JP2003147481A - Non-heatteated high strength and high toughness steel for forging, method of producing the steel, and method of producing forging - Google Patents

Abstract

PROBLEM TO BE SOLVED: To provide a non-heattreated high strength and high toughness steel for forging, a method of producing the steel, and to provide a method of producing a forging. SOLUTION: Steel for forging has a composition in which the contents of C, Si, Mn, Al, N, V, Nb, Mg, Zr, Cr, Ni, Mo, Cu, Ti, B, S, Pb, Ca, and Bi are prescribed to bar steel rolling, and, directly after that, or after reheating, cooling the steel from a temperature of <=1,350 to >=900 deg.C at a cooling rate of 15 to 60 deg.C/sec to form a structure essentially consisting of martensite is subjected to hot forging. In this case, the steel is heated to an Ac3 point or higher to <=900 deg.C and hot forging giving working of >=0.3 in a logarithmic strain is performed in the temperatuer region of an unrecrystallization upper limit temperature or lower to >=700 deg.C for at least one or more times, and the steel is cooled in the temperature region of an Ar3 point or lower to >=500 deg.C at a cooling rate (CR) shown by the following inequality (1) to obtain a fine structure having a mean grain diameter of <=10 μm: 0.1 deg.C/sec<=CR<=(2.5ε+1) deg.C/ sec (1) (ε: the logarithmic strain given in an unrecrystallization temperature region).

Description

Detailed Description of the Invention

[0001]

BACKGROUND OF THE INVENTION 1. Field of the Invention The present invention relates to a forging steel, a method for producing the same, and a method for forging a forged product. More specifically, hot forging is used as a material used for parts of automobiles, construction machines and various industrial machines. The present invention relates to a forging steel having excellent strength and toughness without subsequent heat treatment, a method for producing the same, and a method for forging a forged product.

[0002]

2. Description of the Related Art Conventionally, hot forgings for machine structures are generally made by hot forging a medium carbon steel or low alloy steel material and then reheating, quenching and tempering, that is, tempering treatment. It was provided for use by imparting strength and toughness according to the application. However, the above heat treatment requires a large amount of heat energy cost, and the manufacturing cost is inevitably increased due to an increase in processing steps, an increase in work in process, and the like. Therefore, in recent years, in the production of hot forgings for machine structures, various non-tempered hot forging steels and non-tempered steels have been used to simplify the manufacturing process, in particular, to omit the tempering treatment after hot forging. A method of manufacturing a tempered hot forged product has been proposed. Most of such conventional non-heat treated hot forging steels contain a small amount of V, Nb,
Precipitation hardening type non-heat treated steel to which so-called precipitation hardening type alloying elements such as Ti and Zr are added, in which these are precipitated in the cooling step after hot forging, and high strength is obtained by the precipitation hardening. Is.

For example, in Japanese Examined Patent Publication No. 58-2243, a slight amount of V is added to medium carbon steel, which is heated to a temperature of 1100 ° C. or higher and die-forged, and then 10 to 500 ° C. By air cooling at a cooling rate of 100 ° C / min,
A method for producing a non-heat treated forged product having a ferrite / pearlite structure in which fine V carbonitrides are precipitated in ferrite is described. However, when such a precipitation hardening type non-heat treated steel is used, as described above, 1000 to 1100
It is necessary to heat to a high temperature of ℃ or more,
If ordinary forging is carried out as it is, the crystal grains are significantly coarsened even in the forged product, so that sufficient toughness cannot be obtained.

In order to solve such problems, the addition amount of the precipitation hardening type element is made as small as possible in the material steel and the forging method (for example, JP-A-55-82750).
Low C and high Mn content (for example, JP-A-54-12125)
Japanese Patent Laid-Open Publication No. 56-169723), controlling the type of precipitate (for example, Japanese Patent Laid-Open No. 56-38448), and refining crystal grains by controlled cooling (for example, Japanese Laid-Open Patent Publication No. 56-169723). Although conventionally proposed, it is not easy to obtain a non-heat treated hot forged product excellent in both strength and toughness.

[0005]

SUMMARY OF THE INVENTION It is an object of the present invention to provide a non-heat treated hot forging steel excellent in both strength and toughness, a method for producing the same, and a method for forging a forged product.

[0006]

SUMMARY OF THE INVENTION In order to solve the above problems, the present invention has the following gist. (1) C: 0.1 to 0.8% and Si: 0.
05-2.5%, Mn: 0.2-3%, Al: 0.00
5 to 0.1%, N: 0.001 to 0.02%,
Furthermore, V: 0.05-0.5%, Nb: 0.005-
0.1% of 1 or 2 is contained, the balance is Fe and unavoidable impurities, and martensite is 9 in area ratio.
Non-heat treated high strength characterized by containing 5-100%
High toughness forging steel. (2) In the component (1), in mass%, Mg: 0.0001
-0.005%, Zr: 0.0001-0.005% 1 type or 2 types are added, The non-heat treated high strength and high toughness forging steel. (3) In the component of (1) or (2), in mass%, Cr:
0.05-3%, Ni: 0.05-3%, Mo: 0.0
5 to 3%, Cu: 0.01 to 2%, Ti: 0.003 to
0.05%, B: 0.0005 to 0.005% of 1 type or 2 or more types of non-heat treated high strength and high toughness forging steel. (4) In the component according to any one of (1) to (3), in mass%, S: 0.01 to 0.3%, Pb: 0.03 to.
0.3%, Ca: 0.001 to 0.05%, Bi: 0.
Non-heat treated high strength / high toughness forging steel, characterized by adding one or more of 03 to 0.3%. (5) Immediately after rolling a steel bar made of the component described in any one of (1) to (4) above or 900 to 135
Reheat to 0 ℃, 900 ~ 1350 ℃ to room temperature ~ 30
A method for producing a non-heat treated high strength / high toughness forging steel, which comprises cooling to 0 ° C. at 15 to 60 ° C./sec to obtain steel containing martensite in an area ratio of 95 to 100%. (6) When hot forging the steel according to any one of (1) to (4), the steel is heated to an Ac ₃ point or higher and 950 ° C. or lower to give a work of 0.3 to 3 by logarithmic strain. Non-heat treated high strength / high toughness forging, characterized in that hot forging is performed at least once at a temperature not lower than the upper limit of unrecrystallization temperature of 700 ° C. or more to obtain a structure composed of ferrite and pearlite having an average crystal grain size of 10 μm or less. Method of manufacturing goods. (7) After forging, the temperature range not higher than Ar _{3 and} not lower than 500 ° C. is cooled at a cooling rate (CR) represented by the following formula (1). Manufacturing method of toughness forged products.

0.1 ° C./sec≦CR≦(2.5ε+1)° C./sec (1) (ε: Logarithmic strain given in the non-recrystallization temperature range) (8) Tensile strength is 800 to 1300 MPa The method for producing a non-heat treated high strength / high toughness forged product according to (6) or (7). (9) The non-heat treated high strength according to any one of (6) to (8), characterized in that the yield ratio is 0.7 to 0.95.
Manufacturing method of high toughness forged products.

[0008]

BEST MODE FOR CARRYING OUT THE INVENTION The present invention will be described in detail below. The technical idea that forms the basis of the present invention is as follows. It has been known that in order to obtain a forged product having excellent strength and toughness, the metal structure of the forged product should be made fine. To refine the final structure, γ
(Austenite) is strained by hot forging to make it finer by recrystallization, and dislocations that are normally reduced by recrystallization by lowering the forging temperature and forging at the non-recrystallization temperature until transformation There is a method of making it remain and increasing the nucleation rate. Conventionally, forging in the recrystallization temperature range, that is, forging at a high temperature has a smaller reaction force, and a smaller reaction force makes it easier to improve the forging accuracy. It was a premise to miniaturize. The present inventors have found that by performing forging in a non-recrystallization temperature range which has not been conventionally used for forging, the structure is dramatically refined and the material is improved.

On the other hand, heating at the time of forging is advantageous from the viewpoint of microstructure refinement because low temperature heating has a smaller γ grain size. However, when forging steel with a normal coarse ferrite + pearlite structure is heated at low temperature, due to the influence of the previous structure,
First, since the pearlite portion becomes γ and the ferrite portion becomes γ later, there is a large variation in grain size, and since the diffusion is not sufficiently performed, the concentration deviation remains. As a result, there is a large variation in the structure after forging, and the material does not improve so much.
Further, at low temperature heating, the elements forming carbonitrides do not form a solid solution so much, and the effect of enlarging the non-recrystallization temperature using the solid solution V and Nb as in the present invention cannot be expected. Similarly, strengthening due to work-induced precipitation cannot be expected. As a means of breaking this, the present inventors have found that by adjusting the pre-forging structure to be mainly martensite, a uniform and fine structure can be obtained even at low temperature heating. By making the pre-forging structure a structure mainly composed of martensite, which is a uniform structure with no concentration deviation, a uniform γ can be obtained even at low temperature heating, and it is effective for the refinement of the final structure after forging. Further, in martensite, since the element forming a carbonitride remains in the state of solid solution, when it is heated at a low temperature, the solid solution amount in γ is the conventional ferrite-pearlite structure forging steel. Much more than when heated at low temperature. Therefore, solid solution V, Nb
It is possible to enhance the non-recrystallization temperature by using and to strengthen by processing-induced precipitation. These effects are effective regardless of the structure after forging, even if the structure after forging is ferrite-pearlite, bainite or martensite.

The reasons for limiting the present invention will be described below.

C is an element effective for strengthening steel, but if it is less than 0.1%, sufficient strength cannot be obtained. on the other hand,
If added too much, the toughness will decrease, so the upper limit of the amount added is 0.8%.

Si is effective as a strengthening element for steel, but if it is less than 0.05%, it is not effective. On the other hand, if added too much, the toughness and machinability will decrease, so the upper limit of the addition amount is made 2.5%.

Mn is an element effective for strengthening steel,
If it is less than 0.2%, a sufficient effect cannot be obtained. On the other hand, if added too much, the toughness and machinability deteriorate, so the upper limit of the amount added is 3%.

Al is an effective element for deoxidizing steel and refining crystal grains, but if it is less than 0.005%, it has no effect. On the other hand, if added excessively, the machinability will decrease, so the upper limit of the addition amount is made 0.1%.

N is an element necessary for forming V carbonitrides and Nb carbonitrides for precipitation strengthening, but if less than 0.001%, a sufficient effect cannot be obtained. On the other hand, if added too much, the toughness deteriorates, so the upper limit of the addition is made 0.02%.

V has an effect that solid solution atoms delay dislocation recovery and recrystallization. That is, it is an element that widens the non-recrystallization temperature region to the high temperature side and facilitates the non-recrystallization region forging. Further, after the non-recrystallization rolling, a carbonitride of V is finely precipitated in the entangled portion of dislocations, and the strength is increased by so-called work-induced precipitation, which is an effective element. In order to enjoy these effects, addition of 0.05% or more is necessary. On the other hand, if added excessively, the toughness deteriorates, so the upper limit of the addition amount is made 0.5%.

Like V, Nb is an element necessary for facilitating non-recrystallization and strengthening precipitation, but if it is less than 0.005%, a sufficient effect cannot be obtained. On the other hand, if added excessively, the toughness deteriorates, so the upper limit of the added amount is made 0.1%.

Both Mg and Zr are oxides and sulfides,
Alternatively, it is an element that forms these composites and has an effect of suppressing coarsening of austenite during heating, and also suppresses growth during ferrite transformation, so that it is effective for microstructure refinement. Further, since these oxides become MnS precipitation nuclei, machinability is also improved. In either case, if it is less than 0.0001%, there is no effect, and if it exceeds 0.005%, the toughness deteriorates, so the upper limit of the addition amount is made 0.005%.

Cr, Ni, Mo and Cu are all elements that increase the strength without impairing the toughness when added in appropriate amounts. Cr, Ni, and Mo all have no effect at less than 0.05%, and toughness deteriorates significantly at more than 3%. Therefore, the lower limits of their addition amounts are 0.05% and
The upper limit is 3%. If Cu is less than 0.01%, its effect is not provided, and if it exceeds 2%, toughness is significantly deteriorated. Therefore, the lower limits of the amounts added are 0.01% and 2%, respectively.

Ti produces nitrides and carbides. Since the nitride remains as a solid solution at high temperatures, it is effective in preventing austenite coarsening during heating. Further, the carbide is finely dispersed and is effective for precipitation strengthening. If it is less than 0.003%, these effects do not appear, and if it exceeds 0.05%, the toughness deteriorates, so the lower limit of the addition amount is made 0.003% and the upper limit is made 0.05%.

B is an element that increases hardenability. It is an element effective in increasing the hardenability by increasing the hardenability, preventing the formation of coarse proeutectoid ferrite, and promoting the refinement of the structure. If it is less than 0.0005%, these effects do not appear, and if it exceeds 0.005%, the toughness deteriorates, so the lower limit of the addition amount is made 0.0005% and the upper limit is made 0.005%.

S, Pb, Ca and Bi are all elements that improve machinability. In either case, addition of an excessive amount has no effect, and addition of an excessive amount deteriorates toughness, so S is 0.
01% or more and 0.3% or less, Pb is 0.03% or more and 0.1
The addition amount is limited to 3% or less, Ca to 0.001% to 0.05%, and Bi to 0.03% to 0.3%.

When the area ratio of martensite is less than 95%, the uniformity of the structure and the uniformity of the concentration at the time of reheating are insufficient and the strength and toughness deteriorate. Also, Nb and V during reheating
Since the amount of solid solution is reduced, the effect of the present invention cannot be fully enjoyed.
For the above reasons, the forging steel of the present invention contains martensite in an area ratio of 95% or more. As the other structure, the effect of the present invention can be obtained even if one or more of ferrite, pearlite, and bainite are contained in an area ratio of 5% or less. On the other hand, those containing 100% martensite in area ratio have a uniform structure and concentration upon reheating and can secure a sufficient amount of solid solution of Nb and V.
Most preferred.

Next, a method for manufacturing the forging steel of the present invention will be described.

Quenching from a high temperature is necessary in order for the forging material to have a structure mainly composed of martensite. The heating temperature before the start of quenching was set to 900 ° C. or higher in order to secure the quenching start temperature at 900 ° C. or higher, while the heating temperature was set to 1350 ° C. or lower because it is difficult to heat it further. is there.

The quenching start temperature is set to 1350 ° C. or lower because quenching from a temperature higher than that is difficult. On the other hand, the reason for setting the temperature to 900 ° C. or higher is that if the temperature is lower than that, it is not sufficiently baked and ferrite, pearlite, bainite, etc. are partially formed in an area ratio of 5% or more. On the other hand, the lower limit of the cooling rate at the time of quenching is set to 15 ° C./sec, because at a cooling rate slower than that, ferrite, pearlite, bainite, etc. are partially formed in an area ratio of 5% or more. The upper limit of the cooling rate is set to 60 ° C./sec because it is difficult to cool at a cooling rate higher than that. The cooling method may be water cooling or oil cooling, but is not limited thereto.

The cooling stop temperature is set to room temperature or higher because it is difficult to keep the temperature below room temperature during operation, and is set to 300 ° C. or lower at 300 ° C. or lower because martensite transformation is completed.

Next, the morphology of the structure of the forged product of the present invention will be described.

Usually, in recrystallization γ, dislocations in grains are arranged by recrystallization and the dislocation density is low. Therefore, most of the transformation starts from the γ grain boundary and grows toward the inside of the grain. Further, as long as it is recrystallized γ, the number of transformation nuclei generated per unit area of grain boundary has a substantially constant value. Therefore, the number of grains of the structure after transformation is almost proportional to the area of the γ grain boundary per unit volume, and the smaller the γ grain size after recrystallization, the finer the structure after transformation. On the other hand, in the non-recrystallized γ, the dislocations in the grains are high because the dislocations have not been rearranged by recrystallization. As a result, the transformation starts not only at the grain boundaries but also inside the grains. Further, the effect of processing remains on the grain boundaries, and the number of transformation nucleation per unit area of the grain boundary is larger than that of recrystallization γ. Therefore, a fine transformation structure can be obtained even from coarse γ. The transformation structure obtained by transformation from unrecrystallized γ is ferrite +
It can be roughly classified into pearlite, bainite, and martensite, but all have an average crystal grain size of 10 μm or less. However, depending on the cold speed, it becomes a mixed tissue of these tissues,
Since the toughness is significantly deteriorated, ferrite + pearlite steel is used by the cold speed control described later. The average crystal grain size described here is a crystal grain size serving as a unit of fracture, and refers to the average grain size of ferrite in the case of ferrite + pearlite and the average packet size in the case of bainite and martensite. On the other hand, it is known that when the particle size becomes fine, the strength, toughness, yield ratio, and elongation improve, but when the average particle size is 10 μm or less, these effects become remarkable. If further effect is desired, the average particle size is 5μ
It is desirable to be less than m. On the other hand, although the lower limit of the average particle size is not particularly defined, it is preferably 2 μm or more from the viewpoint of forging cost.

In the present invention, the average particle diameter is 3 to 200 times to 1000 times at the 1/4 t cross section thickness by an optical microscope.
It is defined as a value obtained by observing 5 fields of view and cutting.

The lower limit of the tensile strength is set to 800 MPa in terms of weight reduction of the forged product. On the other hand, when the pressure exceeds 1300 MPa, the toughness is remarkably reduced and the cutting life and the die life are also remarkably reduced. Therefore, the upper limit was made 1300 MPa or less.

Further, the yield ratio is 0.
The lower limit was limited to 7. On the other hand, even if the yield ratio is increased to 0.95 or more, the improvement in fatigue strength saturates, so the upper limit is 0.95.
Limited to.

Next, a method for manufacturing a forged product will be described.

The heating temperature is set to an Ac ₃ point or higher because it is necessary to be in the γ single phase during forging. Further, excessive heating promotes coarsening of γ grains and precipitation of V or Nb during heating, so the upper limit was set to 950 ° C. In addition, Ac ₃ points are (2)
It is defined as the value obtained by the formula.

Ac3 = 910-203 (C) ^{1 /} 2-15.2 (Ni) +44.7 (Si) +104 (V) +31.5 (Mo) +13.1 (W) (2) Unrecrystallized The upper limit temperature is defined as the value obtained by the equation (3). The formula (3) uses V for Nb using a machining master.
It is a regression equation obtained as a result of performing a work hardening test on the contained steel and observing the structure. The expression (3) is a simple expression excluding the term indicating the influence of the workability.

Unrecrystallized upper limit temperature (° C.) = 819 + 61 ((V) +10 (Nb)) ^0.2 (3) The effect of structural refinement by transformation from unrecrystallized γ is the strain given in the unrecrystallized temperature range. Depends on. 0.3 in logarithmic distortion
If the strain is less than 100 μm, the structure cannot be sufficiently refined. Therefore, the lower limit is set to a logarithmic strain of 0.3. If possible, a strain of 0.8 or more is desirable. On the other hand, if the strain is increased, the structure becomes finer, but the effect tends to be saturated. The increase in strain causes an increase in cost due to an increase in forging reaction force and a decrease in die life, so the logarithmic strain is set to 3 or less. In the case of forming by forging multiple times, it may be combined with forging in the recrystallization temperature range. Further, at a forging temperature of 700 ° C. or lower, ferrite is generated before forging, and it becomes work ferrite during forging, which deteriorates toughness, so the lower forging temperature is set to 700 ° C.

The logarithmic distortion described here is the distortion defined by the equation (4). The original thickness average is the average value when the forging direction of the material before forging is the height, and the average finished thickness is the average height after forging. However,
In the case of processing such as extrusion, the formula (5) shall be followed. The average original cross-sectional area is the average cross-sectional area of the surface of the material before forging perpendicular to the forging direction, and the average finished cross-sectional area is the average cross-sectional area after forging.

Logarithmic strain = ln (original thickness height average / finishing thickness height average) (4) Logarithmic strain = ln (original cross-section area average / finishing cross-section average) ... (5) Microstructure by cold speed after forging Although the morphology is different as described above, the cold speed will be described below. In the transformation from unrecrystallized γ, since the nucleation rate is increased, the T-T-T nose shifts to the short time side, and ferrite is easily generated. Therefore, in order to generate ferrite + pearlite, the temperature range of 500 ° C or higher at the Ar ₃ point or less is (1)
It may be cooled at the cooling speed shown in the formula. Expression (1) is an expression obtained from the straight line in FIG. Lower limit of cooling rate is 0.1 ℃ / s
The reason for setting ec is that if the cooling rate is slower than that, sufficient nucleation rate cannot be obtained and the ferrite becomes coarse. On the other hand, the upper limit is set to (2.5ε + 1) ° C./sec because ferrite + pearlite is not generated at a faster cooling rate and a mixed structure with bainite or martensite is formed, resulting in deterioration of toughness. . Further, the reason why the cooling control temperature range is set to the Ar3 point or lower is that it is the temperature at which the transformation starts. On the other hand, the lower limit is set to 500 ° C. because the ferrite + pearlite transformation has already been completed at this temperature.

0.1 ° C./sec≦CR≦(2.5ε+1)° C./sec (1) (ε: Logarithmic strain given in the non-recrystallization temperature range) The Ar ₃ point is determined by the equation (6). Value.

Ar ₃ = 868-396 (C) +24.6 (Si) -58.7 (Mn) -50 (Ni) -35 (Cu) +190 (V) (6)

[0041]

Example From steels having the components shown in Table 1, φ50 × h60
The forging test piece was cut out, heated at high frequency, and subjected to flat plate compression forging in the height direction by applying the method of the present invention and the comparison method shown in Table 2. The distortion in Table 2 was obtained by applying the equation (4). Further, when the method of the present invention was applied and cooled, the particle size, strength, yield ratio and toughness were as shown in Table 2. The temperature control during cooling was performed by a heat-retaining trolley, air blast or water spray cooling. The microstructure of the forged product was photographed at a 1 / 4t position 30 mm away from the center of the forged product, and the average grain size (average packet size) was determined by the cutting method.
The distance from the center is 30 mm to avoid the dead metal part. Mechanical properties are JISA No. 3 tensile test piece and J
The measurement was performed using an IS3 Charpy test piece (width 5 mm). In Table 2, the comparative steels 1, 2 and 14 do not contain the necessary amount of Nb and V which are essential elements of the present invention, so that recrystallization occurs and they have a coarse structure. Therefore, strength, yield ratio, and toughness are low. Since the comparative steels 13 and 15 contain Nb and V more than necessary, the toughness is low. Comparative steel 3,4,5
Indicates that the structure before forging does not contain martensite or its area ratio is low, so the structure after forging becomes coarse and the strength
Yield ratio and toughness are low. Comparative Steel 6 had a too high heating temperature during forging, so the effect of converting the previous structure to martensite was weakened, and the structure became coarse, so the strength, yield ratio, and toughness were low. Since the comparative steel 7 was not forged in the γ single phase at the time of heating because the heating temperature was too low and was forged in the γ + α two-phase state, α was processed and the yield ratio and toughness were low. Comparative Steel 8 has a high working temperature and recrystallization, so that it has a coarse structure and has low strength, yield ratio, and toughness. Since the comparative steel 9 has a low workability, it cannot obtain a sufficient nucleation rate, has a coarse structure, and has low strength, yield ratio, and toughness. Comparative steel 1
In the case of 0, the cold speed after processing was too slow, so a sufficient nucleation rate could not be obtained, resulting in a coarse structure and low strength / yield ratio / toughness. Comparative steel 11 had a low yield rate and toughness because part of bainite was formed because the cold speed after working was too fast. Comparative steel 12 was processed in a state where the processing temperature was too low and α was partially generated during the processing, so α was processed and the yield ratio and toughness were low.

[0042]

[Table 1]

[0043]

[Table 2]

[0044]

[Table 3]

[0045]

According to the present invention, the strength, yield ratio, and
The toughness is improved, and the present invention is effective.

[Brief description of drawings]

FIG. 1 is a diagram showing the influence of logarithmic strain imparted in the non-recrystallized region and cold speed in the temperature range of 500 ° C. to Ar ₃ on the texture formation.

─────────────────────────────────────────────────── ─── Continuation of front page (51) Int.Cl. ⁷ Identification code FI theme code (reference) C22C 38/60 C22C 38/60 (72) Inventor Takashi Fujita 20-1 Shintomi, Futtsu-shi Nippon Steel Corporation Technology Development Headquarters (72) Inventor Toshizo Tarui 20-1 Shintomi, Futtsu City Nippon Steel Corporation Technology Development Headquarters (72) Inventor Naohito Ohno 1 Toyota-cho, Toyota-shi, Aichi Toyota Motor Co., Ltd. (72) Inventor Motohide Mori 1 Toyota Town, Toyota City, Aichi Prefecture Toyota Motor Co., Ltd. (72) Inventor Shigeo Hirota 1 Wano Wari, Arao Town, Tokai City, Aichi Prefecture Aichi Steel Co., Ltd. (72) Invention Person Naoki Iwama 1 Wanowari, Arao-cho, Tokai-shi, Aichi F-term in Aichi Steel Co., Ltd. (reference) 4K032 AA01 AA02 AA03 AA08 AA09 AA12 AA14 AA15 AA16 AA17 AA19 AA20 AA21 AA22 AA23 AA24 AA29 AA31 AA32 AA35 AA36 AA39 CF03 CH04 CH05 CJ02

Claims (9)

[Claims] 1. In mass% C 0.1-0.8% Si 0.05-2.5% Mn 0.2-3% Al 0.005-0.1% N 0.001-0.02 %, And further contains one or two of V 0.05 to 0.5% Nb 0.005 to 0.1%, the balance being Fe and unavoidable impurities, and martensite in area ratio. 95-10
A non-heat treated high strength / high toughness forging steel characterized by containing 0%. 2. The composition according to claim 1, which contains one or two of Mg 0.0001 to 0.005% and Zr 0.0001 to 0.005% by mass.
Non-heat treated high strength / high toughness forging steel. 3. In mass% Cr 0.05 to 3% Ni 0.05 to 3% Mo 0.05 to 3% Cu 0.01 to 2% Ti 0.003 to 0.05% B 0.00055 The non-heat treated high strength / high toughness forging steel according to claim 1 or 2, containing 0.005% of one or more kinds. 4. One or two kinds of S 0.01 to 0.3% Pb 0.03 to 0.3% Ca 0.001 to 0.05% Bi 0.03 to 0.3% in mass%. The non-heat treated high strength / high toughness forging steel according to any one of claims 1 to 3, containing the above. 5. A steel slab comprising the component according to any one of claims 1 to 4 immediately after hot rolling or 900 to 13
Reheat to 50 ℃, 900 ~ 1350 ℃ to room temperature ~ 3
A method for producing a non-heat treated high strength / high toughness forging steel, which comprises cooling to 00 ° C at 15 to 60 ° C / sec to obtain a steel containing martensite in an area ratio of 95 to 100%. 6. When hot forging the forging steel according to any one of claims 1 to 4, heating to an Ac ₃ point or more and 950 ° C. or less, and processing of 0.3 to 3 by logarithmic strain Is a non-heat treated high-strength / high-strength structure characterized by performing hot forging at least once at a temperature not lower than the upper limit of unrecrystallized temperature of 700 ° C. to give a structure consisting of ferrite and pearlite having an average crystal grain size of 10 μm or less. Manufacturing method of toughness forged products. 7. The non-heat treated high strength according to claim 6, wherein after forging, the temperature range not higher than the Ar ₃ point and not lower than 500 ° C. is cooled at a cooling rate (CR) represented by the following formula (1). -Method of manufacturing high toughness forged products. 0.1 ° C./sec≦CR≦(2.5ε+1)° C./sec (1) (ε: logarithmic strain given in the non-recrystallization temperature range) 8. The method for producing a non-heat treated high strength / high toughness forged product according to claim 6, wherein the tensile strength is 800 to 1300 MPa. 9. The method for producing a non-heat treated high strength / high toughness forged product according to claim 6, wherein the yield ratio is 0.7 to 0.95.