KR20040054987A - Method of manufacturing high Si added medium carbon wire rod - Google Patents

Abstract

PURPOSE: A manufacturing method of wire rod is provided which reduces surface decarburizing layer generated in high silicon steel and obtains low tensile strength required during cold forming due to high phase fraction of fine graphite powder by forming ferrite decarburizing layer having low carbon solubility in wire rod heating process. CONSTITUTION: The manufacturing method of high silicon added medium carbon wire rod with less surface decarburization comprises a step of obtaining a billet comprising 0.40 to 0.60 wt.% of carbon, 2.0 to 4.0 wt.% of silicon, 0.1 to 0.8 wt.% of manganese, 0.01 wt.% or less of phosphorus, 0.01 wt.% or less of sulfur, 0.004 to 0.013 wt.% of nitrogen, 0.005 wt.% or less of oxygen, 0.001 to 0.003 wt.% of boron, 0.005 to 0.03 wt.% of titanium, one or more elements selected from the group consisting of 0.3 to 2.0 wt.% of nickel, 0.01 to 0.5 wt.% of vanadium and 0.01 to 0.5 wt.% of niobium, and a balance of Fe and other impurities, wherein the titanium (Ti), nitrogen (N) and boron (B) satisfy the following relational expressions: 0.5<=Ti/N<=2.0, 2<=N/B<=8, and 1.0<=(Ti+5B)/N<=3.5; a step of heating the billet to a temperature of Ac1 transformation point at a heating rate of 20±5 deg.C/min, heating the billet at a heating rate of 9±4 deg.C/min in the temperature range of Ac1 transformation point to Ac3 transformation point, heating the billet to a temperature of 1,050±100 deg.C at a heating rate of 15±5 deg.C/min, maintaining the billet in the temperature range of 1,050±100 deg.C for 30 minutes or more, and wire rod rolling the heated billet; a step of coiling the wire rod at a temperature of 860 to 950 deg.C right after wire rod rolling the heated billet; and a step of obtaining a wire rod by air cooling the second cooled wire rod after first cooling the coiled wire rod to a temperature of 770±30 deg.C at a cooling rate of 1.8±0.5 deg.C/sec and second cooling the first cooled wire rod to a temperature of 620±50 deg.C at a cooling rate of 0.4±0.2 deg.C/sec.

Description

Method of manufacturing high silicon-added medium carbon wire rod with low surface decarburization

The present invention relates to a method for manufacturing wire rods processed by bolts and the like, and more particularly, by forming a ferrite decarburization layer having a low carbon utilization in the wire rod heating process to reduce the surface decarburization layer generated in high silicon steel. The present invention relates to a method for producing a wire rod having a low tensile strength required for cold forming due to a high phase fraction of fine graphite grains.

Wire rod is processed to a certain shape and used for various mechanical parts, for example, bolts, nuts, springs and the like. In order to reduce the weight and performance of such mechanical parts, demand for high strength of wire rods continues to increase.

High-strength materials can cause delayed fracture, in which cracks are propagated by hydrogen under constant load. For example, the bolt has a problem of delayed fracture resistance when the tensile strength is 130 kg / mm ² or more, the current use and range is limited.

Accordingly, the advantages expected when developing bolt steel with excellent delay fracture resistance and high strength are as follows. That is, in terms of steel structures, bolt fastening does not require more skilled techniques than welding joints, and considering the advantages of replacing weak welds, not only increases the stability of the steel structure by strengthening the bolt fastening force, but also decreases the number of bolt fastenings. Steel usage can be reduced. In addition, in terms of automobile parts, it contributes to the weight reduction of parts, and thus, there is an advantage in that design diversification and compactness of the vehicle assembly apparatus are possible. Therefore, the higher the strength without lowering the delayed fracture resistance of the material, it is expected that the advantages of use and the ripple effect on the industry will be considerably greater.

The delayed fracture resistance of high strength materials is known to be lowered because precipitates distributed at grain boundaries act as trapped sites of hydrogen and degrade the strength of grain boundaries. Therefore, if the distribution of the Fe-based precipitates to be distributed to the grain boundary after the heat treatment of the wire to the maximum, it is possible to achieve high strength without deterioration of the delayed fracture resistance.

On the other hand, the wire rod is formed into a bolt of various shapes through the cold forming process, if the strength of the material is high, it is necessary to soften heat treatment of the material before cold forming. It is desirable to secure the tensile strength of 60kg / mm ² or less before cold forming through softening heat treatment. This is to suppress the increase of the die wear rate during cold forming to the maximum. Therefore, in order to use high-strength bolted material, it is necessary to secure not only excellent delayed fracture resistance but also cold forming required for bolt manufacturing.

In general, softening heat treatment for securing bolt cold formability is applied to the spheroidizing heat treatment method, characterized in that the microstructure is composed of ferrite + cementite. However, there is a problem in that the strength of the spheroidized material increases as the amount of alloy element added increases, and thus there is a limit in microstructure configuration to overcome this problem.

The present inventors have found that the use of the graphitized structure in the cold forming of high-strength bolts improves cold forming, and has proposed a high silicon-added medium carbon steel. When the steel is properly blended with graphitization promoting elements, the steel can secure wire tensile strength required for bolt cold forming, and has a fine composite structure with no possibility of precipitation of Fe-based precipitates in crystal grains.

However, when manufacturing high-silicon-added heavy carbon steel as a wire rod, silicon has a great influence on reducing the carbon diffusion coefficient while increasing the activity of carbon in the material during reheating in the wire rod rolling process. As a result, carbon activity increases during reheating in the wire rod rolling process, and the decarburization rate increases on the surface. On the contrary, the carbon supply is not smooth to the surface due to the reduction of diffusion coefficient from the center to the surface, thereby enhancing the concentration of carbon in the surface Done. As the depth of the decarburized layer deepens on the surface, the tightening force of the bolt decreases and the fatigue characteristics of the spring deteriorate. As such, it is well known that the surface decarburization of the high silicon-added steel is because the surface decarburization rate is very fast compared to the low silicon-added steel. Thus, when manufacturing a high silicon-added steel wire, many problems in use may occur when the decarburization control of the material is not appropriate.

Representative examples of decarburization control methods known so far include Korean Patent Nos. 92-24974, 92-24163, 92-24161, Japanese Patent Publications (Pyeong) 2-301514, (Pyeong) 1-31960, (Small) 63 And -216591. Most of these are methods of controlling decarburization by lowering silicon content or adding alloying elements such as lead and tin. However, it is not possible to lower the silicon content in the steel species in which high silicon addition is inevitable for high strength, and when alloying elements such as lead and tin are added, mechanical properties such as impact toughness are deteriorated.

The present invention was devised in the research process to prevent the generation of decarburized layer on the surface of the high-silicon-added medium carbon steel wire heating process, by forming a ferrite decarburized layer with low carbon utilization on the surface of the billet in the wire heating process The present invention provides a method for producing a wire rod having a low tensile strength, which is required for cold forming, while reducing the thickness of the surface decarburized layer by significantly reducing the decarburization rate on the surface thereof.

Wire rod manufacturing method of the present invention for achieving the above object, by weight%, carbon 0.40-0.60%, silicon 2.0-4.0%, manganese 0.1-0.8%, phosphorus 0.01% or less, sulfur 0.01% or less, nitrogen 0.004-0.013 %, Oxygen 0.005% or less, boron 0.001-0.003%, titanium 0.005 to 0.03%, which is selected from the group consisting of nickel 0.3-2.0%, vanadium: 0.01 to 0.5%, niobium: 0.01 to 0.5% Or 2 or more, remaining Fe and other impurities, wherein the titanium (Ti), nitrogen (N), boron (B) is the following relationship, 0.5≤Ti / N≤2.0, 2≤N / B≤8, Obtaining a billet satisfying 1.0 ≦ (Ti + 5B) /N≦3.5,

The billet is heated at a heating rate of 20 ± 5 ° C./min to Ac ₁ transformation point, and is heated at a heating rate of 9 ± 4 ° C./min from Ac ₁ transformation point to Ac ₃ transformation point, which is an ideal range, and 15 to 1050 ± 100 ° C. Heating at a heating rate of ± 5 ° C / min, holding at least 30 minutes in this temperature range, and then rolling the wire;

Winding at 860-950 ° C. immediately after the wire has been rolled, and

Firstly cooling the wound wire up to 770 ± 30 ℃ at a cooling rate of 1.8 ± 0.5 ℃ / sec, and second cooling at 620 ± 50 ℃ at a cooling rate of 0.4 ± 0.2 ℃ / sec and then air cooled It is composed.

Hereinafter, the present invention will be described in detail.

The present inventors temperature between the high (during heating of a wire heating) various angles as a result, billet during heating for the pre-existing rolling research for decarburization reduces the carbon steel wire rod of silicon was added over station (Ac ₃ transformation point in the Ac ₁ transformation point, ferrite and It was found that controlling the heating rate through the austenite phase in which the austenite phase is in equilibrium) can induce a ferrite layer with very low carbon solubility on the surface of the billet. Since the surface ferrite layer has a very low carbon solubility, the diffusion of carbon atoms through this region immediately after precipitation of the ferrite phase becomes very slow, thereby significantly reducing the rate of decarburization reaction and the surface ferrite layer depending on the degree of oxidation. Since the thickness can be adjusted, the billet surface decarburized layer can be remarkably improved during extraction by wire heating.

Furthermore, by controlling the ratio of Ti, N, and B in the alloying elements of the billet, it has been found that wire rods having a low tensile strength can be manufactured by increasing the phase ratio of fine graphite grains during wire fabrication, thereby completing the present invention. The component system of the high silicon-added medium carbon steel which becomes the object of this invention is demonstrated first, and then the method of manufacturing this steel by wire is demonstrated.

[River composition]

Content of carbon (C): 0.40-0.60%

If the carbon content is less than 0.40%, it is difficult to obtain an adequate amount of retained austenite, shape and size in the composite structure after heat treatment for producing a fine composite steel which is effective for delayed fracture resistance. In addition, it is difficult to secure mechanical and thermal stability and sufficient tensile strength and yield strength as high strength bolt steels.

On the other hand, when the carbon content is more than 0.60%, characteristics such as cross-sectional reduction rate, elongation and impact toughness after heat treatment are lowered. In addition, segregation and surface flaws occur in wire rod manufacturing, surface decarburization intensifies during charging, and permanent deformation and static fatigue characteristics deteriorate when bolts are tightened. In addition, the proper shape and size of the microcomposite, and the time required for transformation to secure the complex structure, adversely affect the carbon concentration and interfacial concentration of residual austenite.

Silicon (Si) content: 2.0-4.0%

If the silicon is less than 2.0%, the graphitization heat treatment time for improving the cold forming property is long. In addition, the mechanical and thermal stability of the retained austenite after the ferrite transformation is lowered, it is difficult to secure the ferrite + residual austenite composite structure and the appropriate amount of retained austenite. In addition, it is difficult to secure the strength due to the insufficient solid-solution strengthening effect of ferrite, it affects the delayed fracture resistance, surface corrosion characteristics, impact toughness, permanent deformation during bolting. In addition, it is difficult to secure uniformity and proper thickness of the surface ferrite decarburization layer in the wire heating furnace for controlling the decarburization of wires, which leads to intensification of decarburization, and it is difficult to control the surface scale characteristics by increasing the hardenability during wire cooling.

On the other hand, when the silicon exceeds 4.0%, the above-mentioned effect is saturated and adversely affects the hardenability, the composition of the composite tissue steel, the impact toughness, the fatigue characteristics. In addition, in the production of blooms or billets for wire rods, the microstructures become inhomogeneous due to silicon segregation, thereby degrading quality characteristics in the final product. In addition, it is difficult to control the homogeneous surface decarburization as the thickness of the surface ferrite layer increases during the heat treatment.

More preferred silicone component range in the present invention is 2.8-3.3%. This is because isothermal heat treatment time for producing bainite structure (ferrite + residual austenite), residual austenite fraction, size and shape, and high strength and high toughness of ferrite + residual austenite composite, as well as delayed fracture resistance (diffusion) This is because the amount of hydrogen, precipitation control of grain boundary precipitates, surface decarburization, stress relaxation after bolting or permanent deformation resistance, dynamic and static fatigue characteristics can be improved effectively.

Manganese (Mn) content: 0.1-0.8%

Manganese is an element that forms a solid solution to form a solid to form a solid solution, so that it is very useful for high-strength bolt characteristics. When the content of manganese exceeds 0.8%, local quenchability due to manganese segregation is increased rather than solid solution strengthening effect. Crazy

Content of phosphorus (P) and sulfur (S): 0.01% or less each

Phosphorus segregates at grain boundaries and degrades toughness, so the upper limit is preferably limited to 0.01%.

Sulfur is a low-melting point element segregated to lower the toughness and form an emulsion, which has a detrimental effect on the delayed fracture resistance and stress relaxation characteristics, it is preferable to limit the upper limit to 0.01%.

Nitrogen (N) content: 0.004-0.013%

If the nitrogen content is less than 0.004%, it becomes difficult to form vanadium and niobium-based nitrides that act as non-diffuse hydrogen trap sites. In addition, when the content of nitrogen exceeds 0.013%, the effect is saturated.

Oxygen (O) content: 0.0050% or less

When the oxygen content exceeds 0.0050%, coarse oxide-based nonmetallic inclusions are easily formed, and fatigue life is reduced.

Boron (B, B) content: 0.001 to 0.003%

Boron is a graphitization promoting element and grain boundary strengthening element for improving quenching and delayed fracture resistance. If the boron content is less than 0.0010%, the effect of improving grain boundary strength due to grain boundary strengthening of boron atoms due to grain boundary segregation during heat treatment is insufficient, and the graphitization promoting effect is not sufficient during graphitization treatment for improving cold forming. If the boron content exceeds 0.003%, the effect is saturated, but the precipitation of boron nitride at the grain boundary causes a decrease in grain boundary strength.

Titanium (Ti) content: 0.005-0.03%

If the content of titanium is less than 0.005%, the effect of promoting graphitization and miniaturization of austenite grains is insufficient, and the precipitation distribution of titanium-based carbon and nitride in the grain boundary effective for delayed fracture resistance is insufficient. On the other hand, when the loading of titanium exceeds 0.03%, the addition effect is saturated, and coarse titanium-based carbon and nitride are formed to affect mechanical properties.

The content of Ti, N, and B in the alloy element preferably satisfies the following relationship: 0.5 ≦ Ti / N ≦ 2.0, 2.0 ≦ N / B ≦ 8.0, 1.0 ≦ (Ti + 5B) /N≦3.5. This relationship is to achieve the tensile strength (typically 60kg / mm ² or less) required for cold forming before bolt forming by graphitization during wire rod manufacture and to significantly shorten the time required for wire rod graphitization. Is as follows.

Ti / N: 0.5 to 2.0

If the Ti / N ratio is less than 0.5, the graphene grain formation rate is lowered due to a decrease in nucleation sites of the graphite grains, and when the Ti / N ratio is greater than 2.0, the graphite grains tend to coarsen.

N / B ratio: 2.0-8.0

If the N / B ratio is less than 2.0, the number of BN precipitates required for graphite nucleation is insufficient, and the graphitization rate is lowered. The amount of nitrogen increases and the graphitization rate is rather lowered.

(Ti + 5B) / N ratio: 1.0 to 3.5

If the (Ti + 5B) / N ratio is less than 1.0, the number of TiN and BN precipitates that contribute to the nucleation of graphite grains is insufficient, and if the (Ti + 5B) / N ratio is greater than 3.5, the TiN and The number of BN precipitates is saturated and the amount of nitrogen dissolved in the base metal increases, which adversely affects the graphitization rate.

In addition to the above components, one or two or more of nickel, vanadium and niobium are added.

Nickel (Ni) content: 0.3 ~ 2.0%

Nickel is an element for promoting graphitization, and is an element that forms a nickel enrichment layer on the surface during heat treatment to suppress permeation of external hydrogen and thereby improve delayed fracture resistance. If the nickel content is less than 0.3%, it is difficult to expect the effect of improving the delayed fracture resistance due to the incomplete formation of the surface thickening layer. In addition, the heat treatment time is increased during spheroidization or graphitization treatment to improve decarburization control, toughness and cold formability, and there is no effect of improving cold formability during bolt forming. On the other hand, when the nickel content exceeds 2.0%, the effect is saturated and may adversely affect the proper amount, size and shape of the residual austenite amount.

Vanadium (V) or niobium (Nb) content: 0.01 to 0.5% each

Vanadium and niobium are elements that improve delayed fracture resistance and stress relaxation. If the content of these compounds is less than 0.01%, the distribution of vanadium or niobium-based precipitates in the base material decreases, and thus the role of the non-diffusible hydrogen trap site is insufficient. Therefore, it is difficult to expect the effect of improving the delayed fracture resistance. In addition, it is difficult to expect the precipitation strengthening effect is not enough improvement effect on the stress relaxation resistance, and it is difficult to expect the austenite grain refining, which affects the microstructure of the structure of the ferrite + residual austenite composite structure. On the other hand, if the content of vanadium and niobium exceeds 0.5%, the graphitization heat treatment time is long, and the improvement effect on the delayed fracture resistance and the stress relaxation resistance by the V or Nb-based precipitates is saturated. During the heat treatment of austenite, the amount of coarse alloy carbide that is not dissolved in the base metal increases, which acts like a non-metallic inclusion, leading to a decrease in fatigue characteristics.

[Manufacturing method of wire rod]

In the steel formed as described above, a decarburized layer is formed on the surface in the heating process for manufacturing the wire rod. In order to prevent this, the ferrite decarburized layer is formed on the surface of the billet by appropriately adjusting the heating conditions in the heating process for rolling the billet. It is important to form.

[Wire Rolling Process]

First, according to the present invention, the billet is heated to a heating rate of 20 ± 5 ° C./min to Ac ₁ transformation point. If the heating rate is less than 15 ℃, the charging time is too long for the entire heating, if the temperature exceeds 25 ℃ occurs the temperature difference between the surface and the inside of the billet, Ac ₁ transformation time is different, this time billet can be bent.

Subsequently, heating is performed at a heating rate of 9 ± 4 ° C./minute from the Ac ₁ transformation point in the ideal range to the Ac ₃ transformation point. At this time, if the heating rate is less than 5 ℃ / min, the thickness of the abnormal region ferrite layer on the billet surface of the wire heating furnace is increased more than necessary (≥0.5mm), so it is difficult to obtain the decarburization improvement effect. Charge time) is long. In addition, when the temperature increase rate exceeds 13 ° C./min, it is difficult to form an appropriate surface ferrite layer (0.05-0.5 mm) of billet necessary for suppressing decarburization reaction. More preferably, it heats at the heating rate of 9 +/- 2 degree-C / min. In the present invention, the depth of the ferrite decarburized layer on the surface of the billet is preferably 0.05 to 0.5 mm.

Subsequently, heating is performed at a heating rate of 15 ± 5 ° C / min to 1050 ± 100 ° C, held for at least 30 minutes at this temperature range, and then rolled. After the abnormal zone end temperature, the microstructure of the base material is transformed from the abnormal structure (ferrite + austenite) to the austenite single phase structure. At this time, when the temperature increase rate up to the heating maintenance temperature is less than 10 ℃ / min, This is because when the temperature exceeds 20 ° C / min, the billet warpage is likely to occur due to the deep temperature deviation inside and outside the billet, and it is difficult to control the inside and outside temperature of the billet suitable for hot rolling. More preferably, it is heated at a heating rate of 15 ± 2 ℃ / min.

In addition, when the heating temperature of the wire heating furnace at this time is less than 950 ℃, there is a problem in controlling the proper thickness of the billet surface ferrite layer for decarburization control. In addition, it is not easy to re-use coarse precipitated vanadium-based or niobium-based precipitates during the production of billet, and the workability is poor due to the overload during rolling due to the increase in hot deformation resistance. In addition, when the heating holding temperature exceeds 1150 ° C., a uniform ferrite layer for decarburization control cannot be deposited on the surface. In other words, in order to precipitate the surface ferrite layer having a very low carbon solubility, the surface ferrite layer must remain at the heating and maintaining temperature, but when the heating temperature is higher than 1150 ℃, the surface ferrite layer is transformed into austenite. This is because the decarburization rate increases sharply, which causes deep surface decarburization. If the heating holding time is less than 30 minutes, it is difficult to secure a uniform temperature distribution inside the billet for wire rod rolling.

As described above, the billet is heated to roll the wire. At this time, the decarburization area of the billet extracted from the heating furnace is the same as the decarburization area of the wire after the wire rolling. Therefore, in the present invention, in consideration of this point, it is preferable to roll the wire to a diameter of 30 mm or less.

[Winding process]

Next, the wire is wound, but the winding temperature at this time is preferably 860-950 ℃. The coiling temperature is secured by cooling water spraying immediately after the wire is rolled. This is because the time required for graphitization is insufficient when the coiling temperature is less than 860 ° C., and when the coiling temperature is higher than 950 ° C., the winding failure due to high temperature winding is likely to occur.

[Cooling process]

The wound wire is cooled, first cooling to 1.8 ± 0.5 ° C / sec to 770 ± 30 ° C. The conditions of cooling temperature and cooling rate are taken into account the difference in the degree of cooling of wire rod integrated state, ie overlapping part and non-overlapping part. That is, if the cooling temperature is higher than 800 ℃ at a cooling rate of 1.8 ± 0.5 ℃ / sec, the appropriate graphitization transformation time in the cooling zone is insufficient, the graphitization rate is reduced. In addition, when the cooling temperature is less than 740 ° C, ferrite or perlite transformation is more likely to occur than graphitized tissue.

Next, secondary cooling is performed at 0.4 ± 0.2 ° C / sec to 620 ± 50 ° C.

If the cooling temperature is higher than 670 ℃ the wire graphitization proceeds, if less than 570 ℃ graphitization rate decreases. In addition, if the cooling rate is faster than 0.6 ℃ / sec, the graphitization rate can be reduced, and if it is slower than 0.2 ℃ / sec, it is difficult to secure the proper cooling temperature range of 620 ± 50 ℃ due to the limitation of the cooling equipment, the wire graphitization rate This decreases.

Next, air-cooling, because the transformation is complete and the cooling rate does not affect the tissue.

In the wire rod obtained by cooling, the graphite grain size is 50 µm or less, and the phase fraction is preferably 0.1% or more. If the graphite grain size is greater than 50 μm, surface defects may be caused rather than the effect of improving cold forming property. In addition, when the graphite grain phase fraction is less than 0.1%, there is no tissue softening effect for improving cold forming.

Hereinafter, the present invention will be described in more detail with reference to Examples.

EXAMPLE

The high-silicon-added medium-carbon steel billet having the component system of the invention steel and the comparative steel as shown in Table 1 below was manufactured by 16 mm wire rod by high-speed wire rolling under the conditions of Table 3, and wound and cooled and air-cooled under the conditions of Table 4 immediately after the wire rolling. . The surface decarburized layer depth of the wire rod manufactured as described above was measured by the KS standard (KD D 0216). According to this standard, an optical microscope observation method and a microhardness measuring method are proposed, but the microhardness measuring method was used here. The measured decarburized bed depths are shown in Table 3. The measurement position was measured at 8 equal sections of wire rod section, and the measured value was based on the average value. In addition, the graphitized tissue phase fraction of the wire rods prepared was measured using an image analyzer, wherein the test surface was based on 300 mm ² .

division Chemical composition C Si Mn P S N O Ni B V Nb Ti Invention steel One 0.44 3.03 0.63 0.008 0.009 0.0090 0.0013 - 0.0015 0.05 0.02 0.016 2 0.43 3.24 0.62 0.008 0.008 0.0115 0.0014 - 0.0024 0.06 0.01 0.023 3 0.56 3.23 0.68 0.009 0.009 0.0048 0.0015 0.70 0.0018 0.07 - 0.007 4 0.42 2.35 0.68 0.007 0.009 0.0091 0.0016 - 0.0039 0.01 0.05 0.011 5 0.45 3.99 0.75 0.008 0.009 0.0086 0.0017 - 0.0025 0.04 0.03 0.017 6 0.58 3.12 0.74 0.009 0.008 0.0071 0.0017 - 0.0013 - - 0.013 7 0.57 2.37 0.80 0.009 0.008 0.0056 0.0018 1.10 0.0020 - 0.05 0.009 Comparative steel One 0.44 3.03 0.63 0.008 0.009 0.0090 0.0013 - - 0.05 0.02 - 2 0.44 3.03 0.63 0.008 0.009 0.0090 0.0013 - 0.0005 0.05 0.02 0.026 3 0.44 3.03 0.63 0.008 0.009 0.0090 0.0013 - 0.0055 0.05 0.02 0.007

division Ti / N N / B (Ti + 5B) / N Inventive Steel 1 1.8 6.0 2.6 Inventive Steel 2 2.0 4.8 3.0 Invention Steel 3 1.5 2.7 3.3 Inventive Steel 4 1.2 2.3 3.4 Inventive Steel 5 2.0 3.4 3.4 Inventive Steel 6 1.8 5.5 2.7 Inventive Steel 7 1.6 2.8 3.4 Comparative Steel 1 - - - Comparative Steel 2 2.8 18.0 3.2 Comparative Steel 3 0.8 2.0 3.8

Steel grade used Temperature increase rate up to Ac ₁ (℃ / min) Temperature rise rate to the ideal station (℃ / min) Heating rate up to the heating maintenance temperature after Ac ₃ (℃ / min) Heating holding temperature (℃) Heating holding time (min) Wire stripping depth of 16mm in diameter (mm) Invention 1 Inventive Steel 1 20 5 15 1000 40 0.03 Invention 2 Inventive Steel 1 20 9 15 1050 30 0.04 Invention 3 Inventive Steel 1 20 13 15 1050 30 0.03 Invention 4 Inventive Steel 1 20 9 10 1050 30 0.02 Invention 5 Inventive Steel 1 20 9 15 1050 30 0.04 Invention 6 Inventive Steel 1 20 9 20 1050 30 0.05 Invention 7 Inventive Steel 1 20 9 15 950 40 0.03 Invention Material 8 Inventive Steel 1 20 9 15 1100 30 0.02 Invention 9 Inventive Steel 2 20 9 15 1050 30 0.04 Invention 10 Invention Steel 3 20 9 15 1050 30 0.02 Invention 11 Inventive Steel 4 20 9 15 1050 30 0.03 Invention Material12 Inventive Steel 5 20 9 15 1050 30 0.04 Invention Material 13 Inventive Steel 6 20 9 15 1050 30 0.03 Invention 14 Inventive Steel 7 20 9 15 1050 30 0.05 Comparative Material 1 Inventive Steel 1 20 3 15 1050 30 0.13 Comparative Material 2 Inventive Steel 1 20 17 15 1050 30 0.12 Comparative Material 3 Inventive Steel 1 20 17 5 1050 30 0.16 Comparative Material 4 Inventive Steel 1 20 17 25 1050 30 0.19 Comparative Material 5 Inventive Steel 1 20 17 25 1200 30 0.20 Comparative Material 6 Comparative Steel 1 20 13 15 1050 30 0.03 Comparative Material7 Comparative Steel 2 20 13 15 1050 30 0.03 Comparative Material 8 Comparative Steel 3 20 13 15 1050 30 0.03

Steel grade Winding temperature (℃) Primary cooling temperature (℃) Cooling rate up to the first cooling temperature (℃ / sec) Secondary cooling temperature (℃) Cooling rate up to secondary cooling temperature (℃ / sec) Graphite Grain Percentage (%) Graphite Grain Average Size (㎛) Wire tensile strength (kg / mm ² ) (based on 16mm) Inventive Steel 1 900 740 1.3 670 0.6 1.8 13 54 Inventive Steel 2 860 740 1.8 670 0.4 2.0 8 55 Invention Steel 3 900 740 1.8 640 0.6 1.7 15 53 Inventive Steel 4 950 800 1.8 660 0.6 2.2 8 51 Inventive Steel 5 900 740 1.8 670 0.6 1.8 10 58 Inventive Steel 6 900 800 2.3 570 0.4 2.0 12 54 Inventive Steel 7 950 770 1.8 660 0.6 1.9 11 56 Comparative Steel 1 900 740 1.3 670 0.6 0 0 101 Comparative Steel 2 900 740 1.3 670 0.6 0.3 12 92 Comparative Steel 3 900 740 1.3 670 0.6 0.4 11 91

As shown in Table 4, the comparative steels (1 to 3) exhibited graphite granular phase ratios of 0 to 0.4% and wire tensile strengths of 91 to 101 kg / mm ² , while inventive steels (1-7) were wire-cooled. After showing the graphite grain phase ratio of 1.7 ~ 2.2% and graphite grain size distribution of 8 ~ 15㎛, the wire tensile strength was 51 ~ 58kg / mm ² . . Therefore, it can be seen that the cooling control of the wire using the material composed of the alloy component system of the present invention is very effective in significantly lowering the tensile strength required for forming the bolt cold forming.

As described above, the present invention suppresses the decarburization reaction during heating by providing a uniform and thin ferrite layer on the surface of the billet by reducing the temperature rise rate of the reverse pass when heating the wire rod, thereby providing a graphitized wire rod during wire cooling. Can be.

Claims (5)

By weight%, carbon 0.40-0.60%, silicon 2.0-4.0%, manganese 0.1-0.8%, phosphorus 0.01% or less, sulfur 0.01% or less, nitrogen 0.004-0.013%, oxygen 0.005% or less, boron 0.001-0.003%, titanium 0.005% to 0.03%, and it is composed of one or two or more selected from the group consisting of 0.3-2.0% nickel, 0.01% to 0.5% vanadium, 0.01% to 0.5% niobium, remaining Fe and other impurities, The titanium (Ti), nitrogen (N), boron (B) satisfies the following relationship, 0.5≤Ti / N≤2.0, 2≤N / B≤8, 1.0≤ (Ti + 5B) /N≤3.5 Getting a billet, The billet is heated at a heating rate of 20 ± 5 ° C./min to Ac ₁ transformation point, and is heated at a heating rate of 9 ± 4 ° C./min from Ac ₁ transformation point to Ac ₃ transformation point, which is an ideal range, and 15 to 1050 ± 100 ° C. Heating at a heating rate of ± 5 ° C / min, holding at least 30 minutes in this temperature range, and then rolling the wire; Winding at 860-950 ° C. immediately after the wire has been rolled, and Firstly, the wound wire is cooled to 1.8 ± 0.5 ℃ / sec up to 770 ± 30 ℃ and secondly cooled to 0.4 ± 0.2 ℃ / sec up to 620 ± 50 ℃, followed by air cooling to obtain wire. Method for producing a high silicon-added medium-carbon wire rod with a low surface decarburization comprising a. The method of claim 1, wherein the billet has a heating rate from the Ac ₁ transformation point to the Ac ₃ transformation point in a range of 9 ± 2 ° C./min. The method of claim 1, wherein the heating rate of the billet from the Ac ₃ transformation point to 1050 ± 50 ° C has a range of 15 ± 2 ° C / min. The method for producing a high-silicon-added medium-carbon wire rod having low surface decarburization according to claim 1 or 2, wherein a ferrite decarburization layer having a thickness of 0.05 to 0.5 mm is formed on the surface of the heated billet as described above. The method of claim 1, wherein the wire rod has a graphite grain size of 50 μm or less and a graphite grain phase fraction of 0.1% or more.