.5%. When the ratio of fine austenite is further reduced and the quenching stability is further enhanced, the upper limit of the amount of V may be limited to 0.4%, 0.3%, 0.2%, or 0.1%.

(Ti: 0 to 0.3%)

Like Ti and V, Ti is an element effective for forming carbonitride and preventing remarkable coarsening of austenite grains. When Ti is added to the steel, when the amount of Ti is less than 0.01%, the effect of the addition is not fully manifested, so that the lower limit of the amount of Ti may be set to 0.01%. However, when the amount of Ti exceeds 0.3%, the upper limit of the amount of Ti is set to 0.3% in order to excessively increase the ratio of fine austenite particles in addition to raising the yield ratio (YR). For lowering the alloy cost, the upper limit of the amount of Ti may be limited to 0.2%, 0.1%, or 0.05%.

(Cu: 0 to 0.5%)

Cu is an element incorporated from scrap or the like. If Cu is included in the steel, the workability is lowered or the brittleness in hot is increased. Therefore, the upper limit of the amount of Cu is set to 0.5%. If necessary, the upper limit of the amount of Cu may be limited to 0.4% or 0.3%. The smaller the amount of Cu is, the better, but if the amount of Cu is reduced to less than 0.01%, the refining cost is greatly increased. Therefore, the lower limit of the amount of Cu may be set to 0.01%.

(W: 0 to 0.5%)

W, like Mo, is an element effective for forming carbide and preventing remarkable coarsening of austenite grains. When W is added to the steel, when the W content is less than 0.01%, the effect of the addition is not fully manifested, so the lower limit of the W content may be set to 0.01%. However, when the amount of W exceeds 0.5%, in addition to raising the yield ratio (YR), the upper limit of the amount of W is set to 0.5% in order to excessively increase the ratio of the fine austenite particles. If necessary, the upper limit of the amount of W may be limited to 0.4%, 0.2%, or 0.1%.

(Ta: 0 to 0.5%)

Like Ta, Ta is an element effective in forming carbonitride and preventing remarkable coarsening of austenite grains. In the case where Ta is added to the steel, when the amount of Ta is less than 0.01%, the effect of the addition is not fully manifested, so that the lower limit of the amount of Ta may be set to 0.01%. However, when the amount of Ta exceeds 0.5%, the upper limit of the amount of Ta is set to 0.5% in order to excessively increase the ratio of fine austenite particles in addition to raising the yield ratio (YR). If necessary, the upper limit of the amount of Ta may be set to 0.3%, 0.2%, or 0.1%.

Next, at least one selected from Ni, Mg, Ca, Y, Zr, La and Ce may be added to the steel as a selective element to further strengthen the mechanical properties of the medium carbon steel sheet.

(Ni: 0 to 0.5%)

Ni is an element effective for improving toughness and hardness. When Ni is added to the steel, when the amount of Ni is less than 0.01%, there is no effect of the above addition, so the lower limit of the amount of Ni may be set to 0.01%. However, if the amount of Ni exceeds 0.5%, the effect of addition is saturated and the cost increases, so that the upper limit of the amount of Ni is set to 0.5%. In order to reduce the cost of the alloy, the upper limit of the amount of Ni may be set to 0.3%, 0.2%, or 0.1%.

(Mg: 0 to 0.003%)

Mg is an element effective for controlling the shape of a sulfide by adding a trace amount of Mg, and Mg can be added to the steel as needed. When Mg is added to the steel, when the amount of Mg is less than 0.0005%, the effect can not be obtained. Therefore, the lower limit of the amount of Mg may be set to 0.0005%. In addition, Mg is liable to form oxides, and compounds containing these oxides inhibit the growth of austenite grains. When the amount of Mg is more than 0.003%, Mg is not uniformly distributed in the steel, and a portion where the growth of the austenite grains is suppressed at the quenching heating and a portion where the growth of the austenite grains are not suppressed is unevenly distributed. It becomes difficult to obtain an austenite structure. Therefore, the upper limit of the amount of Mg is set to 0.003%. In order to reduce the cost of the alloy, the upper limit of the Mg amount may be set to 0.002% or 0.001%.

(Ca: 0 to 0.003%)

Ca, like Mg, is an element effective for controlling the shape of a sulfide by adding a trace amount of Ca, and Ca can be added to the steel as needed. When Ca is added to the steel, since the effect is not obtained when the amount of Ca is less than 0.0005%, the lower limit of the amount of Ca may be set to 0.0005%. Further, Ca is liable to form oxides, and compounds containing these oxides inhibit the growth of austenite grains. When the amount of Ca exceeds 0.003%, Ca is not uniformly distributed in the steel, and a place where the growth of the austenite grains during the quenching heating is suppressed and a place where the austenite grains are not restrained is unevenly distributed. It becomes difficult to obtain an austenite structure. Therefore, the upper limit of the amount of Ca is set to 0.003%. In order to reduce the cost of the alloy, the upper limit of the amount of Ca may be set to 0.002% or 0.001%.

(Y: 0 to 0.03%)

Y, like Ca and Mg, is an element effective for controlling the shape of a sulfide by adding a trace amount of Y. If necessary, Y can be added to the steel. When Y is added to the steel, since the effect is not obtained when the Y content is less than 0.001%, the lower limit of the Y content may be 0.001%. In addition, Y is prone to form oxides, and the compounds of these oxides inhibit the growth of austenite grains. When the amount of Y exceeds 0.03%, Y is not uniformly distributed in the steel, and a portion where the growth of the austenite grains during the quenching heating is suppressed and a portion where the growth is not suppressed are unevenly distributed. It becomes difficult to obtain an austenite structure. For this reason, the upper limit of the Y amount is set to 0.03%. For lowering the alloy cost, the upper limit of the Y amount may be set to 0.01% or 0.005%.

(Zr: 0 to 0.03%)

Zr is an element effective for controlling the shape of a sulfide by the addition of a trace amount like Y, Ca and Mg, and Zr can be added in the steel as required. In the case of adding Zr in the steel, since the effect is not obtained when the Zr content is less than 0.001%, the lower limit of the Zr content may be set to 0.001%. Further, Zr is liable to form oxides and carbides, and these oxides and carbide compounds inhibit the growth of austenite grains. When the amount of Zr exceeds 0.03%, Zr is not uniformly distributed in the steel, and a portion where the growth of the austenite grains is suppressed and a portion where the growth of the austenite grains is not suppressed at the time of quenching heating are unevenly distributed. It becomes difficult to obtain an austenite structure. For this reason, the upper limit of the amount of Zr is set to 0.03%. For lowering the alloy cost, the upper limit of the amount of Zr may be set to 0.01% or 0.005%.

(La: 0 to 0.03%)

La, like Zr, Y, Ca and Mg, is an element effective for controlling the shape of a sulfide by adding a trace amount of La, and La can be added to the steel if necessary. When La is added to the steel, the effect can not be obtained when the amount of La is less than 0.001%. Therefore, the lower limit of the amount of La may be set to 0.001%. La is also liable to form oxides, and the compounds of these oxides inhibit the growth of austenite grains. When the amount of La exceeds 0.03%, La is not uniformly distributed in the steel, where the growth of the austenite grains during the quenching heating is suppressed and the portion where the growth is not suppressed is unevenly distributed. Therefore, It becomes difficult to obtain an austenite structure. For this reason, the upper limit of the amount of La is set to 0.03%. In order to reduce the cost of the alloy, the upper limit of the amount of La may be set to 0.02%, 0.01%, or 0.005%.

(Ce: 0 to 0.03%)

As with La, Zr, Y, Ca and Mg, Ce is an element effective for controlling the shape of a sulfide by adding a trace amount of Ce. If necessary, Ce can be added to the steel. When Ce is added to the steel, the effect can not be obtained when the amount of Ce is less than 0.001%. Therefore, the lower limit of Ce amount may be set to 0.001%. Further, Ce is liable to form oxides, and the compounds of these oxides inhibit the growth of austenite grains. When the amount of Ce exceeds 0.03%, Ce is not uniformly distributed in the steel, and a portion where the growth of the austenite grains is suppressed and a portion where the growth of the austenite grains is not suppressed at the time of quenching is unevenly distributed. It becomes difficult to obtain an austenite structure. For this reason, the upper limit of Ce amount is set to 0.03%. In order to reduce the cost of the alloy, the upper limit of Ce amount may be 0.02%, 0.01%, or 0.005%.

(Sn: 0 to 0.03%, Sb: 0 to 0.03%, As: 0 to 0.03%),

Further, when scrap is used as a raw material for medium carbon steel sheet, at least one of Sn, Sb and As is inevitably mixed in the steel in an amount of 0.003% or more. If all of these elements are 0.03% or less, the quenching of the medium carbon steel sheet is not hindered. Therefore, the steel may contain at most 0.03% of Sn, at most 0.03% of Sb, and at most 0.03% of As. The lower limit of the amounts of Sn, Sb and As is not particularly limited, but may be 0.005% or 0.003%, for example, from the viewpoint of refining efficiency when scrap containing a large amount of Sn, Sb and As is used.

As described above, the medium carbon steel sheet according to the present embodiment includes the above-described basic elements, and the balance of the chemical composition including Fe and inevitable impurities or the above-described basic elements and at least one selected from the above- Species, and the remainder portion contains Fe and inevitable impurities.

In the present embodiment, the medium carbon steel sheet has a carbide (cementite) which is dispersed in a microstructure containing ferrite satisfying the above-mentioned chemical composition. Further, Fe atoms and C atoms (for example, Fe ₃ C (cementite)) bonded to Fe atoms and C atoms, and Fe atoms and M atoms substituted for Fe atoms in the bond phase by selective elements M The bonded phase of an atom (carbide containing Fe) is generally referred to as carbide. Further, in the present embodiment, mainly the shape of iron carbide (mainly cementite) is controlled in the carbide, and the carbide can be regarded as iron carbide or cementite.

That is, in the medium carbon steel sheet according to the present embodiment, the ratio of the number of carbides having an average diameter (average carbide diameter) of 0.4 μm or less and a size of 1.5 times or more of the average diameter of the carbides Of the total number of carbides is 30% or less of the total number of carbides, the sintering rate of the carbides is 90% or more, and the average ferrite grain size is 10 占 퐉 or more.

Such a steel sheet is excellent in cold workability and quenching stability when it is formed into a member having a product shape and then quenched.

Fig. 1 shows the relationship between the cold workability and the average carbide diameter and the ratio of the number of coarse carbides. As shown in Fig. 1, when the average carbide diameter is 0.4 mu m or less and the number ratio of coarse carbides is 30% or less, the cold workability can be ensured (see O in Fig. 1).

Cracks during cold working are liable to be generated from coarse carbides (cementites). When the average carbide diameter exceeds 0.4 μm, cold workability can not be ensured. The average diameter of the carbide is 0.4 mu m or less, preferably 0.35 mu m or less or 0.3 mu m or less. The present inventors have also found that the occurrence frequency of cracks is influenced by the particle size distribution of carbides, and in particular, when the number ratio of coarse carbides exceeds 30%, the cold workability can not be secured found.

As described above, the coarse carbides promote cracking during cold working due to deformation concentration. Therefore, it is effective to improve the cold workability by finely dispersing the carbide in the steel as described above and not concentrating the deformation. However, when the strengthening of the dispersion of particles by the carbide particles is suppressed and the cold workability is further enhanced by reducing the deformation resistance, the lower limit of the average carbide diameter may be 0.10 탆 or more. The lower the ratio of the number of coarse carbides, the better the deformability until cracking occurs in the cold working. Therefore, the lower limit of the number ratio of coarse carbides is not particularly limited and may be 0%. On the other hand, in order to make the ratio of the crude carbide to 0%, it is necessary to make the manufacturing process strict. Therefore, in order to reduce the manufacturing cost, the lower limit of the number ratio of coarse carbides may be set to 5%.

Fig. 4 shows the relationship between the cold workability, the spheroidization ratio of carbide, and the mean ferrite grain size. As shown in Fig. 4, when the spheroidization ratio of carbide is 90% or more and the mean ferrite grain size is 10 m or more, the cold workability can be ensured (see O in Fig. 4). In the vicinity of the needle-like carbide, the stress tends to become localized at the time of cold working and tends to be a starting point of cracking. Therefore, it is considered that the higher the spheroidizing rate of the carbide is, the better the improvement in the cold workability.

As shown in Fig. 4, if the spheroidizing rate of the carbide is less than 90%, the locally stress concentration causes the carbide to become a starting point of the crack and degrade the cold workability (see X in Fig. 4). Therefore, in order to obtain sufficient cold workability, the spheroidization ratio of carbide is set to 90% or more. The spheroidizing rate of the carbide may be 91% or more or 92% or more. The higher the spheroidizing ratio of the carbide, the better the cold workability. Therefore, the upper limit of the spheroidizing rate of the carbide is not particularly limited. In this case, ideally, it is most preferable to increase the spheroidization ratio of the carbide to 100%. On the other hand, strict control of the manufacturing process is required to control all the carbides to spherical shape. Therefore, in order to suppress the decrease in the yield, the upper limit of the spheroidizing rate of the carbide may be set to 99.5% or less.

Further, as shown in Fig. 4, when the average diameter (average ferrite grain size) of the ferrite is less than 10 mu m, the deformation is partially concentrated and cracked due to the increase in the yield ratio YR, Can not be ensured. Therefore, the average ferrite grain size is set to 10 탆 or more. The mean ferrite grain size may be 12 占 퐉 or more, 15 占 퐉 or more, or 18 占 퐉 or more. Although the upper limit of the average ferrite grain size is not particularly limited, if the ferrite grain size of the steel sheet becomes too large, surface roughness tends to occur during cold working, thereby deteriorating the appearance of the product. Therefore, in order to further improve the appearance of the product, the upper limit of the average ferrite grain size may be set to 100 탆. The upper limit of the average ferrite grain size is preferably 80 占 퐉 or 60 占 퐉.

3 shows the relationship between the ratio of the number of coarse carbide particles and the number of grains having a grain size number different from the grain size number corresponding to the average grain size by the austenite structure at the time of quenching (that is, the old austenite grains after quenching) (% Area percent of the above austenite)). Here, the ideal austenite is spherical austenite particles having a particle size of not more than 0.5 times or not less than 0.5 times the average spherical austenite particle size, and the area ratio of the ideal austenite is such that the sum of areas of the ideal austenite It is the ratio of the total area.

As shown in Fig. 3, when the average ferrite grain size is less than 10 mu m, the proportion of coarse austenite increases during heating during quenching. When the number ratio of coarse carbides exceeds 30% even when the average ferrite grain size is 10 占 퐉 or more, the proportion of coarse austenite increases during heating during quenching. When the total area of the ideal austenite grains in the austenite structure during heating at the time of quenching exceeds 30% and the degree of sintering becomes high, in the cooling at quenching, the initiation of transformation according to the grain sizes of the respective austenite grains There will be a difference between when and when. As a result, the uniformity of the microstructure of the quenched material is deteriorated, and the heat treatment deformation is increased, resulting in defective shape, and the quenching stability is lowered. Thus, the control of ferrite and carbide is also important for ensuring quenching stability.

Further, in the observation of the structure of the medium carbon steel sheet, a scanning electron microscope was used to measure the visibility of the specimen at a magnification of about 3000 to 10000, and in some cases about 30,000 times, a view including 500 or more carbides (cementite) At least 16 areas are selected, and the area of each carbide contained in the area is measured in detail. Thereafter, the average diameter of the carbide is calculated from the average area per one carbide as the average carbide particle diameter when the grain shape is approximated by a circle. Therefore, in consideration of the measurement method, the lower limit of the average carbide particle diameter may be 0.03 mu m. The carbide having a ratio of the major axis length to the minor axis length of 3 or more is referred to as acicular carbide and the carbide having a ratio of less than 3 is referred to as a spherical carbide to calculate the number of needle carbides and the number of spherical carbides. The value obtained by dividing the number of spherical carbides by the total number of carbides is called the spheroidization ratio of carbide (cementite).

The ferrite grain size of the steel sheet is preferably measured by a scanning electron microscope. Five or more areas are photographed at a magnification that includes 200 or more ferrite particles, and the number of ferrite particles included in the photographed photograph is counted. The photographed photograph is made up of one piece of ferrite particles including all of them and 0.5 piece of ferrite particles including only a part of them. The area per unit area of the ferrite particles is obtained by dividing the photographing area by the number of sent ferrite particles. The square root of the obtained area is referred to as ferrite grain size, and the average value (the average of five or more areas) is referred to as an average ferrite grain size. It is also preferable that bainite and martensite are not mixed in the microstructure.

The plate thickness of the steel sheet may be 1 mm or more, or 1.2 mm or more or 1.8 mm or more. The thickness of the steel sheet may be 12.5 mm or less, or 10 mm or less, 8 mm or less, or 6 mm or less. The tensile strength (TS) is 550 MPa or less. When the tensile strength (TS) is reduced to 550 MPa or less, the ductility is increased, and the molding amount at the time of processing can be sufficiently secured. As described above, when the tensile strength TS is lowered, ductility is improved and workability is improved. The tensile strength TS is preferably 500 MPa or less, 470 MPa or less, or 440 MPa or less. Here, when it is necessary to reduce sagging during punching processing, the tensile strength TS may be increased. In recent years, forging techniques of punching, bending and increasing the thickness are integrated, and therefore, the lower limit of the tensile strength TS may be set to 360 MPa or more or 400 MPa or more depending on the processing method to be applied.

It is also preferable that the yield ratio (YR) (yield point or yield strength / tensile strength) is 60% or less. If the yield ratio (YR) is reduced to 60% or less, concentration of deformation during processing can be avoided. Particularly, in the case of machining a plate, since the material is deformed while being tied to the metal mold, the propagation of the deformation at the time of deformation is uniform so that the material can be uniformly deformed in synchronization with the movement of the metal mold, Uniformity. For this purpose, it has been also found out that it is only necessary to control the work rate to a low yield ratio (YR) so that work hardening can be uniformly started. When the yield ratio (YR) is 60% or less, deformation is likely to propagate. Therefore, at the time of cold working, stress concentration at any one place can be prevented, and as a result, inflow failure and cracks can be further suppressed. It is more preferable to set the yield ratio (YR) to 56% or less or 52% or less. When the yield ratio (YR) is 30% or more, it is possible to prevent scratches on the surface of the steel sheet due to friction with the tool during cold working due to a sufficient yield point or yield strength YP. As described above, since the aesthetic appearance of the member surface obtained from the steel sheet can be further improved, the lower limit of the yield ratio (YR) of the medium carbon steel sheet may be set to 30% or more.

Next, a method for manufacturing a medium carbon steel sheet according to one embodiment of the present invention will be described.

In this embodiment, the pearlite structure of the hot-rolled sheet is optimized by using the material having the chemical composition (chemical composition) range described in the above-described embodiment, and considering the conditions of cooling after finish rolling in hot rolling, (Iron carbide) by cold rolling of the cemented carbide and a short-time annealing at a low temperature (for example, one time cold rolling and one time annealing). Thus, a medium carbon steel sheet having excellent cold workability and excellent quenching stability can be produced. Hereinafter, the manufacturing method according to the present embodiment will be described in detail.

9, the steel having the chemical composition described in the above embodiment is cast (S1), hot-rolled (S2), air-cooled (S3 , Cold rolling (S4), winding (S5), cold rolling (S6), and annealing (S7).

(Hot rolling)

In hot rolling (hot rolling), the cast steel (for example, a continuous casting cast piece) satisfying the range of the chemical composition described in the above embodiment may be directly hot rolled, heated, and then hot rolled. The conditions of hot rolling are not particularly limited, and conditions of general rough rolling and finish rolling (for example, a rolling finish temperature of Ar3 + 50 deg. C or higher) can be applied as conditions of hot rolling.

(Cooling control and winding)

After the hot rolling, cooling control (control of the cooling pattern) is performed to produce pearlite by pearlite transformation, and the pearlite is rolled in a temperature range of 400 to 580 DEG C to obtain a hot rolled steel sheet having an average lamellar thickness of cementite in pearlite suitably controlled Can be obtained. The thinner the lamellar thickness, the more easily the lamellar thickness of the cementite becomes thinner, since the cementite is easily spheroidized at the time of annealing after cold rolling. On the other hand, if the thickness of the lamella is excessively thin, the pearlite becomes excessively hard, and it becomes difficult to introduce deformation into the pearlite during cold rolling. As a result, it is difficult for the ferrite to grow in the subsequent annealing. For this reason, the lower limit of the lamellar thickness may be 0.02 mu m. In addition, if the thickness of the lamella is too thick, spheroidization of the cementite is not promoted at the time of annealing after cold rolling. For this reason, the upper limit of the lamellar thickness may be 0.5 mu m.

In order to obtain such a post-hot-rolled texture, control of the coiling temperature and cooling control after finish rolling are carried out as follows. That is, the steel is air-cooled for 2 seconds to 10 seconds immediately after finish rolling (hot rolling), and then cooled at an average cooling rate of 1080 to 80 占 폚 / s from the end temperature (cold start temperature) Cooled to a pearlite region (finish temperature of arc cooling) at 600 ° C, and then wound in a temperature range of 400 ° C to 580 ° C. By such cooling control and control of the coiling temperature, it is possible to stably obtain the post-hot-rolled texture (for example, the average lamellar thickness of cementite in pearlite is 0.02 to 0.5 mu m). Here, air cooling is cooling with an average cooling rate of less than 10 ° C / s, and the lower limit of the average cooling rate of air cooling is not particularly limited. For example, the average cooling rate of air cooling may be 0 ° C / s.

In the microstructure immediately after finish rolling in the hot state, usually, the recrystallized structure of austenite and the non-recrystallized structure are mixed, and the grain size is often evaluated as uneven. Therefore, if the steel sheet is cooled to a temperature region of pearlite transformation of 480 to 600 ° C at an average cooling rate of 10 ° C / s or more immediately after finish rolling, the recrystallized austenite particles having a small particle size or the non- Pearlite transformation is started, and pearlite having a thick lamellar spacing of cementite is produced. Even if these hot rolled sheets are annealed for 5 to 40 hours in the temperature range of 650 to 720 占 폚 and cold-rolled at a cold rolling rate of 5% or more and less than 30% as described later after acid pickling, the average diameter of the carbide (cementite) (Cementite) is more than 30% or more than 0.4 mu m, or the ratio of crude carbide (cementite) is more than 30%. Therefore, the microstructure of the hot-rolled steel sheet can be made into recrystallized austenite having a uniform grain size by air cooling for 2 seconds or more immediately after the finish rolling in the hot state, and the timing of the transformation after air- And can be aligned in the plate thickness direction. As a result, a hot rolled steel sheet having uniform microstructure including pearlite can be obtained, and cementite and ferrite can be appropriately controlled. The air cooling time may be 3 seconds or more. However, if the air cooling time is longer, it is preferable that the air cooling time is longer, but if it exceeds 10 seconds, the productivity may be lowered, or problems such as occurrence of scratches due to scale growth may occur. Therefore, the upper limit of the air cooling time is set to 10 seconds or less. The air cooling time may be 9 seconds or less or 8 seconds or less.

Here, the microstructure of the hot-rolled steel sheet (steel) can be observed by a scanning electron microscope. Since the lamellar thickness of cermetite in the pearlite after hot rolling is as small as 0.02 to 0.5 占 퐉, pearlite is observed at a magnification of at least 3000 to 10000 to measure the average thickness of the lamella perpendicular to the observation plane. The lamellar thickness is measured by the line segment method, and the measured values obtained from 30 or more measurement points randomly extracted are averaged to calculate the average cementitious lamellar thickness.

(Cold rolling)

(Cold rolling ratio = (plate thickness before cold rolling) - (plate thickness after cold rolling) / (thickness before cold rolling)) of 5% or more and less than 30% after cold rolling the hot- Cold rolling (cold rolling). The number of times of cold rolling is preferably one.

In this cold rolling, deformation is applied to the microstructure of the hot-rolled steel sheet to make the deformation difference occurring between the respective structures remarkable, and the grain growth and recrystallization at the time of annealing are promoted by this deformation difference. In order to obtain the effects of such deformation introduction, the cold rolling rate is set to 5% or more. Further, in this case, the precision of the sheet thickness of the medium carbon steel sheet can be further increased by cold rolling.

Further, if the cold rolling rate is increased, the carbide can be made finer. If the cold rolling ratio is less than 5%, the carbide precipitated by hot rolling (for example, lamellar cementite in the pearlite) does not sufficiently break at the time of cold rolling and remains in the microstructure after annealing, Or the ratio of coarse carbide (cementite) increases or increases. If the cold rolling ratio is 30% or more, the ferrite after recrystallization by annealing becomes fine and the strength becomes high, so that high cold workability can not be achieved. Therefore, the upper limit of the cold rolling ratio is set to less than 30%.

Further, in the technical field of the present embodiment, it is not general that the steel sheet rolled at 400 ° C to 580 ° C is cold-rolled after acid cleaning. The hot-rolled sheet wound in this temperature region has a Vickers hardness of 200 HV or more and is liable to be cracked, so that the load during cold rolling may increase and the productivity may be lowered. Thereby, it was necessary to soften the hot-rolled sheet once by annealing. Further, in the case of using a soft steel having a small amount of C, it is necessary to carry out carburization in a subsequent step (for example, in the production of a quenching member) in order to secure the quenching property. In order to reduce the number of such processes (for example, annealing before cold rolling, cold rolling or carburizing) to reduce costs, it is common to omit cold rolling.

However, the inventors of the present invention have found that even if the hot-rolled sheet is cold-rolled at a cold rolling rate of 5% or more and less than 30% as it is with the recent development of cold rolling technology, the rolling mill can withstand the load during cold rolling . Further, a technology for more appropriately controlling the particle size distribution of cementite by cold rolling and annealing is a newly discovered method for solving the new problem of improving the overall shape accuracy of the quenching member. Furthermore, considering the total cost incurred by the number of steps in the production of the quenching member and the increase in the processing time, the energy can be further reduced.

(Annealing)

The steel sheet subjected to such hot rolling, pickling and cold rolling is annealed to obtain a steel sheet having a structure containing a target carbide. In order to set the strength after annealing to an appropriate level, cold rolling and annealing are preferably performed one cycle at a time.

By the annealing, the pro-eutectoid ferrite is recrystallized at the annealing temperature of Ac1 or less by using the difference in the amount of deformation integrated in each structure, and the grain grows into coarse particles to soften the steel sheet (steel).

In addition, the carbide (cementite) before annealing is uniformly dispersed in the microstructure at a relatively low coiling temperature, and is very fine at the completion of cold rolling by cold rolling. As a result, when the annealing is started, the carbide immediately begins to dissolve and spheroidization is promoted. However, as a result of the above-described production conditions, since a small amount of carbide is dispersed very uniformly, spheroidization of carbide proceeds at once, and very fine spherical carbides are simultaneously produced. In this annealing, it is preferable that the annealing condition is a low temperature in a short time. As a result, the annealing is performed at an annealing time of 5 to 40 hours at a temperature range of 650 to 720 DEG C (annealing temperature). The annealing temperature (lower limit) is preferably 680 DEG C or higher when the speed of spheroidizing the carbide is higher, and preferably 700 DEG C or lower when the coarsening of the carbide is further suppressed. When the spheroidization ratio of the carbide is further increased, the annealing time (lower limit) is preferably 20 hours or more. The optimum annealing conditions are 690 캜 and 20 to 40 hours.

Further, annealing is preferably performed by box annealing from the viewpoint of controllability of the annealing atmosphere. The annealing atmosphere is not particularly limited, but may be a hydrogen concentration of 95% or more, a dew point up to 400 캜 less than -20 캜, and a dew point less than -40 캜 at 400 캜. In this case, the deviation of the characteristic in the width direction of the steel material can be further suppressed. It is also possible to produce a steel material having a desired characteristic even in a nitrogen atmosphere.

In the present embodiment, cold rolling with a cold rolling ratio of 5% or more and less than 30% and annealing after cold rolling are combined as described above. It is preferable that this combination is performed once. Here, after the annealing, it is necessary not to carry out the cold rolling more than 5% of the cold rolling rate again. That is, there is no particular limitation on the number of rolling times (rolling by rollers) during one cold rolling. By combining such cold rolling and annealing after cold rolling, high productivity can be ensured while controlling the shape of carbide.

As described above, by controlling the cooling pattern after hot rolling, the coiling temperature, the cold rolling condition, and the annealing conditions, it is possible to grow ferrite in a state of 10 mu m or more in a state of finely uniformly dispersing the carbides while spherical, The workability can be ensured. At the same time, the ratio of coarsening of the cementite particle size (carbide particle size) can be controlled, and the quenching stability of the obtained medium carbon steel sheet can be enhanced.

A quenching member according to an embodiment of the present invention will be described.

In the quenching member (quenched steel member, steel member obtained by quenching with a product shape) according to the present embodiment, the old austenite particles having an average particle diameter of not more than 0.5 times or not less than twice the average particle diameter of old austenite particles Knit) is within 30% of the total area of the whole of the old austenite grains. The lower limit of the area ratio of the ideal austenite is not particularly limited, and the area ratio may be 0% or more, 1% or more, or 3% or more. The definition of the above-mentioned ideal austenite (grain size of not more than 0.5 times or not smaller than 0.5 times of the average grain size) corresponds to the definition of the grain size number of the austenite (grain size number not less than 2 from the grain size number corresponding to the average grain size) .

The austenite before quenching transforms into various low-temperature structures after quenching, but in a microstructure after quenching, a grain boundary of austenite before transformation is clearly developed by etching or the like. Thus, the particles surrounded by the grain boundaries of austenite, including the low-temperature structure, are evaluated as one old austenite grains.

In addition, the microstructure of the quenching member need not be specified, but the area ratio of the martensite may be 95% or more or 98% or more. The area ratio of the martensite may be 100% or less.

A manufacturing method of a quenching member according to an embodiment of the present invention will be described.

In this embodiment, as shown in Fig. 10, the carbon steel sheet according to the above embodiment is cold worked (S11), heated to a temperature higher than the Ac3 transformation point (S12), cooled (S13) , A quenching member (quenched steel member, steel member obtained by quenching with a product shape) is obtained. Therefore, there are no particular limitations on the conditions except for the heating temperature for transforming ferrite and cementite into austenite (for example, the heating temperature, the heating condition, and the heating condition) The same conditions as for quenching can be applied. For example, it is preferable that the heating (S12) and the cooling (S13) are performed by general high frequency quenching (high frequency heating and quenching). As such a general quenching condition, for example, a medium carbon steel plate (member before quenching) after cold working may be quenched after heating to a temperature range of 900 to 1200 ° C. Further, for example, a medium carbon steel plate (member before quenching) after cold working may be quenched to 200 占 폚 or less at an average cooling rate of 60 to 1000 占 폚 / s by keeping it for 0 to 20 seconds after heating. The microstructure of the finally obtained quenching member does not need to be specified, but the area ratio of the martensite may be 95% or more or 98% or more.

The quenching stability is affected by the degree of agglomeration of austenite during heating for quenching. When the quenching member is produced, the uniformity of the austenite particle size at the time of quenching is preferable. In the case where the austenite grains are mixed, the transformation starts from the austenite grains having a small grain diameter, so that the transformation deformation becomes uneven in the steel and the heat treatment deformation becomes large, such that the quenching stability is impaired. In particular, when the total area of the ideal austenite in the austenite structure at the time of quenching exceeds 30%, the deformation at the time of quenching becomes large. For this reason, a medium carbon steel sheet according to the above embodiment having high quenching stability is used so that deformation at the time of quenching can be sufficiently reduced.

Example

Next, an embodiment will be described.

The level of the embodiment is an example of the execution condition adopted to confirm the feasibility and effect 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.

The steel ingots (steel) having the chemical compositions shown in Tables 1 to 3 (iron and inevitable impurities) were hot-rolled, air-cooled for 2.5 seconds, cooled to 540 占 폚 at an average cooling rate of 30 占 폚 / s Rolled at 500 DEG C, and cooled to room temperature to obtain a hot-rolled sheet. The obtained hot-rolled sheets were cold-rolled at a cold rolling ratio of 15% and then annealed at 680 캜 for 24 hours to obtain various steel sheets having microstructures shown in Table 4, and cold workability was evaluated. Among the results shown in Table 4, hot rolled steel sheets (steel sheets obtained by hot rolling and cooling control) having a good formability were subjected to cold rolling at a cold rolling ratio of 10% and annealing at 680 캜 for 4 hours, Various steel sheets having microstructures were obtained, and cold workability was evaluated. Further, the same hot-rolled sheet was cold-rolled at a cold rolling ratio of 70% and annealed at 680 캜 for 8 hours to obtain various steel sheets having microstructures shown in Table 6, and cold workability was evaluated. A steel sheet was produced from the steel ingot having the chemical composition shown in Tables 1 to 3 under the conditions shown in Tables 7 and 10, and the thickness of the lamellar cementite in the pearlite of the hot rolled steel sheet and the grain size distribution of the carbide of the annealing plate , The ferrite grain size of the annealing plate was measured, and a tensile test was conducted. With respect to the steel sheet after completion of the holding in the annealing, the annealing temperature was lowered to 100 deg. C / hr from the annealing temperature to 400 deg.

In the cold workability test, a disk having a diameter of 100 mm as shown in Fig. 2A was molded at a room temperature in a cup shape having a diameter of 60 mm and a height of 28 mm as shown in Fig. 2B to evaluate the presence or absence of cracks. In the case where cracks were generated, a rating of "No Good", in which the cold workability was inferior, was added. On the other hand, in the case where there is no crack, the cold workability is rated as "Good". This evaluation method is also the same as that of the test method of Fig. 1 described above. "Good" is " Good "

A3, A3, A3, A3, A3, A3, AA4, AB3, AB4, AC3, AC4, AD3, AD4, AE3, AE4, AF3, AF4 6A and 6B in the plate width of 15 mm and the plate length of 150 mm in the case of the plate material having the plate length of 150 mm, AG3, AG4, AH3, AH4, AI3, AI4, AJ3, AJ4, AK3, AK4, AL3, AL4, AM3 and AM4 The core was heated at a frequency of 78 kHz from a room temperature to a heating rate of 100 占 폚 / s, held at 950 占 폚 for 10 seconds, and immediately quenched to room temperature at a cooling rate of 100 占 폚 / s or higher to conduct a high frequency quenching test. As shown in Figs. 6C and 6D, the amount of deformation of the specimen after quenching was evaluated by measuring the bending? From the plate longitudinal direction. Further, the warp? Is preferably less than 3 DEG, more preferably less than 1 DEG.

Further, the distribution of the old austenite grain size of the blank after quenching was measured. In order to measure the distribution of the austenite grain size, the microstructure of the member after quenching was developed in the following solution A, and the microstructure thus obtained was photographed as a digital image. Thereafter, the average grain size of the old austenite The index was obtained. The solution A was prepared by mixing predetermined amounts of sodium hydroxide and picric acid with respect to 1000 ml of distilled water. In the image analysis, the area of each of the old austenite particles was measured, and the circle equivalent diameter was calculated. The average value of the circle equivalent diameters was taken as the average particle diameter. The area ratio obtained by dividing the area occupied by the crystal grains having a grain size not larger than 0.5 times or not smaller than 0.5 times the average grain size by the area of the observation field was determined as the blend index. Fig. 5 shows the relationship between the blended yarn index and the warpage (deformation amount)?. As shown in Fig. 5, when the blend index exceeds 30%, it can be seen that the warpage (deformation)? Is 3 or more. Therefore, when the blend index exceeds 30%, it is determined that the amount of deformation after quenching is large and the quenching stability is poor. Further, by observing the microstructure, it was confirmed that the area ratio of martensite of the quenched material after quenching was 95% or more.

In the steel sheets No. C0, E0, H0, L0 and N0 in Table 4, the average diameter of the carbides is 0.4 占 퐉 or less, the ratio of the number of carbides having a size of 1.5 times or more of the average diameter of the carbides is 30% A spheroidization ratio of 90% or more, and an average ferrite grain size of 10 占 퐉 or more. All of these steel sheet Nos. Exhibited good workability and quenching stability.

In Table 4, the chemical composition of the steel sheet, the average diameter (mu m) of the carbide, the number ratio (%) of the coarse carbides, and the number of the carbides in the steel sheet No. A0, B0, D0, F0, G0, I0, J0, At least one of the spheroidization ratio (%) of carbide and the average ferrite grain size (탆) was not sufficient. As a result, in these steel sheet No.s, the cold workability was not sufficient due to cracks originating from the aggregate of cementite, cracks originating from cementite, cracks starting from MnS, or cracks in the grain.

In the steel sheets No. C1, E1, H1, L1 and N1 shown in Table 5, since the annealing time was less than 5 hours, the sintering rate of the carbide was less than 90%, and the cold workability was not sufficient due to the cracks starting from cementite.

In the steel sheets No. C2, E2, H2, L2 and N2 shown in Table 6, the cold-rolled ratio exceeded 30%, the average ferrite grain size was less than 10 占 퐉 and cracked due to the concentration of deformation, I did.

In Table 7 to Table 12, the steel sheets No. C3, C4, E3, H3, L4, N3, Z4, AA4, AB4, AC4, AD3, AD4, AE4, AF3, AG3, AG4, AH4, AI3, , AL3, AM3 and AM4 were excellent in cold workability and excellent in quenching stability with a blend index of 30% or less.

In the steel sheets No. D3, D4, J3, J4, K3, K4, P3 and P4, the amount of any one of Si, Al, P and Cu was large.

In the steel sheets No. G3, G4, I3, I4, M3, M4, T3, T4, Y3 and Y4, the amount of any one of Mn, N, S, Cr and Zr was large so that the cold workability and the quenching stability were sufficient Did not do it.

Since the amount of C was small in the steel sheets No. A3 and A4, the average ferrite grain size was less than 10 占 퐉 and the cold workability and quenching stability were not sufficient.

In the steel sheets No. B3 and B4, since the amount of Si was small, the number ratio of coarse carbides was higher than 30%, and the cold workability and quenching stability were not sufficient.

In Steel Nos. F3 and F4, since the amount of Mn was small, the average carbide diameter was larger than 0.4 mu m, and cold workability and quenching stability were not sufficient.

In the steel sheets No. O3 and O4, since the amount of C was large, the ratio of the number of coarse carbides and the spheroidizing ratio of carbide could not be properly controlled and the cold workability and quenching stability were not sufficient.

The amount of any one of Nb, Ta, V, W, Ti, and Mo is large in the steel sheets No. Q3, Q4, R3, R4, S3, S4, U3, U4, V3, V4, W3, W4, AN3, , The average ferrite grain size was less than 10 mu m, so that the cold workability and the quenching stability were not sufficient.

In Steel Nos. X3 and X4, since the amount of B was large, the number ratio of coarse carbides was higher than 30%, and the cold workability and quenching stability were not sufficient.

In steel sheet No. AJ3, the annealing temperature was less than 650 占 폚, so that the average ferrite grain size was less than 10 占 퐉, so that the cold workability and quenching stability were not sufficient.

In the steel sheet No. E4, since the annealing temperature was higher than 720 占 폚, the spheroidizing rate of the carbide was less than 90%, and the cold workability was poor.

In steel sheet No. AA3, since the annealing temperature was higher than 720 占 폚, the spheroidization ratio of carbide and the ratio of the number of coarse carbides could not be appropriately controlled and the cold workability and quenching stability were not sufficient.

Since the annealing time was shorter than 5 hours in the steel sheets No. AH3 and AL4, the average ferrite grain size was less than 10 占 퐉 and the cold workability and quenching stability were not sufficient.

In the steel sheet No. H4, since the annealing time was longer than 40 hours, the number ratio of coarse carbides was higher than 30%, and the cold workability and quenching stability were not sufficient.

In the steel sheets No. AE3 and AI4, since the coiling temperature was less than 400 占 폚, the average ferrite grain size was less than 10 占 퐉, so that the cold workability and quenching stability were not sufficient.

Since the coiling temperatures of steel sheets No. L3 and N4 exceeded 580 占 폚, the average carbide diameter, the number ratio of coarse carbides, and the spheroidizing ratio of carbide could not be appropriately controlled and the cold workability and quenching stability were not sufficient I did.

In the steel plate No.AK4, since the cold rolling ratio was less than 5%, the mean ferrite grain size was less than 10 占 퐉, so that the cold workability and quenching stability were not sufficient.

In steel sheet No. AC3, the average ferrite grain size was less than 10 mu m and the number ratio of coarse carbides was higher than 30% because the cold rolling ratio was less than 5% and the annealing time was longer than 40 hr. As a result, cold workability and quenching stability were not sufficient in this steel sheet No..

In the steel sheets No. AB3 and AF4, since the cold rolling ratio exceeded 30%, the average carbide diameter or the ratio of the number of coarse carbides could not be appropriately controlled, and the mean ferrite grain size was less than 10 탆, And the quenching stability was not sufficient.

Further, in Steel No. H shown in Tables 1 to 3, various steel sheets were obtained by variously changing the time of air cooling after finish rolling. As a result, as shown in Fig. 7, the relationship between the air cooling time after finish rolling and the number ratio of coarse carbides was obtained. 7 shows that the number of coarse carbides can be reduced to 30% or less by performing air cooling for 2 seconds or more.

The steel sheet shown in Fig. 7 was produced by the following method. The steel was kept at 1220 占 폚 for 50 minutes, and then the steel was rolled in the hot state from 250 mm plate thickness to 3 mm plate thickness so as to finish finish rolling at 900 占 폚. Immediately after finish rolling, steel was air-cooled for each air-cooling time, cooled to 550 캜 at a cooling rate of 40 캜 / s, and wound at 500 캜 to produce a hot-rolled steel sheet. Each hot-rolled steel sheet was pickled, cold-rolled at a reduction ratio (cold rolling ratio) of 28%, and annealed at 700 ° C for 24 hours.

In addition, in the steel No. H shown in Tables 1 to 3, the reduction rate (cold rolling ratio) at the time of cold rolling was variously changed to obtain various steel sheets. As a result, as shown in Fig. 8, the relationship between the cold rolling ratio and the Vickers hardness was obtained.

The steel sheet shown in Fig. 8 was produced by the following method. The steel was kept at 1220 占 폚 for 50 minutes, and then the steel was rolled in the hot state from 250 mm plate thickness to 8 mm plate thickness so as to finish the finish rolling at 920 占 폚. Immediately after finish rolling, the steel was air-cooled for 3 seconds, cooled to 520 캜 at a cooling rate of 50 캜 / s, and wound at 480 캜 to produce a hot-rolled steel sheet. Each hot-rolled steel sheet was acid-washed, cold-rolled at respective reduction ratios (cold rolling ratio), and annealed at 700 ° C for 24 hours.

Table 13 shows the relationship between the microstructure of the steel sheet and the microstructure of the steel sheet after quenching, for some steel types whose chemical composition is sufficiently adjusted. In Table 13, steel sheets No.C3, C4, E3, H3L4, N3, Z3, Z4, AA4, AB4, AC4, AD3, AD4, AE4, AF3, AG3, AG4, AH4, AI3, AJ4, AK3, In AM3 and AM4, the average carbide diameter, the number ratio of coarse carbides, the spheroidization ratio of carbide, and the average ferrite grain size were all sufficient, and the blend index was 30% or less. In these steel sheet Nos., The steel sheets No. H4, L3, N4, AA3, AB3, N3, N3, and N4, in which at least one of the average diameter of carbide, the number ratio of coarse carbide, the spheroidization ratio of carbide, Martensite of higher area ratio could be obtained than AC3, AE3, AF4, AH3, AI4, AJ3, AK4 and AL4.

It is possible to provide a medium carbon steel sheet excellent in cold workability, a manufacturing method thereof, and a quenching member.

Claims (9)

% By mass,
C: 0.10 to 0.80%
0.01 to 0.3% of Si,
Mn: 0.3 to 2.0%
Al: 0.001 to 0.10%
N: 0.001 to 0.01% and
Nb: 0.12 to 0.5%
≪ / RTI >
P: 0.03% or less,
S: 0.01% or less,
O: 0.0025% or less,
Cr: 1.5% or less,
B: 0.01% or less,
Mo: 0.5% or less,
V: 0.5% or less,
Ti: 0.3% or less,
Cu: 0.5% or less,
W: 0.5% or less,
Ta: 0.5% or less,
Ni: 0.5% or less,
Mg: 0.003% or less,
Ca: 0.003% or less,
Y: 0.03% or less,
Zr: 0.03% or less,
La: 0.03% or less,
Ce: 0.03% or less,
0.03% or less of Sn,
0.03% or less of Sb and
As: not more than 0.03%
, The balance being Fe and inevitable impurities, the average diameter of the carbides being 0.4 占 퐉 or less, and the ratio of the number of carbides having a size of 1.5 times or more of the average diameter of the carbides to the total number of carbides is 30% , A sphericalization ratio of the carbide of 90% or more, an average ferrite grain size of 10 탆 or more, and a tensile strength (TS) of 550 MPa or less. The method according to claim 1,
Wherein the yield ratio (YR) is 60% or less. 3. The method according to claim 1 or 2,
And the plate thickness is 1 to 12.5 mm. % By mass,
C: 0.10 to 0.80%
0.01 to 0.3% of Si,
Mn: 0.3 to 2.0%
Al: 0.001 to 0.10%
N: 0.001 to 0.01% and
Nb: 0.12 to 0.5%
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P: 0.03% or less,
S: 0.01% or less,
O: 0.0025% or less,
Cr: 1.5% or less,
B: 0.01% or less,
Mo: 0.5% or less,
V: 0.5% or less,
Ti: 0.3% or less,
Cu: 0.5% or less,
W: 0.5% or less,
Ta: 0.5% or less,
Ni: 0.5% or less,
Mg: 0.003% or less,
Ca: 0.003% or less,
Y: 0.03% or less,
Zr: 0.03% or less,
La: 0.03% or less,
Ce: 0.03% or less,
0.03% or less of Sn,
0.03% or less of Sb and
As: not more than 0.03%
And the balance having a chemical composition containing Fe and inevitable impurities,
Casting;
Hot-rolled;
Air cooling for 2 to 10 seconds immediately after the end of the hot rolling;
Cooling from a temperature at the end of the air cooling to a temperature range of 480 to 600 캜 at an average cooling rate of 10 to 80 캜 / s;
Winding at a temperature range of 400 to 580 占 폚 and at a temperature lower than the termination temperature of said cooling;
Cold rolling at a cold rolling rate of 5% or more and less than 30%;
Annealing for 5 to 40 hours in a temperature range of 650 to 720 占 폚;
Wherein the carbon steel sheet has a surface roughness of at least 10 mm. 5. The method of claim 4,
Wherein the average lamellar thickness of the cementite in the pearlite included in the steel after the winding is 0.02 to 0.5 占 퐉. A quenching member obtained from the medium carbon steel sheet according to claim 1 or 2,
% By mass,
C: 0.10 to 0.80%
0.01 to 0.3% of Si,
Mn: 0.3 to 2.0%
Al: 0.001 to 0.10%
N: 0.001 to 0.01% and
Nb: 0.12 to 0.5%
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P: 0.03% or less,
S: 0.01% or less,
O: 0.0025% or less,
Cr: 1.5% or less,
B: 0.01% or less,
Mo: 0.5% or less,
V: 0.5% or less,
Ti: 0.3% or less,
Cu: 0.5% or less,
W: 0.5% or less,
Ta: 0.5% or less,
Ni: 0.5% or less,
Mg: 0.003% or less,
Ca: 0.003% or less,
Y: 0.03% or less,
Zr: 0.03% or less,
La: 0.03% or less,
Ce: 0.03% or less,
0.03% or less of Sn,
0.03% or less of Sb and
As: not more than 0.03%
, The balance being Fe and unavoidable impurities,
The ratio of the area occupied by the old austenite grains having the grain size of not more than 0.5 times or not less than twice the average grain size of the old austenite grains is not more than 30%
Wherein the quenching member is a quenched member. The method according to claim 6,
A quenched member characterized in that the area ratio of martensite is 95% or more. Cold-working the medium carbon steel sheet according to any one of claims 1 to 3 as a member;
Heating the member to a temperature higher than the Ac3 transformation point;
Cooling the member
≪ / RTI > Cold-working the medium carbon steel sheet according to claim 3 as a member;
Heating the member to a temperature higher than the Ac3 transformation point;
Cooling the member
≪ / RTI >

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