CN108239726B

CN108239726B - Steel material for high-strength spring excellent in hydrogen embrittlement resistance and method for producing same

Info

Publication number: CN108239726B
Application number: CN201711258437.5A
Authority: CN
Inventors: 全英洙; 李炳甲; 金宽镐; 郑会荣; 金汉辉
Original assignee: Posco Co Ltd
Current assignee: Posco Holdings Inc
Priority date: 2016-12-23
Filing date: 2017-12-04
Publication date: 2020-04-10
Anticipated expiration: 2037-12-04
Also published as: CN108239726A; KR20180074008A

Abstract

One aspect of the present invention provides a steel material for a high-strength spring excellent in hydrogen embrittlement resistance, which comprises, in wt%: 0.45-0.60%, Mn: 0.30% -0.80%, Si: 1.40% -1.80%, P: 0.015% or less, S: 0.015% or less, Cr: 0.2% -0.7%, Mo: 0.05-0.15%, V: 0.05-0.20%, Nb: 0.01% -0.03%, N: less than 0.006%, the balance being Fe and other unavoidable impurities, the prior austenite grain size being of the order of 10.0 or more, and containing carbides (Nb, V, Mo) C having an average size, measured as circle-equivalent diameter, of 15 to 35 nm.

Description

Steel material for high-strength spring excellent in hydrogen embrittlement resistance and method for producing same

Technical Field

The present invention relates to a steel material for high-strength springs excellent in hydrogen embrittlement resistance and a method for producing the same.

Background

Recently, as the amount of fossil fuel used increases, particularly the amount of petroleum fuel used sharply increases, the pollution source resulting from the combustion of petroleum fuel causes serious pollution of the atmosphere, which has become a global problem. In order to avoid the damage caused by petroleum fuel, the technology of reducing the use amount of petroleum fuel as much as possible is studied from various aspects.

Automobiles are heavily using petroleum fuel, and automobile manufacturing companies are continuously performing various attempts and studies to reduce the amount of petroleum fuel used. One of the most traditional ways to reduce the amount of petroleum fuel used is to improve the fuel efficiency of automobiles, which is the mainstream technology currently being developed and applied, and one of the methods to improve the fuel efficiency is to improve the combustion efficiency and power transmission efficiency of engines. Another way to reduce the amount of petroleum fuel used is to reduce the amount of energy required to travel a unit distance by reducing the weight of the vehicle.

In order to reduce the weight of the car body, the car parts can be replaced by light materials, but the application fields of the parts which can replace the advantages of the steel materials per se are not many. Therefore, nowadays, steel materials are still used for manufacturing automobile parts in many cases, and it is also common to try to improve the fuel efficiency of automobiles by reducing the weight of the steel materials.

If only the steel parts are reduced in weight, the load per unit weight is limited, which may bring fatal problems to the safety of automobiles. Therefore, the weight reduction of the component is inevitably possible after the problem of the increase in strength of the component is solved. Further, as the strength of the material increases, a problem of increasing hydrogen embrittlement sensitivity will arise. In particular, calcium chloride or sodium chloride sprinkled on ice surfaces in winter for snow melting and deicing causes corrosion pits in the spring to become stress concentration areas, which further increases the hydrogen embrittlement sensitivity.

In order to solve these problems, patent document 1 discloses a method for increasing hydrogen embrittlement resistance by grain refinement of prior austenite.

Patent document 2 describes that hydrogen embrittlement resistance can be improved by realizing a fine microstructure in which retained austenite is 6% to 15% and the prior austenite grain size class is 10.0 or more.

However, sufficient hydrogen embrittlement resistance cannot be ensured only by realizing a fine microstructure according to patent documents 1 and 2. Therefore, it is necessary to develop a steel material for high-strength springs excellent in hydrogen embrittlement resistance and a method for producing the same.

Documents of the prior art

Patent document 1: korean laid-open patent publication No. 2015-0081366

Patent document 2: korean laid-open patent publication No. 2015-0002848

Disclosure of Invention

Technical problem

The invention aims to provide a high-strength spring steel material with improved hydrogen embrittlement resistance and a manufacturing method thereof.

In addition, technical problems of the present invention are not limited to the above, and can be understood based on the contents throughout the present specification, and it will not be difficult for a person having ordinary skill in the art to understand additional technical problems of the present invention.

Technical scheme

In addition, another aspect of the present invention provides a method for manufacturing a steel material for a high-strength spring excellent in hydrogen embrittlement resistance, including: a wire rod preparation step of preparing a wire rod comprising, in wt%: 0.45-0.60%, Mn: 0.30% -0.80%, Si: 1.40% -1.80%, P: 0.015% or less, S: 0.015% or less, Cr: 0.2% -0.7%, Mo: 0.05-0.15%, V: 0.05-0.20%, Nb: 0.01% -0.03%, N: less than 0.006%, the balance Fe and other unavoidable impurities; a quenching heat treatment step of heating the wire rod to a temperature range of 880 to 1000 ℃ and holding for 10 seconds to 30 minutes, and then cooling at an average cooling rate of 10 ℃/s or more; and a tempering heat treatment step of heating the wire rod subjected to the quenching heat treatment to a temperature range of 380 to 460 ℃ for 10 seconds to 40 minutes, and then cooling the wire rod to 60 ℃ or less at an average cooling rate of 10 ℃/s or more.

In addition, the above technical solutions do not list all features of the present invention. The various features of the present invention, together with the advantages and effects attendant thereto, may be understood in greater detail with reference to the following specific exemplary embodiments.

Effects of the invention

According to the present invention, a steel material having a tensile strength of 1950MPa or more and excellent hydrogen embrittlement resistance can be provided by controlling the prior austenite grain size and carbide (Nb, V, Mo) C, and a method for producing the same.

Drawings

FIG. 1 is a photograph of carbide (Nb, V, Mo) C of invention example 3 taken by a transmission electron microscope.

Fig. 2 is a graph showing the temperature-based hydrogen desorption behavior obtained by the hydrogen thermal desorption experiments of inventive example 3 and comparative example 5.

Fig. 3 is a picture of the microstructures of comparative example 8 and inventive example 1 taken with a scanning electron microscope.

Detailed Description

Preferred embodiments of the present invention are described below. However, the present invention can be modified in various ways, and the scope of the present invention is not limited to the following embodiments. In addition, the embodiments of the present invention are provided to more fully describe the present invention to those of ordinary skill in the art.

The present inventors have considered that it is difficult to prevent hydrogen embrittlement by a method of painting, coating urethane, or the like in order to prevent hydrogen embrittlement of a coil spring, and that it is necessary to improve hydrogen embrittlement resistance of a material itself because a spring is dented by sand or the like involved in a vehicle running, which is an actual use environment. The present inventors have conducted intensive studies in order to solve this problem.

As a result, they have found that hydrogen embrittlement resistance can be improved by realizing a fine microstructure and simultaneously generating hydrogen desorption peaks at 310 to 350 ℃ and 420 to 460 ℃ in a thermal desorption experiment by compositely adding Nb, V, and Mo using an interface of carbide (Nb, Mo, V) C as a non-diffusible hydrogen trapping site, thereby completing the present invention.

High-strength spring steel material excellent in hydrogen embrittlement resistance

The steel material for a high-strength spring excellent in hydrogen embrittlement resistance according to one aspect of the present invention will be described in detail below.

According to one aspect of the present invention, a steel material for a high strength spring excellent in hydrogen embrittlement resistance, which comprises, in wt%: 0.45-0.60%, Mn: 0.30% -0.80%, Si: 1.40% -1.80%, P: 0.015% or less, S: 0.015% or less, Cr: 0.2% -0.7%, Mo: 0.05-0.15%, V: 0.05-0.20%, Nb: 0.01% -0.03%, N: less than 0.006%, the balance being Fe and other unavoidable impurities, the prior austenite grain size being of the order of 10.0 or more, and containing carbides (Nb, V, Mo) C having an average size, measured as circle-equivalent diameter, of 15 to 35 nm.

First, the alloy composition of the present invention will be explained in detail. Unless otherwise specified, the content unit of each element below is% by weight.

C：0.45％～0.60％

C is an element that must be added to ensure the strength of the spring.

If the C content is less than 0.45%, hardenability cannot be ensured, and it is difficult to sufficiently ensure strength required for the steel material for springs. Therefore, the lower limit of the C content is preferably 0.45%, and more preferably 0.48%.

On the contrary, if the C content is more than 0.60%, a sheet-like martensite structure having poor corrosion resistance is formed, and the material is cracked, so that hydrogen embrittlement resistance is remarkably reduced. Further, it is difficult to ensure sufficient toughness for increasing the strength, and it is difficult to suppress the decarburization phenomenon of the material due to the addition of a high content of Si.

Therefore, the C content is preferably 0.45% to 0.60%, and more preferably 0.47% to 0.58%.

Mn：0.30％～0.80％

When Mn is contained in the steel material, hardenability of the steel material is improved, which is advantageous for securing strength. If the Mn content is less than 0.30%, it is difficult to obtain sufficient strength and hardenability required for a material for a high-strength spring, and if it is more than 0.80%, toughness is lowered.

Therefore, the Mn content is preferably 0.30% to 0.80%, and more preferably may be 0.35% to 0.75%.

Si：1.40％～1.80％

Si is dissolved in ferrite to have the effect of strengthening the strength of the base material and improving the deformation resistance. If the Si content is less than 1.4%, the effect of Si solid-dissolved in ferrite to strengthen the base material and improve deformation resistance is not sufficiently exerted. On the contrary, if the Si content is more than 1.8%, surface decarburization is promoted at the time of heat treatment.

Therefore, the Si content is preferably 1.4% to 1.8%, and more preferably 1.43% to 1.75%.

P: less than 0.015%

Since P is an impurity and segregates at grain boundaries to lower the toughness, it is important to control the upper limit, and it is preferably controlled to 0.015% or less. More preferably, it is controlled to 0.012% or less.

S: less than 0.015%

S is an impurity, and is segregated in the grain boundary as a low-melting point element to cause a decrease in toughness, and the formed sulfide adversely affects the spring characteristics, so that the upper limit thereof is important to be controlled, and is preferably controlled to 0.015% or less. More preferably, it is controlled to 0.012% or less.

Cr：0.2％～0.7％

Cr is an element which contributes to ensuring oxidation resistance, temper softening property, hardenability and prevention of surface decarburization. If the Cr content is less than 0.2%, it becomes difficult to sufficiently ensure oxidation resistance, temper softening, hardenability, prevention of surface decarburization, and the like. Conversely, if the Cr content is more than 0.7%, the deformation resistance is lowered, the pitting corrosion is promoted, and C may be formed (V, Cr) to coarsen the carbide size.

Therefore, the Cr content is preferably 0.2% to 0.7%, and more preferably 0.25% to 0.65%.

Mo：0.05％～0.15％

Mo forms carbide, thereby playing the roles of improving precipitation hardening, original austenite grain refinement and hydrogen embrittlement resistance. If the Mo content is less than 0.05%, sufficient carbides cannot be formed, and if the Mo content is more than 0.15%, the production cost increases, the effect of delaying phase transformation is remarkable, and the problem of lowering productivity is caused.

Therefore, the Mo content is preferably 0.05% to 0.15%, and more preferably may be 0.06% to 0.14%.

V：0.05％～0.20％

V forms single or composite carbide, thereby playing the roles of improving precipitation hardening, original austenite grain refinement and hydrogen embrittlement resistance. If the V content is less than 0.05%, sufficient carbides cannot be formed, and if the V content is more than 0.20%, the production cost is increased, and the amount of coarse-grained carbides that are not dissolved during the heat treatment is increased, thereby deteriorating fatigue characteristics and precipitation hardening effects.

Therefore, the V content is preferably 0.05% to 0.20%, and more preferably may be 0.06% to 0.19%.

Nb：0.01％～0.03％

Nb forms composite carbide, thereby playing the roles of improving precipitation hardening, original austenite grain refinement and hydrogen embrittlement resistance. In particular, Nb forms carbides that do not dissolve at high temperatures and does not dissolve again even if the heat treatment time is long, and thus has an excellent effect of refining prior austenite grains. If the Nb content is less than 0.01%, sufficient carbides cannot be formed, and if the Nb content is more than 0.03%, the amount of coarse-grained carbides increases, thereby deteriorating fatigue characteristics and precipitation hardening effects.

Therefore, the Nb content is preferably 0.01% to 0.03%, and more preferably may be 0.013% to 0.027%.

N: less than 0.006%

N is an impurity and is bonded to Nb to form a coarse precipitate that is not redissolved during heat treatment. Therefore, since N offsets the effect of Nb addition, it is important to control the upper limit thereof, and it is preferably controlled to 0.006% or less. More preferably, it is controlled to 0.0055% or less.

The balance of the present invention is iron (Fe). However, impurities from the raw materials or the surrounding environment may be inadvertently and unavoidably mixed in the conventional manufacturing process, and thus it cannot be excluded that these impurities are mixed in. Any person skilled in the art will know these impurities and therefore they will not be described in detail in this specification.

The alloy composition described above can sufficiently ensure the high strength and hydrogen embrittlement resistance desired in the present invention, but may further contain Ni: 0.1% -0.4% and Cu: more than one of 0.1 to 0.3 percent.

Ni：0.1％～0.4％

Ni is an element added to improve hardenability and toughness. If the Ni content is less than 0.1%, the effect of improving hardenability and toughness is insufficient, and if the Ni content is more than 0.4%, the retained austenite amount increases, thereby decreasing fatigue life, and the manufacturing unit price is sharply increased by expensive Ni.

Cu：0.1％～0.3％

Cu is an element advantageous for preventing decarburization and improving corrosion resistance. The decarburized layer significantly reduces fatigue life after spring processing. If the Cu content is less than 0.1%, the above effect is insufficient, and if the Cu content is more than 0.3%, rolling defects due to high temperature brittleness may occur.

The steel material of the present invention has a prior austenite grain size of 10.0 or more, and contains carbides (Nb, V, Mo) C having an average size of 15 to 35nm as measured by an equivalent circle diameter.

The grain size refers to ASTM grain size class, and if the prior austenite grain size is less than 10.0 class, the grain size is large, so that it is difficult to uniformly distribute hydrogen diffused to the grain boundary, and the critical hydrogen concentration causing fracture is reached too early, which causes a problem of poor hydrogen embrittlement resistance.

The carbides (Nb, V, Mo) C refer to NbC, VC, MoC and composite carbides thereof, which pin (pinning) austenite grains during heat treatment to prevent grain growth, thereby functioning to refine the size of the prior austenite grains of the steel for springs and to non-diffusible hydrogen trapping sites.

Carbides in steel may serve as trapping sites for hydrogen that enters during manufacturing processes or hydrogen that enters due to corrosion during use of vehicle springs, and interfaces of carbides (Nb, V, Mo) C serve as non-diffusible hydrogen trapping sites that do not diffuse due to external stress because the activation energy required for hydrogen diffusion is high.

If the average size of the (V, Mo) C carbide measured by circle-equivalent diameter is larger than 35nm, it may rather become a starting point of hydrogen embrittlement, and thus there is a problem that hydrogen embrittlement resistance is deteriorated. In contrast, the lower limit of the average size does not need to be particularly limited, but is difficult to control to less than 15nm, and thus the lower limit thereof is defined to be 15 nm.

In this case, the carbide (Nb, V, Mo) C may be 8 particles/μm²45 pieces/. mu.m²。

If the carbide (Nb, V, Mo) C is less than 8/μm²It is difficult to sufficiently exhibit the function of the non-diffusible hydrogen trapping sites, and if the number of carbide (Nb, V, Mo) C exceeds 45/μm²When the addition amounts of Nb, V and Mo are too large, carbide (Nb, V, Mo) C tends to be coarsened.

Furthermore, the microstructure of the steel material of the present invention may contain 90% to 95% of tempered martensite and 5% to 10% of retained austenite in terms of area fraction.

In addition, for the steel material of the invention, the hydrogen desorption peak can be generated at the same time under the temperature of 310-350 ℃ and 420-460 ℃ during the thermal desorption experiment.

Hydrogen in a steel material is roughly classified into diffusible hydrogen and non-diffusible hydrogen. Diffusible hydrogen is hydrogen that diffuses based on a mechanical driving force or a chemical driving force according to external stress, resulting in hydrogen embrittlement, and non-diffusible hydrogen is hydrogen that does not diffuse due to a driving force. Such diffusible hydrogen and non-diffusible hydrogen can be classified by a thermal desorption experiment in which the amount of hydrogen desorbed from the material when the temperature of the material is raised is observed. The hydrogen desorbed at the temperature rise of 300 ℃ may be defined as diffusible hydrogen. In addition, if the hydrogen trapping site is subjected to a temperature above the activation energy, a peak in the amount of hydrogen desorption occurs at a specific temperature, and the hydrogen trapping site of the material can be derived therefrom. The simultaneous occurrence of hydrogen desorption peaks at 310-350 ℃ and 420-460 ℃ in thermal desorption experiments indicates that sufficient non-diffusible hydrogen trapping sites, namely carbides (Nb, V and Mo) C, exist. There are two peaks indicating the presence of two carbides with different interfacial properties, the peaks at 310 ℃ to 350 ℃ being due to the V, Mo content of carbide (Nb, V, Mo) C interface relatively lower than the Nb content, and the peaks at 420 ℃ to 460 ℃ being due to the V, Mo content of carbide (Nb, V, Mo) C interface relatively higher than the Nb content. The composite carbide has different contents of Nb, V and Mo according to the generation temperature and the re-solid solution fraction, and the composite effect is not particularly limited due to various situations. Nevertheless, the higher the amount of hydrogen trapped at the non-diffusive hydrogen trapping sites, the more excellent the resistance to hydrogen embrittlement.

The tensile strength of the steel material of the present invention may be 1940MPa or more.

If the tensile strength is less than 1940MPa, the effect of weight reduction of the vehicle due to high strength is poor, and it is difficult to satisfy the design shear stress of the coil spring of 1200MPa or more.

Method for producing steel material for high-strength spring excellent in hydrogen embrittlement resistance

The method for manufacturing a steel material for a high strength spring excellent in hydrogen embrittlement resistance according to another aspect of the present invention will be described in detail below.

According to another aspect of the present invention, a method for manufacturing a steel material for a high strength spring excellent in hydrogen embrittlement resistance includes: a wire rod preparation step of preparing a wire rod having the alloy composition; a quenching heat treatment step of heating the wire rod to a temperature range of 880 to 1000 ℃ and holding for 10 seconds to 30 minutes, and then cooling at an average cooling rate of 10 ℃/s or more; and a tempering heat treatment step of heating the wire rod subjected to the quenching heat treatment to a temperature range of 380 to 460 ℃ for 10 seconds to 40 minutes, and then cooling the wire rod to 60 ℃ or less at an average cooling rate of 10 ℃/s or more.

Wire preparation step

A wire rod having the alloy composition described above was prepared.

At this time, the wire rod may be manufactured by a manufacturing method including: heating a steel billet having said alloy composition; performing final hot rolling on the heated steel billet at the temperature range of 810-910 ℃ to obtain a wire rod; and cooling the wire at a cooling rate of 0.15 ℃/s to 1.5 ℃/s.

The heating temperature of the billet is not particularly limited since it is a general temperature, and may be, for example, 1000 to 1150 ℃.

In the wire rod manufacturing process, the fraction of nuclei (nuclei) from which carbides can be precipitated increases as the hot rolling temperature is lower, so that the precipitation amount of carbide elements which are solid-dissolved increases, and the growth rate becomes slow due to the low temperature, so that precipitates tend to become fine. Since the number of nuclei decreases as the rolling temperature increases, the interfaces of precipitated carbides become nucleation sites or precipitates grow by themselves, and therefore, by controlling the hot rolling temperature to be low, the precipitation of coarse carbides can be suppressed.

When the final hot rolling temperature is higher than 910 ℃, coarse carbides may be formed.

On the contrary, when the final hot rolling temperature is lower than 810 ℃, it may cause the generation of surface ferrite decarburization and the generation of a low temperature structure, and the problem of equipment load occurs.

When the cooling rate is higher than 1.5 ℃/s, a low-temperature structure may be generated, and when the cooling rate is lower than 0.15 ℃/s, decarburization of surface ferrite may be generated, and there is a problem that the production rate is lowered.

Quenching heat treatmentStep (ii) of

The wire rod is heated to a temperature range of 880-1000 ℃ and held for 10 seconds to 30 minutes, and then cooled at an average cooling rate of 10 ℃/s or more to perform quenching heat treatment.

The purpose of heating to a temperature in the range of 880-1000 ℃ is to achieve austenitization. When the heating temperature is less than 880 c, undissolved pearlite may remain during heating, and when the heating temperature is more than 1000 c, decarburization and grain growth may be promoted.

When the holding time is less than 10 seconds, austenite reversion may not occur, undissolved pearlite may remain, and when the holding time exceeds 30 minutes, the prior austenite grain size may be less than 10.0.

When the average cooling rate is less than 10 ℃/s, an undesirable structure such as bainite is obtained in addition to a homogeneous martensite structure.

Tempering heat treatment step

The wire rod subjected to the quenching heat treatment is heated to a temperature range of 380-460 ℃ and held for 10 seconds to 30 minutes, and then cooled to 60 ℃ or less at an average cooling rate of 10 ℃/s or more to perform a tempering heat treatment.

When the heating temperature of the tempering heat treatment is less than 380 ℃, the tempering effect of martensite is insufficient to deteriorate the toughness of the spring, and when the heating temperature is more than 460 ℃, it is difficult to secure the strength.

When the holding time is less than 10 seconds, the tempering effect is insufficient, resulting in deterioration of toughness, and when the holding time exceeds 40 minutes, the epsilon martensite migrates to the carburized body, resulting in deterioration of strength and toughness.

When the average cooling rate is less than 10 ℃/s, the structure control during cooling becomes difficult, possibly resulting in a structural deviation of carbide generation in the tempered martensite structure obtained by rapid heating.

The quenching and tempering heat treatment is not particularly limited to heating, and may be, for example, induction heating or heating in a box furnace.

The present invention is described more specifically by examples below. However, the following examples are only for describing the present invention in more detail and are not intended to limit the scope of the present invention. The scope of rights of the present invention depends on the content reasonably derived from the content recited in the claims.

(examples)

Wires having the compositions shown in table 1 below were prepared. The wire rods were 15mm in diameter manufactured by heating a steel slab having the composition shown in table 1 below to 1030 ℃ and performing final hot rolling at 880 ℃ and then cooling at a cooling rate of 0.8 ℃/s.

The wire rods were maintained at the quenching temperatures shown in table 2 below for 10 minutes, and then rapidly cooled (100 ℃/s or more) in oil at 60 ℃ to perform quenching heat treatment. The quenching heat-treated wire rods were kept at the tempering temperature in table 2 for 30 minutes, and then rapidly cooled in oil at 60 ℃ (100 ℃/s or more) to perform tempering heat treatment, thereby producing steel materials.

Tensile strength, grain size, average size and number of carbides (Nb, V, Mo) C, presence or absence of hydrogen desorption peaks at 310 to 350 ℃ and 420 to 460 ℃ in hydrogen thermal desorption experiments, and hydrogen embrittlement breaking time of the produced steel were measured and written in table 2 below.

Tensile strength is measured by ASTM E-8M subsize specification.

For the microstructure, the prior austenite grain size was measured by grinding and corrosion and expressed in ASTM grades. For the average size of carbide (Nb, V, Mo) C, a Transmission Electron Microscope (TEM) replay sample was prepared, the equivalent circle diameter was measured, and the carbide distribution was evaluated by taking five arbitrary sites of stem (scanning transmission electron microscope) images.

For the hydrogen thermal desorption experiment (thermal desorption analysis), the amount of hydrogen desorbed from the steel material by heating at a temperature rising rate of 100 ℃ per hr was observed by a quadrupole mass spectrometer (Quadruple mass spectrometer).

The test was conducted after the steel material was processed into a sheet-like specimen having a size of 1.5mm × 8.0mm × 65mm with respect to the hydrogen embrittlement breaking time. The processed test piece was clamped with a four-point bending jig (4point bending jig)Then, a shear stress of 1400MPa was applied thereto, and the resultant was immersed in distilled water to which 0.01mol/1 of H was added₂SO₄+0.001mol/1 of KSCN, and a potential of-700 mV was applied using a potentiostat (potentiostat), thereby measuring the break time of the sample. In this case, the time for breaking the sample is defined as the hydrogen embrittlement breaking time, and it is acceptable to hold the sample for 600 seconds or more, and the sample having the breaking time shorter than 600 seconds is defined as the presence of the hydrogen embrittlement risk. At this time, the time of 600 seconds is the time until the amount of hydrogen taken in exceeds the hydrogen amount of the critical hydrogen concentration that may cause hydrogen embrittlement in the conventional spring steel material.

[ TABLE 1 ]

A1-A8 are cases where the alloy components given in the present invention are satisfied, and B1-B8 are cases where the alloy components given in the present invention are not satisfied.

[ TABLE 2 ]

In the phase fractions in table 2, TM is tempered martensite, RA is retained austenite, and F is ferrite.

Invention examples 1 to 6 satisfy the alloy composition and production conditions given in the present invention, and the prior austenite grain size is as fine as 10.0 or more, the average size of carbide (Nb, V, Mo) C is 35nm or less, the tensile strength is 1940MPa or more, and the hydrogen peaks appear simultaneously at 310 to 350 ℃ and 420 to 460 ℃, and therefore, the hydrogen embrittlement breaking time of 600 seconds or more is exhibited, and it can be confirmed that the hydrogen embrittlement resistance is excellent. This indicates that even in high-strength steel, hydrogen embrittlement resistance is excellent because hydrogen is dispersed by fine crystal grain size and trapped by fine carbides.

FIG. 1 is a photograph of carbide (Nb, V, Mo) C of invention example 3 taken with a transmission electron microscope, and it can be confirmed that fine carbide of NbC, VC, MoC, (Nb, V, Mo) C is present.

In addition, fig. 2 is a graph showing the temperature-based hydrogen desorption behavior obtained through the thermal desorption experiment of inventive example 3 and comparative example 5, the inventive example 3 having simultaneous hydrogen peaks at 320 ℃ and 430 ℃, and the comparative example 5 having only a hydrogen peak at 430 ℃.

Comparative example 1 in which the C content was 0.44 wt% did not reach the alloy composition of the present invention, and comparative example 2 in which the Si content was 1.30 wt% did not reach the alloy composition of the present invention, and thus the tensile strength desired in the present invention was not reached.

Comparative example 3 has excessive Mn and Cr additions, so grain boundary softening results in an unacceptable hydrogen embrittlement fracture time, and C formation (V, Cr) due to Cr addition also results in an increase in the average size of carbides.

Comparative example 4 was insufficient in Mo addition amount and comparative example 5 was insufficient in V addition amount, and thus (Nb, V, Mo) C composite carbide as an effective non-diffusive hydrogen trapping site could not be formed, resulting in failure in hydrogen embrittlement fracture time. In the thermal desorption experiment, it was confirmed that the non-diffusible hydrogen trapping sites could not be formed efficiently because no hydrogen peak occurred at 310 to 350 ℃.

Fig. 2 is a graph showing the temperature-based hydrogen desorption behavior obtained by the thermal desorption experiment of inventive example 3 and comparative example 5, the inventive example 3 simultaneously showing hydrogen peaks at 320 ℃ and 430 ℃, and the comparative example 5 showing a hydrogen peak only at 430 ℃.

When the diffusible hydrogen trapping site which could cause hydrogen embrittlement was defined as the trapping site of hydrogen desorbed at 300 ℃ or lower, comparative example 5 had a small number of carbides as the non-diffusible hydrogen trapping sites, so that the amount of hydrogen trapped at the diffusible hydrogen trapping site increased, and it was thought that the hydrogen embrittlement breaking time was not acceptable.

Comparative example 6 was insufficient in the amount of Nb added, and thus insufficient in the amount of carbide resulted in failure of the hydrogen embrittlement cracking time. When the steel is held at the austenitizing heat treatment temperature for a long period of time for quenching, high-temperature precipitates such as Nb do not re-dissolve, and (V, Mo) C carbides formed in the Nb-free steel re-dissolve, so that the prior austenite grain size cannot be controlled to be fine, and an effective non-diffusible hydrogen trapping site cannot be formed, resulting in deterioration of hydrogen embrittlement resistance.

Comparative example 7 is a steel material to which no Nb, Mo, V was added, and the fracture time was not acceptable despite the low strength and hydrogen embrittlement. This indicates that even a general steel material having low hydrogen embrittlement sensitivity has a risk of hydrogen embrittlement if the grain size is large and fine carbides are not present.

From comparative example 8, it is understood that if the austenitizing heat treatment temperature is low, undissolved pearlite causes a decrease in toughness, and thus hydrogen embrittlement resistance is also deteriorated.

Fig. 3 is a photograph of the microstructure of comparative example 8 and inventive example 1 taken by a scanning electron microscope, and it can be seen from fig. 3 that comparative example 8 has undissolved pearlite due to a low heat treatment temperature, while inventive example 1 having a high heat treatment temperature in the same composition system forms a tempered martensite structure.

From comparative example 9, it is understood that if the austenitizing heat treatment temperature is high, hydrogen embrittlement resistance is deteriorated due to grain growth.

From comparative example 10, it is understood that the formed coarse carbides cause deterioration of hydrogen embrittlement resistance because the amount of Nb added is high.

The foregoing is a description of the embodiments of the present invention, and those skilled in the art can make various changes and modifications to the present invention without departing from the scope of the present invention, and such changes and modifications also fall within the scope of the appended claims.

Claims

1. A steel material for high-strength springs excellent in hydrogen embrittlement resistance, which comprises, in% by weight, C: 0.45-0.60%, Mn: 0.30% -0.80%, Si: 1.40% -1.80%, P: 0.015% or less, S: 0.015% or less, Cr: 0.2% -0.7%, Mo: 0.05-0.15%, V: 0.05-0.20%, Nb: 0.01% -0.03%, N: less than 0.006%, the balance Fe and other unavoidable impurities,

the prior austenite grain size is more than 10.0 grade, and the prior austenite grain size comprises carbide (Nb, V, Mo) C, the average size of the carbide (Nb, V, Mo) C measured by equivalent circle diameter is less than 35nm, and the microstructure comprises 90-95% of tempered martensite and 5-10% of retained austenite by area fraction.

2. The steel material for high-strength springs excellent in hydrogen embrittlement resistance according to claim 1, wherein:

the steel further contains Ni: 0.1% -0.4% and Cu: more than one of 0.1 to 0.3 percent.

3. The steel material for high-strength springs excellent in hydrogen embrittlement resistance according to claim 1, wherein:

the steel thermal desorption experiment shows that the hydrogen desorption peak value appears at 310-350 ℃ and 420-460 ℃.

4. The steel material for high-strength springs excellent in hydrogen embrittlement resistance according to claim 1, wherein:

the tensile strength of the steel is 1940MPa or more.

5. The steel material for high-strength springs excellent in hydrogen embrittlement resistance according to claim 1, wherein:

the carbide (Nb, V, Mo) C is 8/mum²45 pieces/. mu.m²。

6. A method for producing a steel material for a high-strength spring excellent in hydrogen embrittlement resistance, comprising:

heating a steel slab comprising, in weight percent, C: 0.45-0.60%, Mn: 0.30% -0.80%, Si: 1.40% -1.80%, P: 0.015% or less, S: 0.015% or less, Cr: 0.2% -0.7%, Mo: 0.05-0.15%, V: 0.05-0.20%, Nb: 0.01% -0.03%, N: less than 0.006%, the balance Fe and other unavoidable impurities; performing final hot rolling on the heated steel billet at the temperature range of 810-910 ℃ to obtain a wire rod; and cooling the wire at a cooling rate of 0.15 ℃/s to 1.5 ℃/s;

a quenching heat treatment step of heating the cooled wire rod to a temperature range of 880 to 1000 ℃ for 10 seconds to 30 minutes, and then cooling the wire rod at an average cooling rate of 10 ℃/s or more; and

a tempering heat treatment step of heating the wire rod subjected to the quenching heat treatment to a temperature range of 380 to 460 ℃ for 10 seconds to 40 minutes, and then cooling the wire rod to 60 ℃ or less at an average cooling rate of 10 ℃/s or more.

7. The method for producing a steel material for a high-strength spring excellent in hydrogen embrittlement resistance according to claim 6, wherein:

the wire further comprises, in weight percent, Ni: 0.1% -0.4% and Cu: more than one of 0.1 to 0.3 percent.