JP2015117406A - Middle and high carbon steel sheet excellent in punchability and manufacturing method therefor - Google Patents

Abstract

PROBLEM TO BE SOLVED: To provide a middle and high carbon steel sheet suitable as the raw material of driving system components such as automobile chains, gears, or clutches, and blades, cutting tools or the like, excellent in punchability, especially less punching sag, and a manufacturing method therefor.SOLUTION: There is provided a middle and high carbon hot rolled steel sheet excellent in punchability, containing, by mass%, C:0.10 to 0.70%, Si:0.01 to 1.0%, Mn:0.1 to 3.0%, P:0.001 to 0.025%, S:0.0001 to 0.010%, Al:0.001 to 0.10%, N:0.001 to 0.010% and the balance Fe with inevitable impurities and having an integration degree in a crystal orientation where a (110) face is in parallel accuracy in ±5° to a steel sheet surface in an area from a steel sheet surface layer to 200 μm in a sheet thickness of 2.5 or more.

Description

  The present invention relates to a medium / high carbon hot-rolled steel sheet having excellent punchability and a method for producing the same.

  Medium and high carbon steel plates are used as materials for drive system parts such as automobile chains, gears, and clutches, as well as saws and blades. The original material punched into a predetermined shape with a die from a steel strip cut out from a medium or high carbon steel strip or from a steel strip is formed into a part shape by plastic working such as drawing, hole expansion, thickening, and thinning. The

  In the past, parts were often formed by hot forging. In recent years, there has been a movement to manufacture parts by cold plastic working as described above for cost reduction. After processing, cutting is performed in order to obtain a predetermined dimension, but since the ratio of the cutting cost to the total manufacturing cost is large, the dimensional accuracy required for parts formed by cold plastic processing is increasing. Yes. Since the start of plastic working is a blank plate punched from a steel strip or a cut plate, establishment of a technique for manufacturing a steel plate having high dimensional accuracy by punching is required.

  The improvement of punchability shown in the present invention means reduction of punching blank plate, improvement of properties of the punching end face, and extension of die life as improvement of dimensional accuracy.

  Until now, many proposals have been made on techniques for improving punchability (for example, see Patent Documents 1 to 6).

For example, in Patent Document 1, C: 0.15 to 0.90% by weight of a medium carbon or high carbon steel plate as a precision punching high carbon steel plate that has excellent precision punchability and can be induction hardened, has a spheroidization rate of 80%.
As described above, a structure in which carbide having an average particle size of 0.4 to 1.0 μm is dispersed in the ferrite matrix is adjusted, the notch tensile elongation ElV is set to 20% or more, and the D value [= (3 × ElV2 + 18 × ElV) / TS. , TS: Tensile strength] is 3 or more, an invention for improving the mold life is disclosed, but it relates to the improvement of the punching end face property by controlling the internal structure, and can improve the punching sagging It's not technology.

  Further, Patent Document 2 discloses an invention of a medium- and high-carbon steel sheet that has not only excellent workability but also a soft structure that does not cause wear of the working mold during punching and shaving. In the production of the form, it is necessary to cold-roll the hot-rolled pickled steel sheet once or twice or more, which does not lead to drastic cost reduction.

  Patent Document 3 is a high carbon steel sheet containing 0.70 mass% or more and 0.95 mass% or less of C. By introducing voids into the steel structure by a combination of annealing conditions and cooling conditions, a soft punchable material is disclosed. An invention that aims to improve the speed and punchability is disclosed.

  Patent Document 4 discloses a high carbon cold-rolled steel sheet manufacturing method and a high-carbon cold-rolled steel sheet that can simplify the manufacturing process without deteriorating punching workability, and a high-carbon cold-rolled steel sheet having a predetermined component. Is hot-rolled at a finishing temperature of (Ar3 transformation point−20 ° C.) or more, cooled to a cooling stop temperature of 400 ° C. or more and 550 ° C. or less at a cooling rate exceeding 120 ° C./s, and 600 ° C. or more and Ac1 transformation point or less. After the temperature was raised to the coiling temperature range, coiled in that temperature range, cooled at an average cooling rate of 20 ° C./hr or less until the steel plate temperature reached 400 ° C. to obtain a hot rolled steel plate, pickling Thereafter, an invention is disclosed in which cold rolling is performed at a rolling rate of 30% or more and annealing is performed at an annealing temperature of 600 ° C. or more and Ac1 transformation point or less.

  Patent Document 5 is a high-carbon hot-rolled steel sheet that can be manufactured without using multi-stage annealing that requires a long time and is excellent in stretch flangeability in which cracking of the punched end face is unlikely to occur. An invention is disclosed in which the volume ratio of ferrite grains containing 1 μm or more and less than 1.2 μm and containing no carbide is controlled to 10% or less.

  Patent Document 6 has the same characteristics as a fine and uniform spherical carbide-dispersed steel only in a hot rolling process and a short annealing process, and has the workability that was a problem of the conventional fine carbide-dispersed steel. In particular, an invention relating to a method for producing a high carbon hot-rolled steel sheet that has been advantageously improved in terms of precision punching workability and bending workability is disclosed.

  However, none of these inventions disclose any technique for improving punching sagging that affects the dimensional accuracy of a plastic workpiece.

JP 2009-299189 A JP 2009-215612 A JP 2011-012316 A JP 2007-031761 A JP 2003-013145 A JP-A-8-269541

  In view of the above circumstances, an object of the present invention is to provide a medium / high carbon steel sheet with excellent punchability, particularly with less punching sagging, and a method for producing the same.

  The inventors of the present invention have intensively studied a method for solving the above-described problems. As a result, in order to achieve excellent punchability, particularly suppression of punching sagging, the degree of accumulation of crystal orientations in which the (110) plane of the iron body-centered cubic lattice is within ± 5 ° parallel to the steel plate surface is 2 It was found that it is effective to control to .5 or more. In addition, the hot-rolled sheet annealed sheet has a Vickers hardness of 100 HV or more and 160 HV or less, and the steel sheet has a ferrite grain size of 10 μm or more and 50 μm or less, a cementite particle diameter of 0.1 μm or more and 2.0 μm or less, and cementite spheroidization It has been found that by setting the rate to 85% or more, it is possible to improve the end face properties of the punched parts and to increase the life of the mold.

  In addition, steel sheet manufacturing methods that satisfy these requirements are difficult to manufacture even if the hot rolling conditions and annealing conditions are devised simply, and optimization is achieved through so-called integrated processes such as hot rolling and annealing processes. The fact that it can be produced only by doing this has been found by accumulating various studies, and the present invention has been completed.

  The gist of the present invention is as follows.

(1) In mass%,
C: 0.10 to 0.70%,
Si: 0.01 to 1.0%,
Mn: 0.1 to 3.0%
P: 0.001 to 0.025%,
S: 0.0001 to 0.010%,
Al: 0.001 to 0.10%,
N: 0.001 to 0.010%,
And the balance is a steel plate made of Fe and impurities,
Punchability characterized in that in the region from the steel sheet surface layer to the plate thickness direction 200 μm, the degree of accumulation of crystal orientations in which the (110) plane falls within the parallelism within ± 5 ° with respect to the steel sheet surface is 2.5 or more Medium and high carbon hot-rolled steel sheet with excellent resistance.

(2) The steel sheet is mass% as an additive element,
Ti: 0.01-0.20%,
Cr: 0.01 to 1.50%,
Mo: 0.01 to 0.50%,
B: 0.0001 to 0.010%,
Nb: 0.001 to 0.10%,
V: 0.001 to 0.2%,
Cu: 0.001 to 0.4%
W: 0.001 to 0.5%,
Ta: 0.001 to 0.5%,
Ni: 0.001 to 0.5%,
Mg: 0.001 to 0.03%,
Ca: 0.001 to 0.03%,
Y: 0.001 to 0.03%,
Zr: 0.001 to 0.03%,
La: 0.001 to 0.03%
Ce: 0.001 to 0.030%
The medium / high carbon hot-rolled steel sheet having excellent punchability according to (1) above, comprising one or more of the above.

(3) The steel sheet according to (1) or (2) is
The ferrite particle size is 10 μm or more and 50 μm or less,
The cementite particle size is 0.1 μm or more and 2.0 μm or less,
Having a structure in which the spheroidization rate of cementite is 85% or more,
A medium / high carbon hot-rolled steel sheet excellent in punchability characterized by having a Vickers hardness of 100HV or more and 160HV or less.

  (4) When the continuous cast slab of the component described in (1) or (2) is directly or once cooled and then heated and hot-rolled, the rough bar is heated after the completion of the rough hot rolling, and then 20 to The temperature is raised to 150 ° C., and finish hot rolling is completed in a temperature range of 600 ° C. or more and less than Ae 3-20 ° C., and the hot-rolled steel sheet picked up at 400 ° C. or more and less than 650 ° C. is directly or pickled and annealed A method for producing a medium / high carbon hot-rolled steel sheet having excellent punchability, characterized by being produced.

  (5) Medium / high carbon hot-rolled steel sheet having excellent punching characteristics, characterized by performing annealing at an annealing temperature of 680 ° C. or more and 720 ° C. or less for an annealing time of 3 hours or more and 60 hours or less as the box annealing described in (4). Manufacturing method.

  (6) As the box annealing described in (4) above, after performing the first stage annealing at an annealing temperature of 680 ° C. or more and 720 ° C. or less and an annealing time of 3 hours or more and 60 hours or less, annealing temperature of 730 ° C. or more and 790 ° C. or less Step 2 annealing is performed at a cooling rate of 1 ° C./hr to 20 ° C./hr to 650 ° C., followed by cooling to room temperature. A method for producing medium- and high-carbon hot-rolled steel sheets excellent in punchability characterized by

  ADVANTAGE OF THE INVENTION According to this invention, the medium and high carbon steel plate which suppresses the droop by punching, and its manufacturing method can be provided. In addition, the punched end face is excellent in properties and the mold life can be extended.

It is a figure which shows the relationship between the azimuth | direction accumulation degree to (110) +/- 5 degree, and punching sagging. It is a figure which shows the optimal manufacturing range of hot-rolling finishing temperature and cutting temperature as an influence of manufacturing conditions. It is a figure which shows the optimal form range of a ferrite particle diameter and a cementite particle diameter as an influence of a structure | tissue.

  Hereinafter, the present invention will be described in detail.

  First, the reason why the chemical components of the steel sheet of the present invention are limited will be described. Here, “%” for a component means mass%.

(C: 0.10 to 0.70%)
C is a powerful element that promotes the growth of dislocations in rolling. Addition of 0.10% or more is desirable because it accumulates in the (110) orientation due to growth and accumulation of dislocations of ferrite transformed during rolling. If it is less than 0.10%, the growth of dislocations in hot rolling is suppressed, so the degree of integration is lowered. For this reason, a minimum is made into 0.10%. On the other hand, if it exceeds 0.70%, ferrite hardly transforms in hot rolling, and accumulation in the (110) direction cannot be obtained, so the upper limit is made 0.70%. More preferably, it is 0.15-0.65%.

(Si: 0.01-1.0%)
Si is an element that acts as a deoxidizer and promotes ferrite transformation during hot rolling. If it is less than 0.01%, the effect of addition cannot be obtained, so the lower limit is made 0.01%. On the other hand, if it exceeds 1.0%, the ductility is lowered, so that ear cracks occur during hot rolling, resulting in a decrease in yield and an increase in cost. For this reason, the upper limit is made 1.0%. More preferably, it is 0.05-0.8%, More preferably, it is 0.08-0.35%.

(Mn: 0.1-3.0%)
Mn is an element that promotes the accumulation of dislocations during rolling. Addition of 0.10% or more is desirable because it accumulates in the (110) orientation due to growth and accumulation of dislocations of ferrite transformed during rolling. If the content is less than 0.10%, accumulation of dislocations in hot rolling can be suppressed, and the degree of integration decreases. For this reason, a minimum is made into 0.10%. On the other hand, if it exceeds 3.0%, the ferrite hardly transforms during hot rolling, and accumulation in the (110) direction cannot be obtained, so the upper limit is made 3.0%. More preferably, it is 0.3-2.5%, More preferably, it is 0.5-1.5%.

(P: 0.001 to 0.025%)
P is an element that promotes the growth of dislocations in rolling. If it is less than 0.001%, the effect of addition cannot be obtained, so the lower limit is made 0.001%. On the other hand, if it exceeds 0.025%, the ductility is lowered, so that ear cracks occur during hot rolling, resulting in a decrease in yield and an increase in cost. For this reason, the upper limit is made 0.025%. More preferably, it is 0.002 to 0.02%.

(S: 0.0001 to 0.010%)
S forms nonmetallic inclusions such as MnS and deteriorates the properties of the punched end face, so the upper limit is made 0.010%. However, reducing S to less than 0.0001% causes a significant increase in refining costs, so the lower limit is made 0.0001%. More preferably, it is 0.0003% to 0.007%.

(Al: 0.001-0.10%)
Al is an element that acts as a deoxidizer and promotes ferrite transformation during hot rolling. If it is less than 0.001%, the effect of addition cannot be obtained sufficiently, so the lower limit is made 0.001%. On the other hand, if it exceeds 0.10%, the effect of addition is saturated, the ductility of the steel sheet is lowered, and the cracking of the hot rolled coil and the embrittlement crack in unwinding are caused. For this reason, the upper limit is made 0.10%. More preferably, it is 0.01%-0.08%, More preferably, it is 0.015-0.04%.

(N: 0.001 to 0.010%)
N is an element that works effectively for the growth of dislocations during hot rolling. If it is less than 0.001%, the effect of addition cannot be sufficiently obtained, so the lower limit is made 0.001%. On the other hand, the excessive content causes a reduction in the ductility of the steel sheet, so the upper limit is made 0.010%. More preferably, it is 0.002% to 0.008%.

  Although this invention makes the said component the basic component of a steel plate, the component described below can be selectively contained in order to improve the mechanical characteristic of a steel plate further.

(Ti: 0.01-0.2%)
Ti is an additive element effective for refining austenite grains during hot rolling. Addition of 0.01% or more is desirable because ferrite transformation is promoted by refinement of austenite grains and orientation accumulation in (110) is promoted. If less than 0.01%, the effect cannot be obtained, so the lower limit is made 0.01%. On the other hand, excessive addition leads to a decrease in the ductility of the steel sheet and causes cracking of the hot-rolled steel strip, so the upper limit is made 0.2%. More preferably, it is 0.015% to 0.1%.

(Cr: 0.01-1.50%)
Cr is an additive element that promotes ferrite transformation during hot rolling and improves hardenability during heat treatment of parts. If less than 0.01%, the effect of large addition cannot be obtained, so the lower limit is made 0.01%. On the other hand, if it exceeds 1.50%, the ductility of the steel sheet is reduced and the hot-rolled steel strip is cracked, so the upper limit is made 1.50%. More preferably, it is 0.05% to 1.10%.

(Mo: 0.01 to 0.50%)
Mo is an additive element that promotes ferrite transformation during hot rolling and improves hardenability during heat treatment of parts. If less than 0.01%, the effect of large addition cannot be obtained, so the lower limit is made 0.01%. On the other hand, if added over 0.50%, the ductility of the steel sheet is reduced, and the hot-rolled steel strip is cracked, so the upper limit is made 0.50%. More preferably, it is 0.05% to 0.30%.

(B: 0.0001 to 0.010%)
B is an additive element that refines austenite grains during hot rolling and improves the hardenability during heat treatment of parts. If it is less than 0.0001%, there is no effect of addition, so the lower limit is made 0.0001%. If it exceeds 0.010%, the ductility of the steel sheet will be reduced, and the hot-rolled steel strip will be cracked, so the upper limit is made 0.010%. More preferably, it is 0.0005% to 0.005%.

(Nb: 0.001-0.10%)
Nb is an element that forms nitrides and is effective in increasing the strength of steel. In order to exhibit the effect effectively, it is preferable to contain 0.001% or more. However, if it exceeds 0.10%, the ductility is lowered, and the steel strip is cracked, so the upper limit is made 0.10%. More preferably, it is 0.01% to 0.08%.

(V: 0.001 to 0.2%)
V, like Nb, is an element that forms nitrides and is effective in increasing the strength of steel. In order to exhibit the effect effectively, it is preferable to contain 0.001% or more. If it exceeds 0.2%, the ductility is reduced and the steel strip is cracked, so the upper limit is made 0.2%. More preferably, it is 0.01% to 0.15%.

(Cu: 0.001 to 0.4%)
Cu is an element that increases the strength of the steel material by forming fine precipitates. In order to effectively exhibit the effect of increasing the strength, the content is preferably 0.001% or more. A higher content is desirable for increasing the strength. However, if the content exceeds 0.4%, hot embrittlement of the steel plate is caused and the plate breakage during hot rolling is moved over, so the upper limit is made 0.4%. More preferably, it is 0.01% to 0.35%.

(W: 0.001 to 0.5%)
W is an element effective for improving the hardenability during heat treatment of parts. In order to exhibit the effect effectively, it is preferable to contain 0.001% or more. If it exceeds 0.5%, the ductility is lowered and the steel strip is cracked, so the upper limit is made 0.5%. More preferably, it is 0.01% to 0.4%.

(Ta: 0.001 to 0.5%)
Ta is an element that forms a nitride, increases the strength of the steel sheet, and improves the hardenability. In order to exhibit the effect effectively, it is preferable to contain 0.001% or more. If it exceeds 0.5%, the ductility is lowered and the steel strip is cracked, so the upper limit is made 0.5%. More preferably, it is 0.01% to 0.4%.

(Ni: 0.001 to 0.5%)
Ni is an element effective for improving the toughness of parts and improving hardenability. In order to exhibit the effect effectively, it is preferable to contain 0.001% or more. However, since Ni is an expensive alloy element, the addition of a large amount causes an increase in cost. For this reason, the upper limit is made 0.5%. More preferably, it is 0.01% to 0.4%.

(Mg: 0.001 to 0.03%)
Mg is an element that can control the form of sulfide by adding a small amount, and can be contained as required. If less than 0.001%, the effect cannot be obtained, so the lower limit is made 0.001% or more. On the other hand, if the content is excessive, a coarse Mg-oxide is formed and the properties of the punched end face are lowered, so the upper limit is made 0.03%. More preferably, it is 0.001 to 0.015%.

(Ca: 0.001 to 0.03%)
Ca, like Mg, is an element that can control the form of sulfide by addition of a trace amount, and can be contained if necessary. If less than 0.001%, the effect cannot be obtained, so the lower limit is made 0.001% or more. On the other hand, if the content is excessive, a coarse Ca-oxide is formed and the properties of the punched end face are lowered, so the upper limit is made 0.03%. More preferably, it is 0.001 to 0.015%.

(Y: 0.001 to 0.03%)
Y is an element that can control the form of the sulfide by addition of a trace amount similarly to Mg and Ca, and can be contained if necessary. If less than 0.001%, the effect cannot be obtained, so the lower limit is made 0.001% or more. Further, if the content is excessive, a coarse Y-oxide is formed and the properties of the punched end face are lowered, so the upper limit is made 0.03%. More preferably, it is 0.001 to 0.015%.

(Zr: 0.001 to 0.03%)
Zr is an element that can control the form of sulfide by addition of a trace amount like Mg, Ca, and Y, and can be contained if necessary. If less than 0.001%, the effect cannot be obtained, so the lower limit is made 0.001% or more. Further, if it is excessively contained, coarse Zr-oxide or the like is formed and the properties of the punched end face are lowered, so the upper limit is made 0.03%. More preferably, it is 0.001 to 0.015%.

(La: 0.001 to 0.03%)
La, like Mg, Ca, Y, and Zr, is an element that is effective for controlling the form of sulfides when added in a trace amount, and can be added as necessary. If less than 0.001%, the effect cannot be obtained, so the lower limit is made 0.001% or more. Further, La easily forms oxides, and those compounds serve as starting points for cracking during molding, leading to a decrease in moldability, so the upper limit is made 0.03%.

(Ce: 0.001 to 0.030%)
Ce, like Mg, Ca, Y, Zr, and La, is an element that can control the form of sulfide by addition of a trace amount, and can be added as necessary. If less than 0.001%, the effect cannot be obtained, so the lower limit is made 0.001% or more. Further, if it is excessively contained, coarse Ce-oxide is formed and the properties of the punched end face are lowered, so the upper limit is made 0.030%. More preferably, it is 0.001 to 0.015%.

  In the steel sheet of the present invention, the balance of the components described above is Fe and inevitable impurities. Moreover, even if it is less than the component range of the selective component described above, it does not inhibit the properties of the steel sheet of the present invention. Furthermore, when scrap is used as a raw material of the steel sheet of the present invention, one or more of Sn, Sb, and As are inevitably mixed in by 0.003% or more, but both are 0.03% or less. Then, since the hardenability of the steel sheet of the present invention is not impaired, Sn: 0.003-0.03%, Sb: 0.003-0.03%, and As: 0.003 in the steel sheet of the present invention. Inclusion of one or more of ˜0.03% is allowed as an inevitable impurity.

  In the steel sheet of the present invention, the amount of O is not specified, but if the oxide aggregates and becomes coarser, the formability deteriorates, so O is preferably 0.003% or less. A smaller amount of O is preferable, but since it is technically difficult to reduce it to less than 0.0001%, a content of 0.0001% or more is allowed as an inevitable impurity.

  The steel sheet of the present invention is subjected to optimum hot rolling and annealing in addition to the above-described component composition, and within ± 5 ° from the iron (110) surface of the body-centered cubic lattice in the region from the steel sheet surface layer to 200 μm in the sheet thickness direction. By setting the degree of accumulation of orientation in which the crystal of the crystal is parallel to the steel sheet surface to 2.5 or more, dripping at the time of punching is remarkably suppressed, and the ferrite grain size is 10 μm or more and 50 μm or less as the structure of the steel sheet. The diameter is 0.1 μm or more and 2.0 μm or less, the spheroidization rate of cementite is 85% or more, and the steel sheet hardness is controlled to 100 HV or more and 160 HV or less by Vickers hardness, thereby improving the properties of the punched end face, Improving the life of the mold is a new finding found by the present inventors.

  In the region from the steel sheet surface layer to the plate thickness direction 200 μm, the accumulation of crystal orientations in which the (110) plane falls within the parallelism within ± 5 ° with respect to the steel sheet surface is made 2.5 or more, and punching sagging is remarkably suppressed. Explain what you can do.

  FIG. 1 shows the relationship between the degree of accumulation of crystal orientations in which the (110) plane is oriented within ± 5 ° with respect to the steel sheet surface and punching sagging. There is a clear correlation between the orientation accumulation degree to (110) and punching sagging, and shearing from the steel sheet surface is controlled by controlling the orientation accumulation degree to 2.5 or more in the region from the steel sheet surface to the center of the plate thickness up to 200 μm. It can be seen that the sagging length to the surface is 200 μm or less, and the punching sagging is remarkably suppressed. As described above, the reason why the punching sag can be suppressed when the orientation accumulation degree to (110) is 2.5 or more is not clear, but the shear deformation that the steel sheet receives from the punch and the die at the time of punching is the iron core. Since it is parallel to (110) which is the main slip surface of the cubic lattice, it is surmised that slip in a direction different from shear deformation is suppressed and punching sag is reduced compared to the case of random orientation. A higher degree of orientation integration is better, and more preferably 3.0 or more. Since the higher azimuth accumulation degree is preferable, the upper limit is not particularly limited, but the effect of suppressing punching sagging is saturated when the azimuth accumulation degree is approximately 5.0.

  The orientation accumulation degree is an index representing how much crystal orientation in a specific direction exists. As the crystal orientation, the rotation angle around a specific axis is divided into 5 ° increments, and each orientation existing in the range is regarded as one identical orientation. In addition, the degree of integration of crystals in each orientation is specified as 1 for all orientations in materials with completely random crystal orientations, and the total amount of integration in all orientations is averaged over all orientations. It is obtained by adjusting to be 1. That is, if the degree of integration in one direction increases, the degree of integration in another direction relatively decreases.

  It is desirable to evaluate the orientation integration degree by EBSD (electron beam backscatter diffraction method). In the conventional XRD (X-ray diffraction) evaluation, the accuracy of obtaining crystal orientation information in the region from the steel plate surface to the 200 μm position depends greatly on the preparation accuracy of the thin film sample. This is because the degree of azimuth accumulation measured and calculated by XRD does not necessarily reflect the actual volume ratio state because the perimeter reflection intensity is different.

  In order to obtain the true information of the sample, attention must be paid to the process of EBSD sample preparation and orientation analysis. Samples for evaluation were cut with a discharge wire processing machine from a steel strip, a cut plate cut out from the steel strip, or a punched blank plate from a place where distortion was not applied, and a plane perpendicular to the steel plate surface was used as the observation surface prepare. Since the measurement accuracy of EBSD is affected by the flatness of the observation surface and the strain given by the polishing, the observation surface is finished to a mirror surface by wet polishing and diamond abrasive polishing, and then polishing for distortion removal is performed. The strain relief polishing is performed using a vibration polishing apparatus (Bueller Vibromet 2) under the conditions of an output of 40% and a polishing time of 60 minutes. In addition, in chemical etching in which the sample surface is dissolved by electropolishing or a hydrogen fluoride-based corrosive solution, the corners of the surface layer portion of the observation surface and the steel plate surface are preferentially dissolved, and sagging occurs, so 200 μm from the steel plate surface by EBSD. It cannot be applied to the measurement of crystal orientation in the region up to the position. If it is SEM-EBSD, the device type of SEM or Kikuchi line detector will not be specifically limited. A sample is set so that the observation surface is 70 ° with respect to the detector of the Kikuchi line, and the surface layer of the plate thickness has an area of 200 μm in the plate thickness direction and 900 μm in the plate width direction at a measurement step interval of 0.5 μm. The field of view is measured, a reverse pole figure of (100) is produced from the obtained crystal orientation information, and the degree of integration within ± 5 ° from (110) is obtained. For the analysis of measurement data, OSL analysis software of TSL is good, and in order to eliminate the data influence of measurement error due to noise, analysis is performed without performing cleanup and excluding data with a CI value of 0.1 or less.

  The punchability can be further improved by setting the ferrite grain size to 10 μm or more and 50 μm or less as the structure of the steel sheet before punching. If the ferrite particle size is less than 10 μm, stretcher strain is generated on the surface of the punched sample and the surface appearance is impaired, so the lower limit is made 10 μm or more. On the other hand, if the ferrite grain size exceeds 50 μm, a textured surface is generated on the surface of the punched sample and the surface appearance is impaired, so the upper limit is made 50 μm or less. The ferrite grain size is measured by a line segment method on a photographed image obtained by observing a structure etched with a 3% nitric acid-alcohol solution with an optical microscope or a scanning electron microscope.

  Moreover, the punchability can be further improved by setting the particle diameter of cementite to 0.1 μm or more and 2.0 μm or less and the spheroidization ratio to 85% or more. When the cementite particle diameter is less than 0.1 μm, the fracture surface develops so as to connect fine carbides at the punched end face, so that the shear plane ratio is lowered and the properties are lowered. Therefore, the lower limit is made 0.1 μm or more. . Further, when the cementite particle diameter exceeds 2.0 μm, a serrated fracture surface is formed so as to connect coarse cementite particles on the punched end face, so the upper limit is made 2.0 μm or less.

  Furthermore, the punchability can be further improved by setting the spheroidization rate of cementite to 85% or more. If it is less than 85%, stress concentrates on the needle-like carbides, and the fracture surface develops so as to connect them, so that the properties of the punched end surface are lowered. For this reason, a minimum is made into 85% or more. The higher the spheroidization rate, the better. However, in order to control it to 100%, it is necessary to perform annealing for a very long time, leading to an increase in manufacturing cost. Therefore, the upper limit is preferably less than 100%. As described above, by controlling the structure and hardness of the steel sheet, it is possible to reduce punching sagging, and to achieve better end face properties (aesthetics) and reduced die wear.

The cementite is observed with a scanning electron microscope. A sample for tissue observation is prepared by wet polishing with emery paper and polishing with diamond abrasive grains having a particle size of 1 μm, and then etching with a saturated picric alcohol solution after finishing the mirror surface. The observation magnification is 1000 to 10,000 times, and in the present invention, 16 fields of view containing 500 or more carbides on the tissue observation surface are selected at 3000 times, and a tissue image is acquired. The area of each carbide contained in the region is measured in detail with the image analysis software represented by Mitani Corporation (Win ROOF) on the obtained tissue image. In order to suppress the influence of measurement error due to noise, carbides having an area of 0.01 μm ² or less are excluded from the evaluation target, and the diameter when the average area per piece is approximated by a circle is obtained as the average cementite particle diameter. The case where the ratio of the major axis length to the minor axis length of each carbide is 3 or more is calculated as acicular carbide, and the case where the ratio is less than 3 is calculated as spherical carbide. A value obtained by dividing the number of spherical carbides by the total number of carbides is defined as the spheroidization rate of cementite.

  Moreover, the punchability can be further improved by setting the hardness of the steel plate (sample) before punching to 100 HV or more and 160 HV or less in terms of Vickers hardness. The Vickers hardness is an average value (arithmetic average value) measured at 10 points at 0.1 mm intervals in the thickness direction with a load of 2.94N. If the Vickers hardness is less than 100 HV, pear texture is generated on the surface of the punched sample, and the aesthetic appearance is impaired. For this reason, a minimum is made into 100HV or more. On the other hand, when the Vickers hardness exceeds 160 HV, the life of the mold is reduced and the punching accuracy is likely to be lowered. For this reason, an upper limit shall be 160 HV or less.

  Next, the manufacturing method of this invention steel plate is demonstrated.

  The technical idea of the production method of the present invention is characterized by consistent management of hot rolling and annealing conditions using the materials in the component ranges described above.

  Specific features of the production method of the present invention are as follows.

  Features of hot rolling: After continuous casting of a slab having a predetermined component, it is heated as it is as usual or once after cooling, and when rolling hot, the rough bar is heated after the rough rolling is finished and before the finish rolling is started. Then, the particle size distribution of the austenite grains in the surface layer is adjusted. After the coarse bar is heated, finish rolling is started, finish hot rolling is finished in a temperature range of 600 ° C. or more and less than Ae 3-20 ° C., and the transformed ferrite is rolled in a hot rolling mill. The steel strip after finish rolling is wound into a hot rolled coil in a temperature range of 400 ° C. or higher and lower than 650 ° C. The hot-rolled coil is subjected to box annealing as it is or after pickling to obtain a medium / high carbon steel sheet having excellent punchability.

  Below, the manufacturing method of this invention is demonstrated concretely.

(Hot rolling)
After continuous casting of a slab having a predetermined component, heat it as it is or once after cooling and rolling it hot, after finishing rough rolling, heating the rough bar before starting finish rolling, and then starting finish rolling. The finish hot rolling is finished in a temperature range of 600 ° C. or more and less than Ae 3-20 ° C., and the obtained steel strip is wound in a temperature range of 400 ° C. or more and less than 650 ° C.

  When the heating temperature of the slab exceeds 1150 ° C, it takes a long time to cool down to Ae3-20 ° C and lowers the productivity. In addition, a thick scale is formed on the surface of the slab to promote the generation of wrinkles on the steel strip. To do. For this reason, slab heating temperature sets 1150 degrees C or less as an upper limit. In addition, when the heating temperature is less than 900 ° C., microsegregation and macrosegregation formed by casting cannot be eliminated, and a region where alloy elements such as Si and Mn are concentrated remains in the steel material after rolling and annealing, and punching is performed. Concavities and convexities corresponding to the amount of segregation are formed on the end face of the component, and the lower limit is set to 900 ° C. in order to impair the aesthetic appearance of the end face.

  After completion of the rough hot rolling, the finished hot rolled material has a surface layer of austenite in a mixed grain state in the longitudinal direction of the coarse bar. This is because the temperature difference caused by the presence or absence of contact with the slab support of the hearth in the hot-rolling heating furnace has an influence even after the end of rough rolling, and there is a temperature difference in the longitudinal direction of the coarse bar. If the hot rolling is started without heating the rough bar before finishing hot rolling, the mixed grain structure will not be eliminated, and ferrite transformation will preferentially start from the fine austenite grains during hot rolling. Immediately after finishing rolling, the amount of strain accumulated in the ferrite existing on the surface layer of the steel sheet becomes non-uniform, resulting in variations in orientation accumulation to (110) or a decrease in orientation accumulation degree. Therefore, in order to uniformly control the austenite grain size of the surface layer of the finished hot rolled material over the longitudinal direction of the coarse bar and to align the timing of ferrite transformation in the hot rolling machine, the longitudinal temperature difference is reduced by heating the coarse bar. Is required. The heating temperature of the rough bar is not particularly limited, but it is considered preferable to heat (temperature increase) to a temperature 20 ° C. higher than the temperature after the end of the rough rolling. Although a higher heating temperature is desirable, a heating scale of 150 ° C. or higher (annealing) generates a thick scale on the surface layer of the steel sheet and causes wrinkle formation on the surface of the steel strip, so the upper limit is set to 150 ° C.

  The finish hot rolling is preferably finished at 600 ° C. or more and less than Ae 3-20 ° C. When the finish hot rolling temperature is less than 600 ° C., the rolling load is remarkably increased due to an increase in deformation resistance of the steel material, and the roll wear amount is further increased, resulting in a decrease in productivity. For this reason, a lower limit shall be 600 degreeC or more. On the other hand, when the temperature is Ae3-20 ° C. or higher, only a small amount of ferrite is transformed during hot rolling of the finish, the accumulation in the (110) direction is lowered, and the effect of the present invention cannot be obtained. For this reason, let an upper limit be less than Ae3-20 degreeC.

  The winding temperature is 400 ° C. or higher and lower than 650 ° C. If the coiling temperature is less than 400 ° C., austenite that has not been transformed during finish rolling is transformed into martensite, causing embrittlement of the coil and causing cracking during unwinding, so the lower limit is 400 ° C. or more. To do. Further, when the coiling temperature is 650 ° C. or higher, untransformed austenite is transformed into pearlite having coarse lamellar, and the properties of the punched end face are lowered. For this reason, an upper limit shall be less than 650 degreeC.

  The punchability can be further improved by subjecting the hot-rolled coil manufactured under the above-described conditions as it is or by performing box annealing after pickling.

The Ae3 temperature is preferably obtained by thermodynamic calculation software such as Thermo-Calc, but in the present invention, the following empirical formula (1) is used instead. In addition, all% in a formula is the mass%.
Ae3 (° C.) = 910−203 × (√% C) −15.2 × (% Ni) + 44.7 × (% Si) + 104 × (% V) + 31.5 × (% Mo) + 13.1 × ( % W) (1)
(Annealing conditions in ferrite + cementite region)
The above-mentioned hot-rolled coil is subjected to box annealing for cooling to room temperature after holding at 680 ° C. to 720 ° C. for 3 hours to 60 hours.

  The annealing temperature is preferably 680 ° C. or more and 720 ° C. or less. When the annealing temperature is less than 680 ° C., ferrite grains and cementite particles are not sufficiently coarsened, causing stretcher strain on the punched parts and deterioration of the properties of the punched end face. For this reason, a minimum is made into 680 degreeC or more. Further, when the annealing temperature exceeds 720 ° C., an austenite phase is generated during annealing. When simply cooled to room temperature in this state, pearlite is newly generated from the austenite phase, causing a decrease in the spheroidization rate of cementite and an increase in material hardness. For this reason, an upper limit shall be 720 degrees C or less.

  The annealing holding time is preferably 3 hr (hour) or more and 60 hr or less. If the holding time is less than 3 hours, the ferrite grains and cementite particles are not sufficiently coarsened, causing the occurrence of stretcher strain on the punched parts and the deterioration of the properties of the punched end face. For this reason, a minimum is made into 3 hours or more. When the holding time exceeds 60 hours, ferrite grains and cementite particles are excessively coarsened, which causes the occurrence of satin on the punched parts and the deterioration of the properties of the punched end face. For this reason, the upper limit is set to 60 hr or less.

(Annealing conditions in ferrite + austenite + cementite region)
The above-mentioned hot-rolled coil is subjected to the first stage annealing for holding at 680 ° C. or more and 720 ° C. or less for 3 hours or more and 60 hours, and then further subjected to the second stage annealing for holding 730 ° C. or more and 790 ° C. or less for 1 hr or more and 12 hours or less. Then, it is preferable to set the cooling rate to 650 ° C. to 20 ° C./hr or less and perform annealing to cool to room temperature. In the first stage of annealing, alloy elements such as Mn are concentrated in cementite to enhance thermal stability. In the subsequent second stage high temperature annealing, austenite is generated, cementite which is not concentrated in alloy elements is dissolved in austenite, and then gradually cooled to adsorb C in the austenite to undissolved cementite. This is because the spheroidization rate of cementite can be increased.

  The first stage annealing temperature is set to 680 ° C. or more and 720 ° C. or less. When the annealing temperature is less than 680 ° C., the concentration of the alloy element in the cementite particles is insufficient, and undissolved cementite cannot remain in the second stage annealing, and after the second stage annealing. Even if the cooling rate is controlled to be gradually cooled, generation of new pearlite cannot be suppressed, which causes a decrease in cementite spheroidization rate and an increase in hardness. For this reason, a minimum is made into 680 degreeC or more. In addition, when the first stage annealing temperature exceeds 720 ° C., an austenite phase is generated during annealing. In the place where the austenite phase is generated by the first stage annealing, undissolved cementite cannot remain in the second stage annealing, and even if the cooling rate after the second stage annealing is controlled to be gradually cooled, Since the formation of new pearlite cannot be suppressed, it causes a decrease in the cementite spheroidization rate and an increase in hardness. For this reason, an upper limit shall be 720 degrees C or less.

  The holding time for the first stage annealing is set to 3 hours or more and 60 hours or less. If the holding time is less than 3 hr, the concentration of the alloy element in the cementite particles is not sufficient, and undissolved cementite cannot remain in the second stage annealing, and the cooling rate after the second stage annealing. Even if it is controlled to be gradually cooled, the formation of new pearlite cannot be suppressed, which causes a decrease in the spheroidization rate of cementite and an increase in hardness. For this reason, a minimum is made into 3 hours or more. If the holding time exceeds 60 hours, the ferrite grains become too coarse, and the punched parts are caused to have a satin finish. For this reason, the upper limit is set to 60 hr or less.

  The second stage annealing temperature is set to be 730 ° C. or higher and 790 ° C. or lower. If the annealing temperature is less than 730 ° C., an austenite phase is not generated during annealing. For this reason, a minimum is made into 730 degreeC or more. Further, if the annealing temperature of the second stage exceeds 790 ° C, the dissolution of cementite in the austenite phase is promoted, and it becomes difficult to leave undissolved cementite, so even if it is gradually cooled after the second stage annealing, The formation of new pearlite cannot be suppressed, causing a decrease in cementite spheroidization rate and an increase in hardness. For this reason, an upper limit shall be 790 degrees C or less.

  The holding time for the second stage annealing is 1 hr or more and 12 hr or less. If the holding time is less than 1 hr, the entire coil cannot be made uniform temperature by box annealing, and material fluctuation occurs in the longitudinal direction of the coil and the product power is reduced. Therefore, the lower limit is set to 1 hr or more. When the holding time exceeds 12 hours, the dissolution of cementite in the austenite phase is promoted, and it becomes difficult to leave undissolved cementite. For this reason, an upper limit shall be 12 hr or less.

  The cooling rate to 650 ° C. after the second stage annealing is 20 ° C./hr or less. The slower the cooling rate, the lower the moving speed of the ferrite / austenite interface, and the C adsorption to undissolved cementite is promoted. On the other hand, if the time is too slow, the annealing time becomes longer, leading to a decrease in productivity. Therefore, the cooling rate is preferably 1 ° C./hr or more. On the other hand, when the cooling rate exceeds 20 ° C./hr, the moving speed of the ferrite / austenite interface increases, C adsorption to undissolved cementite becomes insufficient, and pearlite is newly generated from such a location. This causes a decrease in the spheroidization rate of cementite and an increase in the hardness of the material, so the upper limit is made 20 ° C./hr or less. Cooling after the second stage annealing is important for promoting C adsorption to undissolved cementite by controlling the cooling rate to 650 ° C., and then performing step type annealing for cooling to room temperature.

  The atmosphere of the box annealing is not particularly limited, and may be any condition of 95% or more nitrogen atmosphere, 95% or more hydrogen atmosphere, or air atmosphere.

  According to the manufacturing method of the present invention described above, in the region from the steel sheet surface layer to 200 μm in the thickness direction, crystals within ± 5 ° from the (110) plane of iron of the body-centered cubic lattice are parallel to the steel sheet surface. By controlling the degree of integration in a certain direction to 2.5 or more and obtaining a steel sheet that can significantly suppress sagging during punching, and further optimizing the annealing conditions, the ferrite grain size of the material is 10 μm or more and 50 μm or less. The cementite particle size is controlled to 0.1μm or more and 2.0μm or less, the cementite spheroidization ratio is controlled to 85% or more, the Vickers hardness is controlled to 100HV or more and 160HV or less, the properties of the punched end face are improved, and the die life is improved. It is possible to obtain a medium / high carbon steel plate having excellent punchability.

  Next, effects of the present invention will be described with reference to examples.

  The level of the embodiment is an example of execution conditions adopted to confirm the feasibility and effects 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.

  Continuous cast slabs (steel ingots) having the composition shown in Table 1-1 to Table 1-3 are hot-rolled after heating for 1 hr at 1030 ° C., and after hot-rolling a 250 mm slab to 30 mm, finish hot rolling The raw material bar was heated to 38 ° C. and finishing hot rolling was started. After finishing hot rolling at 700 ° C., winding was performed at 510 ° C. to produce a hot rolled coil having a thickness of 3.5 mm. The hot-rolled coil was pickled, annealed at 700 ° C. for 14 hours in a 95% hydrogen-5% nitrogen atmosphere, and a sample for punching evaluation was produced. The punching test was performed under the conditions of a die diameter of 10.7 mm, a punch diameter of 10 mm, and a one-side clearance of 10%. The surface sagging of the sample punched into a circle was measured, and the punching property was evaluated by observing the surface properties. The measurement of the punching sag length and the observation of the surface properties of the punched end face were carried out using a KEYENCE microscope (VHX-1000). The orientation of each sample was measured by the method described in paragraphs (0057) to (0059), the structure in paragraphs (0060) to (0063), and the hardness in the paragraph (0064). Tables 2-1 and 2-2 show the evaluation results of the punchability of the annealed samples manufactured.

  As shown in Tables 2-1 and 2-2, No. 2, 3, 4, 6, 7, 8, 9, 10, 13, 15, 16, 18, 19, 21, 22, 23, 26, 27, 28, 29, 31, 33, 34, 35, 36, 37, 38, 39, 40, 41, 42, 43, 44, 45 are all parallel to the (110) plane within ± 5 ° with respect to the steel sheet surface in the region from the steel sheet surface layer to the plate thickness direction 200 μm. The degree of accumulation of the crystal orientation that fits was 2.5 or more, the amount of punching sagging (mm) was low, and good punching end face properties were shown.

  On the other hand, in Comparative Example 1, since the amount of C in the steel component was low, strain accumulation in hot rolling was suppressed, so the degree of integration was reduced, the amount of punching sagging was high, and the punching end face properties were inferior. . In Comparative Example 5, ear cracks occurred during hot rolling due to a decrease in ductility due to high N. In Comparative Example 11, since the amount of S was high, coarse sulfides were formed, and the quality of the punched end face was lowered. In Comparative Example 12, ear cracks occurred during hot rolling due to a decrease in ductility due to the high Si content. In Comparative Example 14, ear cracks occurred during hot rolling because of a decrease in ductility due to the high amount of P. In Comparative Example 17, ear cracks occurred during hot rolling due to a decrease in ductility due to the high amount of Al. In Comparative Example 20, since the ferrite transformation in hot rolling was suppressed due to the high amount of Mn, the degree of integration was reduced, the amount of punching sagging was high, and the punching end face properties were inferior. In Comparative Example 24, since the ferrite transformation in hot rolling was suppressed due to the high amount of C, the degree of integration was lowered, the amount of punching sagging was high, and the punching end face properties were inferior. In Comparative Example 25, ear cracks occurred during hot rolling due to a decrease in ductility due to the high Ti content. In Comparative Example 30, ear cracks occurred during hot rolling due to a decrease in ductility due to the high Cr content. In Comparative Example 32, due to the formation of coarse Ce-Oxide due to the high amount of Ce, the quality of the punched end face was lowered. In Comparative Example 34, ear cracks occurred during hot rolling because of a decrease in ductility due to the high amount of Mo. In Comparative Example 35, ear cracks occurred during hot rolling because of a decrease in ductility due to a high Nb content.

In order to investigate the influence of hot rolling conditions, 2, 3, 4, 6, 7, 8, 9, 10, 13, 15, 16, 18, 19, 21, shown in Table 1-1 to Table 1-3, Cast a slab having the components 22, 23, 26, 27, 28, 29, 31, 33, 34, 35, 36, 37, 38, 39, 40, 41, 42, 43, 44, 45, Alternatively, after hot cooling at 1060 ° C. for 0.5 to 2 hours, rough hot rolling is started, a rough bar having a thickness of 28 mm is manufactured, and the hot bar is heated before finishing hot rolling, or finish hot rolling without heating. Punching test sample that investigates the influence of manufacturing conditions after finishing hot rolling at various temperatures and winding on a coil, annealing at 700 ° C. for 14 hours in a pickled or unpickled state Was made.
Tables 3-1 and 3-2 show the manufacturing conditions and the evaluation results of the punchability of the manufactured annealing samples.

  As shown in Table 3-1 and Table 3-2, Invention Examples 2-B, 3-A, 4-A, 6-B, 7-A, 8-B, 9-B, 10-B, 13- A, 15-A, 16-B, 18-A, 19-A, 21-B, 22-B, 23-B, 26-B, 27-A, 28-B, 29-B, 31-A, 33-B, 36-A, 36-B, 37-A, 37-B, 38-A, 38-B, 39-A, 39-B, 40-A, 40-B, 41-A, 41- B, 42-A, 42-B, 43-A, 43-B, 44-A, 44-B, 45-A, 45-B are all in the region from the steel plate surface layer to the plate thickness direction 200 μm ( 110) The degree of accumulation of crystal orientations in which the plane falls within a parallelism within ± 5 ° with respect to the steel sheet surface is 2.5 or more, and the amount of punching sag (mm) is as low as 0.15 mm or less, and a good punching end face sex It is shown.

  On the other hand, Comparative Example 2-A is an example in which the steel sheet becomes brittle due to low CT (winding temperature), cracks occur during coil unwinding, and manufacturing conditions are not appropriate. In Comparative Example 3-B, since the surface layer became mixed grains due to the non-heating of the coarse bar, the orientation accumulation degree to (110) was less than 2.5, the amount of punching sag (mm) was high, and the punchability was high It was inferior. In Comparative Example 4-B, the rolling load increased due to the low FT (hot rolling finishing temperature), so that the plate passing property was lowered. Further, the steel plate became brittle due to the low CT, and cracking occurred when the coil was unwound. Since Comparative Example 6-A had a high FT, the orientation accumulation in (110) was less than 2.5, the amount of punching sagging (mm) was high, and the punching end face properties were inferior.

  Since Comparative Example 7-B had a high FT, the orientation accumulation in (110) was less than 2.5, and further, a coarse pearlite lamella was generated by high CT, and the punched end face properties were deteriorated.

  In Comparative Example 8-A, the rolling load increased due to the low FT, so that the plate passing property was lowered.

  Since Comparative Example 9-A had a high FT, the orientation accumulation in (110) was less than 2.5, the amount of punching sagging was high, and the punching end face properties were inferior.

  In Comparative Example 10-A, a rough pearlite lamella was generated by high CT, and thus the punched end face properties were deteriorated.

  Since Comparative Example 13-B had a high FT, the orientation accumulation in (110) was less than 2.5, and the steel sheet became brittle due to low CT, and cracks occurred when the coil was unwound.

  In Comparative Example 15-B, a rough pearlite lamella was generated by high CT, and thus the punched end face properties were deteriorated.

  In Comparative Example 16-A, the rolling load was increased due to the low FT, and thus the plate passing property was lowered.

  In Comparative Example 18-B, a rough pearlite lamella was generated by high CT, and thus the punched end face properties were deteriorated.

  In Comparative Example 19-B, the rolling load increased due to the low FT, so that the sheet passing property decreased. Further, the steel plate became brittle due to the low CT, and cracking occurred when the coil was unwound.

  In Comparative Example 21-A, the surface layer became mixed grains due to the non-heating of the coarse bar, so the orientation accumulation degree to (110) was less than 2.5, the amount of punching sagging was high, and the punching end face property was inferior. It was.

  In Comparative Example 22-A, the steel plate became brittle due to low CT, and cracking occurred when the coil was unwound.

  Since Comparative Example 23-A had a high FT, the orientation accumulation in (110) was less than 2.5, and the steel plate became brittle due to low CT, and cracks occurred when the coil was unwound.

  Since Comparative Example 26-A had a high FT, the orientation accumulation in (110) was less than 2.5, and further, a coarse pearlite lamella was generated by a high CT, so that the punched end face properties were deteriorated.

  In Comparative Example 27-B, the orientation accumulation in (110) was less than 2.5 due to the high FT, the amount of punching sagging was high, and the punching end face properties were inferior.

Since Comparative Example 28-A has a high FT, the orientation accumulation in (110) is less than 2.5, and a coarse pearlite lamella is generated by high CT, so that the punching end face properties are deteriorated. Since the surface layer became mixed grains due to non-heating of the bar, the degree of orientation accumulation in (110) was less than 2.5, the amount of punching sagging was high, and the punching end face properties were inferior.

  In Comparative Example 31-B, a rough pearlite lamella was generated due to high CT, and thus the punched end face properties were lowered.

In Comparative Example 33-A, the rolling load was increased due to the low FT, and thus the sheet passing property was lowered.
As described above, if the manufacturing conditions are not appropriate, it is difficult to manufacture a high-quality steel sheet.

Subsequently, 2, 3, 4, 6, 7, 8, 9, 10, 13, 15, 16, 18, 19, 21, 22 shown in Table 1 to examine the influence of annealing conditions in the ferrite + cementite region. , 23, 26, 27, 28, 29, 31, 33, 34, 35, 36, 37, 38, 39, 40, 41, 42, 43, 44, 45, cast as-is, or After cooling, the hot rolling is started after heating at 1010 ° C. at 0.5 to 2 hours, and the rough bar having a thickness of 35 mm after the hot rolling is heated before the finishing rolling is started, and then the finishing hot rolling is started. Finishing hot rolling at a temperature and winding on a coil were performed, and annealing was performed in a pickled or non-pickled state, and a sample for a punch test for investigating the influence of manufacturing conditions was produced.
Table 4 shows the annealing conditions and the evaluation results of the punchability of the manufactured annealing samples.

  As shown in Table 4, Invention Examples 2-C, 3-C, 4-C, 6-C, 8-C, 9-C, 16-C, 18-C, 19-C, 21-C, 23 -C, 27-C, 29-C, 31-C, 33-C, 36-C, 37-C, 38-C, 39-C, 40-C, 41-C, 42-C, 43-C 44-C and 45-C both have a degree of accumulation of crystal orientations in which the (110) plane falls within the parallelism within ± 5 ° with respect to the steel plate surface in the region from the steel sheet surface layer to the plate thickness direction of 200 μm. The amount of punching sagging (mm) was low, and the punched end face properties were good.

  On the other hand, Comparative Example 7-C had a high annealing temperature, and the punched end face properties were deteriorated due to acicular cementite.

  In Comparative Example 10-C, the annealing temperature was high, and the punched end face properties were lowered due to the acicular cementite.

  In Comparative Example 13-C, the annealing time was too long, pear texture was generated around the sagging surface, and the sagging amount increased due to low hardness.

  In Comparative Example 15-C, the annealing time was too short, and the fracture surface ratio was increased due to the fine cementite, so that the end face properties were lowered, and the mold was damaged due to high hardness.

  In Comparative Example 22-C, since the annealing temperature was low, stretcher strain was generated around the sag surface, and the mold was damaged due to high hardness. In Comparative Example 26-C, since the annealing temperature was low, the fracture surface ratio was increased due to the fine cementite, and thus the punched end face properties were deteriorated.

  In Comparative Example 38-C, the annealing time was too long, the punched end face properties were lowered due to coarse cementite, and the sagging amount was increased due to low hardness.

  Finally, in order to investigate the influence of annealing conditions in the ferrite + austenite + cementite region, 2, 3, 4, 6, 7, 8, 9, 10, 13, 15, shown in Table 1-1 to Table 1-3, Casting slabs with components 16, 18, 19, 21, 22, 23, 26, 27, 28, 29, 31, 33, 36, 37, 38, 39, 40, 41, 42, 43, 44, 45 As it is, or after cooling, it starts rough hot rolling after heating at 1040 ° C. for 0.5 to 2 hours. After heating the rough bar having a thickness of 33 mm after rough rolling before finish rolling, finish hot rolling is started. Samples for punching tests to investigate the effects of manufacturing conditions by performing hot rolling and coiling on coils at various temperatures, annealing in ferrite + austenite + cementite region in pickled or unpickled state MadeTable 5 shows the manufacturing conditions and the evaluation results of the punchability of the manufactured annealed samples.

  As shown in Table 5, Invention Examples 2-D, 3-D, 4-D, 6-D, 7-D, 8-D, 15-D, 18-D, 19-D, 22-D, 28 -D, 29-D, 31-D, 33-D, 36-D, 37-D, 38-D, 39-D, 40-D, 41-D, 42-D, 43-D, 44-D , 45-D has a crystal orientation accumulation degree of 2.5 or more in which the (110) plane falls within the parallelism within ± 5 ° with respect to the steel plate surface in the region from the steel sheet surface layer to the plate thickness direction of 200 μm. Yes, the amount of punching sagging (mm) was low, and good punching end face properties were shown.

  On the other hand, in Comparative Example 9-D, the annealing temperature at the first stage was high, and the punching end face properties were deteriorated due to acicular cementite.

  In Comparative Example 10-D, the annealing temperature at the second stage was high, the punched end face properties were lowered due to acicular cementite, and the mold was damaged due to high hardness.

  In Comparative Example 13-D, the annealing time in the second stage was long, the punched end face properties were lowered due to acicular cementite, and the mold was damaged due to high hardness.

  In Comparative Example 16-D, the annealing time of the first stage was long, and a satin finish was generated around the sag surface, and the aesthetic appearance of the punched member was lowered.

In Comparative Example 21-D, the annealing time of the first stage was short, the punched end face property was lowered due to acicular cementite, and the mold was damaged due to high hardness. Comparative Example 23-D had a high cooling rate. In comparison example 26-D, the die end face was deteriorated due to needle-like cementite, and the mold was damaged due to high hardness. In Comparative Example 26-D, the annealing time of the second stage was short, and material variation occurred in the coil longitudinal direction. did.

  In Comparative Example 27-D, the annealing temperature at the first stage was low, the punched end face properties were lowered due to acicular cementite, and the mold was damaged due to high hardness.

  Table 6-1 and Table 6-2 show the tissue morphology of the samples prepared in Tables 4 and 5. By controlling the ferrite particle size of the material from 10 μm to 50 μm, the cementite particle size from 0.1 μm to 2.0 μm, the spheroidization rate of cementite to 85% or more, and the Vickers hardness from 100 HV to 160 HV, It is clear that the aesthetic appearance (reduction in punching sagging), the properties of the punched end face, and further the extension of the die life can be further improved.

  That is, as shown in Table 6-1 and Table 6-2, Invention Examples 2-C, 3-C, 4-C, 6-C, 8-C, 9-C, 16-C, 18-C, 19-C, 21-C, 23-C, 27-C, 29-C, 31-C, 33-C, 36-C, 37-C, 38-C, 39-C, 40-C, 41- C, 42-C, 43-C, 44-C, 45-C, 2-D, 3-D, 4-D, 6-D, 7-D, 8-D, 15-D, 18-D, 19-D, 22-D, 28-D, 29-D, 31-D, 36-D, 37-D, 38-D, 39-D, 40-D, 41-D, 42-D, 43- D, 44-D and 45-D all have a cementite particle size of 0.1 μm or more and 2.0 μm or less, a cementite spheroidization ratio of 85% or more, a ferrite particle size of 10 μm or more and 50 μm or less, and a Vickers hardness of 100 HV. It becomes less over 160HV, decreased punching sag, properties of the punched end surface was further showed good results in the extension of die life.

  In contrast, in Comparative Example 7-C, the spheroidization rate of cementite was low, and the punched end face properties were lowered due to the acicular cementite.

  In Comparative Example 10-C, the spheroidization rate of cementite was low, and the punched end face properties were lowered due to the acicular cementite.

  In Comparative Example 13-C, a satin texture was generated around the sagging surface, and the sagging amount increased due to low hardness.

  In Comparative Example 15-C, the cementite particle diameter was too small, and the fractured surface ratio was increased due to fine cementite. Therefore, the end face properties were lowered, and the mold was damaged due to high hardness.

  In Comparative Example 22-C, stretcher strain occurred around the sag surface, and the mold was damaged due to high hardness.

  In Comparative Example 26-C, the fractured surface ratio increased due to the fine cementite, and therefore the end face properties decreased.

  In Comparative Example 28-C, the punched end face properties decreased due to coarse cementite, and the amount of sagging increased due to low hardness.

  In Comparative Example 9-D, the spheroidization rate of cementite was low, and the punched end face properties were lowered due to the acicular cementite.

  In Comparative Example 10-D, the spheroidization rate of cementite was low, the punched end face properties were deteriorated due to acicular cementite, and the mold was damaged due to high hardness.

  In Comparative Example 13-D, the spheroidization rate of cementite was low, the punched end face properties were deteriorated due to acicular cementite, and the mold was damaged due to high hardness.

  In Comparative Example 16-D, the ferrite grain size was small, a satin texture was generated around the sag surface, and the appearance of the punched member was lowered.

  In Comparative Example 21-D, the spheroidization rate of cementite was low, the punched end face property was lowered due to acicular cementite, and the mold was damaged due to high hardness.

  In Comparative Example 23-D, the spheroidization rate of cementite was low, the punched end face properties were deteriorated due to acicular cementite, and the mold was damaged due to high hardness.

  In Comparative Example 26-D, material variation occurred in the longitudinal direction of the coil, and the sagging amount increased due to low hardness.

  In Comparative Example 27-D, the spheroidization rate of cementite was low, and the punched end face properties were lowered due to the acicular cementite.

  FIG. 2 shows the optimum range of manufacturing conditions created based on the examples of Tables 3-1 and 3-2. FIG. 2 is a diagram showing the optimum manufacturing range of the hot rolling finishing temperature and the coiling temperature. In the figure, a circle indicates an invention example and a cross indicates a comparative example. Finishing hot rolling in a temperature range of 600 ° C. or more and less than Ae 3-20 ° C., and winding it at 400 ° C. or more and less than 650 ° C. It can be seen that a rolled steel sheet can be obtained.

  FIG. 3 shows an optimum structure map of the medium and high carbon steel plates having excellent punchability created from the examples in Table 6-1 and Table 6-2. FIG. 3 is a diagram showing the optimum form ranges of the ferrite particle size and the cementite particle size. The circles indicate examples with the best punchability (Best), and the crosses indicate examples with the best punchability (Better). In addition to the reduction of punching by controlling the crystal orientation in the region of 200 μm from the surface of the steel sheet, in particular, the structure of the steel sheet has a ferrite particle size of 10 μm to 50 μm and a cementite particle size of 0.1 μm to 2.0 μm. By using the structure (example of ◯), it is possible to reduce the properties of the punched end face and the wear of the mold, and it is understood that the punchability is further improved as compared to the example indicated by X.

Claims (6)

% By mass
C: 0.10 to 0.70%,
Si: 0.01 to 1.0%,
Mn: 0.1 to 3.0%
P: 0.001 to 0.025%,
S: 0.0001 to 0.010%,
Al: 0.001 to 0.10%,
N: 0.001 to 0.010%,
And the balance is a steel plate consisting of Fe and impurities,
Punchability characterized in that in the region from the steel sheet surface layer to the plate thickness direction 200 μm, the degree of accumulation of crystal orientations in which the (110) plane falls within the parallelism within ± 5 ° with respect to the steel sheet surface is 2.5 or more Medium and high carbon hot-rolled steel sheet with excellent resistance. The steel sheet is in mass% as an additive element, and
Ti: 0.01-0.20%,
Cr: 0.01 to 1.50%,
Mo: 0.01 to 0.50%,
B: 0.0001 to 0.010%,
Nb: 0.001 to 0.10%,
V: 0.001 to 0.2%,
Cu: 0.001 to 0.4%
W: 0.001 to 0.5%,
Ta: 0.001 to 0.5%,
Ni: 0.001 to 0.5%,
Mg: 0.001 to 0.03%,
Ca: 0.001 to 0.03%,
Y: 0.001 to 0.03%,
Zr: 0.001 to 0.03%,
La: 0.001 to 0.03%
Ce: 0.001 to 0.030%
The medium / high carbon hot-rolled steel sheet having excellent punchability according to claim 1, wherein one or more of them are contained. The steel sheet according to claim 1 or 2,
The ferrite particle size is 10 μm or more and 50 μm or less,
The cementite particle size is 0.1 μm or more and 2.0 μm or less,
Having a structure in which the spheroidization rate of cementite is 85% or more,
A medium / high carbon hot-rolled steel sheet excellent in punchability characterized by having a Vickers hardness of 100HV or more and 160HV or less.   When the continuous cast slab of the component according to claim 1 or 2 is directly or once cooled and then heated and hot-rolled, the rough bar is heated after the completion of the rough hot rolling to raise the temperature by 20 to 150 ° C. Finishing hot rolling in a temperature range of 600 ° C. or more and less than Ae 3-20 ° C., and hot rolled steel sheet picked up at 400 ° C. or more and less than 650 ° C. as it is or pickled and manufactured by box annealing A method for producing medium and high carbon hot-rolled steel sheets with excellent punchability.   The method for producing a medium / high carbon hot-rolled steel sheet having excellent punchability, characterized by performing annealing at an annealing temperature of 680 ° C. or more and 720 ° C. or less as an annealing time of 3 hr or more and 60 hr or less as the box annealing according to claim 4. .   As the box annealing according to claim 4, after performing the first stage annealing at an annealing temperature of 680 ° C. or more and 720 ° C. or less and an annealing time of 3 hours or more and 60 hours or less, it is 1 hour or more at an annealing temperature of 730 ° C. or more and 790 ° C. or less. It is characterized by performing a second-stage annealing with an annealing time of 12 hr or less, cooling to 650 ° C. at a cooling rate of 1 ° C./hr or more and 20 ° C./hr or less, and then performing step-type annealing to cool to room temperature. A method for producing medium- and high-carbon hot-rolled steel sheets with excellent punchability.

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