JP6737102B2 - Steel material, sliding parts, and method for manufacturing steel material - Google Patents

Description

The present invention relates to a steel material, sliding parts, and a method for manufacturing a steel material.

A sliding component represented by a crankshaft is required to have wear resistance and fatigue strength. In order to increase wear resistance and fatigue strength, many sliding parts are subjected to surface hardening treatment such as induction hardening and cold roll processing.

In addition to the above characteristics, sliding parts are required to have excellent seizure resistance. With the need for lightweight automobiles, there is an increasing demand for narrower shafts and narrower crankshafts. Therefore, the sliding condition between the crankshaft and the bearing is severe. Therefore, the crankshaft is required to have seizure resistance superior to the conventional one.

Japanese Patent No. 4589885 discloses a crankshaft having a thermal conductivity κ of 40 W/mK or more and a surface hardness Hv after induction hardening greater than (2.7×κ+420). Has been done. The same document describes that the dominant factor for seizure is the temperature rise of the sliding surface. It is described that increasing the thermal conductivity and reducing the coefficient of friction are effective for suppressing the temperature rise.

In Japanese Patent No. 4140283, the structure in the region from the surface layer to 200 nm is a mixed structure of proeutectoid ferrite with an area ratio of 5% or less and pearlite with a lamellar spacing of 30 nm or less. A steel crankshaft is disclosed. In the same document, it is the hardness in the region from the surface layer to 200 nm that affects wear resistance, and if the hardness measured in this region using a nanoindentation device is 10 GPa or more, excellent wear resistance is obtained. It is described that sex is obtained.

Japanese Patent No. 4589885 Japanese Patent No. 4140283

As described in the above patent documents, in order to improve the seizure resistance, the hardness (hereinafter referred to as “nano hardness”) in a region within several hundred nm from the surface (sliding surface) is increased. It is effective to do. If the nano hardness is increased, excellent seizure resistance can be obtained even if the inside is soft.

The present inventors have found that the nanohardness of sliding parts changes with long-term sliding. In the above patent documents, the nano hardness after long-term sliding is not examined.

An object of the present invention is to provide a steel material suitable as a material for sliding parts, and to provide a sliding part having excellent seizure resistance.

A steel material according to an embodiment of the present invention has a chemical composition in mass% of C: 0.35 to 0.6%, Si: less than 0.50%, Mn: less than 0.80%, P: 0.02. % To 0.1% or less, S: less than 0.05%, Mo: 0.5 to 3.5%, Al: 0.005 to 0.06%, N: 0.001 to 0.02%, Nb: 0 to 0.05%, balance: Fe and impurities, having a structure in which Mo ₂ C is precipitated in at least the surface layer of the steel material, and the average particle spacing of Mo ₂ C is 1.0 μm or less.

The sliding part according to an embodiment of the present invention has a chemical composition, in mass %, of C: 0.35 to 0.6%, Si: less than 0.50%, Mn: less than 0.80%, P:0. More than 0.02% and 0.1% or less, S: less than 0.05%, Mo: 0.5 to 3.5%, Al: 0.005 to 0.06%, N: 0.001 to 0.02. %, Nb: 0 to 0.05%, the balance: Fe and impurities, the structure of at least the surface layer of the sliding component is martensite, and the average grain size of the former austenite grains of the martensite is 10 μm or less. ..

The sliding part according to an embodiment of the present invention has a chemical composition, in mass %, of C: 0.35 to 0.6%, Si: less than 0.50%, Mn: less than 0.80%, P:0. More than 0.02% and 0.1% or less, S: less than 0.05%, Mo: 0.5 to 3.5%, Al: 0.005 to 0.06%, N: 0.001 to 0.02. %, Nb: 0 to 0.05%, the balance: Fe and impurities, and the structure of at least the surface layer of the sliding component is a high dislocation density structure having a dislocation density of 10 ¹⁵ m ⁻² or more.

In the steel material manufacturing method according to an embodiment of the present invention, the chemical composition is C: 0.35 to 0.6%, Si: less than 0.50%, Mn: less than 0.80%, and P: More than 0.02% and 0.1% or less, S: less than 0.05%, Mo: 0.5 to 3.5%, Al: 0.005 to 0.06%, N: 0.001 to 0. 02%, Nb: 0-0.05%, balance: Fe and a raw material which is an impurity, a step of heating the raw material to 1000 to 1250° C., and a step of hot forging the heated raw material And a step of quenching the hot forged material from a temperature of 850° C. or higher, and a step of tempering the quenched material at a temperature of 500 to 650° C.

In the steel material manufacturing method according to an embodiment of the present invention, the chemical composition is C: 0.35 to 0.6%, Si: less than 0.50%, Mn: less than 0.80%, and P: More than 0.02% and 0.1% or less, S: less than 0.05%, Mo: 0.5 to 3.5%, Al: 0.005 to 0.06%, N: 0.001 to 0. 02%, Nb: 0-0.05%, balance: Fe and a raw material which is an impurity, a step of heating the raw material to 1000 to 1250° C., and a step of hot forging the heated raw material And the step of hot forging includes a warm working step of applying a work with a surface reduction rate of 10% or more at a temperature of 500 to 650° C. to the raw material.

According to the present invention, a steel material suitable as a material for sliding parts and a sliding part having excellent seizure resistance can be obtained.

FIG. 1 is a graph showing the relationship between the added amount of alloying elements and the thermal conductivity. FIG. 2A is a diagram schematically showing a precipitation form of Mo ₂ C. FIG. 2B is a diagram schematically showing a precipitation form of cementite. FIG. 3 is a flow chart showing an example of a method for manufacturing a steel material according to an embodiment of the present invention. FIG. 4 is a flowchart showing another example of the method for manufacturing a steel product according to the embodiment of the present invention. FIG. 5 is a flow chart showing an example of a method for manufacturing a sliding component according to an embodiment of the present invention. FIG. 6 is a flowchart showing another example of the method for manufacturing the sliding component according to the embodiment of the present invention. FIG. 7 is a schematic view of a test material processed for cold roll processing. FIG. 8 is a schematic diagram of a pin-on-disc type wear test. FIG. 9 is a schematic diagram of nanohardness measurement by the nanoindentation method. FIG. 10A is a graph showing the distribution of the nano hardness of the test number 1 after the abrasion test. FIG. 10B is a graph showing the distribution of nano hardness of the test number 2 after the abrasion test. FIG. 10C is a graph showing the distribution of the nano hardness of the test number 3 after the abrasion test. FIG. 10D is a graph showing the distribution of nanohardness of the test number 4 after the abrasion test. FIG. 10E is a graph showing the distribution of nanohardness of the test number 5 after the abrasion test. FIG. 11 is a graph showing the nano hardnesses of Test Nos. 1 to 21 at positions within 20 nm from the surface layer after the abrasion test.

The present inventors have studied the seizure phenomenon, mainly for crankshafts for automobiles. Normally, the crankshaft and the bearing slide in a fluid-lubricated state in which they face each other with an oil film interposed therebetween. However, investigations by the present inventors have revealed that during sliding for a long period of time, there are many situations from fluid lubrication to boundary lubrication. Generally, it has been clarified that when steel is subjected to sliding for a long time in a boundary lubrication environment, work hardening and softening due to frictional heat occur, and the hardness near the sliding surface (within 20 nm from the surface) changes. ..

As a result of observation of the structure, it was found that, on the sliding surface, formation of a processed layer by sliding and deterioration of the structure (tempering, recovery, and recrystallization) due to frictional heat occurred in a superimposed manner. Among these, the formation of the processed layer by sliding is a factor for improving the nanohardness, and the alteration of the structure by frictional heat is a factor for lowering the nanohardness.

Therefore, in order to improve the nano-hardness of the sliding surface, (1) friction heat dissipation by improving the thermal conductivity, (2) surface hardening treatment, and (3) alteration of the structure by improving thermal stability The suppression of is effective.

(1) Dissipation of frictional heat by improving thermal conductivity It has been conventionally known that reducing the Si content is effective for improving thermal conductivity. In order to clarify the effect of other elements, the present inventors measured the thermal conductivity by changing the addition amount of alloying elements based on medium-high carbon steel. The results are shown in Figure 1. As shown in FIG. 1, in addition to C and Si, Mn is an element that greatly affects the thermal conductivity, and it was found that the thermal conductivity can be improved by reducing the Mn content. On the other hand, it was found that Mo and W do not significantly affect the thermal conductivity.

Mn is an element that contributes to the improvement of hardenability, and conventionally, steel used for sliding parts contains 1 to 2% of Mn. In order to improve the thermal conductivity while ensuring the hardenability, it is necessary to contain Mo instead of reducing the Mn content.

(2) Surface hardening treatment In order to improve the wear resistance and seizure resistance of sliding parts, it is necessary to harden the sliding surface. However, simply hardening only makes the surface layer brittle. In order to improve wear resistance and seizure resistance, the work hardening rate must be increased. However, if the amount of work hardening increases, recrystallization tends to occur. Recrystallization partially occurs in the hardened work-hardened structure. Since the recrystallized portion has a soft structure, it easily becomes a starting point of wear. Therefore, it is necessary to select, for the sliding component, a steel material and a steel material structure that have a high work hardening rate and are resistant to recrystallization. As the surface hardening treatment for forming the structure satisfying the above conditions, induction hardening and cold roll working are suitable. The reason is that the martensite produced by induction hardening contains a large amount of mobile dislocations. Also, a large amount of mobile dislocations are contained in the work-hardened structure obtained by cold rolling.

When the steel material is subjected to induction hardening as a surface hardening treatment, at least the surface microstructure thereof becomes fine martensite. The present inventors have found that martensite generated from fine crystal grains has a high work hardening rate. The mechanism is not clear, but it is considered that the martensite variant and the substructures are influencing. The present inventors have also found that crystal grains can be made finer by adding Mo and P together. Although this mechanism is not clear, it is possible that Mo and P form clusters and suppress grain growth. Further, generally, when the crystal grains are made finer, the hardenability is lowered, but the combined addition of Mo and P also contributes to the improvement of the hardenability.

When cold rolling is performed on a steel material as a surface hardening treatment, frictional heat usually reduces the dislocation density in the surface layer portion. In order to improve the work hardening rate, it is preferable that the dislocation density in the surface layer portion can be maintained high. The present inventors have found that the recovery of dislocations can be suppressed by adding Mo and P together.

(3) Suppression of Deterioration of Microstructure by Improving Thermal Stability The present inventors have found that the thermal stability of steel can be improved by adding Mo and P together. Although this mechanism is not clear, it is considered that P hinders the movement of dislocations and suppresses tempering, recovery, and recrystallization. Further, it is possible that Mo and P form clusters by the combined addition of Mo and P to promote this effect.

The present inventors further perform (4) a treatment for precipitating fine Mo ₂ C before the surface hardening treatment (hereinafter, referred to as “fine-graining (precipitation) treatment”), so that after the surface hardening treatment, It was found that the seizure resistance of sliding parts can be further improved.

(4) Fine-graining (precipitation) treatment Mo ₂ C suppresses grain growth. Therefore, if induction hardening is performed as a surface hardening treatment on the steel material that has been subjected to grain refinement (precipitation), the former austenite grains of martensite formed by induction hardening can be significantly refined. More specifically, by setting the average particle spacing of Mo ₂ C to 1.0 μm or less, the average particle size of the prior austenite grains can be set to 10 μm or less. Further, Mo ₂ C has a pinning effect on dislocation movement, and suppresses dislocation recovery. Therefore, a high dislocation density can be imparted to the steel material subjected to the grain refinement (precipitation) by cold rolling as a surface hardening treatment.

Mo ₂ C further suppresses the formation of needle-shaped or plate-shaped coarse cementite. Needle-shaped or plate-shaped coarse cementite is generated by, for example, tempering heat treatment at low temperature performed after induction hardening to remove residual stress and frictional heat during sliding, and reduces wear resistance of sliding parts. As a result, seizure resistance is reduced. By fixing C, Mo ₂ C suppresses the generation of needle-shaped or plate-shaped coarse cementite, and contributes to the improvement of seizure resistance of the sliding component.

The grain refinement promoting process can be realized by the following (A) or (B). In addition, in order to control the average particle spacing of Mo ₂ C, it is necessary to appropriately control the Al content in addition to the C, Mo, and P contents.
(A) A hot forged material is quenched and further tempered at a temperature of 500 to 650°C.
(B) During hot forging, a material having a surface reduction rate of 10% or more is added to the material at a temperature of 500 to 650°C.

The present invention has been completed based on the above findings. Hereinafter, a steel material and a sliding component according to an embodiment of the present invention will be described in detail.

[Steel]
[Chemical composition]
The steel material according to the present embodiment has a chemical composition described below. In the following description, "%" of the content of an element means mass %.

C: 0.35-0.6%
Carbon (C) enhances the hardenability of steel and contributes to the improvement of hardness. If the C content is less than 0.35%, the hardenability of steel is insufficient. On the other hand, if the C content exceeds 0.6%, the machinability of the steel deteriorates. Therefore, the C content is 0.35 to 0.6%. The lower limit of the C content is preferably 0.36%, more preferably 0.38%. The upper limit of the C content is preferably 0.5%, more preferably 0.45%.

Si: less than 0.50% Silicon (Si) is contained in steel as an impurity. Furthermore, positively containing Si has the effect of suppressing the formation of coarse cementite. On the other hand, when the Si content is 0.50% or more, the thermal conductivity of the steel decreases and sufficient seizure resistance cannot be obtained. Therefore, the Si content is less than 0.50%. In order to remarkably obtain the effect of suppressing the formation of coarse cementite, the Si content is preferably 0.05% or more, and more preferably 0.10% or more. The upper limit of the Si content is preferably 0.40%, more preferably 0.30%.

Mn: less than 0.80% Manganese (Mn) is contained in steel as an impurity. Further, if Mn is positively contained, it has the effect of enhancing the hardenability of steel. It also has the effect of suppressing the recovery of dislocations. On the other hand, when the Mn content is 0.80% or more, the thermal conductivity of the steel decreases and sufficient seizure resistance cannot be obtained. Therefore, the Mn content is less than 0.80%. In order to remarkably obtain the effect of enhancing hardenability, the Mn content is preferably 0.05% or more, and more preferably 0.10% or more. The upper limit of the Mn content is preferably 0.60%, more preferably 0.50%, further preferably 0.40%.

P: more than 0.02% and 0.1% or less Phosphorus (P), as a solid solution element, hinders the movement of dislocations and improves the dislocation density. Further, P suppresses the grain growth of austenite, thereby refining the former austenite grains of martensite. As a result, work hardening due to sliding is enhanced, and wear resistance is improved. When the wear resistance is improved, seizure can be suppressed. These effects become remarkable by adding P together with Mo. If the P content is 0.02% or less, these effects cannot be obtained. On the other hand, if the P content exceeds 0.1%, excessive P segregates at the grain boundaries and the fatigue strength of the steel decreases. Therefore, the P content is more than 0.02% and 0.1% or less. The lower limit of the P content is preferably 0.03%, more preferably 0.04%. The upper limit of the P content is preferably 0.08%, more preferably 0.06%.

S: less than 0.05% Sulfur (S) is contained in steel as an impurity. Further, when S is positively contained, a sulfide-based inclusion is formed and the machinability of steel is improved. On the other hand, when the S content is 0.05% or more, the hot workability deteriorates. Therefore, the S content is less than 0.05%. The lower limit of the S content is preferably 0.01%. The upper limit of the S content is preferably 0.04%.

Mo: 0.5-3.5%
Molybdenum (Mo) enhances the hardenability of steel. Further, in both the solid solution state and the state of being precipitated as Mo ₂ C, the movement of dislocations is hindered to improve the dislocation density. Moreover, even if Mo is added to steel, the thermal conductivity of steel is not significantly reduced. Since the steel material according to the present embodiment has a low Mn content in order to increase the thermal conductivity, it is possible to achieve both hardenability and high thermal conductivity by replacing Mn with Mo. When Mo is added together with P, Mo contributes to the improvement of the thermal stability of steel and the suppression of grain growth. If the Mo content is less than 0.5%, these effects cannot be sufficiently obtained. On the other hand, if the Mo content exceeds 3.5%, the machinability of the steel deteriorates. Therefore, the Mo content is 0.5 to 3.5%. The lower limit of the Mo content is preferably 0.8%, more preferably 1.2%, further preferably 2.0%. The upper limit of the Mo content is preferably 3.3%, more preferably 3.2%.

Al: 0.005-0.06%
Aluminum (Al) forms a nitride and contributes to the refinement of austenite grains by the pinning effect of the nitride. It was further revealed that Al contributes to the refinement of Mo ₂ C. If the Al content is less than 0.005%, this effect cannot be sufficiently obtained. On the other hand, if the Al content exceeds 0.06%, the machinability of steel deteriorates. Therefore, the Al content is 0.005 to 0.06%. The lower limit of the Al content is preferably 0.01%, more preferably 0.02%. The upper limit of the Al content is preferably 0.055%, more preferably 0.05%.

N: 0.001-0.02%
Nitrogen (N) forms a nitride or a carbonitride, and contributes to the refinement of austenite grains by the pinning effect of the nitride or the carbonitride. If the N content is less than 0.001%, this effect cannot be sufficiently obtained. On the other hand, if the N content exceeds 0.02%, the toughness of steel decreases. Therefore, the N content is 0.001 to 0.02%. The lower limit of the N content is preferably 0.0015%, more preferably 0.002%. The upper limit of the N content is preferably 0.015%.

The balance of the chemical composition of the steel material according to the present embodiment is Fe and impurities. The impurities referred to here are elements mixed from ores and scraps used as raw materials for steel, or elements mixed from the environment of the manufacturing process.

Nb: 0 to 0.05%
The chemical composition of the steel material according to the present embodiment may contain niobium (Nb) instead of part of Fe. Nb forms a nitride or carbonitride, and contributes to the refinement of austenite grains by the pinning effect of the nitride or carbonitride. On the other hand, if the Nb content exceeds 0.05%, the toughness of steel decreases. Therefore, the Nb content is 0 to 0.05%. The lower limit of the Nb content is preferably 0.005%, more preferably 0.010%. The upper limit of the Nb content is preferably 0.04%, more preferably 0.03%.

[Organization]
The steel material according to the present embodiment has a structure in which Mo ₂ C is deposited on at least the surface layer thereof.

When induction hardening is performed on the steel material as a surface hardening treatment, Mo ₂ C refines the prior austenite grains when martensite is formed by induction hardening and improves work hardenability. Induction hardening generally forms martensite only in the surface layer of steel. Therefore, the steel material according to the present embodiment only needs to have Mo ₂ C precipitated on at least the surface layer thereof. The “surface layer” refers to a region having a depth of 500 μm from the surface of the steel material.

When cold rolling is performed on the steel material as a surface hardening treatment, Mo ₂ C suppresses recovery of dislocations due to frictional heat during cold rolling and contributes to improvement of dislocation density. In this case as well, the steel material according to the present embodiment only needs to have Mo ₂ C precipitated on at least the surface layer thereof.

The matrix of the structure of the steel material in which Mo ₂ C is precipitated is not particularly limited. The texture matrix is, for example, ferrite perlite or tempered martensite.

The average particle spacing of Mo ₂ C is 1.0 μm or less. As the average particle spacing of Mo 2 _C is small, indicating that the fine Mo 2 _C is precipitated at high density. If the average particle spacing of Mo ₂ C is larger than 1.0 μm, the above effects cannot be sufficiently obtained. More specifically, even if the chemical composition of the steel material satisfies the above range, when the steel material is induction hardened, the average particle diameter of the former austenite grains cannot be reduced to 10 μm or less, and the steel material is cold-rolled. In this case, the dislocation density cannot be 10 ¹⁵ m ⁻² or more. The average particle spacing of Mo ₂ C is preferably 0.5 μm or less. Although there is no particular meaning to define the lower limit of the average particle spacing of Mo ₂ C, about 0.05 μm is the practical lower limit.

The average particle spacing of Mo ₂ C is measured as follows.

A sample is taken from the steel material. A field emission type electron probe microanalyzer (FE-EPMA) is used to measure a region of 30 μm×30 μm at an acceleration voltage of 15 kV at 500×500 points at 0.06 μm intervals. The measurement point where the Mo concentration is 50 mass% or more is set to 1, and the point where the Mo concentration is less than 50 mass% is set to 0, and binarized to obtain a two-dimensional distribution image of Mo ₂ C. The average particle spacing of Mo ₂ C is derived by the following procedure.
(1) The number of particles P _A per unit area is measured from a two-dimensional distribution image.
(2) The average particle spacing Δ2 on the two-dimensional plane is derived by the following equation.
Δ2=0.500P _A ^−1/2
(References) Sakuma, Nishizawa: The Japan Institute of Metals, 10, (1971), p. 279-289

[Sliding parts]
The sliding component according to the present embodiment is a sliding component manufactured from the above-mentioned steel material and at least a part of which has been induction hardened or cold rolled. Therefore, the sliding component according to the present embodiment has substantially the same chemical composition as the steel material described above.

The sliding component is, for example, a crankshaft.

At least a part of the sliding component according to the present embodiment is subjected to induction hardening or cold roll processing as a surface hardening treatment. When induction hardening is carried out, at least the surface layer structure becomes martensite. When cold rolling is performed, at least the surface layer structure becomes a high dislocation density structure.

The sliding parts are generally subjected to a surface hardening treatment only on the sliding parts. In the sliding component according to the present embodiment, the surface layer structure does not need to be martensite or high dislocation density structure at all parts of the sliding component, and the surface layer structure in at least the sliding part is martensite or high dislocation density structure. If

[Microstructure when induction hardening is performed as surface hardening treatment]
At least the surface structure of the sliding component subjected to induction hardening as the surface hardening treatment is martensite. Furthermore, the average grain size of the former austenite grains of martensite is 10 μm or less. If the average particle size of the prior austenite particles exceeds 10 μm, sufficient work hardenability cannot be obtained. The average particle size of the prior austenite particles is preferably 5 μm or less. The lower limit of the average grain size of the former austenite grains is not specified, but about 1 μm is the practical lower limit.

The average grain size of the former austenite grains of martensite is measured as follows.

Samples are taken from the sliding parts so that the cross section including the direction perpendicular to the induction hardened surface becomes the observation surface. The observation surface is polished and etched with picric acid to expose old austenite grain boundaries. Then, the average particle diameter of the surface layer of the sliding component is calculated by the intercept method. Specifically, a straight line having a total length L is drawn for a region up to 500 μm from the surface, the number n _{L of} crystal grains that cross the straight line is obtained, and the intercept length (L/n _L ) is obtained. The intercept length (L/n _L ) is determined for five or more straight lines, and the arithmetic mean thereof is taken as the average particle size.

[Microstructure when cold rolling is performed as surface hardening treatment]
In the sliding component that has been subjected to cold roll processing as the surface hardening treatment, at least the surface layer structure is a high dislocation density structure having a dislocation density of 10 ¹⁵ m ⁻² or more. Specifically, the structure having a depth within 100 μm from the surface is a high dislocation density structure having a dislocation density of 10 ¹⁵ m ⁻² or more. If the dislocation density in the surface layer portion is less than 10 ¹⁵ m ⁻² , sufficient surface layer nanohardness cannot be obtained. The dislocation density in the surface layer portion is preferably 1.5×10 ¹⁵ m ⁻² or more, more preferably 2.0×10 ¹⁵ m ⁻² or more.

The dislocation density in the surface layer portion is measured as follows.

A sample including a cold-rolled surface is taken from the sliding component, and the cold-rolled surface is mechanically and electrolytically polished. An X-ray generator using Cu-Kα as a radiation source is used to measure four or more diffraction peaks of the ferrite phase. The background of the measured diffraction peak is removed by dedicated analysis software to separate it into Kα1 and Kα2. The full width at half maximum (full width at half maximum) Δ2θ[rad] at each diffraction angle θ[rad] is measured for the separated Kα1 diffraction peak. The full width at half maximum is corrected by measuring the full width at half maximum derived from the apparatus using a single crystal of LaB ₆ (lanthanum hexaboride) and subtracting it from the measured full width at half maximum.

The dislocation density ρ [m ⁻² ] is calculated from the Modified Williamson-Hall equation shown below.

Here, ΔK _FWHM =Δ2θ cos θ/λ, K=2 sin θ/λ, λ is the wavelength of X-ray (0.154 nm), D is the crystallite size, b is the Burgers vector of ferrite (0.249 nm), and O is the higher-order term. Is.

C is a constant indicating crystal anisotropy as described below.

M is a dimensionless coefficient indicating a dislocation array, and here, M=1.

In the equation (*), ΔK _FHHM and K show a linear relationship when the higher-order terms are ignored. By plotting ΔK _FHHM and K, the crystallite size D can be obtained from the intercept and the dislocation density ρ can be obtained from the slope.

[Common structure when induction hardening is performed and when cold rolling is performed]
The sliding component according to the present embodiment preferably has a structure in which carbide is precipitated in at least the surface layer, and the average major axis of the carbide is 10 nm or less.

The type of carbide is not particularly limited, but examples thereof include Mo ₂ C and cementite. If the average major axis of the carbides exceeds 10 nm, abrasion is promoted during sliding, resulting in a decrease in seizure resistance. Although there is no particular meaning to define the lower limit of the average major axis of the carbides, about 1 nm is a practical lower limit.

The average major axis of carbide is measured as follows.

A sample is taken from the sliding part so that the cross section including the direction perpendicular to the surface that has been induction hardened or cold rolled is the observation surface. Five observation fields are observed by a transmission electron microscope (TEM) at 40,000 times and 100,000 times so as to include the surfaces of the sliding parts. The incident direction of the electron beam is the parent phase ferrite <001> direction. The size of the visual field is, for example, 2.0 μm×1.2 μm in the case of 40,000 times, and 0.8×0.5 μm in the case of 100,000 times.

The type of carbide is specified from the electron diffraction image. FIG. 2A and FIG. 2B are diagrams schematically showing precipitation forms of Mo ₂ C and cementite, respectively. As shown in FIG. 2A, if the incident direction of the electron beam is the [100] direction of the matrix ferrite, Mo ₂ C has a rod-like shape extending in the [100] direction of the matrix ferrite. On the other hand, as shown in FIG. 2B, cementite has a disk shape along [110]. Therefore, it is possible to distinguish between Mo ₂ C and cementite from the bright field image. These directly measure the major axis of each carbide. The major axis of carbide is averaged in each visual field, and the average of 10 visual fields is taken as the average major axis of carbide. However, when obtaining the average, those with a major axis of less than 1 nm are not included. This is because it is difficult to determine whether or not it is a carbide.

[Production method]
Hereinafter, an example of the method for manufacturing the steel material and the sliding component according to the present embodiment will be described. The method for manufacturing the steel material and the sliding component according to the present embodiment is not limited to this.

[Method of manufacturing steel materials]
FIG. 3 is a flow chart showing an example of the method for manufacturing a steel product according to the present embodiment. This manufacturing method includes a step of preparing a material (step A1), a step of heating the material (step A2), a step of hot forging the heated material (step A3), and a step of hot forging the material. It includes a step of quenching (step A4) and a step of tempering the quenched material (step A5).

First, a material having the above-described chemical composition is prepared (step A1). For example, a billet is manufactured by melting steel having the above-described chemical composition, and performing continuous casting and slab rolling.

The material is heated to 1000 to 1250°C (step A2). Then, the heated material is hot forged (step A3). By hot forging, the material is processed into the rough shape of sliding parts. The conditions for hot forging are not particularly limited, but the finishing temperature is 900° C., for example. After hot forging, the material is allowed to cool.

The hot forged material is quenched from a temperature of 850° C. or higher (step A4). More specifically, the hot forged material is heated to a temperature of 850° C. or higher and then oil-cooled or water-cooled.

The quenched material is tempered at a temperature of 500 to 650°C (step A5). The holding time is, for example, 5 to 60 minutes. As a result, a structure in which fine Mo ₂ C is deposited can be obtained.

FIG. 4 is a flow chart showing another example of the method for manufacturing a steel material according to the present embodiment. This manufacturing method includes a step of preparing a raw material (step B1), a step of heating the raw material (step B2), and a step of hot forging the heated raw material (step B3).

First, a material having the above-described chemical composition is prepared (step B1). For example, a billet is manufactured by melting steel having the above-described chemical composition, and performing continuous casting and slab rolling.

The material is heated to 1000 to 1250°C (step B2). Subsequently, the heated material is hot forged (step B3). The step of hot forging the material (step B3) includes a main forging step (step B3-1) and a warm working step (step B3-2).

The conditions for the main forging step (step B3-1) are not particularly limited. In the warm working step (step S3-2), a material having a surface reduction rate of 10% or more is added to the material at a temperature of 500 to 650°C. Due to the strain-induced precipitation by warm working, a structure in which fine Mo ₂ C is precipitated can be obtained. The warm working may be carried out continuously after the main forging step is finished, or after the main forging step is finished, it is allowed to cool and then reheated to a temperature of 500 to 650°C and then the warm working is carried out. May be. The warm working is preferably carried out in the final step of the forging step.

[Manufacturing method of sliding parts]
(1) When Induction Hardening is Performed as Surface Hardening Treatment FIG. 5 is a flow chart showing an example of the method for manufacturing the sliding component according to the present embodiment. This manufacturing method includes a step of preparing a steel material (step C1), a step of cutting the steel material (step C2), a step of induction hardening the cut steel material (step C3), and an induction hardened steel material. A low temperature tempering step (step C4).

First, a steel material having the above-described chemical composition and structure is prepared (step C1). The steel material may be manufactured by the manufacturing method illustrated in FIG. 3 or FIG. 4, or may be manufactured by another manufacturing method.

The steel material is cut into the shape of a sliding part (step C2). Note that the cutting process (step C2) may not be performed depending on the shape of the sliding component.

The steel material is induction hardened (step C3). The ultimate temperature during high frequency heating is preferably 850 to 1000°C. The cooling rate during quenching is preferably 1 to 10°C/sec. The induction hardening may be performed on only a part of the steel material. Specifically, induction hardening may be performed only on the sliding portion.

The induction-hardened steel material is tempered at a low temperature (step C4). The low temperature tempering temperature is not particularly limited, but is, for example, 150 to 200°C. If there is no risk of quench cracking, the low temperature tempering (step C4) may not be performed.

(2) When cold rolling is performed as the surface hardening treatment FIG. 6 is a flowchart showing another example of the method for manufacturing the sliding component according to the present embodiment. This manufacturing method includes a step of preparing a steel material (step D1), a step of cutting the steel material (step D2), and a step of cold-rolling the cut steel material (step D3).

First, a steel material having the above-described chemical composition and structure is prepared (step D1). The steel material may be manufactured by the manufacturing method illustrated in FIG. 3 or FIG. 4, or may be manufactured by another manufacturing method.

The steel material is cut into the shape of a sliding part (step D2). The cutting process (step D2) may not be performed depending on the shape of the sliding component.

Cold rolling is performed on the steel material (step D3). The conditions of cold roll processing are preferably such that the maximum contact surface pressure (Hertz surface pressure) is 2000 to 4000 MPa and the sliding distance is 1000 to 2500 m. Cold rolling may be performed on only a part of the steel material. Specifically, cold rolling may be performed only on the sliding portion.

The embodiments of the present invention have been described above. According to this embodiment, a sliding component having excellent seizure resistance and a steel material suitable for the sliding component can be obtained.

The embodiments described above are merely examples for implementing the present invention. Therefore, the present invention is not limited to the above-described embodiments, and can be implemented by appropriately modifying the above-described embodiments without departing from the spirit thereof.

Hereinafter, the present invention will be described more specifically with reference to Examples. The invention is not limited to these examples.

Steel having the chemical composition shown in Table 1 was melted by a 150 kg vacuum induction melting furnace (VIM) to obtain a raw material. Each material was heated to 1200 to 1250° C. and hot forged to a length of 50 mm and a length of 150 mm to prepare test materials of test numbers 1 to 21 shown in Table 2. In addition, "-" in Table 1 indicates that the content of the corresponding element is at the impurity level.

Part of the test material was produced by performing a grain refining (precipitation) treatment. For the test material described as "warm working" in the column of "pretreatment" in Table 2, as a grain refining (precipitation) treatment, in the final step of hot forging, processing of 15% at 600°C was performed. Was carried out and forged into a 50 mm square and 150 mm length. For the test material described as "Quenching and tempering" in the same column, as a grain refining (precipitation) treatment, after hot forging, it was heated to 950°C, then quenched, and further tempered at 600°C for 30 minutes. did. The test material described as “as forged” in the same column was not subjected to the grain refining (precipitation) treatment, and was allowed to cool to room temperature after hot forging.

A sample was taken from each test material so that the plane perpendicular to the longitudinal direction became the observation plane. The observation surface was observed with an optical microscope to identify the tissue. The results are shown in the column of "structure before hardening treatment" in Table 2. A sample for TEM observation was prepared from the same sample, and the type of carbide (Mo ₂ C or cementite) was specified from the electron beam diffraction image of the observation surface. The results are shown in the column "Type of carbide before hardening treatment" in Table 2. When the ratio of the number of Mo ₂ C to the total number of carbides is less than 10%, “θ” is described in the same column. When the ratio of the number of Mo ₂ C to the total number of carbides is 10% or more, “Mo ₂ C+θ” is described in the same column.

Further, a sample was sampled so that the surface perpendicular to the longitudinal direction was the observation surface, and the average particle spacing of Mo ₂ C was determined by the method described in the above embodiment. It was confirmed that particles having a Mo concentration of 50% by mass or more were Mo ₂ C by comparing with the result of TEM observation. The results are shown in the column of “Average particle spacing of Mo ₂ C before curing treatment” in Table 2.

As shown in Table 2, each of the test materials subjected to the grain refinement (precipitation) treatment had a structure in which Mo ₂ C was precipitated. On the other hand, no Mo ₂ C was precipitated in the structures of the test materials that were not subjected to the grain refining (precipitation) treatment.

Among the test materials on which Mo ₂ C was deposited, in Test Nos. 3, 11, 13, and 15, the average particle spacing of Mo ₂ C was larger than 1.0 μm. The reason why the average particle spacing in Test No. 3 was large is considered to be because the P content of Steel Type B was too small. It is considered that the large average particle intervals in Test Nos. 11 and 13 were due to the P content and Al content of the steel types F and G being too small. The reason why the average particle spacing in Test No. 15 was large is considered to be because the Al content of the steel type H was too small.

One of two types of surface hardening treatments was selected and carried out for each test material. Specifically, (1) heat treatment equivalent to induction hardening or (2) cold roll working was performed.

(1) Induction hardening Quenching was performed on the test materials of Test Nos. 1 to 15 corresponding to induction hardening. Specifically, a test piece of 12 mm square and 60 mm length is cut out from the surface of the forged material, and each test material is heated until the ultimate temperature of 950° C. is reached. Cooled. This is a cooling rate equivalent to oil cooling. After performing heat treatment equivalent to induction hardening, each test material was low temperature tempered at 150 to 200°C.

(2) Cold Roll Working Cold rolling was performed on the test materials of Test Nos. 16 to 21. Specifically, a test material of 50 mm square and 150 mm length was machined into a shape of φ50 mm and 150 mm length as shown in FIG. 7, and further groove processing of R6.5 mm was added. Cold working with a sliding distance of 2000 m was applied to the groove portion with a cold roll made of an R6 pin. As for the roll processing conditions, the rolling load was adjusted so that the maximum contact surface pressure (Hertz surface pressure) was 3500 MPa or more.

"IH" in the "Surface treatment" column of Table 3 indicates that a heat treatment equivalent to induction hardening was performed as the surface hardening treatment. “CR” in the same column indicates that cold roll processing was performed as the surface hardening treatment.

A plurality of blocks having dimensions of 2.7 mm×2.0 mm×2.0 mm were sampled from each of the test materials subjected to the surface hardening treatment. At this time, a block was sampled from the surface of the test material so that the surface of the test material subjected to the surface hardening treatment had a surface of 2.7 mm×2.0 mm. The X-ray full width at half maximum for measuring the dislocation density, and the wear test described later and the measurement of the nano hardness after the wear test are performed on a surface of 2.7 mm×2.0 mm (hereinafter referred to as “test surface”). However, in order to remove the unevenness of the test surface, paper polishing and polishing using colloidal silica were performed. Further, electrolytic polishing was carried out in order to remove the processed layer generated by the paper polishing and the polishing using colloidal silica. The depth of the surface removed by paper polishing, polishing using colloidal silica, and electrolytic polishing was about 3 μm.

Then, the structure is specified for the surface in a vertical relationship with the test surface of the block after the polishing, the carbide is specified, and the average particle diameter of the former austenite grains is measured (only the test material subjected to heat treatment equivalent to induction hardening). The average major axis of the carbide and the Vickers hardness were measured.

The identification of the structure, the measurement of the average grain size of the prior austenite grains, and the measurement of the Vickers hardness were observed at a depth of 500 μm from the test surface after polishing. The identification of the carbide and the measurement of the average major axis of the carbide were performed in a visual field including the test surface after polishing. The tissue was identified by an optical microscope. The results are shown in the "Organization" column of Table 3. The type of carbide (Mo ₂ C or cementite) was specified from an electron diffraction image after preparing a sample for TEM observation. The results are shown in the column "Type of carbide" in Table 3. When the ratio of the number of Mo ₂ C to the total number of carbides is less than 10%, “θ” is described in the same column. When the ratio of the number of cementites to the total number of carbides was less than 10%, “Mo ₂ C” was described in the same column. In cases other than the above, “Mo ₂ C+θ” is described in the same column.

The average grain size of the prior austenite grains, the average major axis of carbides, and the dislocation density were determined by the method described in the above embodiment. The results are shown in Table 3 in the columns of "old γ grain size", "average major axis of carbides", and "dislocation density", respectively. "-" in these columns indicates that the same measurement was not performed. Furthermore, the Vickers hardness was measured. The results are shown in the column of "Vickers hardness" in Table 3.

As shown in Table 3, the structure of the surface layer of the block subjected to the heat treatment corresponding to the induction hardening was all martensite. As shown in Tables 2 and 3, all the blocks (test numbers 3, 5, 7, 9, 11, 13, and 15) subjected to the grain refinement (precipitation) treatment were the average grain size of the old austenite grains. The diameter was 10 μm or less. Among these, the blocks having appropriate chemical compositions (test numbers 5, 7, and 9) did not precipitate cementite and had an average major axis of carbides of less than 10 nm.

Among the blocks subjected to the cold roll processing, the blocks subjected to the grain refinement (precipitation) treatment (Test Nos. 17, 19, and 21) all have a surface layer structure with a dislocation density of 10 ¹⁵ m ⁻² or more. It was a high dislocation density structure.

[Seizure resistance evaluation]
In order to evaluate the nanohardness after sliding for a long time, a wear test was performed on each block and the nanohardness after the wear test was measured as described below.

The pin-on-disc type wear test shown in FIG. 8 was performed on each block.

More specifically, the pin-on-disc type wear test was carried out as follows. The 800th emery paper 20 was stuck on the surface of the rotary disk 10 of the pin-on-disk tester. Then, while the test surface 30a of the test piece 30 as a block was pressed against the emery paper 20 with a surface pressure of 0.3 MPa, the rotary disk 10 was rotated so that the sliding distance became 2000 m.

The above test is carried out in two environments: an environment (lubrication environment) where lubricant oil (5W-20, oil temperature: 130° C.) is applied to the emery paper 20 and an environment where no lubricant oil is applied (non-lubrication environment). did.

After the abrasion test, the nano hardness of the test surface was measured by the nano indexation method. In both the lubricated environment and the non-lubricated environment, the nanohardness at a position within 20 nm from the test surface (hereinafter, referred to as “surface nanohardness”) was set to 14 GPa or more.

The nano hardness was measured using a nano indenter manufactured by Agilent Technology, XP/DCM type. As shown in FIG. 9, the nano indenter probe 40 was brought into contact with the test surface 30 a of the test piece 30 for measurement. More specifically, a Burcovitch type needle made of diamond was pushed into the test surface 30a by the continuous rigidity method (CSM method). The conditions of the continuous stiffness system were that the frequency was 45 Hz, the amplitude was 2 nm, and the indentation depth was 250 nm.

The distributions of the nano hardnesses of the test numbers 1 to 5 after the abrasion test are shown in FIGS. 10A to 10E, respectively.

As shown in FIG. 10A, in Test No. 1, the nano hardness was significantly reduced in the non-lubricated environment as compared to the lubricated environment. It is considered that this is because the Mn content of the steel type A was high, the thermal conductivity was low, and the influence of frictional heat was large.

As shown in FIGS. 10B and 10C, in Test Nos. 2 and 3, although the difference between the lubrication environment and the non-lubrication environment was small, the nanohardness was low as a whole and was less than 14 GPa. It is considered that this is because the thermal stability was low because the P content of steel type B was low.

As shown in FIG. 10D, in Test No. 4, the nano hardness was relatively high, and the difference between the lubrication environment and the non-lubrication environment was small. However, in the non-lubricated environment, the surface nanohardness was less than 14 GPa. It is considered that this is because the coarse carbide, that is, the plate-like or needle-like coarse cementite promoted the wear due to sliding.

As shown in FIG. 10E, in test number 5, the nano hardness was high and the difference between the lubrication environment and the non-lubrication environment was small. Furthermore, the surface layer nanohardness was 14 GPa or more in both the lubricated environment and the non-lubricated environment.

Table 4 and FIG. 11 show surface nanohardnesses of test numbers 1 to 21.

In the test numbers 5, 7, 9, 17, 19, and 21, the surface layer nano-hardness after sliding was 14 GPa or more in both the lubricated environment and the non-lubricated environment. It is considered that these have excellent seizure resistance.

In Test Nos. 1, 2, 6, 8, 10, 12, and 14, the surface nanohardness after sliding was less than 14 GPa in both the lubricated environment and the non-lubricated environment. Test No. 4 had a surface layer nanohardness of less than 14 GPa after sliding in a non-lubricated environment. It is considered that, since the old austenite grains were coarse and needle-like or plate-like coarse cementite was deposited, wear due to sliding was promoted.

In Test Nos. 3, 11, 13, and 15, the surface layer nanohardness after sliding was less than 14 GPa in both the lubricated environment and the non-lubricated environment. It is considered that the former austenite grains were fine, but the acicular or plate-like coarse cementite was deposited, and thus the abrasion due to sliding was promoted.

In the test numbers 16, 18, and 20, the surface layer nanohardness after sliding was less than 14 GPa in both the lubricated environment and the non-lubricated environment. Since these had a low dislocation density, it is considered that wear due to sliding was promoted.

Claims (6)

Steel material for sliding parts ,
The chemical composition is% by mass,
C: 0.35-0.6%,
Si: less than 0.50%,
Mn: less than 0.80%,
P: more than 0.02% and 0.1% or less,
S: less than 0.05%,
Mo: 0.5-3.5%,
Al: 0.005-0.06%,
N: 0.001-0.02%,
Nb: 0 to 0.05%,
Remainder: Fe and impurities,
At least the surface layer of the steel material has a structure in which Mo ₂ C is deposited,
A steel material for sliding parts , wherein the average particle spacing of Mo ₂ C is 1.0 μm or less. The steel material for sliding parts according to claim 1,
The chemical composition is% by mass,
Nb: 0.005-0.05%,
Steel material for sliding parts containing. Sliding parts,
The chemical composition is% by mass,
C: 0.35-0.6%,
Si: less than 0.50%,
Mn: less than 0.80%,
P: more than 0.02% and 0.1% or less,
S: less than 0.05%,
Mo: 0.5-3.5%,
Al: 0.005-0.06%,
N: 0.001-0.02%,
Nb: 0 to 0.05%,
Remainder: Fe and impurities,
At least the surface structure of the sliding component is martensite,
Ri average particle diameter der less 10μm of prior austenite grains of the martensite,
At least the surface layer of the sliding component has a structure in which carbide is precipitated,
A sliding part having an average major axis of the carbide of 10 nm or less . Sliding parts,
The chemical composition is% by mass,
C: 0.35-0.6%,
Si: less than 0.50%,
Mn: less than 0.80%,
P: more than 0.02% and 0.1% or less,
S: less than 0.05%,
Mo: 0.5-3.5%,
Al: 0.005-0.06%,
N: 0.001-0.02%,
Nb: 0 to 0.05%,
Remainder: Fe and impurities,
Wherein at least the surface layer of the tissue of the sliding parts, the dislocation density of Ri 10 ¹⁵ m ^-2 or more high-dislocation-density tissue der,
At least the surface layer of the sliding component has a structure in which carbide is precipitated,
A sliding part having an average major axis of the carbide of 10 nm or less . A method of manufacturing a steel material for sliding parts according to claim 1,
The chemical composition is C: 0.35 to 0.6%, Si: less than 0.50%, Mn: less than 0.80%, P: more than 0.02% and 0.1% or less, S in mass%. : Less than 0.05%, Mo: 0.5 to 3.5%, Al: 0.005 to 0.06%, N: 0.001 to 0.02%, Nb: 0 to 0.05%, balance : A step of preparing a material that is Fe and impurities,
A step of heating the material to 1000 to 1250° C.,
A step of hot forging the heated material,
Quenching the hot forged material from a temperature of 850° C. or higher;
And a step of tempering the quenched material at a temperature of 500 to 650°C, a method for producing a steel material for sliding parts . A method of manufacturing a steel material for sliding parts according to claim 1,
The chemical composition is C: 0.35 to 0.6%, Si: less than 0.50%, Mn: less than 0.80%, P: more than 0.02% and 0.1% or less, S in mass%. : Less than 0.05%, Mo: 0.5 to 3.5%, Al: 0.005 to 0.06%, N: 0.001 to 0.02%, Nb: 0 to 0.05%, balance : A step of preparing a material that is Fe and impurities,
A step of heating the material to 1000 to 1250° C.,
A step of hot forging the heated material,
The hot forging step is a method for manufacturing a steel material for sliding parts , which includes a warm working step in which a surface reduction rate of 10% or more is applied to the material at a temperature of 500 to 650°C.

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