JP4540428B2 - Method for producing hot rolled non-heat treated steel bar - Google Patents

Description

The present invention is a non-tempered material as it is hot-rolled, the tensile strength is 700 MPa or more, and the difference in tensile strength between the surface layer portion and the central portion is 80 MPa or less, the difference in yield strength is 100 MPa or less, Moreover excellent in machinability in high tenacity, a method of manufacturing a hot rolled type non-heat treated bar steel.

  The hot-rolled non-tempered steel bar in the present invention is a steel bar that does not require a tempering treatment after it is processed into a product by precipitating fine precipitates during cooling after rolling the steel bar. Means that.

  Structural components used in transportation equipment and construction machinery, including automobiles, include not only tempered steel that has been tempered and tempered carbon steel for machine structure and alloy steel for machine structure, Non-tempered steel whose strength is secured by adjusting the structure is used.

  Non-tempered steel used for such applications generally has a ferrite-pearlite double-layered structure with V and Nb added. When tensile strength is comparable to tempered steel, yield strength is higher. It has been pointed out that when the drawing value and impact value are low and the yield strength is comparable, the tensile strength, that is, the hardness is excessively increased and the machinability is decreased.

  Under the background described above, Patent Document 1 and Patent Document 2 describe a structure having ferrite, bainitic ferrite, and pseudo martensite in order to obtain a high-strength, high yield ratio, and high toughness non-heat treated steel. A technique comprising aging treatment of a steel material having the above-mentioned after cold working at 600 ° C. or less to precipitate Cu and Ti—Nb-based carbides is disclosed.

  However, in actual production, it is extremely difficult to strictly control the ratio of multiple structures as described above, and when using precipitation strengthening by adding a large amount of Cu, a large amount of expensive Ni is used to prevent hot cracking. Therefore, there is a problem that it is not suitable as a structural part that is consumed in large quantities.

Therefore, the inventors have not reduced the productivity even in actual production, have a tensile strength of 700 MPa or more and a yield ratio of 0.85 or more with an inexpensive component system, and are excellent in hot rolling type non-adjustment. As a technique for obtaining a solid steel bar, a technique for dispersing and precipitating fine precipitates having a particle size of less than 10 nm in a ferrite single phase structure was developed and disclosed in Patent Document 3.
However, the technique of Patent Document 3 was devised for a small-diameter steel bar having a relatively small diameter, as can be seen from the examples, and is not necessarily applicable to a large-diameter steel bar having a large diameter. There was a problem.
In other words, the technique of Patent Document 3 is based on dispersing fine precipitates during cooling after hot rolling, but when the diameter increases and the cooling rate of the steel bar becomes 0.2 ° C./s or less, precipitation occurs. In some cases, the object is not sufficiently fine, and predetermined characteristics cannot be obtained. Moreover, since the cooling rate difference between the surface layer portion and the central portion of the steel bar is not taken into consideration, a low temperature transformation phase such as bainite is generated in the surface layer portion where the cooling rate is relatively fast, and the tensile strength and It has been found that there may be a large difference in yield strength.

JP 2001-123224 JP 2001-131680 A JP 2004-3008 A

The present invention has been developed in view of the above-mentioned actual situation, and can be easily manufactured by an inexpensive component system and by actual manufacturing. Moreover, even if the steel bar has a large diameter of 120 to 400 mm , 700 MPa Hot-rolled non-tempered rod with high tensile strength and high toughness comparable to tempered material, with small difference in tensile strength and yield strength between the surface layer and center, and excellent strength uniformity in the diameter direction The object is to propose an advantageous method for producing steel .

Now, as a result of intensive studies to achieve the above object, the inventors have found that the precipitation behavior of the precipitate and the progress of the ferrite transformation are closely related to the refinement of the precipitate. .
Specifically, the precipitate is effective when the difference between the transformation start temperature of the ferrite transformation that occurs during cooling after rolling and the precipitation start temperature of the precipitate is small and the ferrite transformation and the precipitation of the precipitate compete with each other. The knowledge that it refines automatically is obtained.
The present invention is based on the above findings.

That is, the gist configuration of the present invention is as follows.
1 . % By mass
C: 0.060 to 0.120%,
Si: 0.5% or less,
Mn: the range of the following formula (1),
Al: 0.1% or less,
Ti: 0.03-0.35% and
Mo: 0.05-0.8%
After heating the steel material containing Fe and the inevitable impurities composition to 1100 ° C or higher, finishing temperature: 850 ° C or higher and finishing diameter 120-400mm, and then cooling A method for producing a hot rolled non-heat treated steel bar characterized by the above.
Record
-0.239 × log (CR ₁ ) + 0.889 ≦ Mn ≦ −0.524 × log (CR ₂ ) +1.218 (1)
Where, CR ₁ : Average cooling rate at the center of the steel bar up to 500 ° C after rolling (° C / s)
CR ₂ : Average cooling rate of steel bar surface layer up to 500 ° C after rolling (° C / s)

2 . 1. A method for producing a hot-rolled non-tempered steel bar according to 1 above, wherein the amounts of C, Ti and Mo in the steel satisfy the relationship of the following formula (2).
0.5 ≦ (C / 12) / [(Ti / 48) + (Mo / 96)] ≦ 1.5 (2)

3 . In the above 1 or 2 , the steel material is further in mass%.
Nb: 0.08% or less,
A method for producing a hot-rolled non-tempered steel bar, comprising a composition containing one or more selected from V: 0.15% or less and W: 1.5% or less.

4. 3. The method for producing a hot-rolled non-tempered steel bar according to 3 above, wherein the amounts of C, Ti, Mo, Nb, V and W in the steel satisfy the relationship of the following formula (3).
0.5 ≦ (C / 12) / [(Ti / 48) + (Mo / 96) + (Nb / 93) + (V / 51) + (W / 184)] ≦ 1.5 (3)

According to the present invention, the tensile strength of 700 MPa or more and high toughness comparable to the tempered material, which is an inexpensive component system and can be easily produced by actual production, are further obtained. Further, it is possible to obtain a hot-rolled non-tempered steel bar having a small difference in yield strength and excellent strength uniformity in the diameter direction.
Further, according to the present invention, not only the tempering treatment is unnecessary, but also there is no restriction on the dimensions of the steel bar, which is extremely useful industrially.

The present invention will be specifically described below.
First, the reason why the component composition of steel is limited to the above range in the present invention will be described. Unless otherwise specified, “%” in relation to ingredients means mass%.
C: 0.060 to 0.120%
If C is less than 0.060%, the precipitation amount of fine precipitates is insufficient, and a tensile strength of 700 MPa or more cannot be obtained. Since tensile strength cannot be obtained, the C content is limited to a range of 0.060 to 0.120%.

Si: 0.5% or less
Si is added to improve cold workability, but if the content exceeds 0.5%, the effect is rather impaired, so the Si content is 0.5% or less. More preferably, it is 0.15% or less.

Mn: Range of the following formula (1)
Record
-0.239 × log (CR ₁ ) + 0.889 ≦ Mn ≦ −0.524 × log (CR ₂ ) +1.218 (1)
Where, CR ₁ : Average cooling rate at the center of the steel bar up to 500 ° C after rolling (° C / s)
CR ₂ : Average cooling rate of steel bar surface layer up to 500 ° C after rolling (° C / s)
It is important to control Mn within the range of the above formula (1) in order to precipitate precipitates finely and to prevent the formation of a low-temperature transformation phase such as bainite in the surface layer portion.

Fig. 1 shows the melting of 0.075% C-0.2% Si-0.16% Ti-0.34% Mo-0.015% P-0.015% S-0.045% Al-0.0033% N steel with various Mn contents. After heating to 1200 ° C and hot rolling to a steel bar at a finishing temperature of 900 ° C, the results of investigations on the steel structure and the size of precipitates were investigated. The amount of Mn in the steel and the cooling after hot rolling Shown in relation to speed.
In hot rolling, in order to change the cooling rate after hot rolling, the finishing dimensions of the steel bar are changed in the range of 20mmφ to 400mmφ, and cooling after rolling is air cooling, slow cooling in the heat insulation cover, mist cooling. The following three methods were used. Since it is difficult to measure the cooling rate online, a short steel bar of the same size as the finished dimensions in hot rolling is prepared, and the position of the steel bar 6 mm deep from the surface (hereinafter referred to as the surface layer) and the steel bar A thermocouple is attached to the center (hereinafter referred to as the center), heated to 900 ° C offline, and then air-cooled, gradually cooled in the heat insulation cover, and mist-cooled in the same manner as in the case of the above hot rolling. The average cooling rate to ° C was determined.

As is clear from FIG. 1, the structure of the steel bar and the size of the precipitates change according to the amount of Mn and the cooling rate.
Mn amount is the region between straight line a and straight line b, that is, a value calculated from the cooling rate {−0.239 × log (CR) +0.889} or more and {−0.524 × log (CR) +1.218} or less In some cases, a structure is obtained in which the ferrite is a single phase and fine precipitates having a particle size of less than 10 nm are dispersed in the ferrite.
On the other hand, when the amount of Mn is less than the value calculated from the cooling rate {−0.239 × log (CR) +0.889}, although the ferrite single phase is formed, the grain size of the precipitate becomes 10 nm or more. End up. When the amount of Mn exceeds {−0.524 × log (CR) +1.218}, ferrite exists in a part of the structure, but most of it becomes a low temperature transformation phase such as bainite.

Therefore, in order to obtain a structure in which ferrite is a single phase and fine precipitates having a particle size of less than 10 nm are dispersed in this ferrite, the amount of Mn is related to {−0.239 × log (CR) +0. It is important to adjust to a range of 889} or more and {−0.524 × log (CR) +1.218} or less.
However, the cooling rate of the steel bar is not necessarily uniform, the center portion is the smallest and the surface layer portion is the largest.

Therefore, in the present invention, the cooling rate CR ₁ at the center where the cooling rate is the smallest is applied to the straight line a {= − 0.239 × log (CR) +0.889}, while the straight line b {= − 0.524 × log. For (CR) +1.218}, the cooling rate CR ₂ of the surface layer with the highest cooling rate is applied, and the range in which the Mn content in the steel is sandwiched between these straight lines, that is, the above formula (1) is satisfied. By controlling to the range, a structure in which fine precipitates having a particle size of less than 10 nm were dispersed in the ferrite single phase at any position from the surface layer portion to the central portion of the steel bar was obtained.
It should be noted that since the ferrite transformation and precipitation of precipitates are all finished by 500 ° C., the cooling rate may be considered from 500 ° C. after rolling.

By the way, while proceeding with the examination as described above, the inventors have closely related the precipitation behavior of the precipitate to the progress of the ferrite transformation, and the transformation start temperature and precipitation of the ferrite transformation that occurs during cooling after rolling. It was newly found that precipitates are effectively refined when the difference in precipitation start temperature of the products is small and the ferrite transformation and precipitation compete.
As shown in FIG. 1, the lower limit of the amount of Mn for making precipitates smaller than 10 nm increases as the cooling rate slows. This is also explained as follows from the competitive relationship between ferrite transformation and precipitation. Can do.

In the present invention, precipitation of precipitates starts after ferrite transformation starts during cooling after hot rolling. Here, since the ferrite transformation temperature is higher as the cooling rate is slower, when the cooling rate is slow, after the ferrite transformation starts, precipitation starts after a considerable temperature drop. However, in order to make the precipitate finer, it is necessary to reduce the difference between the start temperature of the ferrite transformation and the start temperature of the precipitation, thereby competing the ferrite transformation and the precipitation.
Therefore, in the present invention, Mn is added to lower the ferrite transformation temperature. In this case, the lower the cooling rate side where the ferrite transformation temperature becomes higher, the lower the allowance for lowering the ferrite transformation temperature. Therefore, it is necessary to increase the amount of Mn added.
Therefore, as shown in FIG. 1, the lower limit of the amount of Mn for refinement of precipitates increases as the cooling rate decreases.

  Further, the upper limit of the Mn amount also increases as the cooling rate is slower. The reason for this is considered to be that when the cooling rate is low, a low-temperature transformation phase such as bainite hardly occurs, so that the amount of addition of Mn, which is a hardenability increasing element, is increased.

Al: 0.1% or less
Al acts as a deoxidizer. However, if the content exceeds 0.1%, the effect is saturated, so the Al content is 0.1% or less. More preferably, it is 0.05% or less.

Ti: 0.03-0.35%
Ti is a useful element that improves the strength by finely depositing precipitates including Ti-based carbides and Ti-Mo-based carbides. Here, in order to ensure a tensile strength of 700 MPa or more, Ti addition of 0.03% or more is necessary. However, if it exceeds 0.35%, precipitates are coarsened and strength and toughness are lowered. Therefore, the Ti amount is limited to the range of 0.03 to 0.35%. More preferably, it is 0.03 to 0.20% of range.

Mo: 0.05-0.8%
Mo is a useful element that improves the strength by finely depositing precipitates including Mo-based carbides and Ti-Mo-based carbides. In addition, Mo has an advantage that the diffusion rate is slow, and when it is precipitated together with Ti, the growth rate of the precipitate is reduced and a fine precipitate is easily obtained. Here, in order to secure a tensile strength of 700 MPa or more, it is necessary to add 0.05% or more of Mo, but if added over 0.8%, a low-temperature transformation phase such as bainite is formed, and precipitation strengthening by fine precipitates is achieved. The amount of Mo was limited to a range of 0.05 to 0.8% because the strength was insufficient. More preferably, it is 0.15 to 0.50% of range.

Moreover, in said component composition, it is advantageous to satisfy the relationship of following Formula (2) regarding especially the amount of C, Ti, and Mo in steel.
0.5 ≦ (C / 12) / [(Ti / 48) + (Mo / 96)] ≦ 1.5 (2)
The parameter shown in the equation (2) affects the size of the precipitate, and when it is 0.5 or more and 1.5 or less, the formation of fine precipitates having a particle size of less than 10 nm is facilitated, which is particularly advantageous. .

  In addition, it was observed that Ti and Mo in carbides have a Ti / Mo atomic ratio of 0.2 to 2.0 in fine carbides and Ti / Mo ratios in 0.7 to 1.5 in finer carbides.

Although the essential components have been described above, in the present invention, one or more selected from Nb, V, and W can be added in order to further improve the strength and the like.
Nb: 0.08% or less
Nb contributes to strength increase by forming fine precipitates with Ti. Moreover, ductility is improved by refining the structure and adjusting the crystal grains. However, if Nb is contained in excess of 0.08%, it will be excessively refined and the ductility will be lowered. Therefore, the addition amount is set to 0.08% or less. More preferably, it is 0.04% or less.

V: 0.15% or less V also forms a fine precipitate with Ti and contributes to an increase in strength. However, if the content exceeds 0.15%, the precipitate becomes coarse, so the content is made 0.15% or less. More preferably, it is 0.10% or less.

W: 1.5% or less W also contributes to strength increase by forming fine precipitates with Ti. However, if the content exceeds 1.5%, the precipitate becomes coarse, so the content is 1.5% or less. More preferably, it is 1.0% or less.

When the above-mentioned Nb, V, W, etc. are added, in order to refine the precipitate, it is advantageous to satisfy the relationship of the following formula (3) with respect to the amount of these elements and C, Ti, Mo. is there.
0.5 ≦ (C / 12) / [(Ti / 48) + (Mo / 96) + (Nb / 93) + (V / 51) + (W / 184)] ≦ 1.5 (3)
The parameter shown in the equation (3) also affects the size of the precipitate as in the equation (2). When this value is 0.5 or more and 1.5 or less, the particle size is finer than 10 nm. Formation of precipitates is facilitated.

  In the case of fine carbides containing one or more of Nb, V, and W, the atomic ratio (Ti + Nb + V) / (Mo + W) of Ti, Mo, Nb, V, and W in the carbide is 0.2 to 2.0, which is finer It was observed that this ratio was 0.7-1.5 for carbides.

Furthermore, in the present invention, in order to improve the machinability at the time of machining the part, S is set to 0.03 to 0.1%, and Pb ≦ 0.2%, Ca ≦ 0.005%, Bi ≦ 0.1% and B ≦ 0.02%. It is advantageous to contain one or more selected from
Here, the reason why the S amount is 0.03 to 0.1% is that if the S amount is less than 0.03%, the machinability cannot be improved, whereas if it exceeds 0.1%, the toughness and ductility deteriorate. Moreover, also about Pb, Ca, Bi, and B, when these elements each exceed an upper limit, toughness and ductility will fall.

In addition, for the purpose of improving strength and ductility, one or more selected from Cr, Ni and Cu can be added in the range of Cr ≦ 0.5%, Ni ≦ 0.5% and Cu ≦ 0.5%.
In order to further improve the toughness of the steel bar, it is desirable to reduce P and N which are inevitable impurities. Specifically, it is preferable to restrict P to 0.03% or less. N is preferably regulated to 0.01% or less, more preferably 0.005% or less.

Next, a description will be given to your Keru steel organization to the present invention.
First, the reason why the steel structure is a ferrite single phase structure in the present invention is that the toughness comparable to the tempered material can be obtained by making the steel structure a ferrite single phase.

  In the present invention, the ferrite single-phase structure means that the area ratio of ferrite is 95% or more, preferably 98% or more in any of the surface layer portion to the central portion of the steel bar in cross-sectional structure observation (observation of optical microscope structure at 200 times). It points to something.

The reason why the fine precipitates precipitated in the ferrite single-phase structure are made as fine precipitates having a particle size of less than 10 nm is as follows.
When the particle size of the precipitate is 10 nm or more, it is difficult to obtain a tensile strength of 700 MPa or more necessary for machine structural parts such as automobiles and other transportation machines and construction machines. In addition, when fine precipitates having a particle size of less than 10 nm are precipitated in the ferrite single phase structure, the yield ratio increases and a high yield ratio comparable to the tempered material can be obtained. When the yield ratio is high, the increase in tensile strength is suppressed with respect to the increase in yield strength, and the hardening of the steel can be reduced, so that machinability comparable to tempered steel is obtained. The fine precipitate is precipitated during cooling after hot rolling.

  The smaller the particle size of the fine precipitates, the more effective the strength increase, and is preferably 5 nm or less, more preferably 3 nm or less. As such fine precipitates, carbides containing a composite of Ti and Mo, and carbides further containing one or more of Nb, V, and W are advantageously suitable.

The particle size of fine precipitates of less than 10nm, when the number is at 1000 / [mu] m ³ or more, it becomes easier to obtain a tensile strength of at least 700 MPa, it is preferable to present 1000 / [mu] m ³ or more.
Although the distribution form of these fine precipitates is not particularly defined, it is preferable to uniformly disperse in the matrix. Moreover, as for the size of the precipitate, if 90% or more of the total precipitate is less than 10 nm, a tensile strength of 700 MPa or more can be obtained. However, a precipitate having an excessively large particle size consumes fine precipitate-forming elements and adversely affects the strength. Therefore, the particle size of the existing precipitate is preferably suppressed to 50 nm or less.

  In addition to the precipitates described above, the effect of the present invention is not impaired even if a small amount of Fe carbide is contained. However, if a large amount of Fe carbide having an average particle size of 1 μm or more is contained, the toughness is inhibited. In the invention, it is desirable that the upper limit of the size of Fe carbide contained is 1 μm and the content is 1% or less of the entire precipitate.

The number of fine precipitates is determined by the following method.
An electron microscope sample is prepared by electropolishing using the twin jet method and observed at an acceleration voltage of 200 kV. At this time, the crystal orientation of the matrix is controlled so that the fine precipitates have a measurable contrast with respect to the matrix, and the focus is shifted from the normal focus in order to minimize the number of precipitates. Observe with focus method. Moreover, the sample thickness of the area | region which measured the deposit particle | grains is evaluated by measuring an elastic scattering peak and an inelastic scattering peak intensity using an electron energy loss spectroscopy.

By this method, the number of particles and the sample thickness can be measured in the same region. The number of particles and the particle diameter are measured at four locations of a 0.5 μm × 0.5 μm region of the sample, and the precipitates distributed per 1 μm ² are calculated as the number for each particle diameter. From this value and the thickness of the sample, the number of precipitates distributed per 1 μm ^{3 of} particle diameter is calculated, and the ratio of the precipitate having a particle diameter of less than 10 nm to the measured total precipitate is calculated.

Next, preferred manufacturing conditions of the present invention will be described.
-Heating temperature In the present invention, in order to precipitate precipitates finely during cooling after hot rolling, it is necessary to once dissolve the precipitates precipitated on the slab before hot rolling in a heating furnace. There is. At that time, if the heating temperature is less than 1100 ° C., Ti—Mo-based carbides and the like are not sufficiently dissolved, so the heating temperature is set to 1100 ° C. or higher.

-Finishing temperature In the present invention, in order to obtain fine precipitates, the amount of Mn is adjusted according to the cooling rate after hot rolling, and the start temperature of ferrite transformation is controlled, thereby competing between ferrite transformation and precipitation. ing. However, when the finishing temperature in hot rolling is low, the strain introduced by rolling changes the starting temperature of the ferrite transformation and affects the appropriate range of the Mn amount. In order to avoid this, the finishing temperature may be set to a high temperature at which the influence of strain does not appear. From this point, the finishing temperature is 850 ° C or higher.

-Cooling rate If the starting temperature of the ferrite transformation is controlled and the ferrite transformation competes with the precipitation, the precipitate precipitates finely. In the present invention, the ferrite transformation start temperature is controlled by adjusting the amount of Mn according to the cooling rate after hot rolling. Moreover, if the amount of Mn is adjusted according to the cooling rate, the generation of a low-temperature transformation phase such as bainite can be prevented.
In this way, in the present invention, the cooling rate is not controlled, but an appropriate structure can be obtained by adjusting the amount of Mn according to the cooling rate . Therefore, it is not necessary to specifically define the cooling rate .

  Steel having the component composition shown in Table 1 was melted, and this was hot-rolled into bar steel of a predetermined size under the conditions described in Table 2 and Table 4. In melting, in addition to the invention steel that satisfies the proper range of the present invention and the comparative steel that deviates from the proper range of the present invention, S45C tempered non-heat treated steel (steel number 10) was also melted as a conventional steel. .

  In hot rolling, the cooling rate was changed by changing the finishing dimensions of the steel bar and the cooling method after rolling, as described in FIG. Since it is difficult to measure the cooling rate online, we prepare short steel bars of the same size as the finished dimensions in hot rolling, attach thermocouples to the surface layer and center, and perform hot rolling offline. The sample was cooled under the same conditions as above, and the average cooling rate up to 500 ° C. was determined. In the present invention, since it is important to optimize the amount of Mn according to the cooling rate after rolling, Table 1 shows the cooling rate after rolling when each steel is hot-rolled, and the calculation based on the cooling rate. The appropriate range of Mn content is also shown.

The steel bar thus obtained was subjected to structure observation, tensile test and impact test.
For tissue observation. After the steel bar cross section was corroded with nital, the surface layer and the center of the steel bar were observed with an optical microscope.
In addition, a thin film sample is prepared by electropolishing from the surface layer and the center of the steel bar and observed with a transmission electron microscope (TEM) to measure the particle size of the precipitate, and an energy dispersive X-ray spectrometer (EDX) was used in combination to identify precipitates.

In the tensile test, a small diameter test piece having a parallel part diameter of 6 mmφ and a parallel part length of 40 mm was used in order to obtain two tensile test values of the surface layer part and the central part.
In the impact test, JIS 3 U-notch impact test pieces were used, and the absorbed energy at a test temperature of 20 ° C. at the center of the steel bar was measured.

Tables 3 and 5 show the results obtained by the above-described structure observation, tensile test, and impact test.
The numbers in the table are used to classify the individual results, and the steel numbers and hot-rolling conditions were combined to indicate the combination of the test steel and hot-rolling conditions (for example, steel numbers). 1 was numbered 1-A when hot rolled under condition A).

Regarding the structure, low temperature transformation phases such as F for ferrite, P for pearlite, bainite, etc. were generated, and the volume fraction was abbreviated as T when the volume fraction exceeded 60% or more.
For the precipitate, the average particle size is described. The variation in particle diameter was ± 1 nm at the maximum for precipitates of less than 10 nm, and ± 3 nm to ± 10 nm for precipitates larger than that. In addition, when a low-temperature transformation phase is generated in the structure, the dislocation density becomes high and observation of precipitates becomes difficult, so measurement of the particle diameter was omitted.

In the table, YS (center part) and TS (center part) are the yield strength and tensile strength of the steel bar center part, and △ YS is the difference between the yield strength of the steel bar surface part and the yield strength of the central part: YS (surface part)- YS (center part), △ TS is the difference between the tensile strength of the steel bar surface layer and the tensile strength of the center part: TS (surface layer part)-TS (center part), and _u E20 is the test temperature of the steel bar center part: 20 ° C Is the absorbed energy.
Here, it is one of the requirements of the present invention that the tensile strength is 700 MPa or more, the absolute value ΔYS is 100 MPa or less, and ΔTS is 80 MPa or less.

As is apparent from Tables 3 and 5, all of the inventive examples satisfying the appropriate range of the present invention in both steel composition and hot rolling conditions can obtain a high tensile strength of 700 MPa or more, and the surface layer portion and the central portion can be obtained. The difference in yield strength ΔYS is 100 MPa or less, and the difference in tensile strength ΔTS between the surface layer portion and the central portion is 80 MPa or less, which shows that the strength in the diameter direction is excellent. Furthermore, the absorbed energy _u E20 also at 100 J / cm ₂ or more, it can be seen the strength is high toughness as compared with the conventional example (No.10-A, 10-B , 10-C) of the same degree.

On the other hand, in the comparative example in which at least one of the steel composition and the hot rolling condition is outside the appropriate range of the present invention, the tensile strength, the difference ΔYS in the yield strength between the surface layer portion and the central portion, and the surface layer portion and the central portion Any of the tensile strength differences ΔTS do not meet the requirements of the present invention.
Since No. 6-A has a low C, the amount of fine precipitates is insufficient, and the tensile strength is low.
In No. 7-A, because C is too high, the precipitates are coarsened. Low tensile strength. Also, the absorbed energy _u E20 is as low as 88 J / cm _2, which is inferior in toughness.
Nos. 8-A and 16-B have a low Mn, so the precipitate in the center is coarsened and the tensile strength is low. Further, both the difference in yield strength ΔYS and the difference in tensile strength ΔTS between the surface layer portion and the central portion are large, and the uniformity in strength in the diameter direction is inferior. Furthermore, the absorbed energy _u E20 is also low and the toughness is also inferior.
No. 29-C also has a low Mn, but in this example, precipitates are coarsened in the surface layer portion and the central portion, and the tensile strength is low. Further, the absorption energy _u E20 also 85 J / cm ₂ and lower, inferior in toughness.
In Nos. 9-A, 17-B, and 30-C having a high Mn, a low temperature transformation phase is generated in the surface layer portion. Further, at least one of the difference in yield strength ΔYS and the difference in tensile strength ΔTS between the surface layer portion and the central portion is large, and the uniformity of the strength in the diameter direction is inferior.
No. 18-B has low tensile strength because Ti is low. On the other hand, in No. 19-B with high Ti, precipitates are coarsened in the surface layer portion and the central portion, the tensile strength is low, and the absorbed energy is also low.
No. 20-B has a low tensile strength because of its low Mo. On the other hand, in No. 21-B having a high Mo, a low temperature transformation phase is generated in both the surface layer portion and the central portion, and precipitation strengthening due to fine precipitates is insufficient, so that the tensile strength is low. Further, the difference in tensile strength ΔTS between the surface layer portion and the central portion is large, and the uniformity of the strength in the diameter direction is poor. Furthermore, the absorbed energy _u E20 is also as low as 51 J / cm ₂ and inferior in toughness.

Since No.1-G has a low heating temperature, precipitates such as Ti—Mo-based carbides precipitated on the slab before hot rolling are not sufficiently dissolved during heating, and the tensile strength is low. In addition, regarding the precipitate, what seems to have precipitated finely during cooling after rolling and what seems to be the undissolved residue of the precipitate deposited on the slab are mixed, and the average particle size of the precipitate is It was over 100nm.
Since No. 1-H had a low finishing temperature, the strain introduced by rolling changed the starting temperature of the ferrite transformation and affected the appropriate range of Mn content. For this reason, the precipitate is coarsened in the center, and the tensile strength is low. Further, both the difference in yield strength ΔYS and the difference in tensile strength ΔTS between the surface layer portion and the central portion are large, and the uniformity in strength in the diameter direction is inferior. Furthermore, the absorbed energy _u E20 is also low and the toughness is also inferior.
In this example, the fine precipitates (particle size <10 nm) observed for the inventive examples are mainly in the case of steel containing Mo, Ti carbide, or Nb, V, or W. Was able to be identified as a carbide containing Mo, Ti and any of Nb, V, and W.

It is the graph which showed the influence of the Mn amount in steel and the cooling rate after hot rolling on the steel structure and the size of precipitates.

Claims (4)

% By mass
C: 0.060 to 0.120%,
Si: 0.5% or less,
Mn: the range of the following formula (1),
Al: 0.1% or less,
Ti: 0.03-0.35% and
Mo: 0.05-0.8%
After heating the steel material containing Fe and the inevitable impurities composition to 1100 ° C or higher, finishing temperature: 850 ° C or higher and finishing diameter 120-400mm, and then cooling A method for producing a hot rolled non-heat treated steel bar characterized by the above.
Record
-0.239 × log (CR ₁ ) + 0.889 ≦ Mn ≦ −0.524 × log (CR ₂ ) +1.218 (1)
Where, CR ₁ : Average cooling rate at the center of the steel bar up to 500 ° C after rolling (° C / s)
CR ₂ : Average cooling rate of steel bar surface layer up to 500 ° C after rolling (° C / s) The method for producing a hot-rolled non-tempered steel bar according to claim 1, wherein the amounts of C, Ti and Mo in the steel satisfy the relationship of the following formula (2).
0.5 ≦ (C / 12) / [(Ti / 48) + (Mo / 96)] ≦ 1.5 (2) The steel material according to claim 1 or 2 , further in mass%.
Nb: 0.08% or less,
A method for producing a hot-rolled non-tempered steel bar, comprising a composition containing one or more selected from V: 0.15% or less and W: 1.5% or less. The method for producing a hot-rolled non-tempered steel bar according to claim 3, wherein the amounts of C, Ti, Mo, Nb, V and W in the steel satisfy the relationship of the following formula (3).
0.5 ≦ (C / 12) / [(Ti / 48) + (Mo / 96) + (Nb / 93) + (V / 51) + (W / 184)] ≦ 1.5 (3)

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