Patent Application
20090126837
The present invention relates to niobium (Nb) and titanium (Ti)-added interstitial free (IF) cold rolled steel sheets that are used as materials for automobiles, household electronic appliances, etc. More specifically, the present invention relates to highly formable IF cold rolled steel sheets whose yield strength is enhanced due to the distribution of fine precipitates, and a process for producing the IF cold rolled steel sheets.
In general, cold rolled steel sheets for use in automobiles and household electronic appliances are required to have excellent room-temperature aging resistance and bake hardenability, together with high strength and superior formability.
Aging is a strain aging phenomenon that arises from hardening caused by dissolved elements, such as C and N, fixed to dislocations. Since aging causes defect, called “stretcher strain”, it is important to secure excellent room-temperature aging resistance.
Bake hardenability means increase in strength due to the presence of dissolved carbon after press formation, followed by painting and drying, by leaving a slight small amount of carbon in a solid solution state. Steel sheets with excellent bake hardenability can overcome the difficulties of press formability resulting from high strength.
Room-temperature aging resistance and bake hardenability can be imparted to aluminum (Al)-killed steels by batch annealing of the Al-killed steels. However, extended time of the batch annealing causes low productivity of the Al-killed steels and severe variation in steel materials at different sites. In addition, Al-killed steels have a bake hardening (BH) value (a difference in yield strength before and after painting) of 10-20 MPa, which demonstrates that an increase in yield strength is low.
Under such circumstances, interstitial free (IF) steels with excellent room-temperature aging resistance and bake hardenability have been developed by adding carbide and nitride-forming elements, such as Ti and Nb, followed by continuous annealing.
For example, Japanese Unexamined Patent Publication No. Sho 57-041349 describes an enhancement in the strength of a Ti-based IF steel by adding 0.4-0.8% of manganese (Mn) and 0.04-0.12% of phosphorus (P). In very low carbon IF steels, however, P causes the problem of secondary working embrittlement due to segregation in grain boundaries.
Japanese Unexamined Patent Publication No. Hei 5-078784 describes an enhancement in strength by the addition of Mn as a solid solution strengthening element in an amount exceeding 0.9% and not exceeding 3.0%.
Korean Patent Laid-open No. 2003-0052248 describes an improvement in secondary working embrittlement resistance as well as strength and workability by the addition of 0.5-2.0% of Mn instead of P, together with aluminum (Al) and boron (B).
Japanese Unexamined Patent Publication No. Hei 10-158783 describes an enhancement in strength by reducing the content of P and using Mn and Si as solid solution strengthening elements. According to this publication, Mn is used in an amount of up to 0.5%, Al as a deoxidizing agent is used in an amount of 0.1%, and nitrogen (N) as an impurity is limited to 0.01% or less. If the Mn content is increased, the plating characteristics are worsened.
Japanese Unexamined Patent Publication No. Hei 6-057336 discloses an enhancement in the strength of an IF steel by adding 0.5-2.5% of copper (Cu) to form ε-Cu precipitates. High strength of the IF steel is achieved due to the presence of the ε-Cu precipitates, but the workability of the IF steel is worsened.
Japanese Unexamined Patent Publication Nos. Hei 9-227951 and Hei 10-265900 suggest technologies associated with improvement in workability or surface defects due to carbides by the use of Cu as a nucleus for precipitation of the carbides. According to the former publication, 0.005-0.1% of Cu is added to precipitate CuS during temper rolling of an IF steel, and the CuS precipitates are used as nuclei to form Cu—Ti—C—S precipitates during hot rolling. In addition, the former publication states that the number of nuclei forming a {111} plane parallel to the surface of a plate increases in the vicinity of the Cu—Ti—C—S precipitates during recrystallization, which contributes to an improvement in workability. According to the latter publication, 0.01-0.05% of Cu is added to an IF steel to obtain CuS precipitates and then the CuS precipitates are used as nuclei for precipitation of carbides to reduce the amount of dissolved carbon (C), leading to an improvement in surface defects. According to the prior art, since coarse CuS precipitates are used during production of cold rolled steel sheets, carbides remain in the final products. Further, since emulsion-forming elements, such as Ti and Zr, are added in an amount greater than the amount of sulfur (S) in an atomic weight ratio, a main portion of the sulfur (S) reacts with Ti or Zr rather than Cu.
On the other hand, Japanese Unexamined Patent Publication Nos. Hei 6-240365 and Hei 7-216340 describe the addition of a combination of Cu and P to improve the corrosion resistance of baking hardening type IF steels. According to these publications, Cu is added in an amount of 0.05-1.0% to ensure improved corrosion resistance. However, in actuality, Cu is added in an excessively large amount of 0.2% or more.
Japanese Unexamined Patent Publication Nos. Hei 10-280048 and Hei 10-287954 suggest the dissolution of. carbosulfide (Ti—C—S based) in a carbide at the time of reheating and annealing to obtain a solid solution in crystal grain boundaries, thereby achieving a bake hardening (BH) value (a difference in yield strength before and after baking) of 30 MPa or more.
According to the aforementioned publications, strength is enhanced by strengthening solid solution or using ε-Cu precipitates. Cu is used to form ε-Cu precipitates and improve corrosion resistance. In addition, Cu is used as a nucleus for precipitation of carbides. No mention is made in these publications about an increase in high yield ratio (i.e. yield strength/tensile strength) and a reduction in in-plane anisotropy index. If the tensile strength-to-yield strength ratio (i.e. yield ratio) of an IF steel sheet is high, the thickness of the IF steel sheet can be reduced, which is effective in weight reduction. In addition, if the in-plane anisotropy index of an IF steel sheet is low, fewer wrinkles and ears occur during processing and after processing, respectively.
It is one object of certain embodiments of the present invention to provide Nb and Ti-added IF cold rolled steel sheets that are capable of achieving a high yield ratio and a low in-plane anisotropy index.
It is another object of certain embodiments of the present invention to provide a process for producing the IF cold rolled steel sheets.
According to the present invention, there is provided a cold rolled steel sheet which has a composition comprising 0.01% or less of C, 0.01-0.2% of Cu, 0.005-0.08% of S, 0.1% or less of Al, 0.004% or less of N, 0.2% or less of P, 0.0001-0.002% of B, 0.002-0.04% of Nb, 0.005-0.15% of Ti, by weight, and the balance of Fe and other unavoidable impurities, wherein the composition satisfies the relationships: 1≦(Cu/63.5)/(S*/32)≦30 and S*═S−0.8×(Ti−0.8×(48/14)×N)×(32/48), and the steel sheet comprises CuS precipitates having an average size of 0.2 μm or less.
According to the present invention, there is provided a cold rolled steel sheet which has a composition comprising 0.01% or less of C, 0.01-0.2% of Cu, 0.01-0.3% of Mn, 0.005-0.08% of S, 0.1% or less of Al, 0.004% or less of N, 0.2% or less of P, 0.0001-0.002% of B, 0.002-0.04% of Nb, 0.005-0.15% of Ti, by weight, and the balance of Fe and other unavoidable impurities wherein the composition satisfies the following relationships: 1≦(Mn/55+Cu/63.5)/(S*/32)≦30 and S*═S−0.8×(Ti−0.833 (48/14)×N)×(32/48), and the steel sheet comprises (Mn,Cu)S precipitates having an average size of 0.2 μm or less.
According to the present invention, there is provided a cold rolled steel sheet which has a composition comprising 0.01% or less of C, 0.01-0.2% of Cu, 0.005-0.08% of S, 0.1% or less of Al, 0.004-0.02% of N, 0.2% or less of P, 0.0001-0.002% of B, 0.002-0.04% of Nb, 0.005-0.15% of Ti, by weight, and the balance of Fe and other unavoidable impurities wherein the composition satisfies the following relationships: 1≦(Cu/63.5)/(S*/32)≦30, 1≦(Al/27)/(N*/14)≦10, S*═S−0.8×(Ti−0.8×(48/14)×N)×(32/48) and N*═N−0.8×(Ti−0.8×(48/32)×S))×(14/48), and the steel sheet comprises CuS and AlN precipitates having an average size of 0.2 μm or less.
According to the present invention, there is provided a cold rolled steel sheet which has a composition comprising 0.01% or less of C, 0.01-0.2% of Cu, 0.01-0.3% of Mn, 0.005-0.08% of S, 0.1% or less of Al, 0.004-0.02% of N, 0.2% or less of P, 0.0001-0.002% of B, 0.002-0.04% of Nb, 0.005-0.15% of Ti, by weight, and the balance of Fe and other unavoidable impurities wherein the composition satisfies the following relationships: 1≦(Mn/55+Cu/63.5)/(S*/32)≦30, 1≦(Al/27)/(N*/14)≦10, S*═S−0.833 (Ti−0.8×(48/14)×N)×(32/48) and N*═N−0.8×(Ti−0.8×(48/32)×S))×(14/48), and the steel sheet comprises (Mn,Cu)S and AlN precipitates having an average size of 0.2 μm or less.
According to the present invention, there is provided a cold rolled steel sheet which has a composition comprising 0.01% or less of C, 0.08% or less of S, 0.1% or less of Al, 0.004% or less of N, 0.2% or less of P, 0.0001-0.002% of B, 0.002-0.04% of Nb, 0.005-0.15% of Ti, at least one kind selected from 0.01-0.2% of Cu, 0.01-0.3% of Mn and 0.004-0.2% of N, by weight, and the balance of Fe and other unavoidable impurities wherein the composition satisfies the following relationships: 1≦(Mn/55+Cu/63.5)/(S*/32)≦30, 1≦(Al/27)/(N*/14)≦10 (with the proviso that the N content is 0.004% or more), S═S−0.8×(Ti−0.8×(48/14)×N)×(32/48) and N*═N−0.8×(Ti−0.8×(48/32)×S))×(14/48), and the steel sheet comprises at least one kind selected from (Mn,Cu)S and AlN precipitates having an average size of 0.2 μm or less.
When the cold rolled steel sheets of the present invention satisfy the following relationships between the C, Ti, Nb, N and S contents: 0.8≦(Ti*/48+Nb/93)/(C/12)≦5.0 and Ti*═Ti−0.8×((48/14)×N+(48/32)×S), they show room-temperature non-aging properties. In addition, when solute carbon (Cs) [Cs═(C−Nb×12/93−Ti*×12/48)×10000 wherein Ti*═Ti−0.8×((48/14)×N+(48/32)×S), provided that when Ti* is less than 0, Ti* is defined as 0], which is determined by the C and Ti contents, is from 5 to 30, the cold rolled steel sheets of the present invention show bake hardenability.
Depending on the design of the compositions, the cold rolled steel sheets of the present invention have characteristics of soft cold rolled steel sheets of the order of 280 MPa and high-strength cold rolled steel sheets of the order of 340 MPa or more.
When the content of P in the compositions of the present invention is 0.015% or less, soft cold rolled steel sheets of the order of 280 MPa are produced. When the soft cold rolled steel sheets further contain at least one solid solution strengthening element selected from Si and Cr, or the P content is in the range of 0.015-0.2%, a high strength of 340 MPa or more is attained. The P content in the high-strength steels containing P alone is preferably in the range of 0.03% to 0.2%. The Si content in the high-strength steels is preferably in the range of 0.1 to 0.8%. The Cr content in the high-strength steels is preferably in the range of 0.2 to 1.2. In the case where the cold rolled steel sheets of the present invention contain at least one element selected from Si and Cr, the P content may be freely designed in an amount of 0.2% or less.
For better workability, the cold rolled steel sheets of the present invention may further contain 0.01-0.2 wt % of Mo.
According to the present invention, there is provided a process for producing the cold rolled steel sheets, the process comprising reheating a slab satisfying one of the compositions to a temperature of 1,100° C. or higher, hot rolling the reheated slab at a finish rolling temperature of the Ar3 transformation point or higher to provide a hot rolled steel sheet, cooling the hot rolled steel sheet at a rate of 300° C./min., winding the cooled steel sheet at 700° C. or lower, cold rolling the wound steel sheet, and continuously annealing the cold rolled steel sheet.
The present invention will be described in detail below.
Fine precipitates having a size of 0.2 μm or less are distributed in the cold rolled steel sheets of the present invention. Examples of such precipitates include MnS precipitates, CuS precipitates, and composite precipitates of MnS and CuS. These precipitates are referred to simply as “(Mn,Cu)S”.
The present inventors have found that when fine precipitates are distributed in Nb and Ti-added IF steels (also referred to simply as “Nb—Ti composite IF steels”), the yield strength of the IF steels is enhanced and the in-plane anisotropy index of the IF steels is lowered, thus leading to an improvement in workability. The present invention has been achieved based on this finding. The precipitates used in the present invention have drawn little attention in conventional IF steels. Particularly, the precipitates have not been actively used from the viewpoint of yield strength and in-plane anisotropy index.
Regulation of the components in the Nb—Ti composite IF steels is required to obtain (Mn,Cu)S precipitates and/or AlN precipitates. If the IF steels contain Ti, Zr and other elements, S and N preferentially react with Ti and Zr. Since the cold rolled steel sheets of the present invention are Nb—Ti composite IF steels, Ti reacts with C, N and S. Accordingly, it is necessary to regulate the components so that S and N are precipitated into (Mn,Cu)S and AlN forms, respectively.
The fine precipitates thus obtained allow the formation of minute crystal grains. Minuteness in the size of crystal grains relatively increases the proportion of crystal grain boundaries. Accordingly, the dissolved carbon is present in a larger amount in the crystal grain boundaries than within the crystal grains, thus achieving excellent room-temperature non-aging properties. Since the dissolved carbon present within the crystal grains can more freely migrate, it binds to movable dislocations, thus affecting the room-temperature aging properties. In contrast, the dissolved carbon segregated in stable positions, such as in the crystal grain boundaries and in the vicinity of the precipitates, is activated at a high temperature, for example, a temperature for painting/baking treatment, thus affecting the bake hardenability.
The fine precipitates distributed in the steel sheets of the present invention have a positive influence on the increase of yield strength arising from precipitation enhancement, improvement in strength-ductility balance, in-plane anisotropy index, and plasticity anisotropy. To this end, the fine (Mn,Cu)S precipitates and AlN precipitates must be uniformly distributed. According to the cold rolled steel sheets of the present invention, contents of components affecting the precipitation, composition between the components, production conditions, and particularly cooling rate after hot rolling, have a great influence on the distribution of the fine precipitates.
The constituent components of the cold rolled steel sheets according to the present invention will be explained.
The content of carbon (C) is preferably limited to 0.01% or less.
Carbon (C) affects the room-temperature aging resistance and bake hardenability of the cold rolled steel sheets. When the carbon content exceeds 0.01%, the addition of the expensive agents Nb and Ti is required to remove the remaining carbon, which is economically disadvantageous and is undesirable in terms of formability. When it is intended to achieve room-temperature aging resistance only, it is preferred to maintain the carbon content at a low level, which enables the reduction of the amount of the expensive agents Nb and Ti added. When it is intended to ensure desired bake hardenability, the carbon is preferably added in an amount of 0.001% or more, and more preferably 0.005% to 0.01%. When the carbon content is less than 0.005%, room-temperature aging resistance can be ensured without increasing the amounts of Nb and Ti.
The content of copper (Cu) is preferably in the range of 0.01-0.2%.
Copper serves to form fine CuS precipitates, which make the crystal grains fine. Copper lowers the in-plane anisotropy index of the cold rolled steel sheets and enhances the yield strength of the cold rolled steel sheets by precipitation promotion. In order to form fine precipitates, the Cu content must be 0.01% or more. When the Cu content is more than 0.2%, coarse precipitates are obtained. The Cu content is more preferably in the range of 0.03 to 0.2%.
The content of manganese (Mn) is preferably in the range of 0.01-0.3%.
Manganese serves to precipitate sulfur in a solid solution state in the steels as MnS precipitates, thereby preventing occurrence of hot shortness caused by the dissolved sulfur, or is known as a solid solution strengthening element. From such a technical standpoint, manganese is generally added in a large amount. The present inventors have found that when the manganese content is reduced and the sulfur content is optimized, very fine MnS precipitates are obtained. Based on this finding, the manganese content is limited to 0.3% or less. In order to ensure this characteristic, the manganese content must be 0.01% or more. When the manganese content is less than 0.01%, i.e. the sulfur content remaining in a solid solution state is high, hot shortness may occur. When the manganese content is greater than 0.3%, coarse MnS precipitates are formed, thus making it difficult to achieve desired strength. A more preferable Mn content is within the range of 0.01 to 0.12%.
The content of sulfur (S) is preferably limited to 0.08% or less.
Sulfur (S) reacts with Cu and/or Mn to form CuS and MnS precipitates, respectively. When the sulfur content is greater than 0.08%, the proportion of dissolved sulfur is increased. This increase of dissolved sulfur greatly deteriorates the ductility and formability of the steel sheets and increases the risk of hot shortness. In order to obtain as many CuS and/or MnS precipitates as possible, a sulfur content of 0.005% or more is preferred.
The content of aluminum (Al) is preferably limited to 0.1% or less.
Aluminum reacts with nitrogen (N) to form fine AlN precipitates, thereby completely preventing aging by dissolved nitrogen. When the nitrogen content is 0.004% or more, AlN precipitates are sufficiently formed. The distribution of the fine AlN precipitates in the steel sheets allows the formation of minute crystal grains and enhances the yield strength of the steel sheets by precipitation enhancement. A more preferable Al content is in the range of 0.01 to 0.1%.
The content of nitrogen (N) is preferably limited to 0.02% or less.
When it is intended to use AlN precipitates, nitrogen is added in an amount of up to 0.02%. Otherwise, the nitrogen content is controlled to 0.004% or less. When the nitrogen content is less than 0.004%, the number of the AlN precipitates is small, and therefore, the minuteness effects of crystal grains and the precipitation enhancement effects are negligible. In contrast, when the nitrogen content is greater than 0.02%, it is difficult to guarantee aging properties by use of dissolved nitrogen.
The content of phosphorus (P) is preferably limited to 0.2% or less.
Phosphorus is an element that has excellent solid solution strengthening effects while allowing a slight reduction in r-value. Phosphorus guarantees high strength of the steel sheets of the present invention in which the precipitates are controlled. It is desirable that the phosphorus content in steels requiring a strength of the order of 280 MPa be defined to 0.015% or less. It is desirable that the phosphorus content in high-strength steels of the order of 340 MPa be limited to a range exceeding 0.015% and not exceeding 0.2%. A phosphorus content exceeding 0.2% can lead to a reduction in ductility of the steel sheets. Accordingly, the phosphorus content is preferably limited to a maximum of 0.2%. When Si and Cr are added in the present invention, the phosphorus content can be appropriately controlled to be 0.2% or less to achieve the desired strength.
The content of boron (B) is preferably in the range of 0.0001 to 0.002%.
Boron is added to prevent occurrence of secondary working embrittlement. To this end, a preferable boron content is 0.0001% or more. When the boron content exceeds 0.002%, the deep drawability of the steel sheets may be markedly deteriorated.
The content of niobium (Nb) is preferably in the range of 0.002 to 0.04%.
Nb is added for the purpose of ensuring the non-aging properties and improving the formability of the steel sheets. Nb, which is a potent carbide-forming element, is added to steels to form NbC precipitates in the steels. In addition, the NbC precipitates permit the steel sheets to be well textured during annealing, thus greatly improving the deep drawability of the steel sheets. When the content of Nb added is not greater than 0.002%, the NbC precipitates are obtained in very small amounts. Accordingly, the steel sheets are not well textured and thus there is little improvement in the deep drawability of the steel sheets. In contrast, when the Nb content exceeds 0.04%, the NbC precipitates are obtained in very large amounts. Accordingly, the deep drawability and elongation of the steel sheets are lowered, and thus the formability of the steel sheets may be markedly deteriorated.
The content of titanium (Ti) is preferably in the range of 0.005 to 0.15%.
Titanium is added for the purpose of ensuring the non-aging properties and improving the formability of the steel sheets. Ti, which is a potent carbide-forming element, is added to steels to form TiC precipitates in the steels. The TiC precipitates allow the precipitation of dissolved carbon to ensure non-aging properties. When the content of Ti added is less than 0.005%, the TiC precipitates are obtained in very small amounts. Accordingly, the steel sheets are not well textured and thus there is little improvement in the deep drawability of the steel sheets. In contrast, when the titanium is added in an amount exceeding 0.15%, very large TiC precipitates are formed. Accordingly, minuteness effects of crystal grains are reduced, resulting in high in-plane anisotropy index, reduction of yield strength and marked worsening of plating characteristics.
To obtain (Mn,Cu)S and AlN precipitates, the Mn, Cu, S, Nb, Ti, Al, N and C contents are adjusted within the ranges defined by the following relationships. The respective components indicated in the following relationships are expressed as percentages by weight.
1≦(Cu/63.5)/(S*/32)≦30 (1)
S═S−0.8×(Ti−0.8×(48/14)×N)×(32/48) (2)
In Relationship 1, S*, which is determined by Relationship 2, represents the content of sulfur that does not react with Ti and thereafter reacts with Cu. To obtain fine CuS precipitates, it is preferred that the value of (Cu/63.5)/(S*/32) be equal to or greater than 1. If the value of (Cu/63.5)/(S*/32) is greater than 30, coarse CuS precipitates are distributed, which is undesirable. To stably obtain CuS precipitates having a size of 0.2 μm or less, the value of (Cu/63.5)/(S*/32) is preferably in the range of 1 to 20, more preferably 1 to 9, and most preferably 1 to 6.
1≦(Mn/55+Cu/63.5)/(S*/32)≦30 (3)
Relationship 3 is associated with the formation of (Mn,Cu)S precipitates, and is obtained by adding a Mn content to Relationship 1. To obtain effective (Mn,Cu)S precipitates, the value of (Mn/55+Cu/63.5)/(S*/32) must be 1 or greater. When the value of Relationship 3 is greater than 30, coarse (Mn,Cu)S precipitates are obtained. To stably obtain (Mn,Cu)S precipitates having a size of 0.2 μm or less, a more preferable value of (Cu/63.5)/(S*/32) is preferably in the range of 1 to 20, more preferably 1 to 9, and most preferably 1 to 6. When Mn and Cu are added together, the sum of Mn and Cu is more preferably 0.05-0.4%. The reason for this limitation to the sum of Mn and Cu is to obtain fine (Mn,Cu)S precipitates.
1≦(Al/27)/(N*/14)≦10 (4)
N═N−0.8×(Ti−0.8×(48/32)×S))×(14/48) (5)
Relationship 4 is associated with the formation of fine (Mn,Cu)S precipitates. In Relationship 4, N*, which is determined by Relationship 5, represents the content of nitrogen that does not react with Ti and thereafter reacts with Al. To obtain fine AlN precipitates, it is preferred that the value of (Al/27)/(N*/14) be in the range of 1-10. To obtain effective AlN precipitates, the value of (Al/27)/(N*/14) must be 1 or greater. If the value of (Al/27)/(N*/14) is greater than 10, coarse AlN precipitates are obtained and thus poor workability and low yield strength are caused. It is preferred that the value of (Al/27)/(N*/14) be in the range of 1 to 6.
The components of the cold rolled steel sheets according to the present invention may be combined in various ways according to the kind of precipitates to be obtained. For example, the present invention provides a cold rolled steel sheet which has a composition comprising 0.01% or less of C, 0.08% or less of S, 0.1% or less of Al, 0.004% or less of N, 0.2% or less of P, 0.0001-0.002% of B, 0.002-0.04% of Nb, 0.005-0.15% of Ti, at least one kind selected from 0.01-0.2% of Cu, 0.01-0.3% of Mn and 0.004-0.2% of N. by weight, and the balance of Fe and other unavoidable impurities, wherein the composition satisfies the following relationships: 1≦(Mn/55+Cu/63.5)/(S*/32)≦30, 1≦(Al/27)/(N*/14)≦10 (with the proviso that the N content is 0.004% or more), S*═S−0.8×(Ti−0.8×(48/14)×N)×(32/48) and N═N−0.8×(Ti−0.8×(48/32)×S))×(14/48), and the steel sheet comprises at least one kind selected from MnS, CuS, MnS and AlN precipitates having an average size of 0.2 μm or less. That is, one or more kinds selected from the group consisting of 0.01-0.2% of Cu, 0.01-0.3% of Mn and 0.004-0.2% of N lead to various combinations of (Mn,Cu)S and AlN precipitates having a size not greater than 0.2 μm.
In the steel sheets of the present invention, carbon is precipitated into NbC and TiC forms. Accordingly, the room-temperature aging resistance and bake hardenability of the steel sheets are affected depending on the conditions of dissolved carbon under which NbC and TiC precipitates are not obtained. Taking into account these requirements, it is most preferred that the Nb, Ti and C contents satisfy the following relationships.
0.8≦(Ti*/48+Nb/93)/(C/12)≦5.0 (6)
Ti*═Ti−0.8×((48/14)×N+(48/32)×S) (7)
Relationship 6 is associated with the formation of NbC and TiC precipitates to remove the carbon in a solid solution state, thereby achieving room-temperature non-aging properties. In Relationship 6, Ti*, which is determined by Relationship 7, represents the content of titanium that reacts with N and S and thereafter reacts with C.
When the value of (Ti*/48+Nb/93)/(C/12) is less than 0.8, it is difficult to ensure room-temperature non-aging properties. In contrast, when the value of (Ti*/48+Nb/93)/(C/12) is greater than 5, the amounts of Nb and Ti remaining in a solid solution state in the steels are large, which deteriorates the ductility of the steels. When it is intended to achieve room-temperature non-aging properties without securing bake hardenability, it is preferred to limit the carbon content to 0.005% or less. Although the carbon content is more than 0.005%, room-temperature non-aging properties can be achieved when Relationship 6 is satisfied but the amounts of Nb and TiC precipitates are increased, thus deteriorating the workability of the steel sheets.
Cs═(C—Nb×12/93−Ti*×12/48)×10000 (8)
(provided that when Ti* is less than 0, Ti* is defined as 0.)
Relationship 8 is associated with the achievement of bake hardenability. Cs, which is expressed in ppm by Relationship 8, represents the content of dissolved carbon that is not precipitated into NbC and TiC forms. In order to achieve a high bake hardening value, the Cs value must be 5 ppm or more. If the Cs value exceeds 30 ppm, the content of dissolved carbon is increased, making it difficult to attain room-temperature non-aging properties.
It is advantageous that the fine precipitates are uniformly distributed in the compositions of the present invention. It is preferable that the precipitates have an average size of 0.2 μm or less. According to a study conducted by the present inventors, when the precipitates have an average size greater than 0.2 μm, the steel sheets have poor strength and low in-plane anisotropy index. Further, large amounts of precipitates having a size of 0.2 μm or less are distributed in the compositions of the present invention. While the number of the distributed precipitates is not particularly limited, it is more advantageous with higher number of the precipitates. The number of the distributed precipitates is preferably 1×105/mm2 or more, more preferably 1×106/mm2 or more, and most preferably 1×107/mm2 or more. The plasticity-anisotropy index is increased and the in-plane anisotropy index is lowered with increasing number of the precipitates, and as a result, the workability is greatly improved. It is commonly known that there is a limitation in increasing the workability because the in-plane anisotropy index is increased with increasing plasticity-anisotropy index. It is worth noting that as the number of the precipitates distributed in the steel sheets of the present invention increases, the plasticity-anisotropy index of the steel sheets is increased and the in-plane anisotropy index of the steel sheets is lowered. The steel sheets of the present invention in which the fine precipitates are formed satisfy a yield ratio (yield strength/tensile strength) of 0.58 or higher.
@ When the steel sheets of the present invention are applied to high-strength steel sheets, they may further contain at least one solid solution strengthening element selected from P, Si and Cr. The addition effects of P have been previously described, and thus their explanation is omitted.
The content of silicon (Si) is preferably in the range of 0.1 to 0.8%.
Si is an element that has solid solution strengthening effects and shows a slight reduction in elongation. Si guarantees high strength of the steel sheets of the present invention in which the precipitates are controlled. Only when the Si content is 0.1% or more, high strength can be ensured. However, when the Si content is more than 0.8%, the ductility of the steel sheets is deteriorated.
The content of chromium (Cr) is preferably in the range of 0.2 to 1.2%.
Cr is an element that has solid solution strengthening effects, lowers the secondary working embrittlement temperature, and lowers the aging index due to the formation of Cr carbides. Cr guarantees high strength of the steel sheets of the present invention in which the precipitates are controlled and serves to lower the in-plane anisotropy index of the steel sheets. Only when the Cr content is 0.2% or more, high strength can be ensured. However, when the Cr content exceeds 1.2%, the ductility of the steel sheets is deteriorated.
The cold rolled steel sheets of the present invention may further contain molybdenum (Mo).
The content of molybdenum (Mo) in the cold rolled steel sheets of the present invention is preferably in the range of 0.01 to 0.2%.
Mo is added as an element that increases the plasticity-anisotropy index of the steel sheets. Only when the molybdenum content is not lower than 0.01%, the plasticity-anisotropy index of the steel sheets is increased. However, when the molybdenum content exceeds 0.2%, the plasticity-anisotropy index is not further increased and there is a danger of hot shortness.
Production of cold rolled steel sheets Hereinafter, a process for producing the cold rolled steel sheets of the present invention will be explained with reference to the preferred embodiments that follow. Various modifications of the embodiments of the present invention can be made, and such modifications are within the scope of the present invention.
The process of the present invention is characterized in that a steel satisfying one of the steel compositions defined above is processed through hot rolling and cold rolling to form precipitates having an average size of 0.2 μm or less in a cold rolled sheet. The average size of the precipitates in the cold rolled plate is affected by the design of the steel composition and the processing conditions, such as reheating temperature and winding temperature. Particularly, cooling rate after hot rolling has a direct influence on the average size of the precipitates.
Hot Rolling Conditions
In the present invention, a steel satisfying one of the compositions defined above is reheated, and is then subjected to hot rolling. The reheating temperature is preferably 1,100° C. or higher. When the steel is reheated to a temperature lower than 1,100° C., coarse precipitates formed during continuous casting are not completely dissolved and remain. The coarse precipitates still remain even after hot rolling.
It is preferred that the hot rolling is performed at a finish rolling temperature not lower than the Ar3 transformation point. When the finish rolling temperature is lower than the Ar3 transformation point, rolled grains are created, which deteriorates the workability and causes poor strength.
The cooling is preferably performed at a rate of 300° C./min or higher before winding and after hot rolling. Although the composition of the components is controlled to obtain fine precipitates the precipitates may have an average size greater than 0.2 μm at a cooling rate of less than 300° C./min. That is, as the cooling rate is increased, many nuclei are created and thus the size of the precipitates becomes finer and finer. Since the size of the precipitates is decreased with increasing cooling rate, it is not necessary to define the upper limit of the cooling rate. When the cooling rate is higher than 1,000° C./min., however, a significant improvement in the size reduction effects of the precipitates is not further shown. Therefore, the cooling rate is preferably in the range of 300-1000° C./min.
Winding Conditions
After the hot rolling, winding is performed at a temperature not higher than 700° C. When the winding temperature is higher than 700° C., the precipitates are grown too coarsely, thus making it difficult to ensure high strength.
Cold Rolling Conditions
The steel is cold rolled at a reduction rate of 50-90%. Since a cold reduction rate lower than 50% leads to creation of a small amount of nuclei upon annealing recrystallization, the crystal grains are grown excessively upon annealing, thereby coarsening of the crystal grains recrystallized through annealing, which results in reduction of the strength and formability. A cold reduction rate higher than 90% leads to enhanced formability, while creating an excessively large amount of nuclei, so that the crystal grains recrystallized through annealing become too fine, thus deteriorating the ductility of the steel.
Continuous Annealing
Continuous annealing temperature plays an important role in determining the mechanical properties of the final product. According to the present invention, the continuous annealing is preferably performed at a temperature of 700 to 900° C. When the continuous annealing is performed at a temperature lower than 700° C., the recrystallization is not completed and thus a desired ductility cannot be ensured. In contrast, when the continuous annealing is performed at a temperature higher than 900° C., the recrystallized grains become coarse and thus the strength of the steel is deteriorated. The continuous annealing is maintained until the steel is completely recrystallized. The recrystallization of the steel can be completed for about 10 seconds or more. The continuous annealing is preferably performed for 10 seconds to 30 minutes.
The present invention will now be described in more detail with reference to the following examples.
The mechanical properties of steel sheets produced in the following examples were evaluated according to the ASTM E-8 standard test methods. Specifically, each of the steel sheets was machined to obtain standard samples. The yield strength, tensile strength, elongation, plasticity-anisotropy index (rm value) and in-plane anisotropy index (Δr value), and the aging index were measured using a tensile strength tester (available from INSTRON Company, Model 6025). The plasticity-anisotropy index rm and in-plane anisotropy index (Ar value) were calculated by the following equations: rm=(r0+2r45+r90)/4 and Δr=(r0−2r45+r90)/2, respectively.
The aging index of the steel sheets is defined as a yield point elongation measured by annealing each of the samples, followed by 1.0% skin pass rolling and thermally processing at 100° C. for 2 hours. The bake hardening (BH) value of the standard samples was measured by the following procedure. After a 2% strain was applied to each of the samples, the strained sample was annealed at 170° C. for 20 minutes. The yield strength of the annealed sample was measured. The BH value was calculated by subtracting the yield strength measured before annealing from the yield strength value measured after annealing.
First, steel slabs were prepared in accordance with the compositions shown in the following tables. The steel slabs were reheated and finish hot-rolled to provide hot rolled steel sheets. The hot rolled steel sheets were cooled at a rate of 400° C./min., wound at 650° C., cold-rolled at a reduction rate of 75%, followed by continuous annealing to produce cold rolled steel sheets. At this time, the finish hot rolling was performed at 910° C., which is above the Ar3 transformation point, and the continuous annealing was performed by heating the hot rolled steel sheets at a rate of 10° C./second to 830° C. for 40 seconds to produce the final cold rolled steel sheets.
First, steel slabs were prepared in accordance with the compositions shown in the following tables. The steel slabs were reheated and finish hot-rolled to provide hot rolled steel sheets. The hot rolled steel sheets were cooled at a rate of 400° C./min., wound at 650° C., cold-rolled at a reduction rate of 75%, followed by continuous annealing to produce cold rolled steel sheets. At this time, the finish hot rolling was performed at 910° C., which is above the Ar3 transformation point, and the continuous annealing was performed by heating the hot rolled steel sheets at a rate of 10° C./second to 830° C. for 40 seconds to produce the final cold rolled steel sheets.
First, steel slabs were prepared in accordance with the compositions shown in the following tables. The steel slabs were reheated and finish hot-rolled to provide hot rolled steel sheets. The hot rolled steel sheets were cooled at a rate of 400° C./min., wound at 650° C., cold-rolled at a reduction rate of 75%, followed by continuous annealing to produce cold rolled steel sheets. At this time, the finish hot rolling was performed at 910° C., which is above the Ar3 transformation point, and the continuous annealing was performed by heating the hot rolled steel sheets at a rate of 10° C./second to 830° C. for 40 seconds to produce the final cold rolled steel sheets.
First, steel slabs were prepared in accordance with the compositions shown in the following tables. The steel slabs were reheated and finish hot-rolled to provide hot rolled steel sheets. The hot rolled steel sheets were cooled at a rate of 400° C./min., wound at 650° C., cold-rolled at a reduction rate of 75%, followed by continuous annealing to produce cold rolled steel sheets. At this time, the finish hot rolling was performed at 910° C., which is above the Ar3 transformation point, and the continuous annealing was performed by heating the hot rolled steel sheets at a rate of 10° C./second to 830° C. for 40 seconds to produce the final cold rolled steel sheets.
First, steel slabs were prepared in accordance with the compositions shown in the following tables. The steel slabs were reheated and finish hot-rolled to provide hot rolled steel sheets. The hot rolled steel sheets were cooled at a rate of 400° C./min., wound at 650° C., cold-rolled at a reduction rate of 75%, followed by continuous annealing to produce cold rolled steel sheets. At this time, the finish hot rolling was performed at 910° C., which is above the Ar3 transformation point, and the continuous annealing was performed by heating the hot rolled steel sheets at a rate of 10° C./second to 830° C. for 40 seconds to produce the final cold rolled steel sheets.
First, steel slabs were prepared in accordance with the compositions shown in the following tables. The steel slabs were reheated and finish hot-rolled to provide hot rolled steel sheets. The hot rolled steel sheets were cooled at a rate of 400° C./min., wound at 650° C., cold-rolled at a reduction rate of 75%, followed by continuous annealing to produce cold rolled steel sheets. At this time, the finish hot rolling was performed at 910° C., which is above the Ar3 transformation point, and the continuous annealing was performed by heating the hot rolled steel sheets at a rate of 10° C./second to 830° C. for 40 seconds to produce the final cold rolled steel sheets.
First, steel slabs were prepared in accordance with the compositions shown in the following tables. The steel slabs were reheated and finish hot-rolled to provide hot rolled steel sheets. The hot rolled steel sheets were cooled at a rate of 400 ° C./min., wound at 650° C., cold-rolled at a reduction rate of 75%, followed by continuous annealing to produce cold rolled steel sheets. At this time, the finish hot rolling was performed at 910° C., which is above the Ar3 transformation point, and the continuous annealing was performed by heating the hot rolled steel sheets at a rate of 10° C./second to 830° C. for 40 seconds to produce the final cold rolled steel sheets.
First, steel slabs were prepared in accordance with the compositions shown in the following tables. The steel slabs were reheated and finish hot-rolled to provide hot rolled steel sheets. The hot rolled steel sheets were cooled at a rate of 400° C./min., wound at 650° C., cold-rolled at a reduction rate of 75%, followed by continuous annealing to produce cold rolled steel sheets. At this time, the finish hot rolling was performed at 910° C., which is above the Ar3 transformation point, and the continuous annealing was performed by heating the hot rolled steel sheets at a rate of 10° C./second to 830° C. for 40 seconds to produce the final cold rolled steel sheets.
First, steel slabs were prepared in accordance with the compositions shown in the following tables. The steel slabs were reheated and finish hot-rolled to provide hot rolled steel sheets. The hot rolled steel sheets were cooled at a rate of 400° C./min., wound at 650° C., cold-rolled at a reduction rate of 75%, followed by continuous annealing to produce cold rolled steel sheets. At this time, the finish hot rolling was performed at 910° C., which is above the Ar3 transformation point, and the continuous annealing was performed by heating the hot rolled steel sheets at a rate of 10° C./second to 830° C. for 40 seconds to produce the final cold rolled steel sheets.
0.55
First, steel slabs were prepared in accordance with the compositions shown in the following tables. The steel slabs were reheated and finish hot-rolled to provide hot rolled steel sheets. The hot rolled steel sheets were cooled at a rate of 400° C./min., wound at 650° C., cold-rolled at a reduction rate of 75%, followed by continuous annealing to produce cold rolled steel sheets. At this time, the finish hot rolling was performed at 910° C., which is above the Ar3 transformation point, and the continuous annealing was performed by heating the hot rolled steel sheets at a rate of 10° C./second to 830° C. for 40 seconds to produce the final cold rolled steel sheets.
0
0.0015
0.13
First, steel slabs were prepared in accordance with the compositions shown in the following tables. The steel slabs were reheated and finish hot-rolled to provide hot rolled steel sheets. The hot rolled steel sheets were cooled at a rate of 400° C./min., wound at 650° C., cold-rolled at a reduction rate of 75%, followed by continuous annealing to produce cold rolled steel sheets. At this time, the finish hot rolling was performed at 910° C., which is above the Ar3 transformation point, and the continuous annealing was performed by heating the hot rolled steel sheets at a rate of 10° C./second to 830° C. for 40 seconds to produce the final cold rolled steel sheets.
0.55
0
0.0027
0.55
First, steel slabs were prepared in accordance with the compositions shown in the following tables. The steel slabs were reheated and finish hot-rolled to provide hot rolled steel sheets. The hot rolled steel sheets were cooled at a rate of 400° C./min., wound at 650° C., cold-rolled at a reduction rate of 75%, followed by continuous annealing to produce cold rolled steel sheets. At this time, the finish hot rolling was performed at 910° C., which is above the Ar3 transformation point, and the continuous annealing was performed by heating the hot rolled steel sheets at a rate of 10° C./second to 830° C. for 40 seconds to produce the final cold rolled steel sheets.
First, steel slabs were prepared in accordance with the compositions shown in the following tables. The steel slabs were reheated and finish hot-rolled to provide hot rolled steel sheets. The hot rolled steel sheets were cooled at a rate of 400° C./min., wound at 650° C., cold-rolled at a reduction rate of 75%, followed by continuous annealing to produce cold rolled steel sheets. At this time, the finish hot rolling was performed at 910° C., which is above the Ar3 transformation point, and the continuous annealing was performed by heating the hot rolled steel sheets at a rate of 10° C./second to 830° C. for 40 seconds to produce the final cold rolled steel sheets.
First, steel slabs were prepared in accordance with the compositions shown in the following tables. The steel slabs were reheated and finish hot-rolled to provide hot rolled steel sheets. The hot rolled steel sheets were cooled at a rate of 400° C./min., wound at 650° C., cold-rolled at a reduction rate of 75%, followed by continuous annealing to produce cold rolled steel sheets. At this time, the finish hot rolling was performed at 910° C., which is above the Ar3 transformation point, and the continuous annealing was performed by heating the hot rolled steel sheets at a rate of 10° C./second to 830° C. for 40 seconds to produce the final cold rolled steel sheets.
209
The preferred embodiments illustrated in the present invention do not serve to limit the present invention, but are set forth for illustrative purposes. Any embodiment having substantially the same constitution and the same operational effects thereof as the technical spirit of the present invention as defined in the appended claims is encompassed within the technical scope of the present invention.
As apparent from the above description, according to the cold rolled steel sheets of the present invention, the distribution of fine precipitates in Nb—Ti composite IF steels allows the formation of minute crystal grains, and as a result, the in-plane anisotropy index is lowered and the yield strength is enhanced by precipitation enhancement.
| Number | Date | Country | Kind |
|---|---|---|---|
| 10-2005-0037183 | May 2005 | KR | national |
| 10-2005-0129236 | Dec 2005 | KR | national |
| 10-2005-0129237 | Dec 2005 | KR | national |
| 10-2005-0129239 | Dec 2005 | KR | national |
| 110-2005-0129235 | Dec 2005 | KR | national |
| Filing Document | Filing Date | Country | Kind | 371c Date |
|---|---|---|---|---|
| PCT/KR06/01670 | 5/3/2006 | WO | 00 | 10/30/2007 |
