KR20000042530A - Method for producing bolt having high strength and high elongation - Google Patents

Abstract

PURPOSE: A method is provided to produce a bolt having high strength and high elongation by improving delayed destruction resistance with high critical delayed destruction strength, suppressing distribution of precipitation and surface decarbonization. CONSTITUTION: A steel billet is composed of 0.40 to 0.60wt% of Carbon, 2.0 to 4.0wt% of Silicon, 0.2 to 0.8wt% of Manganese, 0.2 to 0.8wt% of Chrome, less than 0.01wt% of phosphorus, less than 0.01wt% of sulfur, 0.005 to 0.01wt% of nitrogen, less than 0.005wt% of oxygen and the remain of iron and incidental impurities. Herein, the steel billet contains more than one or two components selected among a group consisting of 0.05 to 0.2wt% of vanadium, 0.005 to.02wt% of niobium, 0.3 to 2.0wt% of nickel, 0.001 to 0.003wt% of boron, 0.01 to 0.5wt% of molybdenum, 0.01 to 0.2wt% of titanium, 0.01 to 0.5wt% of tungsten, 0.01 to 0.2wt% of copper and 0.01 to 0.5wt% of cobalt. The billet is heated and rolled to be a regular shaped bolt. The bolt is heated more than for 20minutes and cooled at 70°C/sec of speed and then cooled by fluid or air to be a bolt having high strength and high elongation.

Description

Manufacturing method of high strength high elongation bolt

The present invention relates to a method for manufacturing bolts used for fastening steel structures and automobile parts, and more particularly, high strength and high elongation bolts having excellent critical delay fracture stress ratio by controlling with a surface ferrite thin film layer and a ferrite + bainite composite structure. It relates to a manufacturing method.

The method of fastening the steel structure members by bolts does not require a skilled technique compared to the welding method through welding, and also increases the stability of the steel structure because there is no fracture phenomenon caused by weak welding. Such bolts are required to have high strength even for stable steel structure member coupling, but for automobile parts and the like, it is necessary to increase the strength of the material even for light weight, multifunction, and high performance. In other words, if the bolt has a high strength, the number of bolted fastenings is reduced, thereby reducing the amount of use thereof, and contributing to the weight reduction of parts in terms of automobile parts, and the weight reduction of such parts diversifies and integrates the design of the automobile assembly apparatus. the possibility of compacting is greater.

Representative technologies for high strength bolts include Japanese Patent Laid-Open Nos. Hei 6-271975, Hei 7-173531, and Japanese Iron and Steel Vol. 82 (1996) No. 4 and the like.

Japanese Patent Laid-Open No. 6-271975 relates to a method for manufacturing a composite tissue steel having excellent resistance to delayed fracture by hydrogen, wherein the weight percentage is 0.05-0.3% C, 0.1-2.5% Si, 0.1-3.0% Mn, 0.05 In steels containing at least one alloy element among -0.1Al, Cu, Ni, Mo, Cr, Nb, V, Ti, and B, the microstructure may be martensite single phase, bainite single phase, or bainite + martensite composite structure. In order to secure the delayed fracture resistance by hydrogen in the microstructure, residual austenite is present in a volume fraction of 1-30%. However, in Japanese Patent Laid-Open No. 6-271975, high elongation was not achieved while securing a critical delay fracture strength of 130 kg / mm ² or more, and a heat treatment process has many disadvantages when manufacturing a composite structure for improving delayed fracture resistance.

Japanese Patent Laid-Open No. 7-173531 has a chemical weight of 0.05-0.3% C, 0.05-2.0% Si, 0.3-5.0% Mn, 1.0-3.0% Cr, 0.01-0.5% Nb, 0.01-0.06% Al The present invention relates to a method for producing bainite + martensite abnormal composite tissue steel by continuously cooling a steel having a composition above a critical cooling rate at which the cornerstone ferrite does not precipitate after hot forming, but having a critical delay fracture strength of 130kg / mm ² or more. Elongation was not achieved, and there are many heat treatment processes in manufacturing composite tissue for improving delayed fracture resistance.

The Japanese Iron and Steel Vol. 82 (1996) No. 4 is based on a conventional tempered martensite structure based on an alloy composition of 0.49% C-0.31% Mn-1.02% Cr-0.68% Mo-0.034% Nb-0.32% V. It consists of -0.009% P-0.004% S critical delayed fracture strength for grain boundary segregation, the reduction of impurities to 130kg / mm ² grade low P, low S, low Mn screen, for preventing the grain boundary precipitation of carbides Ni, and Cr , Mo, V is added, and V, Nb, Ti is added to refine the grains, and the heat treatment is performed at a low tempering temperature. However, in the Japanese Iron and Steel Vol. 82 (1996) No. 4, the problem of delayed fracture resistance is insufficient to use the critical delayed fracture strength of 130kg / mm ² or more and high elongation for improvement of plastic deformation performance is not achieved. I couldn't.

However, as described above, the high-strength bolts used in the prior art are quasi single phase structures in which most of the structures are tempered martensite, and carbide-based precipitates are distributed at the grain boundaries or tempered martensite. And has a bainite complex structure, and sometimes the base material has a characteristic that precipitates are distributed in the ras martensite, resulting in deterioration of delayed fracture resistance due to hydrogen intrusion. Therefore, these materials are currently not available where tensile strength of 130kg / mm ² or more is required or where high strength and high elongation are required at the same time, and thus their use and range are greatly limited.

As such, most of the high-strength bolt materials have a tempered martensite structure in which the precipitates distributed at the grain boundaries act as trapped sites of hydrogen, thereby degrading the strength of the grain boundaries. Therefore, there is a limit to high strength and high elongation at the same time.

Accordingly, the present invention provides a bolt manufacturing method which can obtain a bolt having high strength and high elongation while improving the delayed fracture resistance by increasing the critical delay fracture strength while suppressing the distribution of surface decarburization and grain boundary precipitate to the maximum.

The present invention for achieving the above object is by weight, carbon: 0.40-0.60%, silicon: 2.0-4.0%, manganese: 0.2-0.8%, chromium: 0.2-0.8%, phosphorus: 0.01% or less, sulfur: 0.01 % Or less, nitrogen 0.005-0.01%, oxygen 0.005% or less, including vanadium: 0.05-0.2%, niobium: 0.05-0.2%, nickel: 0.3-2.0%, boron 0.001-0.003%, molybdenum: 0.01 -0.5%, titanium: 0.01-0.2%, tungsten: 0.01-0.5%, copper: 0.01-0.2%, cobalt: 0.01-0.5%, one or more selected from the group consisting of, balance Fe and other The steel billet composed of unavoidable impurities is heated to Ac1 at a heating rate of 15 ± 5 ° C./minute, and the Ac1 to Ac3 in the ideal range is heated at a heating rate of 6 ± 3 ° C./minute, and then 1050 ± 50 ° C. After heating to 10 ± 5 ℃ / min and holding for 30-60 minutes, wire rod is rolled and processed into a bolt of a certain shape to obtain a bolt, and then the obtained bolt is obtained from Ac3- (Ac3-Acl) /1.3 to Ac3-. 20 minutes or more in the range of (Ac3-Acl) /5.5 Heat, quench to Ms + 50 ° C.-Ms + 110 ° C. at a cooling rate of 70 ° C./sec or more, maintain at least 20 minutes at this temperature, and then cool or air-cool the high-strength high elongation bolt. It is about.

Hereinafter, the present invention will be described in detail.

The bolt manufacturing according to the present invention is heated and then rolled the wire while changing the heating rate to suppress the surface decarburization to the steel billet, and the rolled wire is processed into a bolt of a predetermined shape, the processed through the control of heating and cooling conditions By forming a composite structure of ferrite and bainite in the structure of the bolt, the final bolt structure is formed in the surface layer with a decarburized layer in comparison with the existing layer, and the base material has a ferrite + bainite composite structure. .

Next, the reason for limitation of the composition component and the component range of the present invention will be described.

When the carbon (C) is 0.40% or less, it is difficult to secure sufficient tensile strength and elongation as high-strength and high elongation bolt steel after isothermal heat treatment for bainite production, and difficult to secure elongation after heat treatment at 0.60% or more. It is difficult to secure effective bainite structure for high elongation due to too high carbon concentration of austenite due to ferrite formation during abnormal reverse heat treatment, and permanent deformation, fatigue property, denitrification distribution, bainite structure of decarburization and fut fastening in the furnace This is not preferable because it affects bainite transformation time.

The content of silicon (Si) is limited to 2.0-4.0%. Less than 2.0% of the ferrite strengthening effect in the bainite structure is insufficient, making it difficult to secure the strength, and also affect the delayed fracture resistance, elongation, surface corrosion characteristics, impact toughness, bainite structure, and permanent deformation during bolting. This is because it is difficult to properly distribute the surface ferrite decarburization layer in the wire heating furnace for controlling the wire decarburization, and thus, decarburization is intensified, and it is difficult to control the surface scale characteristics by increasing the hardenability during wire rod cooling. Above 4.0%, it is undesirable because the above-mentioned effect saturates and also affects the bainite structure, impact toughness, corrosiveness, and fatigue characteristics in the composite tissue, and is not suitable for the production of bloom or billet for wire rod manufacturing. This is because the quality characteristics in the final product are deteriorated due to the inhomogeneity of the microstructure due to the segregation of silicon at the same time, and it is difficult to control the homogeneous surface decarburization because the thickness of the surface ferrites increases during the heat treatment.

The more preferable component range of this silicon is 2.8-3.3%, isothermal heat treatment time and residual austenite fraction for producing bainite structure, high strength and high elongation of bainite, delayed fracture resistance (diffusion, hydrogen content, grain boundary) Precipitation control), surface decarburization, stress relaxation after bolting or permanent deformation resistance, dynamic and static fatigue characteristics can be improved very effectively.

The reason for the content of manganese (Mn) is 0.2-0.8% is an element that forms a solid solution to form a solid solution to strengthen the solid solution, which is very useful for high-strength bolt characteristics. Tissue heterogeneity caused by simple stones has a more detrimental effect on bolt properties. Macro segregation and micro segregation are more likely to occur depending on the segregation mechanism during steel coagulation. Manganese segregation promotes segregation due to its relatively low diffusion coefficient compared to other elements, and the improvement of hardenability is caused by core martensite. ) Is the main reason for generating When manganese is added below 0.2%, segregation zones are hardly formed by manganese segregation, but it is difficult to expect the effect of stress relaxation by solid solution strengthening. In other words, if the content of manganese is less than 0.2%, the improvement of hardenability and permanent deformation resistance is insufficient due to insufficient employment strengthening effect, and if more than 0.8%, local anisotropy is increased due to manganese segregation and formation of segregation zones intensifies tissue anisotropy. That is, the bolt properties are degraded due to tissue heterogeneity. Therefore, limiting the content of manganese to 0.2-0.8% is a range in consideration of the strength of the base material, the hardenability during heat treatment, the stress relaxation, and the harmful effects of segregation zone formation.

The content of chromium (Cr) is limited to 0.25-0.8%. The reason is that at 0.25% or less, it is difficult to form a surface ferrite layer for surface decarburization control during heat treatment of high silicon-added steel, and thus, there is almost no decarburization inhibitory effect, and it is difficult to expect an improvement in quenchability. This is because it is not preferable because the transformation time of the knight is long, and it is difficult to generate a surface titration ferrite layer when charging the wire rod for wire rod decarburization layer control, which affects homogeneous decarburization control.

The vanadium (V) or niobium (Nb) is an element that improves delayed fracture resistance and stress relaxation resistance, and its content is limited to 0.05-0.2%. The reason is that less than 0.05%, the distribution of vanadium or niyoboom-based precipitates in the base material is less, so the role of the non-diffusible hydrogen trap site is insufficient, so it is difficult to expect the effect of improving the delayed fracture resistance, and the precipitation strengthening is expected. It is difficult because the improvement effect on the stress relaxation resistance is not sufficient, and it is difficult to expect austenite grain refinement, which affects the structure of bainite. In addition, at 0.2% or more, the improvement effect on the delayed fracture resistance and the stress relaxation resistance by the precipitates is saturated, and the amount of coarse alloy carbides which are not dissolved in the base metal during the austenite heat treatment increases, thus acting as a nonmetallic inclusion. This is because the deterioration of characteristics is caused.

The content of oxygen (O) is limited to 0.0015% or less. This is because coarse oxide-based nonmetallic inclusions are easily formed at 0.0015% or more, thereby reducing the fatigue life.

The content of nitrogen (N) is limited to 0.005-0.03%. This is because the formation of vanadium and niobium-based nitride acting as a non-diffusion water trap site is difficult at 0.005% or less, and the effect is saturated at 0.03% or more.

The content of phosphorus (P) and sulfur (S) is limited to 0.01% or less. The reason is that phosphorus segregates at grain boundaries and lowers its toughness, so the upper limit is limited to 0.01%. Sulfur is segregated with low melting point elements to reduce toughness and form an emulsion, which is harmful to delayed fracture resistance and stress relaxation characteristics. It is desirable to limit the upper limit to 0.01% because

The nickel (Ni) is an element that improves the delayed fracture resistance by inhibiting the permeation of external hydrogen caused by the formation of a nickel enriched layer on the surface during heat treatment. If the content is 0.3-2.0%, it is difficult to expect the effect of improving the delayed fracture stability due to the incomplete formation of the surface thickening layer below 0.3%, and the heat treatment time during the graphitization treatment for decarburization control, toughness, and diversification of the manufacturing process. It is longer, because there is no effect of improving the cold formability during cold bolt processing, and the effect is saturated at 2.0% or more.

The boron (boron, B) is a grain boundary strengthening element for improving the hardenability and delayed fracture resistance in the present invention, the boron content of 0.0010-0.003% is less than 0.0010% when the heat treatment, the grain boundary strengthening according to the grain boundary segregation of boron atoms This is because the effect of improving grain boundary strength is inadequate and the graphitization promoting effect is insufficient during the graphitization treatment for improving cold forming property. At 0.003% or more, the effect is saturated, and the precipitation of boron nitride at the grain boundary This is because it causes a decrease.

The content of the molybdenum (Mo) and tungsten (W) is limited to 0.01-0.5%. The reason is that at 0.01% or less, it is difficult to obtain the effect of improving stress relaxation by inhibiting the growth of cementite when cementite transitions from grain carbide to tempering, and it is stable at high temperature by finely distributing molybdenum precipitates during tempering. This is because the organization is difficult to secure. In addition, at 0.5% or more, the effect is saturated, and the increase in hardenability makes it easy to form low-temperature structure (martensite + bainite) during wire rod manufacturing, and the heat treatment time is increased during the graphitization treatment to improve cold formability. Because of this.

The content of copper (Cu) is limited to 0.01-0.2%. The reason for this is that the improvement effect on the corrosion resistance is insufficient at 0.01% or less, and the improvement effect is saturated at 0.2% or more, and the melting point is lowered at the grain boundary segregation. This is because surface grooves are more likely to occur and impact toughness in the final product is lowered.

The content of titanium is limited to 0.01-0.2%. The reason is that at 0.01% or less, the effect of miniaturizing austenite crystal grains is insufficient, and the precipitation distribution of titanium-based carbon and nitride, which is effective for delayed fracture resistance, is insufficient. Therefore, the improvement effect is difficult to be expected. This is because it forms a coarse titanium nitride and is harmful to fatigue properties.

The content of cobalt is limited to 0.01-0.5%. If the content is less than 0.01%, the softening effect is not sufficient in spheroidization or graphitization heat treatment, which is a material soft-nitrification heat treatment for cold forging, and it has no effect on the grain boundary diffusive hydrogen concentration. It is undesirable because it is saturated and the soft nitriding rate is increased greatly during softening heat treatment, resulting in partial microstructure heterogeneity during heat treatment.

Next, using the steel of the composition component as described above, a method for producing a wire rod in detail, to control the heating conditions to reduce the surface desorption.

First, in the present invention, the steel billet having the above composition is heated to Ac1 at a heating rate of 15 ± 5 ° C./min, and from Ac1 to Ac3 in the ideal range is heated at a heating rate of 6 ± 3 ° C./min, after which 1050 Heat to 10 ± 5 ℃ / min to ± 50 ℃, hold for 30-60 minutes, and roll the wire.

If the heating rate up to Ac1 is less than 10 ℃ / minute, the decarburization reaction on the surface of the billet due to the increase in the amount of oxidation and the distribution of heterogeneous oxide layer is difficult to achieve the effect of the present invention, 20 ℃ / It is not preferable to exceed the minutes because the temperature deviation inside and outside the billet is deepened and billet bending occurs.

The heating rate from Ac1 to Ac3 in the above ideal range is performed at 6 ± 3 ° C / min. When the temperature rise rate is less than 3 ° C./min, the thickness of the surface ferrite layer increases more than necessary (for example, ≧ 0.5 mm) when passing through the abnormal temperature range. This is because there is a disadvantage in that the time (charging time) becomes long. In addition, it is because it is difficult to form an appropriate surface ferrite layer (for example, 0.2-0.4 mm) necessary for suppressing decarburization if the temperature increase rate exceeds 9 ° C / min.

The heating rate up to 1050 ± 50 ° C is carried out at 10 ± 5 ° C / min. The reason for this is to limit the temperature increase rate from Ac ₃ temperature, which is the abnormal station termination temperature, to the heating holding temperature. After the abnormal station termination temperature, the microstructure of the base material is transformed from the abnormal structure (ferrite + austenite) to the austenite single phase structure. Done. Therefore, at this time, if the temperature increase rate is less than 5 ℃ / min, the material time increases, and if it exceeds 15 ℃ / min, the heating rate is 10 ± 5 ℃ / min, because the billet bends easily occur due to the deep temperature deviation inside and outside the billet. It is limiting.

The heating rate of 10 ± 5 ℃ / min is performed to the heating temperature by the wire rod heating, the heating temperature of the wire heating is limited to 1050 ± 50 ℃ range. This is a problem in controlling the proper thickness of the billet surface ferrite layer for the decarburization control when the heating temperature is less than 1000 ℃ by the wire heating, it is not easy to re-use the coarse precipitated vanadium-based or niobium-based precipitates, hot This is because the workability becomes poor due to the overload during rolling due to the increase in the deformation resistance, and if the ferrite layer for the decarburization control cannot be deposited on the surface when it exceeds 1100 ° C. That is, it is possible to have a surface ferrite layer with a very low carbon solubility. However, if the heating temperature exceeds 1100 ° C, the surface ferrite layer is rapidly decarburized because the surface ferrite layer transforms into austenite, thereby increasing the surface decarburization. There is a problem.

The heating holding time at the heating temperature of the wire heating furnace as described above is limited to the range of 30-60 minutes, because in less than 30 minutes, it is difficult to secure a uniform temperature distribution inside and outside the billet for rolling the wire, and exceeds 60 minutes. This is because the amount of oxidation rapidly increases, making it difficult to secure an appropriate ferrite layer for billet surface decarburization control.

In this case, the ferrite layer is more preferably composed of 0.01-0.05mm of the surface in the bolt manufacturing.

By such a method, a surface ferrite layer having an appropriate thickness can be obtained, and decarburization can be reduced as compared with the conventional method.

Next, a method of manufacturing a bolt having high strength and high elongation by appropriately controlling heat treatment conditions using a wire rod manufactured by the above method will be described in detail.

In the present invention, by using a wire rod formed with the surface ferrite layer, to obtain a bolt by processing into a bolt of a certain shape, and then obtained bolts from Ac3- (Ac3-Acl) / 1.3 to Ac3- (Ac3-Acl) / 5.5 After heating for 20 minutes or more in the range and quenching to Ms + 50 ° C-Ms + 110 ° C at a cooling rate of 70 ° C / sec or more, and maintaining at this temperature for at least 20 minutes, followed by oil cooling or air cooling.

The abnormal reverse heat treatment condition is limited to Ac3- (Ac3-Ac1) /1.3 to Ac3- (Ac3-Ac1) /5.5 temperature range, and the reason for limitation is as follows.

Here, Ac3 is the austenite transformation temperature when heated, Ac1 is the transformation temperature in the abnormal region (ferrite + austenite) when heating, Ac3, Ac1 transformation temperature is different depending on the alloy component system according to the alloy component system.

This is because it is not preferable that the amount of ferrite produced when the reverse heat treatment is higher than 25% at the temperature of Ac3- (Ac3-Ac1) /1.3 is lower than Ac3- (Ac3-Ac1) /, as mentioned above. This is because at 5.5 or higher, the ferrite amount required for grain boundary discontinuity exceeds 5%, so the effect is difficult to expect.

The heat treatment may be completed by 20 minutes or more to complete the desired transformation, after the heat treatment is quenched at a cooling rate of 70 ℃ / sec or more.

The isothermal heat treatment conditions for the fabrication of bainite tissue after the heating is limited to the martensite transformation temperature (Ms) directly to Ms + (80 ° C. ± 30 ° C.).

If the isothermal heat treatment is carried out below Ms + 50 ° C., the bainite transformation time is long and the martensite is likely to occur when an isothermal heat treatment temperature deviation occurs, which is not desirable, and the elongation and impact toughness are reduced. If it exceeds + 110 ℃, there is a problem in securing proper yield strength due to the drastic reduction of yield ratio (yield strength / tensile strength ratio) .Therefore, there is a problem in that stress relaxation is poor when bolting, and fracture resistance is reduced by reducing impact toughness. This is because it is harmful to and also affects the critical delay strength and fatigue characteristics.

The present invention applying the above method obtains high strength and high elongation, because this effect is controlled to 5-25% of the ferrite fraction in the composite structure of bainite and ferrite. In other words, the high-strength high elongation effect of the present invention is to reduce the distribution of precipitates at grain boundaries, and for this purpose, it is preferable that the microstructure fraction of ferrite is in the range of 5-25%. For this reason, if the ferrite structure fraction is less than 5%, the amount of ferrite is too small to discontinue the grain boundaries of austenite, and the effect is insufficient. If the ferrite structure is more than 25%, the ferrite structure is maintained like the parent material due to the excessive ferrite fraction. This is because there is a lack of bainite structure effective at high elongation due to the decrease in yield strength and the increase of carbon concentration in austenite due to the increase of ferrite fraction.

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

Example

Ingot of 50 kg having a composition as shown in Table 1 was cast for each steel type.

C Si Mn Cr V Nb Mo Ti W B P S N ₂ Inventive Steel 1 0.45 3.03 0.29 0.58 0.54 - - - - - 0.005 0.004 0.008 Inventive Steel 2 0.40 3.42 0.31 0.79 0.2 - - 0.01 - 0.0013 0.006 0.005 0.014 Invention Steel 3 0.60 2.99 0.32 0.33 0.05 0.54 - - 0.02 - 0.007 0.009 0.007 Inventive Steel 4 0.45 2.0 0.77 0.51 0.11 - 0.2 0.03 - - 0.006 0.008 0.009 Inventive Steel 5 0.44 3.96 0.23 0.27 0.06 - - - 0.2 0.0015 0.008 0.008 0.008 Inventive Steel 6 0.53 3.01 0.35 0.55 - - 0.05 0.05 0.07 0.0010 0.004 0.009 0.004 Inventive Steel 7 0.58 2.56 0.60 0.29 - 1.10 0.13 0.10 - - 0.005 0.006 0.005

The ingot prepared as shown in Table 1 was heated at a rate under the conditions shown in Table 2 below, and the heated billet was rolled to prepare a wire having a diameter of 13 mm.

The surface of the decarburized layer was measured according to the standard of KS D 0216 of the manufactured wire rod, and the results are shown in Table 2 below. At this time, the measurement position was measured at the position where the wire rod section was divided into 8 and the measured value was based on the average value.

division Wire heating furnace heating pattern Wire Stripping Depth (mm) Heating rate up to Ac ₁ (℃ / min) Ideal temperature rise rate (℃ / min) Temperature rise rate after Ac ₃ (℃ / min) Heating temperature (℃) Holding time (min) Total decarburization (mm) Ferrite Layer Thickness (mm) Invention One 15 4 11 1050 30 0.04 0.03 2 15 5 13 1100 60 0.04 0.03 3 15 6 12 1000 30 0.05 0.04 4 16 7 12 1050 60 0.06 0.05 5 17 8 12 1050 50 0.05 0.04 6 15 10 11 1050 50 0.06 0.04 7 15 10 13 1100 60 0.07 0.05 Comparative material One 16 12 12 1050 50 0.10 0.08 2 16 15 12 1050 40 0.12 0.10 3 16 15 12 1050 60 0.13 0.11 4 15 17 11 1100 30 0.17 0.14 5 15 17 11 1050 50 0.18 0.15 6 15 19 11 1100 30 0.19 0.15 7 15 19 11 1050 50 0.20 0.16

As can be seen in Table 2, the wire rod surface decarburization depth of the invention material (1-7) that satisfies the conditions of the present invention shows a range of 0.04-0.06mm, while the comparative material outside the conditions of the present invention (1- 7) showed wire decarburization in the range of 0.10-0.20 mm. That is, in the present invention, as a result of using a wire having a significantly reduced total depth of the ferrite decarburized layer, as shown in the following results, it is very advantageous for characteristics such as delayed fracture of the bolt.

On the other hand, using the wires of the Table 2 prepared as described above was processed into bolts, and then subjected to a heat treatment test for the processed bolts under the same conditions as Table 3 below. In the results of Table 2, Ac ₃ and Ac ₃ transformation temperatures for determining the microstructure phase fraction and the abnormal range were measured using a thermal analyzer (dilatometry), and the results of the invention are shown in Table 3 together. It was.

Heating temperature (℃) Ac ₃ -[(Ac ₃ -Ac ₁ ) / X] Heating time (minutes) Isothermal heating temperature (C) (Ms + X) Isothermal holding time (min) Ferrite Percentage (%) Transformation temperature (C) Ac ₃ Ac ₁ Ms Inventive Example 1 X = 2 40 80 40 12 915 818 272 Inventive Example 2 X = 2 30 80 40 17 955 833 290 Inventive Example 3 X = 2 70 80 40 10 883 803 229 Inventive Example 4 X = 2 80 80 40 8 880 782 260 Inventive Example 5 X = 2 30 80 40 20 961 842 288 Inventive Example 6 X = 2 40 80 40 12 899 817 250 Inventive Example 7 X = 2 120 80 40 8 857 775 208

Tensile characteristics, impact characteristics, and delayed fracture characteristics of the bolts manufactured as described above were evaluated, and the results are shown in Tables 4 and 5, respectively.

Tensile test pieces were used in No. 4 test specimens according to the standard of KS B 0801, the tensile test was tested at a cross head speed (5 mm / min).

In addition, the impact test piece was manufactured according to No. 3 test piece of KS B 0809, wherein the notch direction was machined from the side (LT direction) of the wire rod rolling direction. The microstructure fraction was investigated using a point counting method, which is a general optical microscopy method, wherein the test surface was 1000 mm ² .

In addition, the delayed fracture resistance evaluation is applied to the commonly used constant load method. This evaluation method is a general method for evaluating delayed fracture resistance by the time of additional stress or the time required to break under a specific stress. In the delayed fracture test, the test stress was determined based on the notched tensile strength.

The delayed fracture tester used a constant loading type delayed fracture testing machine. The delayed fracture specimens were prepared with a specimen diameter of 6 mm, a notch diameter of 4 mm, and a notch root radius of 0.1 mm. The test piece atmosphere solution was performed at room temperature (25 ± 5C) at a pH of 2 ± 0.5 with a wolpole buffer solution (HCl + CH ₃ COONa).

The critical delay fracture strength is the tensile strength at which the required time from failure to break for a specific stress ratio (load stress / notch tensile strength ratio) is not more than 150 hours. Tooth section area). The number of specimens for setting the critical delay fracture strength was based on the case where more than 13 specimens were not broken.

Tensile Strength (kg / mm ² ) Yield strength (kg / mm ² ) Elongation (%) Cross section reduction rate (%) Critical Delay Break Strength (kg / mm ² ) Critical Delay Failure Stress Ratio Invention Example 1 142 104 34 53 140 0.65 Invention Example 2 149 119 32 56 150 0.6 Invention Example 3 163 127 29 61 150 0.6 Invention Example 4 154 119 32 49 150 0.65 Invention Example 5 151 121 31 51 150 0.6 Invention Example 6 155 123 30 55 150 0.6 Invention Example 7 164 128 30 49 150 0.65

Tensile Strength (kg / mm ² ) Yield strength (kg / mm ² ) Elongation (%) Cross section reduction rate (%) Critical Delay Break Strength (kg / mm ² ) Critical Delay Failure Stress Ratio Comparative Example 1 130 95 33 50 145 0.5 Comparative Example 2 136 108 33 55 145 0.5 Comparative Example 3 148 115 28 60 145 0.5 Comparative Example 4 140 108 33 48 145 0.5 Comparative Example 5 138 111 30 49 145 0.5 Comparative Example 6 142 112 32 54 145 0.5 Comparative Example 7 150 116 29 47 145 0.5

As can be seen in Tables 4 and 5, the elongation of the present invention is 21-35% and the critical delay fracture strength is 120-150kg / mm ² level, but the comparative materials have an elongation of 8-18% and the critical delay fracture strength 95 The present invention using the composite structure consisting of ferrite + bainite at -140kg / mm ² level can be seen that the elongation and the critical delay strength is significantly improved compared to the comparative materials.

As described above, the present invention proposes an alloy composition system and heat treatment conditions of a ferritic composite steel for excellent elongation and resistance to delayed fracture resistance, thereby achieving high strength of bolts and simultaneously ensuring excellent critical delay stress ratio. High strength high elongation bolts are obtained.

Claims (2)

In weight%, carbon: 0.40-0.60%, silicon: 2.0-4.0%, manganese: 0.2-0.8%, chromium: 0.2-0.8%, phosphorus: 0.01% or less, sulfur: 0.01% or less, nitrogen 0.005-0.01%, Oxygen 0.005% or less is contained, including vanadium: 0.05-0.2%, niobium: 0.05-0.2%, nickel: 0.3-2.0%, boron 0.001-0.003%, molybdenum: 0.01-0.5%, titanium: 0.01-0.2 %, Tungsten: 0.01-0.5%, copper: 0.01-0.2%, cobalt: 0.01-0.5% of the steel billet containing one or two or more selected from the group consisting of balance Fe and other unavoidable impurities, Heat up to Ac1 at a heating rate of 15 ± 5 ° C / minute, and heat up from Ac1 to Ac3, an ideal range, at a heating rate of 6 ± 3 ° C / minute, and then heat up to 1050 ± 50 ° C at 10 ± 5 ° C / minute After heating for 30-60 minutes, the wire is rolled, processed into a bolt of a certain shape to obtain a bolt, and then the obtained bolt is transferred from Ac3- (Ac3-Acl) /1.3 to Ac3- (Ac3-Acl) /5.5. Heating for more than 20 minutes in the range, Ms at a cooling rate of 70 ℃ / sec or more After quenching to + 50 ° C.-Ms + 110 ° C., the mixture is maintained at this temperature for at least 20 minutes, and then oil-cooled or air-cooled. The method of claim 1, The wire is a method of producing a high strength high elongation bolt, characterized in that the surface has a ferrite layer of 0.01-0.05mm