JP2005133199A - High carbon cold rolled steel sheet, and method for manufacturing the same - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a high carbon cold rolled steel sheet in which cracks at the blanked edge faces are hard to be generated, and which has excellent stretch flanging property or ductility as well. <P>SOLUTION: The high carbon cold rolled steel sheet comprises 0.20 to 0.58% C, ≤0.1% Si, 0.20 to 0.60% Mn, ≤0.02% P, ≤0.01% S, ≤0.1% sol.Al, ≤0.005% N, 0.001 to 0.005% B and 0.05 to 0.3% Cr, and has a structure in which the average grain size of ferrite is ≤6 μm, the average grain size of carbide is 0.1 to <1.20 μm, and the volume ratio of the ferrite grains substantially including no carbide is ≤15%. In the method of manufacturing a high carbon cold rolled steel sheet, the steel having the above composition is subjected to hot rolling under the conditions where the finishing temperature is ≥(an Ar<SB>3</SB>point-10°C), the cooling velocity is >120°C/s, the cooling stop temperature is ≤620°C, and the coiling temperature is ≤600°C, is thereafter subjected to cold rolling at a draft of ≥30%, and is annealed at 640°C to an Ac<SB>1</SB>transformation point. <P>COPYRIGHT: (C)2005,JPO&NCIPI

Description

  The present invention relates to a high-carbon cold-rolled steel sheet that is used for structural parts of automobiles and the like and has excellent stretch flangeability having a strength of a cold-rolled steel sheet that is 440 MPa or more, or further excellent in ductility, and a method for producing the same.

  High carbon steel sheets used for tools or automobile parts (gears, missions) and the like are subjected to heat treatment such as quenching and tempering after punching and forming. One of the requirements of users who perform these parts processing is to improve hole expansion processing (burring) in forming after punching. This hole expansion workability is evaluated as stretch flangeability as press formability. Therefore, a material excellent in stretch flangeability is desired. In addition, when forming into a complex shape, it is also required that the elongation characteristic, which is an index of ductility, is good.

  Several techniques have been studied for improving the stretch flangeability of such a high-carbon steel sheet. For example, Patent Document 1 proposes a method for producing a medium / high carbon steel sheet having excellent stretch flangeability in a process after cold rolling. This technique consists of steel containing C: 0.1 to 0.8% by mass, the metal structure is substantially a ferrite + pearlite structure, and the pro-eutectoid ferrite area ratio is determined by the C content (% by mass) as necessary. Hot rolling steel sheets with a perlite lamellar spacing of 0.1μm or more are given a cold rolling of 15% or more, and then subjected to three-stage or two-stage annealing for a long time in a three-stage or two-stage temperature range. It is to give.

  Patent Document 2 discloses that the heat of pro-eutectoid ferrite + pearlite structure is made of steel containing C: 0.1 to 0.8% by mass, and the pro-eutectoid ferrite area ratio (%) is not less than a predetermined value determined by the C content. A technique is disclosed in which a steel sheet is subjected to annealing in which the first stage of heating and holding and the second stage of heating and holding are continuously performed.

Further, Patent Document 3 proposes a high carbon hot-rolled steel sheet excellent in stretch flangeability. This is because steel containing 0.2 to 0.7% by mass of C is hot-rolled at a finishing temperature (Ar ₃ transformation point -20 ° C) or higher and then cooled at a cooling rate exceeding 120 ° C / second and a cooling stop temperature of 650 ° C or lower. And then winding at a coiling temperature of 600 ° C. or lower, pickling, and annealing at an annealing temperature of 640 ° C. or higher and an Ac ₁ transformation point or lower. The metal structure is characterized by controlling the average particle size of carbide to 0.1 μm or more and less than 1.2 μm and the volume fraction of ferrite grains not containing carbide to 10% or less.

Patent Document 4 proposes a high carbon cold-rolled steel sheet having excellent stretch flangeability. This is because steel containing 0.2 to 0.7 mass% of C is hot-rolled at a finishing temperature (Ar ₃ transformation point -20 ° C) or higher and then cooled at a cooling rate of over 120 ° C / second and a cooling stop temperature of 650 ° C or lower. And then winding at a coiling temperature of 600 ° C. or lower, pickling, cold rolling at a cold pressure ratio of 30% or higher, and annealing at an annealing temperature of 600 ° C. or higher and an Ac ₁ transformation point or lower. The metal structure is characterized by controlling the average particle size of carbide to 0.1 μm or more and less than 2.0 μm, and the volume fraction of ferrite grains not containing carbide to 15% or less.
Japanese Patent Laid-Open No. 11-269552 Japanese Patent Laid-Open No. 11-269553 Japanese Patent Laid-Open No. 2003-13145 Japanese Patent Laid-Open No. 2003-13144

  In the techniques described in Patent Documents 1 and 2, since the ferrite structure is composed of pro-eutectoid ferrite and does not substantially contain carbide, it is soft and excellent in ductility, but stretch flangeability is not necessarily good. This is considered as follows. That is, during the punching process, the pro-eutectoid ferrite portion is greatly deformed in the vicinity of the punched end face, and therefore the deformation amount differs greatly between the pro-eutectoid ferrite and the ferrite containing the spheroidized carbide. As a result, stress concentrates in the vicinity of the grain boundaries of the grains having greatly different deformation amounts, and voids are generated at the interface between the spheroidized structure and the ferrite. Since this grows into a crack, it is considered that the stretch flangeability is deteriorated as a result.

  As a countermeasure against this, it is conceivable to soften the whole by strengthening the spheroidizing annealing. However, in that case, the spheroidized carbides become coarse and become the starting point of void generation during processing, and the carbides are difficult to dissolve in the heat treatment stage after processing, leading to a decrease in quenching strength.

  In recent years, demands for processing levels from the viewpoint of productivity improvement have become stricter than ever before. Therefore, also in the hole expanding process of the high carbon steel sheet, the punched end face is likely to be cracked due to an increase in the degree of processing. Therefore, a high stretch steel sheet is also required to have high stretch flangeability.

  In view of such circumstances, the present inventors have been able to manufacture without using multi-stage annealing requiring a long time, and for the purpose of providing a high carbon steel sheet excellent in stretch flangeability in which cracking of the punched end face is unlikely to occur. The technologies described in Patent Documents 3 and 4 were developed. These technologies have made it possible to produce high carbon hot rolled steel sheets or high carbon cold rolled steel sheets having excellent stretch flangeability.

  Recently, for applications such as driveline parts, the strength of non-heat treated parts has been increasing with integral molded parts from the viewpoint of high durability and light weight, and the tensile strength (TS) of the steel plate is 440 MPa. The above strength has been demanded.

  In addition, in the integral molding, there are more than 10 pressing processes, and not only burring, but also molding is performed by complex combinations of molding modes such as overhang and bending, so stretch flangeability and ductility Both of these characteristics are required at the same time.

  However, with the techniques described in Patent Documents 3 and 4 above, when trying to achieve TS ≧ 440 MPa (73 points or more in terms of HRB hardness, 135 points or more in terms of Hv hardness), sufficient stretch flangeability is not necessarily obtained. It was. That is, stretch flangeability is evaluated by the hole expansion ratio (λ). In high-carbon cold-rolled steel sheets, λ ≧ 80%, preferably λ ≧ 85% is desired. At the same time, it was not possible to secure a stable demand. Also, no mention was made of ductility.

  In view of such circumstances, the present invention can be manufactured without using a multi-stage annealing that requires a long time, cracking of the punched end face is unlikely to occur, the tensile strength is 440 MPa or more, and the hole expansion ratio λ ≧ 80%, preferably Is intended to provide a high carbon cold-rolled steel sheet with excellent stretch flangeability satisfying λ ≧ 85%, and moreover, when particularly excellent ductility is required, ductility and stretch flange satisfying 35% elongation. It aims at providing the high carbon cold-rolled steel plate excellent in property.

  The above problem is solved by the following invention. The invention is, in mass%, C: 0.20 to 0.58%, Si: 0.1% or less, Mn: 0.20 to 0.60%, P: 0.02% or less, S: 0.01% or less, sol. Al: 0.1% or less, N: 0.005% or less, B: 0.001 to 0.005%, Cr: 0.05 to 0.3%, the composition of the balance iron and inevitable impurities, the average ferrite grain size of 6 μm or less, the average carbide A high carbon cold-rolled steel sheet having a structure in which a grain size is 0.1 μm or more and less than 1.20 μm, and a volume fraction of ferrite grains substantially not containing carbide is 15% or less.

  In the invention of the high-carbon cold-rolled steel sheet, a high-carbon cold-rolled steel sheet characterized by having an average carbide particle size of 0.5 μm or more and less than 1.20 μm can also be provided.

  In the invention of the high carbon cold-rolled steel sheet, a high-carbon cold-rolled steel sheet characterized in that the volume fraction of ferrite grains substantially free of carbides is 10% or less.

As an invention of the manufacturing method, a steel material having the above composition is hot-rolled at a finishing temperature equal to or higher than the finishing temperature (Ar ₃ transformation point −10 ° C.), and then the cooling rate exceeds 120 ° C./second and the cooling stop temperature is 620 ° C. Cooling is performed as follows, and then winding is performed at a coiling temperature of 600 ° C. or less, cold rolling is performed at a reduction rate of 30% or more, and then annealing is performed at an annealing temperature of 640 ° C. or more and an Ac ₁ transformation point or less. It is a manufacturing method of a carbon cold-rolled steel sheet.

As an invention of the manufacturing method, in the above-mentioned invention, it is also possible to provide a manufacturing method of a high carbon cold-rolled steel sheet characterized in that the annealing is performed at an annealing temperature of 680 ° C. or more and an Ac ₁ transformation point or less.

  Further, as an invention of a manufacturing method, in the above-described invention, further, the cooling is performed at a cooling stop temperature of 600 ° C. or less, and the winding is performed at a winding temperature of 500 ° C. or less. It can also be a method.

Further, as an invention of a manufacturing method, in the above-mentioned invention, after the winding, and before the cold rolling, further annealing is performed at an annealing temperature of 640 ° C. or more and an Ac ₁ transformation point or less. It can also be set as the manufacturing method of a steel plate.

  These inventions were made in the course of diligent research on the effects of composition and microstructure on stretch flangeability and ductility of high carbon steel sheets. In the process, it has been found that factors affecting the stretch flangeability and ductility of the steel sheet have a great influence not only on the composition and shape and amount of carbide, but also on the dispersion state of carbide.

  Further, by controlling the carbide average particle size as the shape of the carbide and the volume fraction of ferrite grains substantially free of carbide as the carbide dispersion state, the stretch flangeability of the high carbon cold-rolled steel sheet is improved. I understood it. Furthermore, it has been found that by controlling the composition and ferrite particle size, it is possible to achieve both stable and high levels of stretch flangeability and strength, and to stably increase the elongation by further defining and controlling the carbide particle size. It was. And based on this knowledge, the manufacturing method for controlling said structure | tissue was examined, and the high carbon cold-rolled steel plate excellent in stretch flangeability, or also excellent in ductility, and its manufacturing method were established.

  Hereinafter, the components of the present invention will be described.

C content: 0.20 to 0.58% (mass%, the same applies hereinafter)
C is an important element that forms carbides and imparts hardness after quenching. However, if the C content is less than 0.20%, the formation of pro-eutectoid ferrite becomes remarkable in the structure after hot rolling, the number of ferrite grains substantially not containing carbides increases, and the distribution of carbides becomes uneven. Further, the ferrite grains are also coarsened. Furthermore, in that case, sufficient strength cannot be obtained as a machine structural component even after quenching. On the other hand, when the C content exceeds 0.58%, stretch flangeability and ductility are low even after annealing. Therefore, the C content is set to 0.20% or more and 0.58% or less.

Si: 0.1% or less Since Si is an element that improves hardenability and increases the strength of the material by solid solution strengthening, it is preferably contained in an amount of 0.005% or more. However, if the content exceeds 0.1%, pro-eutectoid ferrite is likely to be generated, and ferrite grains that do not substantially contain carbides increase, and stretch flangeability deteriorates. Therefore, the Si content is limited to 0.1% or less.

Mn: 0.20 to 0.60%:
Mn is an element that improves the hardenability and increases the strength of the material by solid solution strengthening as in the case of Si. Moreover, it is an important element which fixes S as MnS and prevents the hot crack of a slab. And it is known that the content of Mn has a great influence on the hardenability. Therefore, the influence of the amount of Mn on the hardenability in the B and Cr-added steel of the present invention was investigated.

  C: 0.36%, Si: 0.03%, Mn: 0.10 to 0.90%, P: 0.01%, S: 0.003%, sol. After melting steel composed of Al: 0.03%, N: 0.0040%, B: 0.0025%, Cr: 0.25%, hot rolling was performed at a heating temperature of 1250 ° C, a hot rolling finishing temperature of 880 ° C, and a winding temperature of 560 ° C. . Next, cold rolling was performed at a reduction rate of 50%, and annealing was performed at 710 ° C. for 40 hours, thereby producing a steel plate having a thickness of 2.5 mm. The obtained steel sheet was cut into a size of 50 × 100 mm, heated to 820 ° C. in a heating furnace, held for 60 seconds, and then quenched into oil at about 60 ° C. Hardness of the test piece after quenching was measured on a Rockwell C scale (HRc) using the sample surface as a measurement surface at 10 points to evaluate the hardenability. Evaluation made the 10-point measured average hardness (HRc) 50 or more favorable. The obtained results are shown in FIG.

  FIG. 1 is a graph showing the relationship between the amount of Mn and the average hardness after quenching. From Fig. 1, when the Mn content is 0.20% or more, an average hardness (HRc) of 50 or more is secured, and when the Mn content is 0.35% or more, the average hardness (HRc) reaches 55, and a higher quenching hardness is stabilized. It turns out that it is obtained.

  Also, since the strength of the material is increased, S is fixed as MnS, and hot cracking of the slab is prevented, when the Mn content is less than 0.20%, these effects are reduced and the formation of proeutectoid ferrite is promoted. And coarsening the ferrite grains.

  On the other hand, if it exceeds 0.60%, the tensile strength can be obtained, but the formation of a manganese band which is a segregation band becomes remarkable, and the stretch flangeability and elongation deteriorate.

  Accordingly, the Mn content is 0.20% or more and 0.60% or less, preferably 0.35% or more and 0.60% or less.

P: 0.02% or less P is an element that must be reduced in order to segregate at grain boundaries and reduce toughness. However, since the P content is acceptable up to 0.02%, the P content is limited to 0.02% or less.

S: 0.01% or less S is an element that must be reduced in order to form Mn and MnS and degrade stretch flangeability. However, since the S content is acceptable up to 0.01%, the S content is limited to 0.01% or less.

sol.Al: 0.1% or less Al is used as a deoxidizer, and is added in the steelmaking stage to improve the cleanliness of the steel. Usually, the steel contains 0.005% or more of sol.Al. On the other hand, even if Al is added so that the sol.Al content exceeds 0.1%, the effect of improving the cleanliness is saturated and the cost increases. Therefore, the sol.Al content in the steel is 0.1% or less.

N: 0.005% or less N is an element that must be reduced because BN is formed to reduce the amount of solid solution B effective for hardenability and lower hardenability. However, since the N content is acceptable up to 0.005%, the N content is limited to 0.005% or less. More preferably, it is limited to 0.0050% or less.

B: 0.001 to 0.005%
B is an important element that suppresses the formation of pro-eutectoid ferrite during cooling after hot rolling, improves stretch flangeability, and at the same time enhances hardenability. However, if the B content is less than 0.001%, a sufficient effect cannot be obtained. On the other hand, if it exceeds 0.005%, the effect is saturated and the hot rolling load is increased and the operability is lowered. Therefore, the B content is set to 0.001% or more and 0.005% or less. More preferably, it is 0.0010% or more and 0.0050% or less.

Cr: 0.05-0.3%
Cr, like B, is an important element that suppresses the formation of pro-eutectoid ferrite during cooling after hot rolling, improves stretch flangeability, and at the same time enhances hardenability. However, if the Cr content is less than 0.05%, a sufficient effect cannot be obtained. On the other hand, if the content exceeds 0.3%, the hardenability is improved, but the effect of suppressing the formation of pro-eutectoid ferrite is saturated and the cost is increased. Therefore, the Cr content is 0.05% or more and 0.3% or less. More preferably, it is 0.05% or more and 0.30% or less.

  Next, the structure of the steel sheet of the present invention will be described.

Average ferrite particle diameter: 6 μm or less The ferrite particle diameter is an important factor governing stretch flangeability and material strength, and is an important requirement in the present invention. By refining the ferrite grains, it is possible to increase the strength without deteriorating stretch flangeability. That is, by setting the average ferrite grain size to 6 μm or less, excellent stretch flangeability can be obtained while securing the tensile strength of the material to 440 MPa or more. On the other hand, when the particle size is less than 1.0 μm, the strength is remarkably increased and the load during press working may increase. Therefore, it is preferably 1.0 μm or more. The ferrite grain size can be controlled by manufacturing conditions, particularly the hot rolling finishing temperature and the cooling stop temperature.

Carbide average particle size: 0.1 μm or more and less than 1.20 μm Carbide particle size greatly affects the workability in general and the generation of voids in hole expansion processing. When the carbide becomes finer, the generation of voids can be suppressed. However, when the carbide average particle size is less than 0.1 μm, the ductility decreases with an increase in hardness, and the stretch flangeability also decreases. On the other hand, the workability generally improves as the average carbide particle size increases, but if it exceeds 1.20 μm, the stretch flangeability deteriorates due to the generation of voids in the hole expanding process. Therefore, the carbide average particle size is controlled to be 0.1 μm or more and less than 1.20 μm. More preferably, it is controlled to be 0.1 μm or more and less than 1.2 μm. Furthermore, by controlling the carbide average particle size to 0.5 μm or more and less than 1.20 μm, an increase in strength can be suppressed, and at the same time, elongation can be increased and excellent elongation characteristics can be obtained. Therefore, it is preferably 0.5 μm or more and less than 1.20 μm. More preferably, it is controlled to 0.5 μm or more and less than 1.2 μm. The average carbide grain size is the manufacturing conditions, especially the cooling stop temperature after hot rolling, the coiling temperature, and the annealing temperature in the annealing treatment of the cold-rolled steel sheet, or further annealing in the annealing treatment of the hot-rolled steel sheet before cold rolling. It can be controlled by temperature. Here, regarding the particle size of the carbide, the average of the major axis and the minor axis of the carbide is the particle size of each carbide, and the average value of the particle size of each individual carbide is the average particle size of the carbide.

Dispersion state of carbide: The volume fraction of ferrite grains substantially free of carbide is 15% or less. By making the dispersion state of carbide uniform, the stress concentration on the punched end face is alleviated as described above. And generation of voids can be suppressed. By making ferrite particles substantially free of carbides to have a volume ratio of 15% or less, the dispersion state of carbides can be made uniform, and stretch flangeability is remarkably improved. Accordingly, the volume fraction of ferrite grains substantially free of carbides is set to 15% or less. Furthermore, by making the volume fraction of ferrite grains substantially free of carbides 10% or less, the dispersion state of the carbides can be made more uniform, and extremely excellent stretch flangeability can be obtained. Therefore, it is preferably 10% or less. On the other hand, considering that this component system is hypoeutectoid steel and it is difficult to completely suppress pro-eutectoid ferrite, the lower limit of the volume fraction of ferrite grains substantially free of carbide is about 1%. It is preferable to do this. The carbide dispersion state, that is, the volume fraction of ferrite grains substantially free of carbide, is controlled by production conditions, particularly the finishing temperature of hot rolling, the cooling rate after rolling, the cooling stop temperature, and the coiling temperature. be able to.

  Here, the ferrite grain substantially free of carbide means a ferrite grain in which carbide is not detected by observation of a metal structure with a normal optical microscope, and means a ferrite grain in which carbide is not detected at a low magnification even with a scanning electron microscope. To do. That is, the ferrite grains substantially free of carbides in the present invention is a ferrite grain in which the thickness of a steel sheet sample is polished, corroded with a night, and no carbide is detected even when observed with a scanning electron microscope at 1000 times And Such a ferrite grain is a portion generated as pro-eutectoid ferrite after hot rolling, and it can be said that no carbide is observed in the grain even after annealing, that is, a ferrite grain substantially free of carbide.

  Next, the reason for limiting the production conditions suitable for producing the high carbon cold rolled steel sheet of the present invention will be described.

Hot rolling finishing temperature: (Ar ₃ transformation point -10 ° C) or higher If the finishing temperature during hot rolling of steel is less than (Ar ₃ transformation point -10 ° C), ferrite transformation proceeds in part, so carbide As a result, ferrite grains that do not contain substantially increase and stretch flangeability deteriorates. Further, the coarsening of the ferrite grains becomes remarkable and the average ferrite grain diameter exceeds 6 μm, so that the strength decreases with stretch flangeability. Therefore, the finishing temperature of the hot rolling finish rolling is set to (Ar ₃ transformation point −10 ° C.) or higher. Thereby, uniform refinement | miniaturization of a structure | tissue can be achieved and the stretch flangeability and intensity | strength can be aimed at. On the other hand, the upper limit of the finishing temperature is not particularly defined. However, when the temperature is higher than 1000 ° C., a scale defect is likely to occur. The Ar ₃ transformation point (° C.) can be calculated by the following formula.

Ar ₃ = 930.21-394.75C + 54.99Si-14.40Mn + 5.77Cr (1)
Here, the element symbol in a formula represents content (mass%) of each element.

Cooling conditions after hot rolling: Cooling rate> 120 ° C./second In the present invention, rapid cooling (cooling) is performed after rolling in order to reduce the volume fraction of ferrite grains after transformation. If the cooling method after hot rolling is slow cooling, the degree of supercooling of austenite is small and a large amount of proeutectoid ferrite is generated. When the cooling rate is 120 ° C./second or less, pro-eutectoid ferrite is remarkably produced, and ferrite grains substantially not containing carbides exceed 15%, and stretch flangeability deteriorates. Therefore, the cooling rate of the cooling after rolling is set to more than 120 ° C./second. On the other hand, the upper limit of the cooling rate is about 700 ° C./second based on the current facility capacity.

  Here, the cooling rate is an average cooling rate from the start of cooling after finish rolling to the stop of cooling. In addition, after finishing rolling, cooling is usually started within about 3 seconds, but the finer precipitates such as ferrite crystal grains and pearlite after transformation, from the point of further improving workability, after finishing rolling, It is preferable to start cooling within a time period exceeding 0.1 second and less than 1.0 second.

Cooling stop temperature: 620 ° C. or less When the cooling stop temperature for cooling after hot rolling is high, ferrite is generated during cooling until winding, and pearlite colonies and lamella spacing increase. Therefore, the ferrite grains become coarse after cold rolling-annealing, and at the same time, fine carbides cannot be obtained, the strength is lowered, and the stretch flangeability is deteriorated. When the cooling stop temperature is higher than 620 ° C., ferrite grains substantially not containing carbides exceed 15%, and stretch flangeability deteriorates. Therefore, the cooling stop temperature for cooling after hot rolling is set to 620 ° C. or lower. Further, when the ferrite grains substantially not containing carbides are made 10% or less, the cooling stop temperature is preferably made 600 ° C. or less. On the other hand, the lower limit of the cooling stop temperature is not particularly defined, but the shape of the steel sheet deteriorates as the temperature becomes lower, and therefore it is preferably set to 200 ° C. or higher.

Winding temperature: 600 ° C or less The steel sheet is wound after cooling is stopped. The higher the winding temperature, the larger the pearlite lamella spacing. Therefore, the carbide after cold rolling-annealing becomes coarse, and when the coiling temperature exceeds 600 ° C., stretch flangeability deteriorates. Accordingly, the coiling temperature is set to 600 ° C. or lower. Furthermore, by setting the coiling temperature to 500 ° C. or lower, the dispersion state of carbides becomes even more uniform and extremely excellent stretch flangeability can be obtained. On the other hand, the lower limit of the coiling temperature is not particularly defined, but the shape of the steel sheet deteriorates as the temperature is lowered, and therefore it is preferably set to 200 ° C. or higher.

  In order to further uniformize the dispersion state of carbides and obtain excellent stretch flangeability, it is preferable to cool the cooling stop temperature to 600 ° C. or lower and wind it at a winding temperature of 500 ° C. or lower.

  Furthermore, the hot-rolled steel sheet after winding is preferably subjected to pickling for scale removal before cold rolling. In particular, when annealing a hot-rolled steel sheet, it is preferable to perform pickling before the annealing to remove the influence of the scale on the steel sheet surface. Pickling may be performed according to a conventional method.

Annealing temperature of hot-rolled steel sheet: When annealing, 640 ° C or higher and Ac ₁ transformation point or lower After hot rolling, cold rolling is performed, but before that, annealing (primary annealing) is preferably performed in order to spheroidize the carbide. . The primary annealing at this time may be either box annealing or continuous annealing. When the annealing temperature of primary annealing is less than 640 ° C., the effect of annealing cannot be obtained. On the other hand, when the annealing temperature exceeds the Ac ₁ transformation point, a part is austenitized and pearlite is generated again during cooling, so that the annealing effect cannot be obtained. Therefore, the annealing temperature when performing the primary annealing is set to 640 ° C. or higher and Ac ₁ transformation point or lower. In order to obtain excellent stretch flangeability, the annealing temperature is preferably 680 ° C. or higher.

Cold rolling reduction: 30% or more Cold rolling improves the stretch flangeability by finely dispersing carbides uniformly. However, if the rolling reduction of the cold rolling is less than 30%, not only the effect is not obtained, but also the non-recrystallized portion remains after annealing, and the stretch flangeability is deteriorated. Also, the elongation decreases. Therefore, the rolling reduction of cold rolling is set to 30% or more. The upper limit of the rolling reduction is not particularly limited, but is preferably 80% or less from the viewpoint of rolling load.

Annealing temperature of cold-rolled steel sheet: 640 ° C. or higher and Ac ₁ transformation point or lower After cold rolling, annealing is performed to promote recrystallization and spheroidization of carbides. When the annealing temperature is less than 640 ° C., the spheroidization of the carbide is insufficient or the average particle size of the carbide is less than 0.1 μm, and the stretch flangeability deteriorates. On the other hand, when the annealing temperature exceeds the Ac ₁ transformation point, a part is austenitized, and pearlite is generated again during cooling, so the stretch flangeability deteriorates. Elongation also deteriorates. From the above, the annealing temperature of the cold-rolled steel sheet is set to 640 ° C. or more and Ac ₁ transformation point or less. Furthermore, by setting the annealing temperature to 680 ° C. or higher, the average carbide particle size becomes 0.5 μm or higher, and high elongation characteristics can be obtained. Therefore, it is preferably 680 ° C. or higher and Ac ₁ transformation point or lower. The Ac ₁ transformation point (° C.) can be calculated by the following formula.

Ac ₁ = 754.83-32.25C + 23.32Si-17.76Mn + 17.13Cr (2)
Here, the element symbol in a formula represents content (mass%) of each element.

  According to the present invention, in order to improve stretch flangeability, not only the control of the component composition and production conditions, but also the ferrite particle size, carbide particle size, and the dispersion state of the carbide are controlled, so that the end face at the time of punching Generation of voids can be suppressed, and crack growth in the hole expanding process can be slowed. As a result, it is possible to provide a high-carbon cold-rolled steel sheet having a tensile strength of 440 MPa or more and extremely excellent stretch flangeability or even excellent ductility.

  Either a converter or an electric furnace can be used for preparing the components of the high carbon steel of the present invention. The high carbon steel whose components are prepared in this way is used as a steel slab, which is a steel material, by ingot-bundling rolling or continuous casting. The steel slab is hot-rolled. At this time, the slab heating temperature is preferably 1300 ° C. or lower in order to avoid deterioration of the surface state due to generation of scale.

  In addition, rough rolling may be omitted during hot rolling and finish rolling may be performed, or direct feed rolling may be performed in which a continuously cast slab is rolled as it is or for the purpose of suppressing temperature reduction. In order to secure the finishing temperature, the rolled material may be heated by a heating means such as a bar heater during hot rolling. In order to promote spheroidization or reduce hardness, the coil may be kept warm by means such as a slow cooling cover after winding.

  After winding up, it is pickled according to a conventional method in some cases. Next, annealing is performed after cold rolling. Note that the hot-rolled steel sheet may be subjected to primary annealing before cold rolling and annealing after cold rolling. As for annealing after cold rolling, either box annealing or continuous annealing may be used. After annealing after cold rolling, temper rolling is performed as necessary. Since this temper rolling does not affect the hardenability, there is no particular limitation on the conditions.

  From the above, a high carbon cold-rolled steel sheet having excellent stretch flangeability or further excellent ductility can be obtained. The above shows one embodiment of the production method of the present invention, and the present invention is not limited to this.

  The reason why the high carbon cold-rolled steel sheet obtained in this way has excellent stretch flangeability is considered as follows. Stretch flangeability is greatly affected by the internal structure of the punched end face. In particular, when there are many ferrite grains substantially free of carbides (part corresponding to pro-eutectoid ferrite after hot rolling), it has been confirmed that cracks are generated from the grain boundary with the part of the spheroidized structure.

  Looking at the behavior of the microstructure, voids due to stress concentration become prominent at the carbide interface during punching. This stress concentration increases as the size of the carbide increases and as the number of ferrite grains substantially free of carbide increases. During the hole expanding process, these voids are connected to form a crack.

  Thus, by controlling not only the production conditions but also the average particle size of carbides and the proportion of ferrite grains substantially free of carbides, stress concentration can be reduced and void generation can be reduced. it can.

  Continuous casting slab of steel with chemical components shown in Table 1 is heated at 1250 ° C, hot-rolled finishing temperature 880 ° C, time after finish rolling to start cooling 0.7 seconds, cooling rate after hot-rolling 150 ° C / second, cooling stopped Hot rolling was performed at a temperature of 610 ° C. and a winding temperature of 560 ° C. to obtain a hot-rolled steel sheet. Thereafter, pickling, cold rolling at a reduction rate of 50%, and box annealing for 40 hours at 710 ° C. were performed to produce a steel sheet with a thickness of 2.5 mm. Here, Steel Nos. A to E are invention examples in which the chemical components (compositions) are within the scope of the present invention, and Steel Nos. F to M are comparative examples in which the composition is outside the scope of the present invention.

  Samples were taken from these steel plates, and the ferrite average particle size, carbide average particle size and carbide dispersion state measurement, stretch flangeability evaluation, and tensile test were performed. Each test and measurement method and conditions are shown below.

(i) Average ferrite particle diameter, carbide average particle diameter and dispersion state After the plate thickness cross section of the sample was polished and subjected to night corrosion, the microstructure was photographed with a scanning electron microscope, and the characteristic values shown were measured.

  First, the average ferrite grain size was measured according to the cutting method in the ferrite grain size test method defined in JIS standard G0552, with respect to the structure photograph taken at 1000 times with the scanning electron microscope.

The average carbide grain size, using tissue photographs were similarly taken at 3000 times, the range of actual area 0.01 mm ^2, pull the twenty segments of 100mm in the thickness direction, the carbides cross the line segments The major axis and minor axis were measured, the average value of both was taken as the grain size of the carbide, and the average of the grain sizes of all the measured carbides was determined as the average grain size of the carbide.

  As for the dispersion state of carbide, for the structure photograph taken at a magnification of 1000 times, the area ratio of ferrite grains in which carbide is not observed is measured, and this is used as the volume ratio of ferrite grains substantially free of carbide. It was used as an indicator of the dispersion state.

(ii) Evaluation of stretch flangeability The sample was punched with a punching tool having a punch diameter d ₀ = 10 mm and a die diameter 11 mm (clearance 20%), and then a hole expansion test was performed. The hole expansion test is performed by pushing up with a cylindrical flat bottom punch (50mmφ, 5R (shoulder radius 5mm)), and the hole diameter db (mm) at the time when a plate thickness penetration crack occurs at the hole edge is measured. The hole expansion rate λ (%) defined by the equation was obtained.

λ = 100 × (db−d ₀ ) / d ₀ (3).
(iii) Tensile test A JIS No. 5 test piece was taken along the 90 ° direction (C direction) with respect to the rolling direction, a tensile test was performed at a tensile speed of 10 mm / min, and tensile strength and elongation were measured.

  Table 2 shows the ferrite average particle diameter, carbide average particle diameter, carbide dispersion state, stretch flangeability, tensile strength, and elongation obtained from the above test results. Here, the stretch flangeability was evaluated by the hole expansion ratio λ of the above formula (3). In the present invention, the target for the tensile strength TS is 440 MPa or more, and the hole expansion rate λ is 80% or more (plate thickness 2.5 mm). In addition, when the excellent ductility is required, the target is 35% or more.

  In Table 2, steel Nos. A to E have chemical components (compositions) within the scope of the present invention, the average ferrite particle size is 6 μm or less, the average carbide particle size is 0.1 μm or more and less than 1.20 μm, and the carbide is substantially This is an invention example in which the volume fraction of ferrite grains not included is 15% or less. These achieve the objectives of the present invention with a tensile strength (TS) of 440 MPa or more and a hole expansion ratio λ of 80% or more. In addition, since the average particle size of carbide is 0.5 μm or more, the elongation is 35% or more.

  On the other hand, Steel Nos. F to M in Table 2 are comparative examples in which chemical components (compositions) are out of the scope of the present invention. Steel No. F has low C, ferrite average particle size, carbide average particle size, volume fraction of ferrite grains substantially free of carbides is beyond the scope of the present invention, tensile strength is less than 440 MPa, hole expansion rate Is lower than the target. Steel No. G has a high C and the structure is within the scope of the invention, but the hole expansion rate is lower than the target. Also, the elongation is low. Steel No. H is high in Si and P, and Steel No. L and M are low in B and Cr respectively. Therefore, in each case, a large amount of pro-eutectoid ferrite is formed, and the volume fraction of ferrite grains substantially free of carbides is high. The upper limit of 15% of the present invention range is exceeded and the hole expansion rate is lower than the target.

  Steel No. I in the comparative example has a low Mn, so that a large amount of pro-eutectoid ferrite is generated, the volume fraction of ferrite grains substantially free of carbides is higher than the range of the present invention, and the average ferrite grain size exceeds 6 μm. Strength and hole expansion rate are lower than target. Steel No. J has a high Mn and a band structure, so the hole expansion rate is lower than the target. Also, the elongation is low. Steel No. K has high S, MnS increases, and the hole expansion rate is greatly reduced.

  Among the steels shown in Table 1 above, after heating the continuous casting slabs of the steel Nos. A and C of the inventive examples to 1250 ° C, they were hot-rolled into hot-rolled steel sheets under the conditions shown in Table 3 Thereafter, pickling, cold rolling and annealing were performed to produce a steel plate having a thickness of 2.5 mm. Some hot-rolled steel sheets were subjected to primary annealing after pickling. Here, steel plates Nos. 1 to 12 are invention examples whose manufacturing conditions are within the scope of the present invention, and steel plates Nos. 13 to 19 are comparative examples whose manufacturing conditions are outside the scope of the present invention.

  Samples were collected from these steel plates, and in the same manner as in Example 1, the ferrite average particle size, the carbide average particle size and the dispersion state of the carbide, the stretch flangeability measurement, and the tensile test were performed. The results are shown in Table 4.

  From Table 4, steel sheets No. 1 to 12 whose production conditions are within the scope of the present invention are ferrite average particle size of 6 μm or less, carbide average particle size of 0.1 μm or more and less than 1.20 μm, and substantially free of carbides. The volume ratio of the grains is 15% or less, which is a steel sheet of the invention example. The steel plates of these inventive examples achieve the objectives of the present invention with a tensile strength (TS) of 440 MPa or more and a hole expansion ratio λ of 80% or more.

  Among them, especially steel plates No. 3, 4, 5, 6, 11, and 12 have a cooling stop temperature of 600 ° C. or lower, a coiling temperature of 500 ° C. or lower, and steel plates No. 5, 6, 9, 10, 11, No. 12 is an example in which primary annealing is performed, and each is within the preferable range of the production conditions of the present invention. These have a high hole expansion rate (85% or more). Steel plates No. 1, 3, 5, 7, 9, and 11 have an annealing temperature after cold rolling of 680 ° C. or higher, and they have achieved an elongation of 35% or higher.

  In contrast, steel plates Nos. 13 to 19 in Table 4 are comparative examples in which the production conditions (Table 3) are outside the scope of the present invention. Steel plate No. 13 has a rolling end temperature of hot rolling lower than the range of the present invention, the ferrite average particle size, the volume fraction of ferrite grains substantially not containing carbide exceeds the upper limit of the range of the present invention, and the tensile strength And the hole expansion rate is lower than the target. Steel plate No. 14 has a cooling rate after rolling lower than the range of the present invention, the volume fraction of ferrite grains substantially free of carbides also exceeds the upper limit of the range of the present invention, and the hole expansion rate is lower than the target.

  Steel plate No. 15 of the comparative example has a cooling stop temperature higher than the range of the present invention, and the ferrite average particle size, carbide average particle size, and the volume fraction of ferrite grains substantially free of carbides also exceed the upper limit of the range of the present invention. The tensile strength and hole expansion rate are lower than the target. Steel plate No. 16 as a comparative example has a coiling temperature higher than the range of the present invention, an average carbide particle size exceeding the upper limit of the range of the present invention, and a hole expansion rate lower than the target.

  Steel plate No. 17 has a cold rolling reduction ratio lower than the scope of the present invention, an unrecrystallized structure remains, ferrite grains do not become finer, tensile strength is higher, and elongation and hole expansion ratio are higher than the target. Low. Steel plate No. 18 has an annealing temperature after cold rolling higher than the range of the present invention, the average particle size of carbide, the volume fraction of ferrite grains substantially free of carbide exceeds the upper limit of the range of the present invention, The rate is lower than the target. In addition, the elongation is also decreasing. Steel plate No. 19 has an annealing temperature after cold rolling lower than the range of the present invention, carbide spheroidization is insufficient and accurate particle size measurement is impossible, but the average particle size of carbide is clearly 1.2 μm. The hole expansion rate is lower than the target. Also, the elongation is low.

  By using the high-carbon cold-rolled steel sheet of the present invention, it is possible to increase the degree of processing in processing of transmission parts and the like typified by gears. As a result, the manufacturing process is omitted and parts and the like are manufactured at low cost. It becomes possible to do.

It is a figure which shows the relationship between the amount of Mn, and the average hardness after hardening.

Claims (7)

  In mass%, C: 0.20 to 0.58%, Si: 0.1% or less, Mn: 0.20 to 0.60%, P: 0.02% or less, S: 0.01% or less, sol. Al: 0.1% or less, N: 0.005% or less, B: 0.001 to 0.005%, Cr: 0.05 to 0.3%, the composition of the balance iron and unavoidable impurities, the average ferrite particle size is 6 μm or less, and the carbide average A high carbon cold-rolled steel sheet having a structure in which a grain size is 0.1 µm or more and less than 1.20 µm, and a volume fraction of ferrite grains substantially not containing carbide is 15% or less.   In mass%, C: 0.20 to 0.58%, Si: 0.1% or less, Mn: 0.20 to 0.60%, P: 0.02% or less, S: 0.01% or less, sol. Al: 0.1% or less, N: 0.005% or less, B: 0.001 to 0.005%, Cr: 0.05 to 0.3%, the composition of the balance iron and unavoidable impurities, the average ferrite particle size is 6 μm or less, and the carbide average A high carbon cold-rolled steel sheet having a structure in which a grain size is 0.5 μm or more and less than 1.20 μm, and a volume ratio of ferrite grains substantially not containing carbide is 15% or less.   The high carbon cold-rolled steel sheet according to claim 1 or 2, wherein the volume fraction of ferrite grains substantially free of carbides is 10% or less. The steel having the composition according to claim 1 is hot-rolled at a finishing temperature of (Ar ₃ transformation point −10 ° C.) or higher, and then cooled at a cooling rate exceeding 120 ° C./second and a cooling stop temperature of 620 ° C. or less. Next, it is wound at a coiling temperature of 600 ° C. or lower, cold-rolled at a rolling reduction of 30% or higher, and then annealed at an annealing temperature of 640 ° C. or higher and below the Ac ₁ transformation point. Method. The steel having the composition according to claim 2 is hot-rolled at a finishing temperature of (Ar ₃ transformation point −10 ° C.) or higher, and then cooled at a cooling rate exceeding 120 ° C./second and a cooling stop temperature of 620 ° C. or less. Next, it is wound at a coiling temperature of 600 ° C. or lower, cold-rolled at a rolling reduction of 30% or higher, and then annealed at an annealing temperature of 680 ° C. or higher and an Ac ₁ transformation point or lower. Method.   The method for producing a high-carbon cold-rolled steel sheet according to claim 4 or 5, wherein cooling is performed at the cooling stop temperature of 600 ° C or lower and winding is performed at the winding temperature of 500 ° C or lower. The prior between after winding the cold rolling, further method for producing a high carbon cold rolled steel sheet according to claims 4 to 6, characterized in that annealing in the following _{transformation} point annealing temperature 640 ° C. or higher Ac.

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