Patent Application
20150368739
Aspects of the present invention relate to a high-strength hot-rolled steel sheet having high burring formability and a method for manufacturing the high-strength hot-rolled steel sheet. A high-strength hot-rolled steel sheet according to aspects of the present invention may be used in automotive body components, for example, structural parts, such as members and frames of automotive bodies, and chassis parts, such as suspensions. However, the present invention is not limited to these applications.
In recent years, for weight saving of automotive bodies, high-strength steel sheets have been actively used as materials for automotive parts. High-strength steel sheets are widely used as automotive structural parts. For further weight saving of automotive bodies, there is a strong demand for application of high-strength steel sheets not only to structural parts but also to chassis parts in which hot-rolled steel sheets are generally used.
Most of automotive parts made of steel sheets are formed into predetermined shapes by press forming or burring forming. However, in general, higher-strength steel sheets have lower workability. Thus, high-strength steel sheets for use in automotive parts must have high workability as well as desired strength. For example, because chassis parts are formed by severe processing, both high strength and workability must be satisfied. In particular, applicability of high-strength steel sheets to these parts and mass productivity of these parts often depend on burring formability.
Various microstructure control and reinforcement methods have been used to improve the workability of high-strength hot-rolled steel sheets. For example, these methods include use of a complex structure of ductile ferrite and hard martensite, use of a bainite microstructure, and precipitation strengthening of a ferrite microstructure. However, high-strength hot-rolled steel sheets having sufficient workability that can be applied to parts to be subjected to severe burring forming, such as chassis parts, cannot be manufactured in the related art. Thus, there is a demand for high-strength hot-rolled steel sheet having high workability.
Patent Literature 1 describes a hot-rolled steel sheet that has a composition including, on a weight percent basis, C: 0.05% to 0.2%, Si: 0.01% to 0.5%, Mn: 0.01% or more and less than 0.5%, P: 0.05% or less, S: 0.01% or less, Al: 0.005% to 0.1%, N: 0.007% or less, and Ti: 0.05% to 0.3%, and has a microstructure including a limited amount of cementite precipitation. In the technique proposed in Patent Literature 1, a decrease in austenite former Mn and extension of an a region promote TiC precipitation after hot rolling and before coiling, thereby securing the strength of the steel sheet due to precipitation strengthening of TiC. Furthermore, a decrease in the amount of cementite significantly improves the hole expandability of the steel sheet. As a result, 400 to 800 N/mm2 high-strength hot-rolled steel sheets having high workability can be manufactured.
Patent Literature 2 describes a technique for manufacturing a hot-rolled steel sheet having a structure consisting essentially of ferrite having an average grain size of 5 μm or less by heating, rolling, and cooling steel having a composition including, on a mass percent basis, C: 0.01% to 0.10%, Si: 1.0% or less, Mn: 2.5% or less, P: 0.08% or less, S: 0.005% or less, Al: 0.015% to 0.050%, and Ti: 0.10 to 0.30%, and coiling the steel sheet at a coiling temperature outside the temperature range in which TiC is coherently precipitated in a matrix phase. In the technique proposed in Patent Literature 2, a single phase structure of ferrite having controlled grain size and morphology can impart high stretch-flangeability to the hot-rolled steel sheet without reducing the high strength of the hot-rolled steel sheet. Furthermore, coherent precipitation of TiC in the main phase matrix impairs ductility or stretch-flangeability.
Patent Literature 3 describes a hot-rolled steel sheet that has a composition including, on a mass percent basis, C: 0.005% or more and 0.050% or less, Si: 0.2% or less, Mn: 0.8% or less, P: 0.025% or less, S: 0.01% or less, N: 0.01% or less, Al: 0.06% or less, Ti: 0.05% or more, and 0.10% or less and has a matrix consisting essentially of ferrite and a microstructure including finely precipitated Ti carbide. This technique is based on the knowledge that solute strengthening elements Mn and Si adversely affect stretch-flangeability. Thus, the Mn and Si contents are minimized, and fine Ti carbide is utilized to secure the strength. The technique proposed in Patent Literature 3 can be used to manufacture a high-strength hot-rolled steel sheet having a tensile strength of 590 MPa or more and high stretch-flangeability. Furthermore, the addition of B to the composition can suppress coarsening of Ti carbide.
In the technique proposed in Patent Literature 1, a decrease in Mn content results in a high ferrite transformation temperature and coarsening of TiC precipitated in the hot-rolled steel sheet. In the manufacture of the hot-rolled steel sheet, TiC is mainly formed during austenite→ferrite transformation in cooling and coiling steps after hot rolling. Thus, a high ferrite transformation temperature results in precipitation of TiC in a high-temperature region and tends to result in coarsening of TiC. When coarsening of TiC occurs in the hot-rolled steel sheet, high burring formability cannot be achieved.
In the technique proposed in Patent Literature 2, the steel sheet is coiled at a temperature at which TiC is not coherently precipitated in the main phase matrix in the hot-rolled steel sheet manufacturing process. Fine TiC that contributes to increased strength of a steel sheet is not precipitated in a hot-rolled steel sheet manufactured under such conditions. Thus, high strength and burring formability cannot be simultaneously satisfied.
With respect to the technique proposed in Patent Literature 3, because of the low Mn content, it is difficult to uniformly decrease the ferrite transformation temperature. This results in low production stability and makes it impossible to finely control the size of Ti carbide precipitated in the hot-rolled steel sheet. It is also stated that the addition of B can suppress coarsening of TiC. However, the addition of B tends to elongate ferrite grains, and it is impossible to achieve high ductility. Thus, it is difficult to manufacture a hot-rolled steel sheet having high strength and burring formability.
Burring formability required for mass production of automotive parts is not referred to in these techniques.
Burring formability of steel sheets has been principally evaluated in a hole-expanding test by a method according to the Japan Iron and Steel Federation standard. However, it is difficult to say that the hole-expanding test accurately simulates a punching process and a hole-expanding process in mass production of automotive parts in actual production lines. Thus, there is a problem that steel sheets that are experimentally shown to have good burring formability according to the standard often suffer from processing defects in mass production of automotive parts.
For example, evaluation of workability in a laboratory alone is insufficient for mass production of parts. It is necessary to ensure workability of materials also in consideration of variations in processing conditions in mass production. Such problems are not investigated in the related art. Thus, the resulting high-strength hot-rolled steel sheets do not necessarily have desired strength and workability required for mass production of automotive parts, particularly burring formability (hereinafter also referred to as mass production burring formability). Techniques that utilize Ti carbide in known structures consisting essentially of ferrite, for example, the techniques proposed in Patent Literatures 1 to 3, cannot realize high production stability and mass production burring formability of high-strength hot-rolled steel sheets.
As described above, many studies have been made on hot-rolled steel sheets having high stretch-flangeability (burring formability). However, high-strength hot-rolled steel sheets that satisfy mass production burring formability, that is, severe burring formability required in actual automotive part production lines are not necessarily manufactured in the related art.
It is an aim of aspects of the present invention to advantageously provide a high-strength hot-rolled steel sheet having a tensile strength (TS) of 540 MPa or more and having high burring formability, particularly high mass production burring formability, and a method for manufacturing the high-strength hot-rolled steel sheet.
The “mass production burring formability” herein is evaluated as a burring ratio measured in a hole-expanding test using a 60-degree conical punch after punching with a 50-mmφ punch (clearance of stamping: 30%) and is different from burring formability evaluated as a λ value determined by a known hole-expanding test method, for example, a hole-expanding test method according to the Japan Iron and Steel Federation standard.
The present inventors studied methods for evaluating Mass production burring formability. Burring formability has been evaluated as a λ value, for example, measured by a hole-expanding test method according to the Japan Iron and Steel Federation standard. In this case, the punch diameter is 10 mmφ. However, the present inventors found that burring formability in actual mass production settings for parts is not correlated with the λ value measured in laboratories according to the Japan Iron and Steel Federation standard. It was found that burring formability evaluated in a new hole-expanding test that includes hole-expanding using a 60-degree conical punch after punching with a 50-mmφ punch (clearance of stamping: 30%) is closely correlated with mass production punchability and mass production burring formability.
The present inventors also extensively studied various factors that contribute to high strength and workability, particularly mass production burring formability, of hot-rolled steel sheets by evaluating mass production burring formability in the new hole-expanding test.
Moreover, extensive studies have been made on means for improving mass production burring formability of a hot-rolled steel sheet based on a structure consisting essentially of a ductile ferrite phase while reinforcing the hot-rolled steel sheet, with consideration given to all the precipitates that can be precipitated in the hot-rolled steel sheet, such as nitrides, sulfides, carbides, and complex precipitates thereof (for example, carbonitride).
It was found that a hot-rolled steel sheet having a tensile strength of 540 MPa or more that satisfies severe mass production burring formability required in actual automotive part production lines can be manufactured by optimizing the balance between the amount of C in the hot-rolled steel sheet and the amount of Ti (Ti*) that contributes to the formation of carbide and increasing the percentage of carbide having a grain size of less than 9 nm in carbide precipitated in the hot-rolled steel sheet. It was also found that mass production burring formability can be further improved by controlling the size of not only carbide but also all the precipitates that can be precipitated in a hot-rolled steel sheet (nitrides, sulfides, carbides, and complex precipitates thereof).
The present inventors also studied means for achieving a desirable size of precipitates that are precipitated in a hot-rolled steel sheet (nitrides, sulfides, carbides, and complex precipitates thereof), that is, a size required to impart desired strength (e.g., tensile strength of 540 MPa or more) and good mass production burring formability to the hot-rolled steel sheet. As a result, it was found that it is desirable to properly control the Mn content and C, S, N, and Ti contents of a hot-rolled steel sheet and optimize the hot rolling conditions and the cooling and coiling conditions after hot rolling.
The following are non-limiting embodiments of the present invention.
[1] A high-strength hot-rolled steel sheet, containing: on a mass percent basis, C: 0.013% or more and less than 0.08%, Si: less than 0.5%, Mn: more than 0.8% and less than 1.2%, P: 0.05% or less, S: 0.005% or less, N: 0.01% or less, Al: 0.1% or less, and Ti: 0.03% or more and 0.15% or less such that C, S, N, and Ti satisfy the following formulae (1) and (2), the remainder being Fe and incidental impurities, wherein the high-strength hot-rolled steel sheet has a microstructure in which a ferrite phase fraction is more than 90%, a carbide containing Ti is precipitated, and 70% or more of the carbide has a grain size of less than 9 nm.
0.05≦Ti*<0.1 (1)
C×(48/12)−0.16<Ti* (2)
wherein Ti*=Ti−N×(48/14)−S×(48/32), and C, S, N, and Ti denote the amounts (% by mass) of the corresponding elements.
[2] The high-strength hot-rolled steel sheet according to [1], wherein 50% by mass or more of Ti is precipitated as precipitates containing Ti having a grain size of less than 20 nm.
[3] The high-strength hot-rolled steel sheet according to [1] or [2], further containing at least one of V: 0.002% or more and 0.1% or less and Nb: 0.002% or more and 0.1% or less on a mass percent basis.
[4] The high-strength hot-rolled steel sheet according to any one of [1] to [3], further containing at least one of Cu: 0.005% or more and 0.2% or less, Ni: 0.005% or more and 0.2% or less, Cr: 0.002% or more and 0.2% or less, and Mo: 0.002% or more and 0.2% or less on a mass percent basis.
[5] The high-strength hot-rolled steel sheet according to any one of [1] to [4], further containing B: 0.0002% or more and 0.003% or less on a mass percent basis.
[6] The high-strength hot-rolled steel sheet according to any one of [1] to [5], further containing at least one of Ca: 0.0002% or more and 0.005% or less and REM: 0.0002% or more and 0.03% or less on a mass percent basis.
[7] A method for manufacturing a high-strength hot-rolled steel sheet, including: heating steel having a composition according to any one of [1] and [3] to [6] to 1100° C. or more, hot-rolling the steel at a finish-rolling temperature of (Ar3+20° C.) or more and at a total reduction ratio of 60% or less at last two finish rolling stands, cooling the hot-rolled steel sheet at an average cooling rate of 40° C./s or more, and coiling the hot-rolled steel sheet at a coiling temperature in the range of 560° C. to 720° C.
[8] A method for manufacturing a high-strength hot-rolled steel sheet, comprising: heating steel having a composition according to any one of [1] and [3] to [6] to 1100° C. or more, hot-rolling the steel at a finish-rolling temperature of (Ar3+20° C.) or more and at a total reduction ratio of 60% or less at last two finish rolling stands, cooling the hot-rolled steel sheet at an average cooling rate of 40° C./s or more, coiling the hot-rolled steel sheet at a coiling temperature in the range of 500° C. to 660° C., annealing the hot-rolled steel sheet at a soaking temperature of 750° C. or less after pickling, and plating the hot-rolled steel sheet by immersing the hot-rolled steel sheet in a molten zinc bath.
[9] The method for manufacturing a high-strength hot-rolled steel sheet according to [8], wherein the plating treatment is followed by alloying treatment.
Aspects of the present invention provide a high-strength hot-rolled steel sheet having a tensile strength of 540 MPa or more and high burring formability such that the high-strength hot-rolled steel sheet can be subjected to processing in mass production of automotive parts. Thus, in one embodiment, a high-strength hot-rolled steel sheet can be applied to structural parts, such as members and frames of automotive bodies, and chassis parts, such as suspensions. Embodiments of the present invention may contribute greatly to weight saving of these parts.
Aspects of the present invention can provide a hot-rolled steel sheet having a tensile strength of 540 MPa or more and good mass production burring formability. Thus, the high-strength hot-rolled steel sheet can be applied not only to automotive parts but also to other applications. Thus, the present invention has industrially advantageous effects.
Non-limiting embodiments of the present invention will be further described below.
A high-strength hot-rolled steel sheet according to one embodiment contains, on a mass percent basis, C: 0.013% or more and less than 0.08%, Si: less than 0.5%, Mn: more than 0.8% and less than 1.2%, P: 0.05% or less, S: 0.005% or less, N: 0.01% or less, Al: 0.1% or less, and Ti: 0.03% or more and 0.15% or less such that C, S, N, and Ti satisfy the following formulae (1) and (2), the remainder being Fe and incidental impurities, wherein the high-strength hot-rolled steel sheet has a microstructure in which a ferrite phase fraction is more than 90%, a carbide containing Ti is precipitated, and 70% or more of the carbide has a grain size of less than 9 nm.
0.05≦Ti*<0.1 (1)
C×(48/12)−0.16<Ti* (2)
wherein Ti*=Ti−N×(48/14)−S×(48/32), and C, S, N, and Ti denote the amounts (% by mass) of the corresponding elements.
First, examples of the composition of a hot-rolled steel sheet according to aspects of the present invention will be described below. Unless otherwise specified, the percentages of the components are on a mass percent basis.
C: 0.013% or More and Less than 0.08%
C is an important element that forms an appropriate carbide in a hot-rolled steel sheet and secures the strength of the steel sheet. In order to achieve the desired tensile strength (540 MPa or more), the C content is 0.013% or more. However, a C content of 0.08% or more results in poor workability and undesired burring formability of a hot-rolled steel sheet. Thus, the C content is 0.013% or more and less than 0.08%, preferably 0.03% or more and 0.07% or less.
Si: Less than 0.5%
A Si content of 0.5% or more results in very low surface quality of a hot-rolled steel sheet, which adversely affects fatigue characteristics, chemical conversion treatability, and corrosion resistance. Si increases the ferrite transformation temperature and thereby adversely affects the formation of fine precipitates. Thus, the Si content is less than 0.5%, preferably 0.001% or more and less than 0.1%, more preferably 0.001% or more and less than 0.05%.
Mn: More than 0.8% and Less than 1.2%
Mn is an important element. Mn significantly influences precipitation of a carbide containing Ti, through control of austenite-to-ferrite transformation temperatures.
In the case of a hot-rolled steel sheet containing Ti, a carbide containing Ti is mainly precipitated by austenite→ferrite transformation in cooling and coiling steps after finish rolling in a hot-rolled steel sheet manufacturing process. Among carbides precipitated in a hot-rolled steel sheet, fine carbide contributes to high strength of the hot-rolled steel sheet, but coarse carbide does not contribute to high strength and adversely affects the workability of the hot-rolled steel sheet.
A high austenite-ferrite transformation temperature results in precipitation of a carbide containing Ti in a high-temperature region and consequently coarsening of the carbide containing Ti. Thus, in order to decrease the size of the carbide containing Ti, it is preferable to decrease the austenite-ferrite transformation temperature.
Mn is an element that has an effect of decreasing the austenite-ferrite transformation temperature. A Mn content of 0.8% or less results in an insufficient decrease in the austenite-ferrite transformation temperature. As a result, it is difficult for a carbide containing Ti to have a desirable size. Thus, it is difficult to provide a high-strength hot-rolled steel sheet having high mass production burring formability. On the other hand, a Mn content of 1.2% or more results in saturation of the effect and increased costs. An excessively high Mn content of 1.2% or more also results in increased Mn segregation in the central portion in the thickness direction. This center segregation impairs a punched surface before burring forming and is therefore responsible for low mass production burring formability. Thus, the Mn content is more than 0.8% and less than 1.2%, preferably more than 0.8% and less than 1.0%.
P promotes low workability of a hot-rolled steel sheet due to segregation. Thus, the P content is 0.05% or less, preferably 0.001% or more and 0.03% or less. In the case of a galvanized steel sheet formed by galvanizing treatment of a hot-rolled steel sheet, the P content is preferably 0.005% or more, more preferably 0.01% or more, still more preferably 0.015% or more, in terms of platability.
S forms a sulfide and decreases the workability of a hot-rolled steel sheet. Thus, the S content is 0.005% or less, preferably 0.0001% or more and 0.003% or less, more preferably 0.0001% or more and 0.0015% or less.
An excessively high N content of more than 0.01% results in the formation of a large amount of nitride in a hot-rolled steel sheet manufacturing process, low hot ductility, and very low burring formability of a hot-rolled steel sheet due to coarsening of nitride. Thus, the N content is 0.01% or less, preferably 0.0001% or more and 0.006% or less, more preferably 0.0001% or more and 0.004% or less.
Al is a useful element as a deoxidizing agent for steel. However, an Al content of more than 0.1% makes casting of steel difficult and results in a large amount of residual inclusion in steel and low surface quality and workability of a hot-rolled steel sheet. Thus, the Al content is 0.1% or less, preferably 0.001% or more and 0.06% or less.
Ti is an important element in the present invention. Ti forms fine carbide and contributes to increased strength of a hot-rolled steel sheet. In order to achieve the desired strength of a hot-rolled steel sheet (tensile strength of 540 MPa or more), the Ti content is 0.03% or more. However, a Ti content of more than 0.15% tends to result in the remaining coarse carbide in a hot-rolled steel sheet. Coarse carbide has no strength increasing effect and greatly impairs the workability, toughness, and weldability of a hot-rolled steel sheet. Thus, the Ti content is 0.03% or more and 0.15% or less, preferably 0.04% or more and 0.12% or less.
A hot-rolled steel sheet according to one embodiment contains C, S, N, and Ti in the ranges described above so as to satisfy the formulae (1) and (2). The formulae (1) and (2) are desirably satisfied in order to achieve high strength and good mass production burring formability of a hot-rolled steel sheet and are beneficial indicators. In the formulae (1) and (2), Ti*=Ti−N×(48/14)−S×(48/32), and C, S, N, and Ti denote the amounts (%) of the corresponding elements.
0.05≦Ti*<0.1 (1)
As described below, in one embodiment, a predetermined amount of Ti is added to steel, and carbide in steel is dissolved by heating before hot rolling. A carbide containing Ti is mainly precipitated during coiling after hot rolling. However, Ti added to steel does not entirely contribute to the formation of carbide and is partly consumed by forming nitride or sulfide. This is because Ti is likely to form nitride or sulfide rather than carbide in a higher temperature region than the coiling temperature. Thus, Ti forms nitride or sulfide before the coiling step in the production of a hot-rolled steel sheet. Thus, the minimum amount of Ti that can contribute to the formation of carbide out of Ti added to steel can be represented by Ti* (=Ti−N×(48/14)−S×(48/32)).
A hot-rolled steel sheet cannot have the desired strength (tensile strength of 540 MPa or more) at Ti* of less than 0.05. Thus, in one embodiment, Ti* is 0.05 or more, preferably 0.055 or more. However, a hot-rolled steel sheet has a tensile strength of 700 MPa or more at Ti* of 0.1 or more. A hot-rolled steel sheet having such excessively high strength has poor workability. Furthermore, an excessively high Ti* results in coarsening of precipitates containing Ti, such as Ti carbide, Ti carbonitride, Ti nitride, and Ti sulfide, which results in low mass production burring formability. Thus, in in one embodiment, Ti* is less than 0.1.
C×(48/12)−0.16<Ti* (2)
The formula (2) represents the relationship between the amount of Ti* and the amount of C. An excessively higher amount of C than the amount of Ti* results in coarsening of Ti carbide and Ti carbonitride, precipitation of coarse cementite or pearlite, and a significant reduction in the workability of a hot-rolled steel sheet, such as mass production burring formability. In one embodiment, C×(48/12)−0.16<Ti*, preferably C×(48/12)−0.15<Ti*. However, an excessively high amount of Ti than the amount of C tends to result in low toughness and weldability of a hot-rolled steel sheet. Thus, Ti*<C×(48/12)+0.08 is preferred. More preferably, Ti*<C×(48/12)+0.06.
The following are base components of a hot-rolled steel sheet according to non-limiting embodiments of the present invention. A hot-rolled steel sheet according to one embodiment may contain at least one of V: 0.002% or more and 0.1% or less and Nb: 0.002% or more and 0.1% or less, if desirable.
V and Nb are effective in decreasing the size of crystal grains and improving the toughness of a hot-rolled steel sheet. Thus, V and Nb may be added as desired. Part of added V and/or Nb is precipitated together with Ti as fine complex carbide or complex precipitates and thereby contributes to precipitation strengthening. In order to produce such an effect, the V content is preferably 0.002% or more, and the Nb content is preferably 0.002% or more. However, a V or Nb content of more than 0.1% is not worth the cost. Thus, the V content is preferably 0.002% or more and 0.1% or less, more preferably 0.002% or more and 0.08% or less. The Nb content is preferably 0.002% or more and 0.1% or less, more preferably 0.002% or more and 0.08% or less.
A hot-rolled steel sheet according to one embodiment may contain at least one of Cu: 0.005% or more and 0.2% or less, Ni: 0.005% or more and 0.2% or less, Cr: 0.002% or more and 0.2% or less, and Mo: 0.002% or more and 0.2% or less, if desirable.
Cu and Ni are elements that contribute to high strength of a hot-rolled steel sheet and may be added, if desirable. In order to produce such an effect, the Cu content is preferably 0.005% or more, and the Ni content is preferably 0.005% or more. However, a Cu or Ni content of more than 0.2% may result in surface layer cracking during hot rolling in the production of a hot-rolled steel sheet. Thus, the Cu content is preferably 0.005% or more and 0.2% or less, more preferably 0.005% or more and 0.1% or less. The Ni content is preferably 0.005% or more and 0.2% or less, more preferably 0.005% or more and 0.15% or less.
Cr and Mo are carbide formation elements, contribute to high strength of a hot-rolled steel sheet, and may be added, if desirable. In order to produce such an effect, the Cr content is preferably 0.002% or more, and the Mo content is preferably 0.002% or more. However, a Cr or Mo content of more than 0.2% is not worth the cost. Thus, the Cr content is preferably 0.002% or more and 0.2% or less, more preferably 0.002% or more and 0.1% or less. The Mo content is preferably 0.002% or more and 0.2% or less, more preferably 0.002% or more and 0.1% or less.
A hot-rolled steel sheet according to one embodiment may contain B: 0.0002% or more and 0.003% or less, if desirable.
B is an element that retards austenite-ferrite transformation of steel. B decreases the precipitation temperature of a carbide containing Ti by suppressing austenite-ferrite transformation and contributes to a reduction in the size of the carbide. In order to produce such an effect, the B content is preferably 0.0002% or more. However, a B content of more than 0.003% results in a strong bainite transformation effect of B, making it difficult for a hot-rolled steel sheet to have a structure consisting essentially of a ferrite phase. The B content is preferably 0.0002% or more and 0.003% or less, more preferably 0.0002% or more and 0.002% or less.
A hot-rolled steel sheet according to one embodiment may contain at least one of Ca: 0.0002% or more and 0.005% or less and REM: 0.0002% or more and 0.03% or less, if desirable.
Ca and REM are elements that are effective in morphology control of an inclusion in steel and contribute to improved workability of a hot-rolled steel sheet. In order to produce such an effect, the Ca content is preferably 0.0002% or more, and the REM content is preferably 0.0002% or more. However, a Ca content of more than 0.005% or a REM content of more than 0.03% may result in an increased inclusion in steel and low workability of a hot-rolled steel sheet. Thus, the Ca content is preferably 0.0002% or more and 0.005% or less, more preferably 0.0002% or more and 0.003% or less. The REM content is preferably 0.0002% or more and 0.03% or less, more preferably 0.0002% or more and 0.003% or less.
The remainder may be Fe and incidental impurities. Examples of the incidental impurities include W, Co, Ta, Sn, Sb, Zr, and O. The amount of each of the incidental impurities may be 0.1% or less.
Next, the microstructure of a hot-rolled steel sheet according to non-limiting embodiments of the present invention will be described below.
A hot-rolled steel sheet according to one embodiment has a microstructure in which the ferrite phase fraction is more than 90%, a carbide containing Ti is precipitated, and 70% or more of the carbide has a grain size of less than 9 nm. A hot-rolled steel sheet preferably has a microstructure in which 50% by mass or more of Ti is precipitated as precipitates having a grain size of less than 20 nm.
Ferrite Phase Fraction: More than 90%
The burring formability of a hot-rolled steel sheet can be effectively improved when the hot-rolled steel sheet has a microstructure including a ductile ferrite phase. In order to achieve a preferable high mass production burring formability, the ferrite fraction in the microstructure of a hot-rolled steel sheet is more than 90%, preferably more than 94%, more preferably more than 96% by area. It is desirable that the ferrite grains have a polygonal shape from the perspective of burring formability. It is also desirable that the ferrite grain size be as small as possible. The hot-rolled steel sheet preferably has a single phase structure of ferrite in terms of burring formability. In order to improve punchability, the ferrite fraction is preferably 99% by area or less.
A hot-rolled steel sheet according to one embodiment may have a microstructure other than the ferrite phase, such as cementite, pearlite, bainite, martensite, and/or retained austenite. Although these microstructures in the hot-rolled steel sheet impair burring formability, these microstructures may constitute approximately less than 10%, preferably less than 6%, more preferably less than 4% by area in total.
In one embodiment, the desired strength (tensile strength of 540 MPa or more) of a hot-rolled steel sheet is achieved by precipitation of a carbide containing Ti in the hot-rolled steel sheet. The carbide containing Ti is mainly precipitated carbide resulting from austenite→ferrite transformation in the cooling and coiling steps after finish rolling in a hot-rolled steel sheet manufacturing process.
In order to make the maximum use of the precipitation strengthening effect and optimize the balance between strength and workability (mass production burring formability), it may be preferable to reduce the size of a carbide containing Ti precipitated in a hot-rolled steel sheet. As a result of extensive studies, the present inventors found that 70% by number or more, preferably 80% or more, of a carbide containing Ti has a grain size of less than 9 nm in order to achieve the desirable characteristics. The “carbide containing Ti” includes complex carbides containing Ti and at least one of Nb, V, Cr, and Mo as well as Ti carbide.
The size of precipitates containing Ti may be controlled to further improve the mass production burring formability of a hot-rolled steel sheet.
As described above, in the case of a hot-rolled steel sheet made of steel containing Ti, in addition to carbide (carbide containing Ti) that contributes to high strength of a hot-rolled steel sheet, nitride, carbonitride, and sulfide containing Ti are precipitated. In the production of a hot-rolled steel sheet, these nitride, carbonitride, and sulfide are precipitated faster than carbide containing Ti. Thus, nitride, carbonitride, and sulfide containing Ti are precipitated in a higher temperature range than carbide and are therefore easily coarsened and tend to impair mass production burring formability.
As a result of extensive studies, the present inventors found that the control of the amount and size of these precipitates is very effective in improving mass production burring formability. Preferably, 50% by mass or more, more preferably 60% by mass or more and 85% by mass or less, of Ti in a hot-rolled steel sheet is preferably precipitated as precipitates containing Ti having a grain size of less than 20 nm. The precipitates containing Ti having a grain size of less than 20 nm are mostly carbide containing Ti and also include nitride, carbonitride, and sulfide containing Ti.
The precipitates containing Ti may be precipitates of Ti carbide, Ti nitride, Ti sulfide, and/or Ti carbonitride, and/or complex precipitates, such as complex carbide, complex nitride, complex sulfide, and/or complex carbonitride containing Ti and at least one of Nb, V, Cr, and Mo.
Even when precipitates having a grain size of 20 nm or more out of precipitates containing Ti are precipitated, it is surmised that a proper amount of precipitates contribute to improved punchability before burring forming and consequently contribute to improved burring formability.
Formation of a coated layer on a surface of a hot-rolled steel sheet in order to impart corrosion resistance does not reduce the advantages of the present invention. The type of coated layer formed on a surface of a hot-rolled steel sheet is not particularly limited and may be galvanic electroplating or hot-dip plating. The hot-dip plating may be hot-dip galvanization. The coated layer may also be galvannealed steel, which was subjected to alloying treatment after plating.
Methods for manufacturing a hot-rolled steel sheet will be described below with reference to non-limiting embodiments.
One embodiment includes heating steel having the composition described above to 1100° C. or more, hot-rolling the steel at a finish-rolling temperature of (Ar3+20° C.) or more and at a total reduction ratio of 60% or less at last two finish rolling stands, cooling the hot-rolled steel sheet at an average cooling rate of 40° C./s or more, and coiling the hot-rolled steel sheet at a coiling temperature in the range of 560° C. to 720° C.
The steel may be melted by any method, for example, in a converter, electric furnace, or induction furnace. After that, secondary smelting is preferably performed with vacuum degassing equipment. Subsequent casting is preferably performed in a continuous casting process in terms of productivity and quality. A blooming method can also be used. A slab (steel) to be casted may be a general slab having a thickness in the range of approximately 200 to 300 mm or a thin slab having a thickness of approximately 30 mm. In the case of a thin slab, rough rolling may be omitted. A slab after casting may be subjected to hot direct rolling or may be subjected to hot rolling after reheating in a furnace.
Steel thus produced is subjected to hot rolling. In one embodiment, it is desirable to heat the steel (slab) before hot rolling and redissolve carbide in the steel. At a steel heating temperature of less than 1100° C., carbide is not redissolved in the steel, and desirable fine carbide cannot be formed in the cooling and coiling steps after hot rolling. Thus, the steel heating temperature is 1100° C. or more, preferably 1200° C. or more, more preferably 1240° C. or more.
However, an excessively high steel heating temperature results in excessively accelerated oxidation of the surface of a steel sheet and very poor surface quality and adversely affects the workability of a hot-rolled steel sheet. Thus, the steel heating temperature is preferably 1350° C. or less.
After heating of steel, the steel is subjected to hot rolling, which is composed of rough rolling and finish rolling. The rough rolling conditions are not particularly limited. As described above, when steel is a thin slab, rough rolling may be omitted. In the finish rolling, the finish-rolling temperature is (Ar3+20° C.) or more, and the total reduction ratio at last two stands of a finish rolling mill is 60% or less.
At a finish-rolling temperature of less than (Ar3+20° C.), austenite→ferrite transformation in the cooling and coiling steps after hot rolling is ferrite transformation from unrecrystallized austenite grains. In such a case, desired fine carbide cannot be formed, and a hot-rolled steel sheet cannot have a desirable strength (e.g., tensile strength of 540 MPa or more). Thus, the finish-rolling temperature is (Ar3+20° C.) or more, preferably (Ar3+40° C.) or more. However, an excessively high finish-rolling temperature results in coarsening of crystal grains and adversely affects the punchability of a hot-rolled steel sheet. Thus, the finish-rolling temperature is (Ar3+140° C.) or less.
The Ar3 transformation point herein refers to a transformation temperature at a change point of a thermal expansion curve measured in a thermecmastor test (thermo-mechanical simulation test) at a cooling rate of 5° C./s.
When the total reduction ratio at last two finish rolling stands exceeds 60%, this results in increased residual strain and accelerates ferrite transformation from unrecrystallized austenite grains. Thus, the total reduction ratio at last two stands of a finish rolling mill is 60% or less, preferably 50% or less.
When the average cooling rate in cooling after hot rolling is less than 40° C./s, this results in a high ferrite transformation temperature. As a result, carbide is precipitated in a high temperature region, desired fine carbide cannot be formed, and a hot-rolled steel sheet cannot have the desired strength (e.g., tensile strength of 540 MPa or more). Thus, the average cooling rate is 40° C./s or more, preferably 50° C./s or more. However, at an excessively high average cooling rate, it may be difficult to achieve the desired ferrite microstructure. Thus, the average cooling rate is 150° C./s or less.
The average cooling rate herein refers to the average cooling rate between the finish-rolling temperature and the coiling temperature.
In one embodiment, a carbide containing Ti is precipitated in a period from immediately before coiling to the beginning of the coiling step by decreasing the ferrite transformation temperature so as to be close to the coiling temperature at the average cooling rate. This can prevent precipitation and coarsening of the carbide containing Ti in a high temperature region. Thus, the resulting hot-rolled steel sheet can contain precipitated fine carbide.
As described above, in one embodiment, fine carbide containing Ti is mainly precipitated in a period from immediately before coiling to the beginning of the coiling step. Thus, in order to precipitate a large amount of fine carbide containing Ti, the coiling temperature is controlled in a temperature range suitable for precipitation of the carbide containing Ti. At a coiling temperature of less than 560° C. or more than 720° C., fine carbide that contributes to high strength of steel is not sufficiently precipitated, and the hot-rolled steel sheet may not have a desirable strength. For these reasons, the coiling temperature ranges from 560° C. to 720° C., preferably 600° C. to 700° C.
In one embodiment, a hot-rolled steel sheet after coiling may be subjected to pickling and annealing treatment and then to plating treatment by immersion in a molten zinc bath. After the plating treatment, the hot-rolled steel sheet may be subjected to alloying treatment. When the plating treatment is performed, the coiling temperature ranges from 500° C. to 660° C., and the soaking temperature for the annealing treatment is 750° C. or less.
A higher coiling temperature facilitates the formation of an internal oxidation layer in a hot-rolled steel sheet. The internal oxidation layer may promote plating defects. For example, a coiling temperature of more than 660° C. results in low plating quality. On the other hand, a low coiling temperature is preferred in order to prevent plating defects. However, a coiling temperature of less than 500° C. results in insufficient precipitation of a carbide containing Ti, and a hot-rolled steel sheet may not have a desirable strength. Thus, when plating treatment is performed after coiling, the coiling temperature ranges from 500° C. to 660° C., preferably 500° C. to 600° C.
As described above, when the coiling temperature is lowered for plating treatment, fine carbide that contributes to high strength of a hot-rolled steel sheet (carbide containing Ti) may be insufficiently precipitated during coiling. Thus, in one embodiment, the desired strength (tensile strength of 540 MPa or more) of a hot-rolled steel sheet after plating treatment is achieved by precipitating fine carbide (carbide containing Ti) during annealing treatment before the plating treatment. When the soaking temperature for annealing treatment exceeds 750° C., precipitated carbide (carbide containing Ti) is coarsened, and a hot-rolled steel sheet has low strength. Thus, the soaking temperature for annealing treatment is 750° C. or less, preferably 720° C. or less. In order to promote precipitation of fine carbide (carbide containing Ti), the soaking temperature for annealing treatment is preferably 600° C. or more. The holding time at the soaking temperature preferably ranges from 10 to 1000 seconds.
After the annealing treatment, the steel sheet is immersed in a hot-dip galvanizing bath to form a hot-dip galvanized layer on the surface of the steel sheet. After immersion in the hot-dip galvanizing bath, the steel sheet may be subjected to alloying treatment. The annealing treatment and plating treatment are preferably performed in a continuous hot-dip galvanizing line.
The type of plating is not limited to hot-dip galvanization or galvannealing described above and may be electrogalvanizing.
The plating treatment conditions, the alloying treatment conditions, and other manufacturing conditions are not particularly limited and may be general conditions.
Steel slabs (Nos. A to P) containing the components listed in Table 1 and having the Ar3 transformation point listed in Table 1 were heated to a temperature in the range of 1180° C. to 1290° C., and hot-rolled steel sheets (Nos. 1 to 22) were formed under the hot-rolling conditions listed in Table 2. The hot-rolled steel sheets had a thickness in the range of 2.0 to 4.5 mm. The Ar3 transformation points listed in Table 1 were determined as described above. Part of the hot-rolled steel sheets (Nos. 3, 4, 9 to 11, 16, 18, and 19) were subjected to pickling and were then subjected to annealing treatment at a soaking temperature listed in Table 2 and hot-dip galvanizing treatment in a hot-dip galvanization line. In the hot-dip galvanizing treatment, each of the hot-rolled steel sheets subjected to the annealing treatment was immersed in a galvanizing bath (0.1% by mass Al—Zn) at 480° C., and a hot-dip galvanized layer was formed on both faces of the steel sheet at 45 g/m2. Part of the hot-rolled steel sheets (Nos. 9 to 11, 16, 18, and 19) were subjected to the hot-dip galvanizing treatment and alloying treatment. The alloying treatment temperature was 520° C.
Test specimens were taken from the hot-rolled steel sheets (Nos. 1 to 22) and were subjected to microstructure observation, a tensile test, and a hole-expanding test. The microstructure observation method and various test methods were as follows:
Scanning electron microscope (SEM) test specimens were taken from the hot-rolled steel sheets. A vertical cross section of each of the test specimens parallel to the rolling direction was polished and was subjected to nital etching. SEM photographs were taken in 10 visual fields at a quarter thickness in the depth direction and at a magnification ratio of 3000. A ferrite phase and a non-ferrite phase were separated by image analysis. The fraction of each of the phases (area fraction) was determined.
Thin film samples were prepared from the hot-rolled steel sheets (at a quarter thickness in the depth direction). Photographs were taken in 10 visual fields with a transmission electron microscope at a magnification ratio of 200,000.
The number of carbide grains containing Ti (N0) was determined from the photographs. The grain size of each carbide grain containing Ti was determined as an equivalent circle diameter by image processing. The number of carbide grains having a grain size of less than 9 nm (N1) out of the carbide grains containing Ti was determined. The ratio of the number of carbide grains having a grain size of less than 9 nm to the number of carbide grains containing Ti (N1/N0×100(%)) was calculated from these numbers (N0 and N1).
Precipitates were extracted from the hot-rolled steel sheets by constant-current electrolysis using an AA electrolyte solution (an ethanol solution of acetylacetone-tetramethylammonium chloride). The extract was passed through a filter having a pore size 20 nm. Precipitates having a size of less than 20 nm were separated in this manner, and the amount of Ti in the precipitates having a size of less than 20 nm was measured by inductively-coupled plasma optical emission spectrometry (ICP). The ratio (percentage) of Ti in the precipitates having a size of less than 20 nm was determined by dividing the amount of Ti in the precipitates having a size of less than 20 nm by the amount of Ti in the hot-rolled steel sheet.
Three JIS No. 5 test pieces for tensile test were taken from each of the hot-rolled steel sheets such that the tensile direction is a direction perpendicular to the rolling direction. The tensile strength and total elongation were measured in a tensile test (strain rate: 10 mm/min) according to JIS Z 2241 (2011). Each of the hot-rolled steel sheets were subjected to measurements three times in the tensile test. The tensile strength (TS) and total elongation (El) were averages of the three measurements.
(iii) Hole-Expanding Test (Evaluation of Mass Production Burring Formability)
Test specimens (size: 150 mm×150 mm) were taken from the hot-rolled steel sheets. A hole having an initial diameter d0 was formed in each of the test specimens by punching with a 50-mmφ punch (clearance of stamping: 30%). The hole was then expanded by inserting a conical punch having a vertex angle of 60 degrees into the hole from the punched side. A hole diameter d1 at which a crack penetrated the steel sheet (test specimen) in the thickness direction was measured. The burring ratio (%) was calculated using the following equation.
Burring ratio (%)={(d1−d0)/d0}×100
A burring ratio of 60% or more was considered to be high mass production burring formability.
Table 3 shows the results.
K
0.31
L
1.37
M
N
0.68
785
840
75
10
455
11
30
17
K
18
L
19
M
20
N
98(B: 2)
55
40
50
10
65
11
25
99(B: 1)
17
18
93(B: 7)
19
25
20
55
The hot-rolled steel sheets according to examples (Nos. 1 to 3, 5, 6, 9, 12 to 16, 21, and 22) had the desired tensile strength (540 MPa or more) and good mass production burring formability. By contrast, the hot-rolled steel sheets according to comparative examples (Nos. 4, 7, 8, 10, 11, and 17 to 20), which were outside the scope of the present invention, did not have a desired high strength or sufficient burring ratios.
| Number | Date | Country | Kind |
|---|---|---|---|
| 2013-016455 | Jan 2013 | JP | national |
This is the U.S. National Phase application of PCT International Application No. PCT/JP2014/000336, filed Jan. 23, 2014, which claims priority to Japanese Patent Application No. 2013-016455, filed Jan. 31, 2013, the disclosures of each of these applications being incorporated herein by reference in their entireties for all purposes.
| Filing Document | Filing Date | Country | Kind |
|---|---|---|---|
| PCT/JP2014/000336 | 1/23/2014 | WO | 00 |
