JP5245997B2 - High strength hot forged non-tempered steel with excellent toughness and method for producing the same - Google Patents

Description

  The present invention is a high-strength hot-forged non-tempered steel that is suitable for parts such as automobiles and construction machinery and has excellent toughness without being subjected to heat treatment such as quenching or tempering after hot forging. And a manufacturing method thereof.

  Structural parts used in automobiles and construction machinery are made of carbon steel for mechanical structures and alloy steels for mechanical structures. In order to adjust strength and toughness, they are quenched and tempered after hot forging. Is common. In recent years, for the purpose of cost reduction, tempering processes such as quenching and tempering have been omitted. For example, non-tempered steels, such as carbon steel for machine structure specified in JIS G 4051 and manganese steel for machine structure specified in JIS G 4106, with addition of precipitation hardening elements such as V and Nb, etc. have been developed. It is used for automobile engine parts, undercarriage parts and construction machine parts.

  On the other hand, undercarriage parts of large vehicles such as trucks are required to have a high elastic limit and low temperature toughness. 0.05% proof stress, which is an index of elastic limit, and 2 mmU notch at −50 ° C., which is an index of low temperature toughness. Improvement of the absorbed energy during the Charpy test has been an issue. In addition, Pb-free steel materials are required as environmental measures.

  Conventional tempered steel mainly composed of ferrite and pearlite structure has an advantage that a high elastic limit is easily obtained. For example, a technique for preventing toughening of austenite grains by precipitated BN and improving toughness by making ferrite pearlite fine after transformation has been proposed (see, for example, Patent Document 1). However, non-heat treated steel mainly composed of ferrite and pearlite has a demerit that absorbed energy at low temperature is inferior to non-heat treated steel mainly composed of bainite. Therefore, the non-tempered steel proposed in Patent Document 1 has insufficient low temperature toughness of -50 ° C.

  Moreover, the refinement of prior austenite is achieved by adding Nb, and the structure after hot forging is made of one or both of martensite and bainite, thereby providing a non-heat treated steel that simultaneously satisfies high strength and high toughness. It has been proposed (for example, see Patent Document 2). However, this non-heat treated steel is subjected to water cooling after hot forging for the purpose of making the structure martensite or bainite. For this reason, there are problems such as distortion and burning cracks caused by a rapid temperature change, and there remains room for improvement.

  Further, bainite steel is proposed in which the amount of C is reduced and Mn and Cr are added to ensure hardenability (see, for example, Patent Document 3). If the amount of C is too large, the toughness is lowered and the machinability is lowered. However, a large amount of Mn and Cr must be added to make the structure bainite low in the amount of C, untempered, and bainite. For this reason, the alloy cost has increased, and the cost merit due to non-tempering cannot be fully enjoyed.

JP-A-62-74055 JP-A-63-312949 JP-A-8-277437

In the present invention, 0.05% proof stress is 500 MPa or more, and U-notch Charpy impact value at −50 ° C. is 20 J / cm ² or more without performing tempering treatment after hot forging and without increasing the alloy cost. Therefore, an object of the present invention is to provide a high-strength hot-forged non-tempered steel excellent in toughness having a machinability equivalent to that of a conventional material even if Pb-free and a method for producing the same.

The gist of the present invention is as follows.
(1) By mass%, C: 0.25 to 0.5%, Si: 0.05 to 1.0%, Mn: 1.0 to 2.3%, S: 0.04 to 0.2% Cr: 0.2-1.0%, V: 0.05-0.2%, B: 0.0005-0.004%, Ti: 0.001-0.1%, S and The Mn content satisfies 0.02 ≦ S / Mn ≦ 0.12, the balance is Fe and inevitable impurities, the area ratio of bainite is 95% or more, and the 0.05% proof stress is 500 MPa or more. A high-strength hot-forged non-tempered steel excellent in toughness, characterized in that the U-notch Charpy absorbed energy at −50 ° C. is 20 J / cm ² or more.
(2) The toughness described in (1) above, further containing, by mass%, one or both of Al: 0.005 to 0.06% and Nb: 0.01 to 0.3%. Excellent high strength hot forged non-tempered steel.
(3) Further, high strength hot forging non-tempered with excellent toughness according to the above (1) or (2), characterized by containing N: 0.001 to 0.007% by mass% steel.
(4) Further, by mass%, Cu: 0.01 to 0.5%, Ni: 0.01 to 2.0%, Mo: 0.01 to 1.0%, or one or more of them are contained. The high-strength hot-forged non-tempered steel excellent in toughness according to any one of (1) to (3) above.
(5) The above (1) to (4), further comprising one or both of Zr: 0.0001 to 0.02% and Mg: 0.0001 to 0.02% by mass%. A high-strength hot-forged non-tempered steel excellent in toughness according to any one of the above.
(6) The high-strength hot-forged non-tempered steel according to any one of the above (1) to (5), wherein the prior austenite grain size is 200 μm or less and the bainite lath width is 5 μm or less. .
(7) The method for producing high-strength hot forged non-heat treated steel according to any one of (1) to (6) above, which is described in any one of (1) to (5) above. High strength hot forged non-tempered steel with excellent toughness characterized by hot forging a steel material having the following components and cooling the temperature range from 800 ° C to 300 ° C at 0.1 to 5 ° C / s Manufacturing method.

According to the present invention, a high strength heat excellent in toughness having a 0.05% proof stress of 500 MPa or more and a -50 ° C. U-notch Charpy impact value of 20 J / cm ² or more without performing a tempering treatment after hot forging. It becomes possible to provide a non-tempered non-tempered steel and a manufacturing method thereof, and the industrial contribution is extremely remarkable. Moreover, since the hardenability is improved by increasing the amount of C, the alloy cost can be reduced.

  Unlike ferritic and pearlite steels, the bainite steel has a round type (stress-strain curve that gradually changes from an elastic region to a plastic region) without an upper yield point. Therefore, bainite steel has a 0.05% yield strength lower than that of ferritic pearlite steel having the same strength. Therefore, it is necessary to increase the strength in order to ensure 0.05% yield strength.

  Moreover, in order to ensure the low temperature toughness of bainite steel, it is necessary to refine the prior austenite and the bainite lath width. Refinement of prior austenite is achieved by utilizing precipitated MnS and suppressing grain growth of austenite after hot forging and before the start of transformation. Furthermore, in order to suppress the growth of austenite grain boundaries, precipitated particles such as VN and TiN, more preferably precipitated particles such as AlN are also used. The bainite lath width is refined by lowering the transformation temperature, that is, improving hardenability. As a hardenability improving element, C is mainly used, and Mn, Cr, V and the like are added in a necessary amount.

  However, when the hardenability is mainly increased by C, ferrite is generated in the form of a film at the grain boundary, resulting in a significant decrease in toughness. Therefore, B is added to suppress the formation of grain boundary ferrite. Further, it is solute B that is effective in suppressing grain boundary ferrite, and BN precipitates promote the generation of grain boundary ferrite. Therefore, fixation of N by the generation of TiN and AlN is also effective in suppressing the generation of BN.

  Hereinafter, the present invention will be described in detail. First, the reasons for limiting chemical components will be described.

C: 0.25 to 0.5%
C is an element effective for securing the material hardenability, and is an important element for securing the strength while making the structure bainite in the cooling process after hot forging. Also, the alloy cost is low, and if a large amount of C can be added, the alloy cost can be reduced. From the viewpoint of securing strength, it is necessary to contain 0.25% or more of C, but if it is too much, the toughness and machinability deteriorate, so 0.5% was made the upper limit. In addition, the preferable lower limit of C amount is 0.30%, and the preferable upper limit is 0.4%.

Si: 0.05-1.0%
Si is an element used for deoxidation at the time of melting, and is also an element effective for ensuring necessary strength. In particular, Si is solid-dissolved in a matrix, thereby improving proof stress and fatigue strength. For this purpose, 0.05% or more of Si needs to be added. However, if more than 1.0% is added, machinability deteriorates and grain boundary ferrite is generated, so 1.0% was made the upper limit. In addition, the preferable lower limit of Si amount is 0.1%, and the preferable upper limit is 0.5%.

Mn: 1.0 to 2.3%
Mn is an effective element as a deoxidizing element at the time of melting similarly to Si. Moreover, it is effective for improving hardenability and is an important element for making the structure bainite in the cooling process after hot forging. Furthermore, it combines with S to form MnS and pin the austenite grain boundaries, thereby preventing austenite coarsening during hot forging. In order to obtain these effects, it is necessary to add 1.0% or more of Mn. On the other hand, when Mn exceeds 2.3% Mn, machinability deteriorates and 0.05% proof stress and low temperature toughness are deteriorated, and martensite and austenite composite (MA) are produced. In order to increase, the upper limit was 2.3%. In addition, the preferable lower limit of the amount of Mn is 1.5%, and a preferable upper limit is 2.0%.

S: 0.04 to 0.2%
S is an element that combines with Mn to generate MnS, and is effective in preventing austenite coarsening during hot forging. MnS has crystallized MnS (produced in the liquid phase) and precipitated MnS (produced in the solid phase), but many MnS are crystallized MnS, which are coarse and not very effective for pinning austenite grains. Absent. In the present invention, solidification is completed after crystallization of coarse MnS, and solid solution Mn and solid solution S remaining in the solid phase are finely precipitated in a particle size of 10 nm to 1 μm and used for pinning austenite grains. 0.04% or more of S needs to be added. On the other hand, if adding more than 0.2% S, the toughness deteriorates, so the upper limit was made 0.2%. In addition, the preferable lower limit of S amount is 0.05%, and the preferable upper limit is 0.1%. Furthermore, since crystallization MnS and precipitated MnS are effective for improving machinability, this steel is Pb-less and can secure machinability that is higher than that of conventional steel.

Cr: 0.2 to 1.0%
Cr, like Mn, is an effective element for improving hardenability, and is an important element for making the structure a bainite structure in the cooling process after hot forging. For this reason, addition of 0.2% or more of Cr is necessary. However, when 1.0% or more of Cr is added, island martensite and MA increase, and 0.05% yield strength and low temperature toughness deteriorate. Therefore, the upper limit of Cr content is set to 1.0%. In addition, the preferable lower limit of Cr amount is 0.3%, and the preferable upper limit is 0.7%.

V: 0.05-0.2%
V is an element effective for forming carbides or nitrides to refine austenite crystal grains. Further, in the high temperature range where hot forging is performed, a part of V is dissolved, and the hardenability is improved. Therefore, it is effective for making the structure bainite in the cooling process after hot forging. Furthermore, V dissolved in the hot forging precipitates at a low temperature and improves the strength. In order to exert such an effect, it is necessary to add 0.05% or more of V. However, the addition exceeding 0.2% increases the deformation resistance of the steel material during hot forging, and improves the forgeability. In order to inhibit, the upper limit was made 0.2%.

B: 0.0005 to 0.004%
Solid solution B is an important element for improving hardenability and for making the structure bainite in the cooling process after hot forging. For this purpose, it is necessary to add 0.0005% or more, but if added over 0.004%, the low temperature toughness deteriorates, so the upper limit was made 0.004%. In addition, the preferable lower limit of B amount is 0.0015%, and a preferable upper limit is 0.003%.

Ti: 0.001 to 0.1%
Ti has the effect of generating nitrides as TiN and making the austenite crystal grains finer. Moreover, it is necessary to add nitrogen in order to fix nitrogen in austenite and improve the hardenability by adding B. For this purpose, it is necessary to add 0.001% or more of Ti, but if it exceeds 0.1%, the machinability is remarkably lowered, so the upper limit was made 0.1%. In addition, the preferable lower limit of the amount of Ti is 0.005% or more, and the preferable upper limit is 0.05% or less.

  Further, one or more of Al, Nb, N, Cu, Ni, Mo, Zr, and Mg may be added.

Al: 0.005-0.06%
Al is a deoxidizer, and is an element effective for forming nitrides and refining austenite crystal grains. In order to exert such an effect, it is preferable to add 0.005% or more. On the other hand, if the amount of Al is too large, the transformation temperature rises and the bainite lath may become rough, so it is preferable to set the upper limit to 0.06%. A preferred lower limit is 0.01%, and a preferred upper limit is 0.04%.

Nb: 0.01 to 0.3%
Nb is an effective element for forming carbides or nitrides to refine austenite crystal grains. In order to exert such an effect, it is preferable to add 0.01% or more of Nb. Even if Nb is added in excess of 0.3%, the effect is saturated, so it is preferable to set 0.3% as the upper limit.

N: 0.001 to 0.007%
N is an element effective for refining austenite crystal grains by forming nitrides by combining with elements such as Nb, Ti, and A1. In order to exhibit the effect, addition of 0.001% or more is preferable. However, if the amount of N exceeds 0.007%, it combines with B to form BN and inhibit the effect of improving the hardenability of B, so 0.007% is preferable as the upper limit.

Cu: 0.01 to 0.5%
Cu is an element effective for ensuring good hardenability, and addition of 0.01% or more is preferable. On the other hand, if adding more than 0.5% Cu, the castability of the steel may deteriorate, so the upper limit is preferably made 0.5%.

Ni: 0.01 to 2.0%
Ni is an element that ensures good hardenability and is also effective in improving toughness, and is preferably added in an amount of 0.01% or more. On the other hand, even if Ni over 2.0% is added, the effect is saturated, so it is preferable to set the upper limit to 2.0%.

Mo: 0.01 to 1.0%
Mo, like Ni, is an element that ensures good hardenability and is effective in improving toughness, and it is preferable to add 0.01% or more. On the other hand, even if Mo is added in excess of 1.0%, the effect is saturated, so it is preferable to set 1.0% as the upper limit.

Zr: 0.0001 to 0.02%
Zr is an element effective in improving machinability and forming oxides to refine austenite crystal grains. In order to exert the effect, addition of 0.0001% or more of Zr is preferable. On the other hand, if the amount is too large, the toughness may be lowered, so the upper limit of Zr is preferably 0.02%.

Mg: 0.0001 to 0.02%
Like Zr, Mg is an element effective for improving machinability and refining austenite crystal grains, and 0.0001% or more is preferably added. On the other hand, if the amount of Mg is too large, the toughness may decrease, so the upper limit is preferably made 0.02%.

0.02 ≦ S / Mn ≦ 0.12
In the present invention, since crystallization MnS and precipitated MnS are used for preventing austenite grain coarsening, the mass ratio between the S content and the Mn content is important. When S / Mn is 0.02 or less, sufficient MnS cannot be obtained, austenite grains become coarse, and low temperature toughness cannot be ensured. On the other hand, when S / Mn is larger than 0.12, coarse crystallized MnS increases and becomes a starting point of fracture, so that low temperature toughness cannot be secured. For this reason, S / Mn was 0.02-0.12 by mass ratio. In addition, the preferable lower limit of S / Mn is 0.04 or more, and a preferable upper limit is 0.1 or less.

  Next, the structure of the non-heat treated steel of the present invention will be described.

Bainite area ratio of 95% or more The bainite area ratio of the structure obtained by air cooling after hot forging is set to 95% or more. This is because it becomes difficult to ensure low temperature toughness when the area ratio of bainite is less than 95%. The remaining structures other than bainite are island-like martensite, MA, and grain boundary ferrite. When the total area ratio of these structures exceeds 5%, the toughness is significantly reduced. The total area ratio of bainite and the remaining ferrite, island-shaped martensite, and MA is averaged by image analysis by taking a photomicrograph after mirror-polishing the longitudinal section (L section) and mirror-polishing. Find the value.

Prior austenite grain size of 200 μm or less Refinement of prior austenite grain size is effective in improving toughness. When the prior austenite grain size exceeds 200 μm, the toughness decreases, so the upper limit is set to 200 μm. The grain boundaries of the prior austenite appear after corroding the L cross-section after being polished with picric acid. The prior austenite particle size is measured by taking a photomicrograph and cutting.

The width of the bainite lath is 5 μm or less Refining the width of the bainite lath after transformation is effective in improving toughness. When the width of the bainite lath exceeds 5 μm, the toughness decreases, so the upper limit is set to 5 μm. As for the width of the bainite lath, the L cross-section is mirror-polished and then nitrite-corroded to appear. The width of the bainite lath is obtained by taking a scanning electron microscope (SEM) photograph, measuring the lath width of a portion where each lath is clearly visible in each photograph, and obtaining an average value.

  Next, the manufacturing method of the non-heat treated steel of this invention is demonstrated. The non-heat treated steel of the present invention heats steel material, performs hot forging, and cools without performing accelerated cooling such as water cooling. Also, tempering such as quenching and tempering is not performed. The heating temperature of the steel material and the end temperature of hot forging are not particularly limited, but the following conditions are preferable.

  When the heating temperature of the steel material exceeds 1300 ° C., the scale becomes thick and the yield decreases. Moreover, since the austenite coarsens, the effect of pinning particles such as MnS, TiN, and AlN becomes insufficient and the crystal grains may become coarse, the heating temperature of the steel is preferably 1300 ° C. or lower.

  When the end temperature of hot forging exceeds 1250 ° C., coarsening of crystal grains and local abnormal grain growth may occur. Therefore, it is preferable to set the upper limit of the end temperature of hot forging to 1250 ° C. or less. On the other hand, when the end temperature of hot forging is less than 900 ° C., the deformation resistance increases and the die life is shortened. Therefore, the end temperature of hot forging is preferably 900 ° C. or higher.

Average cooling rate from 800 ° C to 300 ° C 0.1-5 ° C / s
The non-tempered steel of the present invention is manufactured without performing a tempering treatment such as quenching or tempering after hot forging. Usually, after completion of hot forging, it may be allowed to cool, but since the cooling rate varies depending on the size, mist cooling or the like is performed as necessary. In the temperature range from 800 ° C. to 300 ° C., a bainite structure is generated during cooling. Therefore, when the average cooling rate is less than 0.1 ° C./s, grain boundary ferrite is generated, and the lath width of the bainite is 5 μm or more. Decreases. On the other hand, when the average cooling rate from 800 ° C. to 300 ° C. exceeds 5 ° C./s, cracks and strains are likely to occur after cooling, as in the case of non-heat treated steel produced by water cooling. Therefore, after hot forging, the average cooling rate in the temperature range from 800 ° C. to 300 ° C. needs to be 0.1 to 5 ° C./s.

By hot forging the steel material having the essential component range of the present invention and cooling the temperature range from 800 ° C. to 300 ° C. at 0.1 to 5 ° C./s, the area ratio of bainite is 95% or more. High-strength hot steel having a prior austenite grain size of 200 μm or less, a bainite lath width of 5 μm or less, a 0.05% proof stress of 500 MPa or more, and a U-notch Charpy absorbed energy at −50 ° C. of 20 J / cm ² or more. Forged non-tempered steel can be manufactured.

  Steel having the chemical composition shown in Table 1 was melted in a 300 kg vacuum melting furnace and then rolled to φ70 mm. Further, a 200 mm long test piece was cut out from the round bar, heated to 1200 ° C. in a heating furnace, and forged into a 50 mm square bar. The forging end temperature was 1100 ° C., and air cooling was performed after forging. The cooling rate from 800 ° C. to 300 ° C. is 0.38 ° C./s.

  The test piece for structure | tissue observation was extract | collected from the 1/4 thickness part of the L direction of the obtained forged material. For the prior austenite grain size, the L section of the specimen for texture observation was mirror-polished and then corroded with picric acid, and 200 times optical micrographs were taken for each three fields of view, and the prior austenite grain size was measured by a cutting method. The average value was obtained. The bainite lath width was corroded after re-polishing and 10 times each of 2000 magnified SEM photographs were taken, and the lath width was measured for the portion where the lath was clearly visible in each photograph, and the average value was obtained.

  The area ratio of ferrite and the remaining area ratio (total of one or both of island-like martensite and MA) were subjected to repeller etching after re-polishing the specimen, and five optical microscope photographs of 500 times were taken for each image. The average value was obtained by analysis. It was confirmed that the structure other than ferrite, island martensite, and MA was bainite, and the area ratio of bainite was calculated from the sum of ferrite and the remainder.

  Furthermore, JIS Z 2201 No. 4 tensile test piece and JIS Z 2242 2 mmU notch Charpy test piece were collected from the 1/4 thickness part in the L direction of the obtained forging. The tensile test was performed at room temperature, and the Charpy test was performed at -50 ° C.

The results are shown in Table 2. Production No. satisfying the component composition of the present invention. Nos. 1 to 10 have a high yield strength and a high toughness with a 0.05% proof stress exceeding 500 MPa and an absorbed energy of ² mmU notch Charpy test at −50 ° C. of 20 J / cm ² or more.

  On the other hand, production No. 11 has a smaller amount of C, more Mn, less S, and a smaller S / Mn ratio than the present invention. For this reason, the area ratio of bainite is reduced, the effect of refining austenite with MnS is weak, the prior austenite grain size becomes coarse, and both the 0.05% proof stress and the −50 ° C. impact value are lowered.

  Production No. No. 12 has a large amount of C. No. 15 has a large amount of Mn. No. 16 has a large amount of Cr. Since 18 has a large amount of V, the area ratio of island martensite or MA is increased, and the area ratio of bainite is decreased. These have a 0.05% proof stress exceeding 500 MPa, but have a low -50 ° C impact value.

  Production No. No. 13 has a small amount of S and a small S / Mn. In addition, production No. No. 17 does not add V. Therefore, the prior austenite grain size becomes coarse and the 0.05% yield strength exceeds 500 MPa, but the -50 ° C impact value is low.

  Production No. No. 14 has a large amount of S and a large S / Mn. Therefore, although the prior austenite grain size is made fine by MnS, the hardenability is lowered and the lath of bainite is coarsened. Production No. Since 19 does not contain B, hardenability is insufficient, resulting in coarsening of the bainite lath width and generation of intergranular ferrite. These have a 0.05% proof stress exceeding 500 MPa, but have a low -50 ° C impact value.

Production No. No. 20 has a large amount of B, BN and M ₂₃ CB ₆ are precipitated, the hardenability is lowered, and the bainite lath width is coarsened. Furthermore, due to the formation of island martensite and MA, the area ratio of bainite is decreased, and the 0.05% proof stress and the −50 ° C. impact value are both low.

  Production No. No. 21 has an S / Mn of 0.02 or less, so the effect of refining austenite by MnS is weak, the prior austenite grain size becomes coarse, and the area ratio of bainite decreases due to the formation of island martensite and MA. However, the impact value at −50 ° C. is lowered.

  Production No. No. 22 has an S / Mn value of 0.2 or more, so that the prior austenite grain size is reduced by MnS, but the hardenability is reduced, the bainite lath is coarsened, and film-like ferrite is formed at the grain boundaries. Is generated. For this reason, both 0.05% yield strength and -50 ° C impact value are low values.

Steel No. 1 in Table 1 A 200 mm long test piece was cut out from the rolled material (φ70 mm) of A, the heating temperature was set to 1250 ° C., the forging end temperature was set to 1150 ° C., and a 50 mm square bar was molded. It cooled at the cooling rate shown. Cooling rate from 800 ° C to 300 ° C, 8 ° C / s is oil-cooled by immersion in an oil bath of 60 ° C (production No. 26), 3 ° C / s is water mist cooling (production No. 23) , 0.5 ° C./s is air-cooled (Production No. 24), 0.2 ° C./s is forged and cooled in a heat insulating material (Production No. 25), and 0.05 ° C./s is furnace cooled (Production No. .27)

  The prior austenite grain size and bainite lath width were measured in the same manner as in Example 1. The area ratio of ferrite and martensite and the area ratio of the remainder (the sum of one or both of island-shaped martensite and MA) were polished and repeller etched as in Example 1, and an optical micrograph was taken. The average value was obtained by analysis. The area ratio of bainite was calculated from the sum of ferrite, martensite and the balance. The tensile test and the Charpy test were performed in the same manner as in Example 1.

The results are shown in Table 3. As is apparent from Table 3, when the average cooling rate between 800 and 300 ° C. is in the range of 0.1 to 5 ° C./s, the bainite area ratio is 95% or more, 0.05% proof stress is 500 MPa or more, −50 The impact toughness value at 20 ° C. satisfies 20 J / cm ² .

  When the average cooling speed between 800 and 300 ° C. is 8 ° C./s, the martensite area ratio becomes 80% and the strength increases, so the impact toughness value at −50 ° C. is low. On the other hand, when the average cooling rate between 800 and 300 ° C. is 0.05 ° C./s, grain boundary ferrite, island martensite, and MA are generated, the area ratio of bainite is reduced, and 0.05% proof stress is low. It is a value.

Claims (7)

% By mass
C: 0.25 to 0.5%,
Si: 0.05 to 1.0%,
Mn: 1.0 to 2.3%
S: 0.04 to 0.2%,
Cr: 0.2 to 1.0%
V: 0.05-0.2%
B: 0.0005 to 0.004%,
Ti: 0.001 to 0.1%
And the contents of S and Mn are
0.02 ≦ S / Mn ≦ 0.12
The balance is composed of Fe and inevitable impurities, the area ratio of bainite is 95% or more, the 0.05% proof stress is 500 MPa or more, and the U-notch Charpy absorbed energy at −50 ° C. is 20 J / cm ² or more. High strength hot forged non-heat treated steel with excellent toughness characterized by Furthermore, in mass%,
Al: 0.005 to 0.06%,
Nb: 0.01 to 0.3%
The high-strength hot-forged non-tempered steel excellent in toughness according to claim 1, wherein one or both of the above are contained. Furthermore, in mass%,
N: 0.001 to 0.007%
The high-strength hot-forged non-tempered steel excellent in toughness according to claim 1 or 2, characterized by comprising: Furthermore, in mass%,
Cu: 0.01 to 0.5%,
Ni: 0.01 to 2.0%,
Mo: 0.01 to 1.0%
The high-strength hot-forged non-tempered steel excellent in toughness according to any one of claims 1 to 3, characterized by containing one or more of the following. Furthermore, in mass%,
Zr: 0.0001 to 0.02%,
Mg: 0.0001 to 0.02%
One or both of these are contained, The high strength hot forging non-heat-treated steel excellent in toughness of any one of Claims 1-4 characterized by the above-mentioned.   The high-strength hot-forged non-tempered steel according to any one of claims 1 to 5, wherein the prior austenite grain size is 200 µm or less and the width of the bainite lath is 5 µm or less.   It is a manufacturing method of the high intensity | strength hot forging non-tempered steel of any one of Claims 1-6, Comprising: Hot forging the steel material which has a component of any one of Claims 1-5. And the manufacturing method of the high strength hot forging non-heat-treated steel excellent in toughness characterized by cooling the temperature range from 800 degreeC to 300 degreeC at 0.1-5 degreeC / s.