CN116783316A - Wire and part with improved delayed fracture resistance and method of making same - Google Patents

Wire and part with improved delayed fracture resistance and method of making same Download PDF

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Publication number
CN116783316A
CN116783316A CN202180090890.8A CN202180090890A CN116783316A CN 116783316 A CN116783316 A CN 116783316A CN 202180090890 A CN202180090890 A CN 202180090890A CN 116783316 A CN116783316 A CN 116783316A
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China
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delayed fracture
present disclosure
content
fracture resistance
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CN202180090890.8A
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Chinese (zh)
Inventor
全英洙
李相润
崔锡欢
崔明洙
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Posco Holdings Inc
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Posco Co Ltd
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    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/18—Hardening; Quenching with or without subsequent tempering
    • C21D1/25—Hardening, combined with annealing between 300 degrees Celsius and 600 degrees Celsius, i.e. heat refining ("Vergüten")
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/26—Methods of annealing
    • C21D1/32—Soft annealing, e.g. spheroidising
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/06—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of rods or wires
    • C21D8/065—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of rods or wires of ferrous alloys
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/0093—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for screws; for bolts
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/52—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for wires; for strips ; for rods of unlimited length
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/52—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for wires; for strips ; for rods of unlimited length
    • C21D9/525—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for wires; for strips ; for rods of unlimited length for wire, for rods
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/001—Austenite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/008—Martensite

Abstract

Wires and components having improved delayed fracture resistance and methods of making the same are disclosed. Wires with improved delayed fracture resistance according to the present disclosure comprise in weight%: 0.15% to 0.30% of C, 0.15% to 0.25% of Si, 0.95% to 1.35% of Mn, 0.030% or less of P, 0.030% or less of S, 0.015% to 0.030% of Ti, 0.0010% to 0.0040% of B, 0.0010% to 0.0080% of N, and Fe and unavoidable impurities as the balance, and the wire satisfies formula 1.[ formula 1] 2.0.ltoreq.5.5xSi+[ Mn ].ltoreq.2.4 (in formula 1, [ Si ] and [ Mn ] represent the content (wt.%) of the corresponding element).

Description

Wire and part with improved delayed fracture resistance and method of making same
Technical Field
The present disclosure relates to wires and parts having improved delayed fracture resistance and methods of manufacturing the same, and more particularly, to wires and parts that can be used for fastening bolts and the like of automobiles and structures exposed to various stress and corrosive environments and methods of manufacturing the same.
Background
With weight reduction and miniaturization of automobiles and structures, high strength is required for wires for fastening bolts and the like of automobiles and structures. In general, cold working, grain refinement, martensite strengthening, precipitation strengthening, and the like are used to increase the strength of steel.
However, dislocations, grain boundaries, martensite lath boundaries, fine precipitate boundaries for strengthening lead to poor delayed fracture properties by acting as hydrogen traps in the steel. For this reason, delayed fracture properties become poor in high-strength bolts having a tensile strength of 1GPa or more.
In order to solve this problem, cr-Mo alloy steel added with Mo is used in steel for high-strength bolts having a tempered martensite structure with a tensile strength of 1GPa or more. However, in response to the demand for cost reduction with the development of bolt manufacturing technology, cr-B steel has been tried instead of Cr-Mo steel. Therefore, cost reduction is achieved by using Cr-B steel for bolts used in the structure without significant impact on safety. Then, after the safety thereof is determined, cr-B steel is used for some fastening bolts of automobiles.
In addition, in the automotive industry, there is a need to develop materials for bolts that can reduce costs more than Cr-B steel. In response to this demand, a technical development has recently been made to apply Mn-B steel using Mn cheaper than Cr to a high strength bolt of 1GPa or more.
However, since Mn causes high solid solution strengthening in a ferrite matrix as compared with Cr, mn-B steel may cause cracking in a threaded portion of a bolt. Therefore, steel of high content Mn added for manufacturing a high strength bolt of 1GPa or more is difficult to be used for the high strength bolt because delayed fracture may occur in a threaded portion of the bolt due to cracking.
Disclosure of Invention
Technical problem
In one aspect, the present disclosure is directed to providing a wire rod for a high strength bolt with improved delayed fracture resistance, a bolt, and a method of manufacturing the same by optimizing a solid solution strengthening effect of Mn-B steel and improving formability via controlling an alloy element.
Technical proposal
A wire having improved delayed fracture resistance according to one exemplary embodiment of the present disclosure comprises in weight%: 0.15% to 0.30% of C, 0.15% to 0.25% of Si, 0.95% to 1.35% of Mn, 0.030% or less of P, 0.030% or less of S, 0.015% to 0.030% of Ti, 0.0010% to 0.0040% of B, 0.0010% to 0.0080% of N, and Fe and unavoidable impurities as the balance, and the wire satisfies formula 1.
[ 1]2.0 is less than or equal to 5.5 xSiMn is less than or equal to 2.4
In formula 1, [ Si ] and [ Mn ] represent the content (wt%) of the corresponding element.
In another exemplary embodiment of the present disclosure, the wire may satisfy formula 2.
[ 2]1.0 < [ Ti ]/3.42[ N ] < 2.0
In formula 2, [ Ti ] and [ N ] represent the content (wt%) of the corresponding element.
In another exemplary embodiment of the present disclosure, the size of the TiN inclusion may be 15 μm or less.
A method for manufacturing a wire rod having improved delayed fracture resistance according to an exemplary embodiment of the present disclosure includes: a step of finish rolling a steel material at 880 ℃ to 980 ℃, the steel material comprising in weight-%: 0.15% to 0.30% of C, 0.15% to 0.25% of Si, 0.95% to 1.35% of Mn, 0.030% or less of P, 0.030% or less of S, 0.015% to 0.030% of Ti, 0.0010% to 0.0040% of B, 0.0010% to 0.0080% of N, and Fe and unavoidable impurities as the balance, and the steel satisfies formula 1;
and a step of winding at 830 to 930 ℃.
[ 1]2.0 is less than or equal to 5.5 xSiMn is less than or equal to 2.4
In formula 1, [ Si ] and [ Mn ] represent the content (wt%) of the corresponding element.
In another exemplary embodiment of the present disclosure, the steel may satisfy formula 2.
[ 2]1.0 < [ Ti ]/3.42[ N ] < 2.0
In formula 2, [ Ti ] and [ N ] represent the content (wt%) of the corresponding element.
A method for manufacturing a part with improved delayed fracture resistance according to one exemplary embodiment of the present disclosure includes: a step of drawing a wire rod manufactured according to the present disclosure; a step of spheroidizing heat treatment of the drawn wire rod at 745 to 770 ℃; a step of heating the spheroidized heat-treated drawn wire rod at 870 ℃ to 940 ℃; a step of quenching the spheroidized heat-treated drawn wire rod at 50 to 80 ℃; and tempering the quenched component at 400 ℃ to 600 ℃.
A part with improved delayed fracture resistance according to one exemplary embodiment of the present disclosure comprises in weight-%: 0.15% to 0.30% of C, 0.15% to 0.25% of Si, 0.95% to 1.35% of Mn, 0.030% or less of P, 0.030% or less of S, 0.015% to 0.030% of Ti, 0.0010% to 0.0040% of B, 0.0010% to 0.0080% of N, and Fe and unavoidable impurities as the balance, and the component satisfies formula 1.
[ 1]2.0 is less than or equal to 5.5 xSiMn is less than or equal to 2.4
In formula 1, [ Si ] and [ Mn ] represent the content (wt%) of the corresponding element.
In another exemplary embodiment of the present disclosure, the component satisfies formula 2.
[ 2]1.0 < [ Ti ]/3.42[ N ] < 2.0
In formula 2, [ Ti ] and [ N ] represent the content (wt%) of the corresponding element.
In another exemplary embodiment of the present disclosure, the component comprises 0.3% to 2% by volume fraction of the retained austenite structure and the retained tempered martensite structure.
Advantageous effects
A high strength bolt component with improved delayed fracture resistance according to one exemplary embodiment of the present disclosure improves formability during working of a threaded portion of an Mn-B steel bolt. Accordingly, delayed fracture of the 1 GPa-order high-strength bolt can be suppressed by preventing cracks in the threaded portion of the bolt.
Drawings
Fig. 1 is an image of the threaded portion of comparative example 3 before evaluating the delayed fracture resistance.
Detailed Description
Not all elements of the exemplary embodiments are described in the present specification, and descriptions of general contents or overlapping contents between the exemplary embodiments in the technical field to which the present disclosure pertains are omitted.
Furthermore, when a component is described as "comprising" a certain component, it is intended that it may also include, rather than exclude, additional components, unless otherwise indicated.
Unless the context clearly indicates otherwise, singular expressions include plural expressions.
Hereinafter, the present disclosure is described in detail.
The following exemplary embodiments are provided to fully convey the concept of the present disclosure to those of ordinary skill in the art to which the present disclosure pertains. The present disclosure is not limited to the exemplary embodiments presented, but may also be embodied in other forms.
The inventors of the present disclosure found that by controlling the content of Si and Mn, it is possible to improve formability by optimizing the solid solution strengthening effect while securing strength, and therefore, since cracking caused by poor formability of a threaded portion is suppressed, it is possible to improve delayed fracture resistance.
Furthermore, they found that by controlling the contents of Ti and N and controlling the size of TiN inclusion, a fine grain size can be obtained, and thus formability can be improved and delayed fracture resistance can be ensured.
A wire having improved delayed fracture resistance according to one exemplary embodiment of the present disclosure comprises in weight%: 0.15 to 0.30% of C, 0.15 to 0.25% of Si, 0.95 to 1.35% of Mn, 0.030% or less of P, 0.030% or less of S, 0.015 to 0.030% of Ti, 0.0010 to 0.0040% of B, 0.0010 to 0.0080% of N, and Fe and unavoidable impurities as the balance.
The reason for limiting the content of the alloy element will be described in detail. Unless otherwise indicated, the amounts are in weight percent.
The content of carbon (C) is 0.15% to 0.30%.
C is an element added to ensure the strength of the product. If the carbon content is less than 0.15%, it is difficult to ensure the target strength. If it exceeds 0.30%, delayed fracture characteristics may become poor because the formation of retained austenite having excellent mechanical stability is hindered by hydrostatic pressure formed at lath martensite during quenching. Therefore, in the present disclosure, the C content is limited to 0.15% to 0.30%.
The content of silicon (Si) is 0.15% to 0.25%.
Si is an element used not only for deoxidizing steel but also for securing strength by solid solution strengthening. If the Si content is less than 0.15%, deoxidization of the steel and strength improvement by solid solution strengthening may be insufficient. In addition, if it exceeds 0.25%, formability and impact properties may be poor due to solid solution strengthening. Therefore, in the present disclosure, the Si content is limited to 0.15% to 0.25%.
The content of manganese (Mn) is 0.95% to 1.35%.
Mn is an element that improves hardenability. Which is a very useful element to provide a solid solution strengthening effect by forming a substitutional solid solution in the matrix tissue. If the Mn content is less than 0.95%, it is difficult to secure the strength desired in the present disclosure due to the solid solution strengthening effect and insufficient hardenability. If the Mn content exceeds 1.35%, formability may be poor due to the solid-solution strengthening effect. Therefore, in the present disclosure, the Mn content is limited to 0.95% to 1.35%
The content of phosphorus (P) is 0.030% or less (excluding 0%).
P is an element that segregates in grain boundaries and reduces toughness and delayed fracture resistance. Therefore, in the present disclosure, the upper limit of the P content is limited to 0.030%.
The content of sulfur (S) is 0.030% or less (excluding 0%).
Like P, S segregates in grain boundaries and reduces toughness. Furthermore, it hinders hot rolling by forming a low melting emulsion. Therefore, in the present disclosure, the upper limit of the S content is limited to 0.030%.
The content of titanium (Ti) is 0.015% to 0.030%.
Ti is an element that combines with N introduced into steel to form titanium carbonitride (TiN). In the present disclosure, tiN can prevent cracking due to poor formability of parts and improve delayed fracture resistance by reducing grain size. Further, since Ti forms TiN, it can prevent free N from combining with B to form BN that deteriorates formability. If the Ti content is less than 0.015%, tiN is not sufficiently formed, and free N forms BN. Therefore, the hardening effect of B cannot be utilized. If it exceeds 0.03%, the delayed fracture resistance may become poor due to the formation of coarse carbonitrides. Therefore, in the present disclosure, the Ti content is limited to 0.015% to 0.03%.
The content of boron (B) is 0.0010% to 0.0040%.
B is an element that improves hardenability. If the B content is less than 0.0010%, it is difficult to expect improvement of hardenability. In addition, if it exceeds 0.0040%, fe is formed in the grain boundary 23 (CB) 6 In the case of carbide, delayed fracture resistance becomes poor due to embrittlement of austenite grain boundaries, and formability becomes poor due to formation of BN. Therefore, in the present disclosure, the B content is limited to 0.0010% to 0.0040%.
The nitrogen (N) content is 0.0010% to 0.0080%.
N is an element that forms carbonitrides. If the N content is less than 0.0010%, tiN precipitates of reduced grain size may be formed insufficiently. If it exceeds 0.0080%, toughness and ductility of steel may become poor due to an increase in the content of dissolved nitrogen, and free N may combine with B to form BN deteriorating formability. Therefore, in the present disclosure, the N content is limited to 0.0010% to 0.0080%.
The remainder of the alloy composition is iron (Fe). However, the wire rods of the present disclosure with improved delayed fracture resistance may contain other impurities that may be contained in common industrial steel production processes. These impurities are well known to those of ordinary skill in the art to which the present disclosure pertains, and the type and content thereof are not particularly limited in the present disclosure.
Wires having improved delayed fracture resistance according to one exemplary embodiment of the present disclosure satisfy formula 1:
[ 1]2.0 is less than or equal to 5.5 xSiMn is less than or equal to 2.4
In formula 1, [ Si ] and [ Mn ] represent the content (wt%) of the corresponding element.
In the present disclosure, the contents of Si and Mn are controlled so that formability and delayed fracture resistance of a wire rod can be improved by suppressing excessive solid solution strengthening while securing strength by a solid solution strengthening effect. Equation 1 is a formula for optimizing the solid solution strengthening effect. In equation 1, if the value of 5.5× [ Si ] + [ Mn ] is less than 2.0, the strength desired by the present disclosure cannot be ensured. If the value of 5.5x [ Si ] + [ Mn ] exceeds 2.4, delayed fracture may be caused due to cracking caused by poor formability during the forming of the high-strength member due to excessive solid solution strengthening. Therefore, in the present disclosure, in order to improve the delayed fracture resistance, the value of 5.5× [ Si ] + [ Mn ] is limited to 2.0 to 2.4.
Furthermore, the wire rod having improved delayed fracture resistance according to one exemplary embodiment of the present disclosure satisfies formula 2.
[ 2]1.0 < [ Ti ]/3.42[ N ] < 2.0
In formula 2, [ Ti ] and [ N ] represent the content (wt%) of the corresponding element.
The present disclosure aims to improve the delayed fracture resistance of a wire by reducing the grain size and improving the formability. The inventors of the present disclosure found that the grain size can be reduced by controlling the size of TiN inclusions, and that formability and delayed fracture resistance can be ensured by suppressing the formation of BN. Equation 2 is a formula obtained to control the size of TiN inclusions and suppress the formation of BN. In formula 2, if the value of [ Ti ]/3.42[ N ] is 1.0 or less, formability may become poor due to BN or the like formed by free N not bonded with Ti. If the value of [ Ti ]/3.42[ N ] is 2.0 or more, the size of TiN inclusion increases due to excessive Ti, and the grain refining effect cannot be achieved. Thus, in the present disclosure, the value of [ Ti ]/3.42[ N ] is limited to satisfy 1.0 < [ Ti ]/3.42[ N ] < 2.0.
In the present disclosure, the size of TiN inclusions used to reduce the grain size may be 15 μm or less. If the size of the TiN inclusion exceeds 15 μm, it is difficult to secure delayed fracture resistance by grain refinement.
Parts with improved delayed fracture resistance made from wires according to the present disclosure contain 0.3% to 2% residual austenitic structure and residual tempered martensitic structure by volume fraction. If the fraction of the retained austenite structure is less than 0.3%, it is difficult to expect the effect as a barrier to delay hydrogen diffusion. Further, if it exceeds 2%, residual austenite is thickly formed not only in the lath boundaries but also in the austenite grain boundaries and the like, which makes it difficult to retard the diffusion of hydrogen and reduce the effect of improving the delayed fracture resistance.
Next, a method for manufacturing wires and parts having improved delayed fracture resistance according to an exemplary embodiment of the present disclosure will be described.
Wires and components having improved delayed fracture resistance according to the present disclosure may be manufactured by various methods without particular limitation. As an exemplary embodiment, it may be manufactured by the following method.
Wires with improved delayed fracture resistance according to the present disclosure may be manufactured by a process comprising: a step of finish rolling a steel material at 880 ℃ to 980 ℃, the steel material comprising in weight-%: 0.15% to 0.30% of C, 0.15% to 0.25% of Si, 0.95% to 1.35% of Mn, 0.030% or less of P, 0.030% or less of S, 0.015% to 0.030% of Ti, 0.0010% to 0.0040% of B, 0.0010% to 0.0080% of N, and Fe and unavoidable impurities as the balance; and a step of winding at 830 to 930 ℃.
First, a steel material satisfying the above alloy composition is prepared and finish-rolled into a wire rod at 880 to 980 ℃. The rolled wire is then wound into a coil shape at 830 ℃ to 930 ℃.
If the wire rod rolling temperature is lower than 880 c or if the winding temperature is lower than 830 c, it is possible to form the decarburized ferrite layer on the surface by transformation because the surface layer is quasi two phases, and the delayed fracture resistance may become poor because the decarburized ferrite layer is also formed on the surface of the bolt during the heat treatment. Further, since the prior austenite grain size of the bolt decreases and the fraction of the residual austenite increases, the delayed fracture resistance may become poor. If the wire rod finish rolling temperature exceeds 980 ℃ or if the winding temperature exceeds 930 ℃, a decarburized ferrite layer may be formed on the surface as decarburization is accelerated by hydrogen, and the delayed fracture resistance may become poor due to an increase in the prior austenite grain size.
Then, the wound wire may be drawn, spheroidized, heat treated, coated, and formed into a bolt; austenitizing, quenching, and then tempering to obtain the final bolt component. For example, it can be prepared by the following method.
A method for manufacturing a component for a bolt according to an exemplary embodiment of the present disclosure includes: a step of drawing a wire rod manufactured according to the present disclosure; a step of spheroidizing heat treatment of the drawn wire rod at 745 to 770 ℃; a step of heating the spheroidized heat-treated drawn wire rod at 870 ℃ to 940 ℃; a step of quenching the spheroidized heat-treated drawn wire rod at 50 to 80 ℃; and a step of tempering at 400 to 600 ℃.
The spheroidizing heat treatment may be performed at 745 to 770 ℃. If the heat treatment temperature is lower than 745 deg.c or exceeds 770 deg.c, the degree of spheroidization may be reduced, which may result in an increase in hardness, poor formability of the threaded portion of the bolt after forming, and cracking of the threaded portion.
The austenitizing heat treatment may be performed at 870 ℃ to 940 ℃. If the heat treatment temperature is below 870 ℃, toughness may become poor due to insufficient reverse austenite transformation, with non-uniform formation of a martensitic structure after quenching. If the heat treatment temperature exceeds 940 ℃, the delayed fracture resistance may become poor due to an increase in the prior austenite grain size.
Quenching may be performed at 50 to 80 ℃. If the quenching temperature is lower than 50 ℃, fine quench cracks may occur in the threads of the bolt due to thermal deformation, which may cause delayed fracture. If it exceeds 80 ℃, in addition to mechanically stable retained austenite formed in the lath due to insufficient quenching, retained austenite may be formed in the prior austenite grain boundaries, and delayed fracture may be caused due to accumulation of hydrogen.
Tempering may be performed at 400 to 600 c in order to provide strength and toughness depending on the use and purpose of the final product. If the tempering temperature is lower than 400 ℃, brittleness may be caused due to tempering. Further, if it exceeds 600 ℃, it is difficult to achieve the strength desired by the present disclosure.
Parts with improved delayed fracture resistance made in accordance with the present disclosure include 0.3% to 2% residual austenitic structure and residual tempered martensitic structure by volume fraction.
Hereinafter, the present disclosure is described in more detail by way of examples. However, the following examples merely illustrate the present disclosure, and the present disclosure is not limited by the examples. The scope of the present disclosure is to be determined by the appended claims and the reasonable inferences to be drawn therefrom.
Examples
The wires of examples 1 to 9 and comparative examples 1 to 7 satisfying the alloy compositions of table 1 were prepared into final bolts under the manufacturing conditions of the present disclosure for performing the test. Specifically, the steel member satisfying the alloy composition of table 1 was finish rolled into a wire rod at 880 to 980 ℃ and wound into a coil shape at 830 to 930 ℃. The wound wire is subjected to spheroidizing heat treatment at 745 to 770 ℃. Then, the spheroidized heat-treated wire rod is formed into a bolt, austenitized at 870 ℃ to 940 ℃, quenched at 50 ℃ to 80 ℃, and then tempered at 400 ℃ to 600 ℃ to ensure a tensile strength of 1050±16 MPa.
TABLE 1
After evaluating the maximum size of TiN precipitates and delayed fracture cracks of the bolts of examples 1 to 6 and comparative examples 1 to 5, the values of formulas 1 and 2 were determined. The results are given in table 2. Maximum size of TiN precipitates by 160mm at 30 fields by analysis via extremum 2 Is defined by the largest dimension of the inclusions observed in the L-section (longitudinal direction) of the bolt. The results are also given in table 2.
The delayed fracture resistance was tested by tightening the bolts with a clamping force corresponding to the yield strength according to the delayed fracture simulation method and immersing them in a solution of 5% hydrochloric acid +95% distilled water for 10 minutes. The presence of cracks in the thread (which is the part where stress is concentrated) was observed. X represents no cracking, and O represents cracking.
TABLE 2
As seen from table 2, for examples 1 to 6 in which the requirements of the alloy composition, the formula, and the TiN size set forth in the present disclosure were satisfied, no delayed fracture crack was shown in the threaded portion of the bolt before and after evaluating the delayed fracture resistance. In contrast, for comparative example 1 in which the [ Ti ]/3.42[ N ] value exceeded the upper limit of 2.0 set forth in the present disclosure as 2.506, coarse TiN was formed and delayed fracture cracking occurred.
For comparative example 2 in which the [ Ti ]/3.42[ N ] value exceeded the upper limit of 2.0 set forth in the present disclosure as 3.070, coarse TiN was formed and delayed fracture cracking occurred.
For comparative example 3 in which the Si content exceeded the upper limit of 0.25% and the 5.5× [ Si ] + [ Mn ] value exceeded the upper limit of 2.4 as 0.26%, delayed fracture cracking occurred because the formability of the threaded portion of the bolt after spheroidizing heat treatment was poor due to excessive solid solution strengthening. Fig. 1 shows an image of the threaded portion of comparative example 3 before evaluating the delayed fracture resistance. As can be seen from fig. 1, for comparative example 3, which did not meet the requirements set forth in the present disclosure, delayed fracture cracking occurred, indicating that delayed fracture resistance was not achieved.
For comparative example 4 in which the Mn content exceeded the upper limit of 1.35% as 1.45% and the 5.5× [ Si ] + [ Mn ] value exceeded the upper limit of 2.4 as 2.61 as the disclosure, delayed fracture cracking occurred because formability of the threaded portion of the bolt after spheroidizing heat treatment was poor due to excessive solid solution strengthening.
For comparative example 5 in which the C content exceeds the upper limit of 0.30% set forth in the present disclosure as 0.33%, the formation of the retained austenite structure with excellent mechanical stability was prevented and delayed fracture cracking occurred.
Further, final bolt samples of example 3 and comparative examples 6-1 to 6-6 were prepared under the conditions described in table 3.
TABLE 3
For example 3 in which the finish rolling temperature, the winding temperature, the spheroidizing heat treatment temperature, and the austenitizing temperature were satisfied, no delayed fracture crack occurred. In contrast, for comparative example 6-1 in which the rolling temperature exceeded the upper limit 980 ℃ set forth in the present disclosure as 990 ℃ and the winding temperature exceeded the upper limit 930 ℃ set forth in the present disclosure as 940 ℃, delayed fracture cracking occurred due to the increase in the prior austenite grain size in the wire rod and in the bolt.
For comparative example 6-2 in which the rolling temperature did not reach the lower limit 880 ℃ set forth in the present disclosure as 870 ℃ and the winding temperature did not reach the lower limit 830 ℃ set forth in the present disclosure as 820 ℃, the fraction of retained austenite increased and delayed fracture cracking occurred due to the increase in the prior austenite grain size in the wire rod and in the bolt.
For comparative examples 6-3 in which the austenitizing heat treatment temperature exceeded the upper limit of 940 c set forth in the present disclosure as 950 c, delayed fracture cracking occurred due to the increase in the prior austenite grain size of the bolts.
For comparative examples 6 to 4 in which the austenitizing heat treatment temperature was 860 ℃ and did not reach the lower limit 870 ℃ proposed in the present disclosure, undissolved ferrite was formed because QT heat treatment was performed in a state where the bolt was not sufficiently austenitized, and thus, delayed fracture cracks occurred.
For comparative examples 6 to 5 in which the spheroidization temperature was 740 ℃ and did not reach the lower limit 745 ℃ proposed by the present disclosure, and for comparative examples 6 to 6 in which the spheroidization temperature was 775 ℃ and exceeded the upper limit 770 ℃ proposed by the present disclosure, formability was poor and delayed fracture cracking occurred due to low spheroidization degree.
Although exemplary embodiments of the present disclosure have been described, the present disclosure is not limited thereto, and those of ordinary skill in the art will appreciate that various changes and modifications may be made without departing from the concept and scope of the appended claims.

Claims (9)

1. A wire having improved delayed fracture resistance comprising in weight percent: 0.15% to 0.30% of C, 0.15% to 0.25% of Si, 0.95% to 1.35% of Mn, 0.030% or less of P, 0.030% or less of S, 0.015% to 0.030% of Ti, 0.0010% to 0.0040% of B, 0.0010% to 0.0080% of N, and Fe and unavoidable impurities as the balance, and the wire satisfies formula 1:
[ 1]2.0 is less than or equal to 5.5 xSiMn is less than or equal to 2.4
Wherein [ Si ] and [ Mn ] represent the content (wt%) of the corresponding element.
2. The wire rod of claim 1, which satisfies formula 2:
[ 2]1.0 < [ Ti ]/3.42[ N ] < 2.0
Wherein [ Ti ] and [ N ] represent the content (wt%) of the corresponding element.
3. The wire rod of claim 1, wherein the TiN inclusion has a size of 15 μιη or less.
4. A method for manufacturing a wire having improved delayed fracture resistance, comprising:
a step of finish rolling a steel material at 880 ℃ to 980 ℃, the steel material comprising in weight-%: 0.15% to 0.30% of C, 0.15% to 0.25% of Si, 0.95% to 1.35% of Mn, 0.030% or less of P, 0.030% or less of S, 0.015% to 0.030% of Ti, 0.0010% to 0.0040% of B, 0.0010% to 0.0080% of N, and Fe and unavoidable impurities as the balance, and the steel satisfies formula 1; and
a winding step at 830 to 930 ℃,
[ 1]2.0 is less than or equal to 5.5 xSiMn is less than or equal to 2.4
Wherein [ Si ] and [ Mn ] represent the content (wt%) of the corresponding element.
5. The method for manufacturing a wire rod according to claim 4, wherein the steel material satisfies formula 2:
[ 2]1.0 < [ Ti ]/3.42[ N ] < 2.0
Wherein [ Ti ] and [ N ] represent the content (wt%) of the corresponding element.
6. A method for manufacturing a part with improved delayed fracture resistance, comprising:
a step of drawing the wire rod manufactured according to claim 4 or 5;
a step of spheroidizing heat treatment of the drawn wire rod at 745 to 770 ℃;
a step of heating the spheroidized heat-treated drawn wire rod at 870 ℃ to 940 ℃;
a step of quenching the spheroidized heat-treated drawn wire rod at 50 to 80 ℃; and
tempering the quenched component at 400 ℃ to 600 ℃.
7. A part having improved delayed fracture resistance comprising in weight percent: 0.15% to 0.30% of C, 0.15% to 0.25% of Si, 0.95% to 1.35% of Mn, 0.030% or less of P, 0.030% or less of S, 0.015% to 0.030% of Ti, 0.0010% to 0.0040% of B, 0.0010% to 0.0080% of N, and Fe and unavoidable impurities as the balance, and the component satisfies formula 1:
[ 1]2.0 is less than or equal to 5.5 xSiMn is less than or equal to 2.4
Wherein [ Si ] and [ Mn ] represent the content (wt%) of the corresponding element.
8. The component of claim 7, the component satisfying equation 2:
[ 2]1.0 < [ Ti ]/3.42[ N ] < 2.0
Wherein [ Ti ] and [ N ] represent the content (wt%) of the corresponding element.
9. The component of claim 7, comprising 0.3% to 2% residual austenitic structure and residual tempered martensitic structure by volume fraction.
CN202180090890.8A 2020-12-18 2021-12-14 Wire and part with improved delayed fracture resistance and method of making same Pending CN116783316A (en)

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PCT/KR2021/018977 WO2022131752A1 (en) 2020-12-18 2021-12-14 Wire rod and parts with improved delayed fracture resistance, and methods for manufacturing same

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