JP5512494B2 - High-strength, high-toughness non-tempered hot forged parts and manufacturing method thereof - Google Patents

Description

  The present invention relates to a high-strength, high-toughness non-heat treated hot forged part used as a component such as a connecting rod, a crankshaft, a hub, etc. constituting an engine and undercarriage of a transport machine such as an automobile or a ship, and a manufacturing method thereof .

  2. Description of the Related Art Conventionally, hot forged parts used as components of transport aircraft have been required to increase the yield strength of the parts in order to reduce the weight of the parts. And, in order to increase the strength of hot forged parts while maintaining non-tempered steel, the application of steel in which a predetermined amount of V is added to medium carbon steel is in progress. V precipitates as carbides and carbonitrides during cooling and contributes to yield strength improvement by precipitation strengthening. The precipitation strengthening amount increases as the size of the carbides and carbonitrides becomes finer and as the volume ratio increases.

  For example, Patent Document 1 has a composition in which V: 0.05 to 0.5 mass% and Ca: 0.0005 to 0.02 mass% are added to medium carbon steel, and the structure after hot forging is Ferrite-pearlite, and the area ratio of sulfides having a Ca content exceeding 40% by mass is 30% or less, and the area ratio of sulfides having a Ca content of 0.3 to 40% by mass is 10% or more. Non-tempered hot forged parts are described. However, in Patent Document 1, after forging, air cooling is not performed and sufficient cooling control is not performed. Therefore, the precipitation amount of fine precipitates does not increase, and the precipitation strengthening amount does not improve. As a result, there is a problem that the yield strength of the non-tempered hot forged part does not increase.

  Patent Document 2 discloses a steel having a composition obtained by adding V: more than 0.30% by mass to 0.70% by mass and Ti: 0.003 to 0.050% by mass to medium carbon steel at a sufficiently high temperature. A non-tempered hot forged part and a method for producing the same are described, which are characterized by being forged at a temperature of 300 ° C. and then cooled to an average cooling rate of 0.05 ° C./second or more and less than 2 ° C./second. However, even in Patent Document 2, since only a monotonous cooling pattern is taken, the precipitation strengthening amount of fine precipitates in ferrite is not appropriately controlled, and ferrite having a low precipitation strengthening amount is locally formed. As a result, there is a problem that the yield strength of non-tempered hot forged parts does not increase because ferrite with a low precipitation strengthening rate determines the yield strength.

  Further, in hot forged parts used as components of transport aircraft, high toughness is also required as an important characteristic. Generally, as steel becomes stronger, the toughness decreases.

  In order to solve the above problem, for example, in Patent Documents 3 and 4, one or two kinds of V: 0.05 to 0.5 mass% and Nb: 0.005 to 0.1 mass% are added to the medium carbon steel. Is manufactured, and the steel is heated to Ac3 point or higher and 950 ° C. or lower (Patent Document 3) or Ac3 point or higher and 1350 ° C. or lower (Patent Document 4), and the true strain amount (logarithmic strain amount) is obtained. Hot forging that gives a working of 0.3 to 3 is at least 1 at non-recrystallization upper limit temperature (lower than 900 ° C. in examples) 700 ° C. or higher (Patent Document 3) or lower than 800 ° C. and 700 ° C. or higher (Patent Document 4). The temperature range of Ar 3 point or less and 500 ° C. or more was cooled at 0.1 ° C./sec≦cooling rate (CR) ≦ (2.5ε + 1) ° C./sec (ε: logarithmic strain), and the average particle size Non-refining characterized by obtaining a ferrite-pearlite structure of 10 μm or less Manufacturing method between the forged parts are described.

JP-A-11-350065 Japanese Patent Application Laid-Open No. 07-3386 Japanese Patent Laid-Open No. 2003-147481 JP 2003-147482 A

  However, Patent Documents 3 and 4 have a problem that since low-temperature forging is performed, deformation resistance becomes large and forging becomes difficult, and productivity of non-tempered hot forged parts decreases. Moreover, in Patent Documents 3 and 4, in order to obtain a yield strength of 1000 MPa or more, the amount of true strain during low-temperature forging needs to be 2.3 or more (see Examples), and forging becomes difficult. Productivity of tempered hot forged parts decreases. In addition, since the cooling rate after forging (cooling rate) is 3 ° C./sec (see Examples) at the maximum, the ferrite transformation temperature is not sufficiently lowered, and fine precipitates (carbides) precipitated with the transformation are coarse. And precipitation strengthening cannot be fully utilized. As a result, there is a problem that the yield strength of the non-tempered hot forged part does not increase.

  The present invention has been made in view of the above problems, and its object is to provide a non-tempered hot forged part having high yield strength and toughness without lowering the forging temperature and a method for producing the same. There is.

In order to solve the above-mentioned problem, the high-strength and high-toughness non-tempered hot forged part according to the present invention has C: 0.20 to 0.80 mass%, Si: 0.50 mass% or less, Mn: 0 .40 to 1.00 mass%, P: 0.050 mass% or less, S: 0.050 mass% or less, V: 0.20 to 0.80 mass%, Nb: 0.02 to 0.30 mass% N: 0.0100% by mass or less, with the balance being Fe and inevitable impurities, consisting of a ferrite-pearlite structure in which the austenite grain has a Gf particle size number of 7 or more, in ferrite and pearlite Diameter: The number density of (V, Nb) C precipitates of 10 nm or less is 5000 pieces / μm ² or more.

  According to the said structure, since it has the steel composition which added Nb in addition to V to medium carbon steel, since precipitation NbC pins the coarsening of the austenite recrystallized grain after forging, Gf grain size number is predetermined. A ferrite-pearlite structure having fine prior austenite grains exceeding the value can be obtained. As a result, the toughness of the non-tempered hot forged part is increased. Moreover, when the number density of fine (V, Nb) C precipitates having a predetermined diameter or less precipitated in ferrite is equal to or higher than a predetermined value, the Vickers hardness of the non-tempered hot forged part becomes 400 Hv or higher, yielding. Strength increases.

Further, the method for producing a high-strength, high-toughness non-tempered hot tempered part according to the present invention is a method for producing the high-strength, high-toughness non-tempered hot tempered part , wherein the steel material comprising the composition is A heat treatment step of heat-treating so as to be _TNbC calculated by the following formula (1), and hot-forging the heat-treated steel material at a forging temperature of 1050 ° C. or higher and lower than the heating temperature, the heat A hot forging process in which the amount of true strain during hot forging is 0.3 or more, and after completion of hot forging, rapid cooling is performed at a cooling rate of 3.0 ° C./second or more, and rapid cooling during the rapid cooling is performed. A rapid cooling process with an end temperature of 550 to 720 ° C., and after the rapid cooling, slowly cool at a cooling rate of 0.1 ° C./second or more and 1.5 ° C./second or less, and finish the slow cooling in the slow cooling And a slow cooling step in which the temperature is 400 ° C. or lower.
−14000 / (log ([% C] · 0.01) −7.58) −273 ≦ T _NbC (° C.) ≦ −14000 / (log ([% C] · ([% Nb] −0.01) ) -7.58) -273 (1)
(However, in the formula (1), [% C] and [% Nb] are the contents (mass%) of C and Nb.)

According to the above procedure, in the heat treatment step, by heating to the range of _TNbC , while completely dissolving V, a part (for example, 0.01% by mass or more) of Nb is dissolved. At the same time, the remaining Nb (for example, 0.01% by mass or more) Nb is not dissolved yet, and by securing NbC for pinning recrystallized austenite grains generated during forging, high Vickers hardness (high strength) The refinement of prior austenite grains (higher toughness) can be realized.

  And, in the hot forging process, by performing hot forging so that the true strain amount is 0.3 or more at 1050 ° C. or more and below the heating temperature of the heat treatment process, coarse (V, Nb) Ferrite having a prior austenite grain with a Gf grain size number greater than or equal to a predetermined value, by which the precipitation of C can be suppressed, and a fine recrystallized austenite structure is obtained by pinning of NbC that has been precipitated from the heating stage. A pearlite structure can be obtained.

  Then, in the cooling step, rapid cooling is performed so that the average cooling rate to the quenching stop temperature (550 to 720 ° C.) is 3.0 ° C./s or higher, and after the rapid cooling is finished, the average cooling rate to 400 ° C. or lower is obtained. By performing slow cooling so as to be 0.1 ° C./s or more and 1.5 ° C./s or less, the generation of coarse (V, Nb) C in austenite is suppressed, and ferrite-pearlite is reduced at a low temperature. By transforming, phase interface precipitates can be finely dispersed, and fine (V, Nb) C precipitates can be obtained with a number density of a predetermined value or more, and high strength can be realized. In addition, since it forges at high temperature, there is little deformation resistance and it can manufacture easily.

  The non-tempered hot forged part according to the present invention can have high Vickers hardness (yield strength) and fine prior austenite grains (high toughness). Further, the method for producing a non-tempered hot forged part according to the present invention is a non-tempered material having high Vickers hardness (yield strength) and fine prior austenite grains (high toughness) without lowering the forging temperature. Hot forged parts can be manufactured and productivity is improved.

1 shows the steel structure of a high-strength, high-toughness non-tempered hot forged part according to the present invention. 4 is an optical micrograph of the ferrite-pearlite structure in FIG. 15 is an optical micrograph of a ferrite-pearlite structure in FIG. 1 shows the steel structure of a high-strength, high-toughness non-tempered hot forged part according to the present invention. 4 is a TEM photograph of (V, Nb) C precipitate in FIG. 8 is a TEM photograph of (V, Nb) C precipitate in FIG. It is a flowchart which shows the manufacturing method of the high strength and toughness non-tempered hot forged part which concerns on this invention.

  The high strength and high toughness non-tempered hot forged part (hereinafter referred to as non-tempered hot forged part) according to the present invention will be described.

  The non-tempered hot forged part according to the present invention contains a predetermined amount of C, Si, Mn, P, S, V, Nb, and N, and the balance is Fe and inevitable impurities. It consists of a ferrite-pearlite structure having prior austenite grains, and the number density of fine (V, Nb) C precipitates in the ferrite and pearlite is not less than a predetermined value. The reasons for limiting the composition and structure will be described below.

(C: 0.20 to 0.80 mass%)
C is combined with V and Nb to precipitate (V, Nb) carbide and (V, Nb) carbonitride (both are referred to as (V, Nb) C precipitates), and increases the amount of precipitation strengthening. It is an element that improves the hardness of pearlite and contributes to increasing the yield strength of non-tempered hot forged parts. When the C content is less than 0.20% by mass, the precipitation strengthening amount of the (V, Nb) C precipitate is decreased, and the yield strength is decreased. On the other hand, when the amount of C exceeds 0.80% by mass, ferrite transformation and pearlite transformation are suppressed, so that bainite is formed and the yield strength is lowered. Moreover, C amount becomes like this. Preferably it is 0.30-0.60 mass%, More preferably, it is 0.40-0.50 mass%.

(Si: 0.50 mass% or less)
Si is an element that contributes to increasing the yield strength of non-tempered hot forged parts by solid solution strengthening. When the amount of Si exceeds 0.50% by mass, the hardenability is increased, bainite is formed, and the yield strength of the non-tempered hot forged part is reduced. In addition, 0 mass% may be sufficient.

(Mn: 0.40 to 1.00% by mass)
Mn is an element that contributes to increasing the yield strength of non-tempered hot forged parts by solid solution strengthening. When the amount of Mn is less than 0.40% by mass, the amount of solid solution is small and the yield strength of the non-tempered hot forged part is lowered. On the other hand, when the amount of Mn exceeds 1.00 mass%, hardenability will become high, a bainite will be formed and it will become a factor which the yield strength of a non-tempered hot forging part falls.

(P: 0.050 mass% or less)
P is an element contained as an unavoidable impurity, but if it exceeds 0.050% by mass, it causes embrittlement of the non-tempered hot forged part. In addition, 0 mass% may be sufficient.

(S: 0.050 mass% or less)
S is an element that forms MnS and contributes to the improvement of the machinability of the non-tempered hot forged part. However, when the content exceeds 0.050% by mass, the ductility and toughness of the non-tempered hot forged part are increased. Reduce. In addition, 0 mass% may be sufficient.

(V: 0.20 to 0.80 mass%)
V is an element that strengthens ferrite and pearlite by precipitating as carbide and carbonitride together with Nb in lamellar ferrite in ferrite and pearlite, and contributes to increasing the yield strength of non-tempered hot forged parts. is there. Although V addition has been used in the past, if the V addition amount is increased monotonically without cooling control after forging, phase interface precipitation of V carbide and V carbonitride is less likely to occur during ferrite transformation, yield strength. Decreases.

  When the amount of V is less than 0.20% by mass, the precipitation strengthening amount of V carbide and V carbonitride is low, and the number density of fine (V, Nb) C precipitates is a predetermined value in a non-tempered hot forged part. The strength is lowered. On the other hand, if the amount of V exceeds 0.80% by mass, ferrite-pearlite transformation is suppressed, so that bainite is formed or the transformation temperature becomes too low and phase interface precipitation in ferrite occurs. On the contrary, the strength is reduced. Moreover, V amount becomes like this. Preferably it is 0.35-0.80 mass%, More preferably, it is 0.45-0.80 mass%.

(Nb: 0.02-0.30 mass%)
Nb is an element that strengthens ferrite and pearlite by precipitating as carbide and carbonitride together with V in lamellar ferrite in ferrite and pearlite, and contributes to increasing the yield strength of non-tempered forged parts. Further, Nb partially exists as NbC in an insoluble state even in a high temperature range where V is completely dissolved. NbC acts as pinning particles for the austenite recrystallized grains, contributing to refinement of the recrystallized austenite structure. As a result, a tough ferrite-pearlite structure is obtained.
When the content of Nb is less than 0.02% by mass, in the non-tempered forged part, the Gf particle size number of the prior austenite grains becomes less than a predetermined value, and the toughness decreases. On the other hand, when the Nb content exceeds 0.30 mass%, the yield strength and toughness improving effects are saturated.
Therefore, the Nb content is 0.02 to 0.30 mass%.
The Nb content is preferably 0.05 to 0.25% by mass, and more preferably 0.15 to 0.23% by mass.

(N: 0.0100 mass% or less)
N is an element that forms carbonitride by bonding with V or Nb at a high temperature. When the amount of N exceeds 0.0100% by mass, coarse carbonitride is formed. Since the amount of V and Nb that can be finely precipitated by phase interface precipitation is reduced by the amount of N combined in the high temperature region, the analysis strengthening amount is reduced and the yield strength is reduced. Moreover, N amount becomes like this. Preferably it is 0.0090 mass% or less, More preferably, it is 0.0080 mass% or less. In addition, 0 mass% may be sufficient.

(Inevitable impurities)
Inevitable impurities are elements such as Sn and Sb, for example, and are contained within a range that does not hinder the effects of the present invention.

(Old austenite grains: Gf grain number 7 or more)
The prior austenite grains become finer as the Gf grain size number is larger. And the toughness of a non-tempered hot forged part improves, so that a prior austenite grain is fine. Therefore, the Gf particle size number of the prior austenite grains is 7 or more, preferably the Gf particle size number is 8 or more, more preferably the Gf particle size number is 9 or more.

(Ferrite-pearlite structure)
When V and Nb are sufficiently added, the yield strength of ferrite and pearlite formed by diffusion transformation occurring at high temperatures can be greatly increased by precipitation strengthening of (V, Nb) C precipitates precipitated at the phase interface. On the other hand, in the case of a transformation phenomenon occurring at a low temperature such as bainite and martensite, although the tensile strength is improved, the (V, Nb) C precipitate cannot be precipitated at the phase interface, and therefore the yield strength is decreased. For this reason, when bainite or martensite is mixed, the yield strength cannot be secured, so it is necessary to have a ferrite-pearlite structure. Further, by adopting a ferrite-pearlite structure, there is no risk of delayed fracture occurring in non-tempered hot forged parts. With a ferrite-pearlite structure, 95% or more of the area ratio is a ferrite-pearlite structure.

(Number density of (V, Nb) C precipitates in ferrite and pearlite: 5000 / μm ² or more having a diameter of 10 nm or less)
The (V, Nb) C precipitate deposited at the phase interface during the ferrite-pearlite transformation strengthens the structure and contributes to the improvement of the yield strength. The amount of precipitation strengthening increases as the volume fraction occupied by the precipitate increases and as the precipitate particle size decreases. Therefore, since the (V, Nb) C precipitates are preferably as fine as possible, the number density of (V, Nb) C precipitates having a diameter of 10 nm or less is set to 5000 / μm ² or more. The number density is preferably 6000 / μm ² or more of (V, Nb) C precipitates having a diameter of 10 nm or less, more preferably 7000 / μm ² or more of (V, Nb) C precipitates having a diameter of 10 nm or less. It is.

A method for manufacturing a non-tempered hot forged part according to the present invention will be described with reference to FIG.
The method for manufacturing a non-tempered hot forged part according to the present invention includes a heat treatment step S1, a hot forging step S2, a rapid cooling step S3, and a slow cooling step S4. Hereinafter, each step will be described.

(Heat treatment step S1)
Heat treatment process S1 is a process of heating the steel raw material which consists of the said composition so that it may become the range of _TNbC .
Here, T _NbC is −14000 / (log ([% C] · 0.01) −7.58) −273 ≦ T _NbC (° C.) ≦ −14000 / (log ([% C] · ([% Nb] -0.01))-7.58) -273 (1). However, in said Formula (1), [% C] and [% Nb] are taken as each content (mass%) of said C and said Nb. By this process, the total amount of V of the steel material and a part of Nb are dissolved. When the heating temperature is less than the lower limit of the range of _TNbC , Nb does not dissolve in 0.01% by mass or more, and the number density of fine (V, Nb) C composite precipitates is predetermined in the non-tempered forged part. It becomes less than the value, yield strength decreases. On the other hand, when the heating temperature exceeds the upper limit of the range of _TNbC , undissolved Nb becomes 0.01% by mass or less, NbC for pinning the recrystallized austenite grains is not sufficiently secured, and the recrystallized austenite grains are coarse. And toughness decreases.
Moreover, in addition to the range of _TNbC , the upper limit of the heating temperature is preferably less than the melting temperature of steel, and more preferably about 1300 ° C. from the capacity of the equipment.
Here, _TNbC is a temperature derived by equation transformation from the solubility product of NbC (Yoshiyuki Imai, Tetsuo Shono, Iron and Steel, 1966, p.110), and is a temperature below the lower limit of the temperature range. Is heated to a temperature exceeding the upper limit of the temperature range, and non-solid Nb is reduced to 0.01% by mass or less.
Note that the temperature in the heat treatment step is the highest temperature of the workpiece in the heat treatment step. Further, the steel material is produced by, for example, casting and / or forging and extrusion.

(Hot forging process S2)
Hot forging step S2 is a step of hot forging a heat-treated steel material at a forging temperature not lower than 1050 ° C. and not higher than the heating temperature, and the true strain amount during hot forging is not less than 0.3. To do. In the hot forging step S2, a tough microstructure is obtained. Generally, in order to obtain a high toughness microstructure, it is necessary to perform forging in a low temperature range. However, in the present invention, NbC precipitated in a high temperature range acts as pinning particles. Recrystallized austenite grains are obtained, and the Gf grain size number of the prior austenite grains of the non-tempered hot forged parts is 7 or more. In addition, deformation resistance is reduced due to high temperature forging.

  When the forging temperature is lower than 1050 ° C., (V, Nb) C precipitates are precipitated in austenite. In particular, Nb tends to precipitate. When the (V, Nb) C precipitate is precipitated in austenite, the precipitation amount of (V, Nb) C precipitate that can be finely precipitated in the ferrite and pearlite during ferrite and pearlite transformation decreases. Decreases. As a result, the number density of fine (V, Nb) C precipitates in the ferrite and pearlite in the non-tempered hot forged part becomes a predetermined value or less, and the strength decreases. In addition, the upper limit value of the forging temperature is set to be equal to or lower than the heating temperature in the heat treatment step S <b> 1 from the capacity of the equipment.

In the heat treatment step S1, since the steel material is heated at a high temperature in the range of _TNbC , the austenite grains of the steel material before hot forging are coarsened. Therefore, in order to obtain fine recrystallized austenite grains by pinning NbC in the hot forging step S2, it is necessary to perform hot forging with a true strain amount of 0.3 or more to promote recrystallization of austenite grains. . Preferably, hot forging with an amount of true strain of 0.4 or more is performed, more preferably hot forging with an amount of true strain of 0.5 or more. The true strain amount is preferably 5.0 or less so that the deformation resistance during hot forging does not become too high.
The true strain amount is a strain amount calculated by ln [(cross-sectional area of the workpiece before hot forging) / (cross-sectional area of the workpiece after hot forging)].

(Rapid cooling step S3)
The rapid cooling step S3 is a step of rapid cooling at the cooling rate of 3.0 ° C./second or more after completion of hot forging, and the rapid cooling end temperature at the time of rapid cooling is set to 550 to 720 ° C. In the rapid cooling step S3, the amount of (V, Nb) C precipitates precipitated and coarsened in the austenite region before ferrite transformation decreases, and the amount of precipitation at the interface increases, thereby maximizing the amount of precipitation strengthening. To do.

  After the hot forging is completed, if it is slowly cooled (cooling rate: less than 3.0 ° C./second), (V, Nb) C precipitates are precipitated and coarsened in the austenite region before the start of ferrite transformation. Therefore, the amount of precipitation at the V / Nb phase interface decreases, and the amount of precipitation strengthening decreases. In addition, when the cooling is slow, the temperature at which the ferrite transformation starts increases. In particular, as a result of the recrystallized austenite grains being refined by the pinning effect of NbC, the austenite grain interface increases and the ferrite transformation is promoted. When the ferrite transformation temperature is high, the precipitation driving force of (V, Nb) C precipitates when phase interface precipitation occurs decreases, so that the size of nucleated (V, Nb) C precipitates becomes coarse, Interfacial precipitation itself does not occur, and ferrite with a small precipitation strengthening amount is locally formed. As a result, the yield strength of non-tempered hot forged parts cannot be improved.

  In order to improve the yield strength, the cooling rate at which the ferrite is not formed in the temperature range where the ferrite with a small precipitation strengthening amount is formed, that is, the temperature range from the hot forging end temperature to the rapid cooling end temperature (550 to 720 ° C.). It may be cooled at (3.0 ° C./second or more). The preferable range of the rapid cooling rate is 5.0 ° C./second or more, more preferably 10 ° C./second or more. Further, the preferable range of the rapid cooling end temperature is 570 to 680 ° C, more preferably 590 to 660 ° C.

  When the rapid cooling end temperature exceeds 720 ° C., the number density of fine (V, Nb) C precipitates becomes less than a predetermined value, and the yield strength of the non-tempered hot forged part cannot be improved. On the other hand, when the rapid cooling end temperature is less than 550 ° C., bainite is formed and the number density of fine (V, Nb) C precipitates is less than a predetermined value, and the yield strength of the non-tempered hot forged part. Cannot be improved.

(Slow cooling step S4)
The slow cooling step S4 is a step of slow cooling at a cooling rate of 0.1 ° C./second or more and 1.5 ° C./second or less after the end of the rapid cooling, and the slow cooling end temperature during the slow cooling is 400 ° C. or less To do. In the slow cooling step S4, ferrite transformation is promoted, and the number density of (V, Nb) C precipitates precipitated by phase interface precipitation increases.

  When V is added, nucleation of ferrite and pearlite is strongly suppressed. Therefore, when the temperature range from the rapid cooling end temperature to the slow cooling end temperature is cooled at a cooling rate exceeding 1.5 ° C./second, the ferrite-pearlite transformation Is not completed, and bainite and martensite are formed. As a result, in the non-tempered hot forged part, the number density of fine (V, Nb) C precipitates becomes less than a predetermined value, and the yield strength decreases. Therefore, it is necessary to cool at a cooling rate of 1.5 ° C./second or less at which bainite and martensite are not formed. On the other hand, when the cooling rate is as low as less than 0.1 ° C./second, a special slow cooling / warming facility is required or productivity is extremely reduced. Further, during the slow cooling, the (V, Nb) C precipitate is coarsened and the precipitation strengthening amount is reduced. As a result, in the non-tempered hot forged part, the number density of fine (V, Nb) C precipitates becomes less than a predetermined value, and the yield strength decreases. On the other hand, when the end temperature of the slow cooling exceeds 400 ° C., the ferrite-pearlite transformation is not completed, the number density of (V, Nb) C precipitates is less than a predetermined value, and the yield strength is insufficient.

  The method for producing a non-tempered hot forged part according to the present invention is a non-tempered step after a step other than the above-described steps, for example, a step of preparing a steel material into a predetermined shape before the heat treatment step, or a slow cooling step. It may be configured to include a step of cooling the hot forged part.

The steel having the chemical composition shown in Table 1 prepared by vacuum melting was heated at 1250 ° C. for 30 minutes, and then hot-forged into a round bar of φ50 mm and air-cooled.
A test piece of φ8 × 12 mm was cut out from D / 4 part of the round bar, and this was subjected to thermomechanical treatment under the conditions shown in Table 2 with a cermek master Z manufactured by Fuji Electric Koki Co., to obtain a forged material. The forging rate is 10 s ⁻¹ in terms of true strain rate, and the maximum load at that time is shown in Table 2.

The structure / precipitate and Vickers hardness of the produced forging were evaluated as follows.
The forged material is equally divided into two in parallel with the length direction before forging, and the cut surface is mirror-polished using one of them and corroded with 3% nital liquid, and at the center in the original length direction, D For the / 4 position, optical microscope observation was performed at magnifications of 100 and 400 to identify the tissue. Moreover, the Gf particle size number of the prior austenite (γ) grains was measured. The Gf grain size number was measured in accordance with JIS G0551 “Austenite grain size test method for steel”, and the appearance of old austenite grain boundaries was in accordance with “a) slow cooling method” in Appendix 1 of JIS G0551. Five fields of view were observed, and the average value was calculated. The results are shown in Table 2. Moreover, the hardness test was performed by the load of 10 kgf using the Vickers hardness tester for the optical microscope observation position vicinity, and the Vickers hardness was measured. Five points were measured and the average value was calculated. The results are shown in Table 2. In addition, the Vickers hardness set 400Hv or more as the pass.

  An extraction replica sample was prepared from the same position as that observed with the optical microscope, and a transmission electron microscope (TEM) image with a magnification of 100,000 was observed with respect to ferrite having a size of approximately 900 nm × 770 nm and five visual fields in pearlite. , Nb) The number density of C precipitates was measured. The results are shown in Table 2.

As shown in Table 2, Example No. 1 satisfying the requirements of the present invention. Nos. 2 to 4, 16, and 17 have high Vickers hardness, a large Gf particle size number, and fine prior austenite grains. In addition, FIG. 4 shows an optical micrograph of the ferrite-pearlite structure of FIG. 4 shows a TEM photograph of 4 (V, Nb) C precipitates.
On the other hand, comparative example No. which does not satisfy the requirements of the present invention. 1, 5 to 15 are inferior in at least one of Vickers hardness and Gf particle size number.

  Specifically, Comparative Example No. In No. 1, since the amount of V is less than the lower limit of the claims, the number density of fine (V, Nb) C precipitates is small, and the Vickers hardness is inferior. Comparative Example No. In No. 5, since the heating temperature and the forging temperature are less than the lower limit values of the claims, the number density of fine (V, Nb) C precipitates is small and the Vickers hardness is inferior. Comparative Example No. In No. 6, since the forging temperature is less than the lower limit of the claims, the Gf grain size number is small, and fine prior austenite grains are not obtained. Further, the number density of fine (V, Nb) C precipitates is small, and the Vickers hardness is inferior. Comparative Example No. In No. 7, since the true strain amount is less than the lower limit value of the claims, the Gf grain size number is small, and fine prior austenite grains are not obtained.

  Comparative Example No. Since No. 8 was not rapidly cooled, the number density of fine (V, Nb) C precipitates was small, and Vickers hardness was inferior. In addition, FIG. 8 shows a TEM photograph of 8 (V, Nb) C precipitates. Comparative Example No. In No. 9, since the rapid cooling rate is less than the lower limit of the claims, the number density of fine (V, Nb) C precipitates is small, and the Vickers hardness is inferior. Comparative Example No. No. 10 has a rapid cooling end temperature exceeding the upper limit of the claims, so that the number density of fine (V, Nb) C precipitates is small and Vickers hardness is inferior. Comparative Example No. No. 11 has a rapid cooling rate less than the lower limit of the claims, so that bainite is formed, the number density of fine (V, Nb) C precipitates is small, and Vickers hardness is inferior.

  Comparative Example No. No. 12, since the slow cooling rate is less than the lower limit of the claims, the number density of fine (V, Nb) C precipitates is small, and the Vickers hardness is inferior. Comparative Example No. In No. 13, since the slow cooling rate exceeds the upper limit of the claims, bainite is formed, the number density of fine (V, Nb) C precipitates is small, and the Vickers hardness is inferior. Comparative Example No. No. 14, since the amount of V exceeds the upper limit of the claims, bainite is formed, the number density of fine (V, Nb) C precipitates is small, and the Vickers hardness is inferior. Comparative Example No. Since No. 15 does not contain Nb, the Gf particle size number is small, and fine old austenite grains are not obtained. Further, in FIG. An optical micrograph of 15 ferrite-pearlite structures is shown.

S1 Heat treatment process S2 Hot forging process S3 Rapid cooling process S4 Slow cooling process

Claims (2)

C: 0.20 to 0.80 mass%,
Si: 0.50 mass% or less,
Mn: 0.40 to 1.00% by mass,
P: 0.050 mass% or less,
S: 0.050 mass% or less,
V: 0.20-0.80 mass%,
Nb: 0.02 to 0.30 mass%,
N: 0.0100% by mass or less, with the balance being Fe and inevitable impurities,
It consists of a ferrite-pearlite structure whose Gf particle size number of the prior austenite grains is 7 or more,
A high-strength, high-toughness, non-tempered hot forged part characterized in that the number density of (V, Nb) C precipitates having a diameter of 10 nm or less in ferrite and pearlite is 5000 pieces / μm ² or more. A method for producing a high-strength, high-toughness non-tempered hot forged part according to claim 1 ,
A heat treatment step of heat-treating the steel material having the above composition so as to be _TNbC calculated by the following formula (1);
A hot forging process in which the heat-treated steel material is hot forged at a forging temperature of 1050 ° C. or higher and lower than the heating temperature, and the true strain amount during the hot forging is 0.3 or more;
After completion of hot forging, rapid cooling is performed at a cooling rate of 3.0 ° C./second or more, and a rapid cooling process at a rapid cooling end temperature of 550 to 720 ° C.,
Including a slow cooling step in which the slow cooling is performed at a cooling rate of 0.1 ° C./second or more and 1.5 ° C./second or less after the end of the rapid cooling, and the slow cooling end temperature is 400 ° C. or less. A manufacturing method of high strength and high toughness non-tempered hot forged parts characterized by
−14000 / (log ([% C] · 0.01) −7.58) −273 ≦ T _NbC (° C.) ≦ −14000 / (log ([% C] · ([% Nb] −0.01) ) -7.58) -273 (1)
(However, in the formula (1), [% C] and [% Nb] are the contents (mass%) of C and Nb.)