JP3823627B2 - Method for producing 60 kg grade non-tempered high strength steel excellent in weldability and toughness after strain aging - Google Patents

Description

【００６２】
【表２】

【００６３】
【表３】

【００６４】
【発明の効果】
本発明によれば、歪時効後の靭性に優れると共に、溶接性に優れる６０キロ級非調質鋼の製造方法の提供が可能で、産業上その効果は極めて大きい。[0001]
BACKGROUND OF THE INVENTION
This invention is a 60 kg class structural steel used for hydraulic iron pipes, pressure vessels, line pipes, marine structures, etc., and has excellent low temperature toughness even after cold working such as bending. The present invention relates to a method for producing excellent non-tempered high-tensile steel.
[0002]
[Prior art]
When steel is plastically deformed cold, a phenomenon called strain aging embrittlement occurs that deteriorates toughness. Research on strain aging embrittlement has been carried out mainly on thin steel sheets for automobile bodies, but in recent years, the demand for structural reliability has increased, and even thick steel sheets are processed not only at the material stage. The toughness after plastic deformation due to accidents or accidents has become a problem.
[0003]
As a test for evaluating strain aging embrittlement, a strain aging Charpy test in which 5% tensile pre-strain is applied and a Charpy test after aging treatment at 250 ° C. for 1 hour is known. Recently, it has been requested as one of material evaluation tests. Increasing cases are being conducted.
[0004]
As a technique for suppressing strain aging embrittlement for thick steel plates, there are JP-A-5-320820, JP-A-59-182915, JP-A-56-127750, and the like. It is not a technology for meat steel plates.
[0005]
Japanese Patent Laid-Open No. 5-320820 discloses a low yield point quenched steel for a spherical bow having a tensile strength of 400 MPa. The grain size of the steel material structure is adjusted to prevent toughness deterioration after strain aging, but the amount of C is 0.002 to 0.03%, and the component composition containing almost no other reinforcing elements is the target. It cannot be applied to 60kg class steel.
[0006]
Japanese Unexamined Patent Publication No. 59-182915 discloses a production method for suppressing strain aging embrittlement in TMCP type 500 MPa class steel. Paying attention to the fact that embrittlement after cold working occurs when TMCP 50 kg steel is cold worked, strain is concentrated in the ferrite phase of ferrite and bainite structure, cooling solid solution N and solid solution C in ferrite It is a technology that reduces the ferrite phase by embrittlement by controlling the stop temperature. For this reason, it cannot be applied to 60 kg grade steel which is cooled to near room temperature and becomes a quenched structure.
[0007]
Japanese Patent Application Laid-Open No. 56-127750 describes a technology for suppressing strain aging embrittlement of 600 MPa class steel. This technology is applied to strain aging embrittlement caused by VN precipitation type steel containing 0.01% or more of N. It can be suppressed by adding Ca or Mg. However, this technology is effective only for VN steel manufactured by as roll or normal, and the steel of the examples also has a high C content of 0.12% or more and Pcm of 0.25% or more. Steels that are inferior in weldability are described and do not meet the demands of current general customers.
[0008]
[Problems to be solved by the invention]
As described above, a technique for suppressing embrittlement after plastic deformation with a 60 kg thick steel material excellent in welding workability has not yet been completed. The present invention provides a method for producing a 60 kg grade non-tempered high strength steel having excellent weldability and excellent toughness even after strain aging. Specifically, the fracture surface transition of a strain aging Charpy test is provided. A method for producing a 60 kg grade non-tempered high strength steel having a temperature vTs (aged) of −40 ° C. or lower is provided.
[0009]
[Means for Solving the Problems]
Conventionally, regarding the toughness degradation mechanism after strain aging, in the case of thin steel sheets, the movement of dislocations is hindered by the interaction between C and N dissolved in the steel and the dislocations due to strain application, and the yield point increases. And it is known to become brittle. However, as a result of the snake peak measurement by the internal friction measurement method for the solid solutions C and N of steels having different toughness deterioration after strain aging, the inventors of the present invention have found that solid steel is not solid. The amount of dissolved C and N is less than 3 ppm, and the positive influence of the amount of dissolved C and N on the toughness deterioration after strain aging of thick steel is not recognized. Since steel is a bainite-based structure, it was assumed that it was due to structural differences.
[0010]
Therefore, the present inventors varied the composition of the steel material and the heat treatment conditions in various ways, examined in detail the influence of the structure on the strain aging characteristics, and described the strain aging characteristics peculiar to the Nb-containing thick steel described below. I figured it out.
[0011]
1. Decreasing the C content reduces toughness deterioration after strain aging.
[0012]
2. The toughness after strain aging is affected by the manufacturing process, and the quenching and tempering steel deteriorates as the tempering temperature increases, and the toughness after strain aging deteriorates. The degree of toughness deterioration after aging is small.
[0013]
3. The structure changes by tempering, and cementite precipitates relatively finely as the tempering temperature is lowered. The precipitation sites were observed to be dispersed within the prior austenite grain boundaries, the bainite packet boundaries, and the prior austenite grains. As the tempering temperature was increased, cementite was agglomerated and coarsened, and most of the precipitation sites became the former austenite grain boundaries and bainite packet boundaries. On the other hand, in the cooling stop material, cementite precipitates relatively finely regardless of the cooling stop temperature, and the precipitation sites are dispersed in the prior austenite grain boundaries, bainite packet boundaries, and the prior austenite crystal grains. Observed.
[0014]
4). The finer the austenite crystal grain size during quenching and the smaller the bainite packet size, the less the toughness deterioration after strain aging.
[0015]
5). When film-like ferrite having a thickness of several μm or less is formed at the prior austenite grain boundaries, the grain boundary area is substantially increased, and cementite is refined.
[0016]
6). The fracture unit of the brittle fracture surface in the Charpy impact test after strain aging corresponds to the bainite packet size.
[0017]
These results show that the toughness after strain aging of Nb-containing thick-walled steel is governed by the size and amount of cementite precipitated at the boundary between the prior austenite grain boundaries and the bainite packet, and the toughness after strain aging is improved. In order to achieve this, the size of the precipitated cementite is reduced and the amount thereof is reduced, which shows that the tempering omission type cooling stop process is excellent. And as a factor that affects the size and precipitation amount of cementite, the former has the effect of directly increasing the precipitation site with respect to the C amount and tempering temperature, indirectly with respect to a certain amount of cementite, and reducing the precipitation size. Austenite grains, bainite packet size, and film-like ferrite precipitated on the prior austenite grain boundaries were observed.
[0018]
That is, the main production conditions affecting the toughness after strain aging are the amount of C, the slab heating temperature that affects the prior austenite grain size and the bainite packet size, the rolling method in the recrystallization region, and the ferrite precipitation amount. It becomes the cumulative rolling reduction in the non-recrystallization temperature range that affects.
[0019]
The present invention has been made based on the above findings and further studies.
[0020]
1. In mass% , C: 0.04 to 0.09%, Si: 0.1 to 0.5%, Mn: 1.2 to 1.8%, Nb: 0.01 to 0.05%, sol . Al: 0.002 to 0.07%, N: 0.001 to 0.004%, and satisfy Pcm ≦ 0.20%, Ceq (WES) ≦ 0.42% , remaining Fe and inevitable The steel made of impurities is heated to 1150 ° C. or lower, heated and then rolled at 1 d / hm ≧ 1.0 in the temperature range of 900 to 1000 ° C. for one or more passes, and then cumulatively reduced in the temperature range of Ar 3 to less than 900 ° C. 60kg class excellent in weldability and toughness after strain aging, characterized by cooling to a temperature range of 300-600 ° C at a cooling rate of 2 ° C / second or higher than Ar3 after rolling at a rate of 10-60% Production method of non-tempered high-tensile steel.
[0021]
However, Pcm = C + Mn / 20 + Si / 30 + Cu / 20 + Ni / 60 + Cr / 20 + Mo / 15 + V / 10 + 5B,
Ceq (WES) = C + Mn / 6 + Si / 24 + Ni / 40 + Cr / 5 + Mo / 4 + V / 14.
[0022]
Id: Projected contact arc length Id = (R · (hi−ho)) ^1/2
hm: Average plate thickness hm = (hi + 2ho) / 3
R: roll radius, hi: plate thickness before rolling, ho: plate thickness after rolling The method for producing a 60 kg grade non-tempered high strength steel excellent in weldability and toughness after strain aging according to 1, further comprising Cr: 0.1 to 0.5% by mass as a steel composition.
[0023]
3. The weldability and toughness after strain aging according to 1 or 2, wherein the steel composition further contains one or two of Mo: 0.02-0.3% and Cu: 0.1-0.6% in mass% . Of excellent 60kg grade non-tempered high strength steel.
[0024]
4). The steel composition further contains Ni: 0.1 to 0.5% by mass% . 60 kg class non-heat treated high excellent in weldability and toughness after strain aging according to any one of claims 1 to 3. Tensile steel manufacturing method.
[0025]
5). 60 kg class non-tempered high-tensile steel excellent in weldability and toughness after strain aging according to any one of 1 to 4, further comprising, as a steel composition, V: 0.01 to 0.08% by mass% Manufacturing method.
[0026]
6). As a steel composition, weldability and strain according to any one of 1 to 5 further containing one or two of Ti: 0.005 to 0.02% and Ca: 0.001 to 0.004% by mass%. A method for producing 60 kg grade non-tempered high strength steel with excellent toughness after aging.
[0028]
DETAILED DESCRIPTION OF THE INVENTION
The component composition and production conditions in the present invention will be described below.
[0029]
1. Production condition C: 0.04% or more and 0.09% or less C is added to ensure a predetermined strength. If it is less than 0.04%, it is difficult to secure a tensile strength of 60 kg in the case of a thick material. If it exceeds 0.09%, the toughness after strain aging deteriorates, so 0.04% or more and 0.0. Add 09% or less.
[0030]
Si: 0.1% or more and 0.5% or less Si is added to ensure predetermined strength and toughness. If the content is less than 0.1%, the effect is not sufficient. If the content exceeds 0.5%, the effect is saturated and the toughness of the weld heat affected zone is remarkably deteriorated. Therefore, 0.1% to 0.5% is added.
[0031]
Mn: 1.2% or more and 1.8% or less Mn is added to ensure a predetermined strength. If it is less than 1.2%, it is difficult to secure a tensile strength of 60 kg in the case of a thick material, and if it exceeds 1.8%, the toughness of the weld heat-affected zone is significantly deteriorated, so 1.2% or more 1 Add 8% or less.
[0032]
Nb: 0.01% or more and 0.05% or less Nb suppresses recrystallization of austenite during rolling, activates austenite grain boundaries during direct quenching, and facilitates the formation of film ferrite. Moreover, it precipitates as Nb carbide at the time of tempering and is added because it is effective for increasing the strength. If it is less than 0.01%, these effects are insufficient, and if it exceeds 0.05%, the toughness deteriorates due to significant precipitation strengthening of Nb carbide, so 0.01% or more and 0.05% or less are added.
[0033]
sol. Al: 0.002% or more and 0.07% or less Al is added for deoxidation. sol. If the amount of Al is less than 0.002%, the effect is not sufficient, and if added over 0.07%, surface flaws of the steel material are likely to occur, so 0.002% or more and 0.07% or less are added. .
[0034]
N: 0.001% or more and 0.004% or less N is combined with Al or Ti at the time of rolling and heating to produce AlN and TiN, thereby refining austenite. If the content is less than 0.001%, the effect is not sufficient. If the content exceeds 0.004%, significant strain aging embrittlement occurs due to solute N even after quenching and tempering, so 0.001% or more and 0.004% or less. To do.
[0035]
Pcm ≦ 0.20, Ceq (WES) ≦ 0.42
Pcm, Ceq (WES) is an index of weld cold cracking property and toughness of the heat affected zone. If Pcm exceeds 0.20%, cold cracking may occur in welding without preheating. When WES) exceeds 0.42, the heat-affected zone toughness of high heat input welding is remarkably deteriorated, so Pcm ≦ 0.20 and CeqWES ≦ 0.42. Here, Pcm = C + Mn / 20 + Si / 30 + Cu / 20 + Ni / 60 + Cr / 20 + Mo / 15 + V / 10 + 5B, Ceq (WES) = C + Mn / 6 + Si / 24 + Ni / 40 + Cr / 5 + Mo / 4 + V / 14.
[0036]
The above is the basic component composition of the steel of the present invention, but Cr, Mo, Cu, Ni, V, Ti, and Ca can be added alone or in combination in order to improve desired characteristics.
[0037]
Cr: 0.1% to 0.5% Cr is added to ensure strength and toughness. If it is less than 0.1%, the effect is insufficient. If it exceeds 0.5%, the weldability and the toughness of the weld heat affected zone are remarkably deteriorated, so 0.1% or more and 0.5% or less are added.
[0038]
Mo: 0.02% or more and 0.3% or less, Cu: 0.1% or more and 0.6% or less of one or two kinds of Mo improve strength, and are added because they are particularly effective for thick materials. If it is less than 0.02%, the effect is not sufficient, and if it exceeds 0.3%, the weldability and the toughness of the heat affected zone are significantly deteriorated, so 0.02% or more and 0.3% or less. Cu is added to improve the strength.
[0039]
If it is less than 0.1%, the effect is not sufficient, and if it exceeds 0.6%, there is a concern about Cu cracking, so 0.1% or more and 0.6% or less.
[0040]
Ni: 0.1% or more and 0.5% or less Ni is added to improve toughness. If it is less than 0.1%, the effect is not sufficient, and if it exceeds 0.5%, the steel material cost is remarkably increased.
[0041]
V: 0.01% or more and 0.08% or less V is added to precipitate as carbide during tempering and to improve strength. If it is less than 0.01%, the effect is not sufficient, and if it exceeds 0.08%, the toughness deteriorates due to significant precipitation strengthening of V carbide, so the content is made 0.01% or more and 0.08% or less.
[0042]
Ti: 0.005% or more and 0.02% or less, Ca: 0.001% or more and 0.004% or less of one or two types of Ti and Ca are added to improve the base material toughness and the toughness of the heat affected zone. . Ti produces TiN at the time of rolling and heating or welding to refine the austenite grain size.
[0043]
If less than 0.005%, the effect is not sufficient, and if added over 0.02%, TiNb composite carbide precipitates during rolling, and the precipitation amount of Nb carbide during tempering becomes insufficient, resulting in a decrease in strength. Therefore, the content is made 0.005% or more and 0.02% or less.
[0044]
Ca is present in the steel as Ca sulfide and refines the austenite grain size at the time of rolling heating or welding. If it is less than 0.001%, the effect is not sufficient, and if added over 0.004%, the cleanness is remarkably deteriorated by a large amount of Ca-sulfuric acid, so 0.001% or more and 0.004% or less.
[0045]
Furthermore, in the present invention, it is desirable to regulate B, O, P, and S within the following ranges.
[0046]
B: 0.0002% or less, O: 0.001% or more and 0.004% or less B is treated as an impurity element in the present invention. If it is present as solute B during direct quenching, the formation of film-like ferrite at the prior austenite grain boundaries is suppressed, so the content is restricted to 0.0002% or less by selecting the melting raw material. O is an unavoidable impurity, but if it is less than 0.001%, the production cost becomes expensive, and if it exceeds 0.004%, a large amount of Ca sulfate is collected and the cleanliness deteriorates, so 0.001% Above 0.004%.
[0047]
P ≦ 0.010%, S ≦ 0.002%
P and S are impurity elements. When P ≦ 0.010% and S ≦ 0.002%, central segregation is reduced, and the toughness and weldability at the center of the plate thickness are improved.
[0048]
2. Rolling under rolling conditions of 900 to 1000 ° C .: 900 to 1000 which is a low temperature range of the recrystallization temperature of Nb-containing steel in order to suppress toughness deterioration after strain aging of one or more passes for rolling satisfying ld / hm ≧ 1.0. Rolling that satisfies ld / hm ≧ 1.0 at 1 ° C. is performed at 1 ° C. for at least one pass, and processing strain is effectively introduced to the center of the plate thickness in a temperature range where recovery is fast.
[0049]
When the rolling temperature is less than 900 ° C., recrystallization is not sufficient, and when it exceeds 1000 ° C., the recrystallized grain size becomes large, so the temperature is set to 900 to 1000 ° C. If ld / hm is less than 1.0, sufficient processing strain to excite recrystallization up to the center of the plate thickness is not added, so 1.0 or more.
[0050]
Rolling in a temperature range of Ar 3 or more and less than 900 ° C .: In order to suppress toughness deterioration after rolling strain aging with a cumulative reduction rate of 10 to 60%, accumulation of work strain due to cumulative reduction in the non-recrystallization temperature range of Nb-containing steel To promote ferrite transformation.
[0051]
If the rolling temperature is less than Ar3, the hardenability deteriorates due to the progress of ferrite transformation at the start of direct quenching, the predetermined strength cannot be obtained, and if it is 900 ° C. or higher, processing strain does not accumulate and ferrite transformation becomes insufficient due to recrystallization. Therefore, Ar3 is set to be lower than 900 ° C. If the cumulative rolling reduction is less than 10%, the ferrite transformation is not promoted. If the cumulative rolling reduction exceeds 60%, the effect is saturated and the anisotropy of the steel material increases rapidly, so the cumulative rolling reduction is set to 10 to 60%. Ar3 is obtained, for example, as Ar3 = 910-310C-80Mn-20Cu-15Cr-55Ni-80Mo.
[0052]
3. Heat treatment condition Cooling start temperature: Ar3 or higher When the cooling start temperature is lower than Ar3, the fraction of the austenite phase that transforms into a bainite phase decreases due to accelerated cooling, and a predetermined strength cannot be obtained.
[0053]
Cooling rate: 2 ° C./second or more The higher the cooling rate, the more effective the strength rises.
[0054]
Cooling stop temperature: 300-600 ° C
The cooling stop temperature affects the strength-toughness balance, and if it is less than 300 ° C., the toughness is not recovered, and if it exceeds 600 ° C., the cooling effect becomes insufficient and a predetermined strength cannot be obtained.
[0055]
If the slab heating temperature exceeds 1150 ° C., the austenite crystal grains become coarser, which makes it difficult to refine by subsequent rolling and may deteriorate the toughness after strain aging. Is preferred.
[0056]
【Example】
Table 1 shows the chemical components of the test steel used in the examples (the remainder not shown is substantially composed of Fe and inevitable impurities). The slab having these chemical components was heated and then rolled to 25 to 75 mm. After rolling, the cooling start temperature, the cooling rate and the cooling stop temperature were variously changed, and the characteristics were investigated. Table 2 shows the manufacturing conditions, and Table 3 shows the characteristics of the steel sheet.
[0057]
As mechanical properties, strength, toughness, and toughness after strain aging were determined. The tensile test was a test using a collected JIS14A (14φ) test piece from 1/4 t.
[0058]
The impact test was a test using a 2 mm V notch Charpy impact test piece (JIS No. 4 standard test piece) taken so that the longitudinal direction was perpendicular to the rolling direction from 1/4 t. The toughness after strain aging was tested by giving a 5% tensile pre-strain to a plate-shaped test piece, and after aging treatment at 250 ° C. for 1 hour, a 2 mm V notch Charpy impact test piece was taken in the tensile direction.
[0059]
Hereinafter, an Example and a reference example are explained in full detail.
Steel types A, B, C, D, and F in Table 1 are steels having a composition that satisfies the invention according to any one of claims 1 to 6, and steel type G has a C content and Pcm is outside the scope of the invention. . Moreover, the steel type E of Table 1 shows the steel of the component composition corresponding to a reference example. Steel numbers 1 to 5 and 7 in Tables 2 and 3 are production examples using steel types A, B, C, D and F , respectively, and are examples of the invention according to any one of claims 1 to 6 . Moreover, steel number 6 in Tables 2 and 3 shows a production example using steel type E.
[0060]
The vTs after strain aging is −40 ° C. or less, the change in vTs before and after strain aging is small, and good strain aging embrittlement resistance is obtained. Steel numbers 8 and 9 are production examples using steel types A and B, but the cooling conditions are outside the scope of the present invention. Steel No. 8 has a low cooling start temperature, and Steel No. 9 has a high cooling stop temperature and a low strength. Steel No. 10 has a slab heating temperature higher than 1150 ° C. and slightly inferior in strain aging embrittlement resistance. Steel No. 11 did not undergo rolling with an ld / hm of 1.0 or more in the recrystallization temperature range, so the degree of degradation of vTs due to strain aging is large. Steel No. 12 has a low cumulative rolling reduction and a large degree of degradation of vTs due to strain aging. Steel No. 13 has a too low cooling stop temperature and is inferior in toughness after strain aging. Steel No. 14 has a slow cooling rate and a decrease in strength. Although it is a manufacture example by the steel type F, the steel number 15 is a manufacture example by the steel type G, and a component composition is outside the range of this invention, and it is inferior to the strain aging embrittlement resistance.
[0061]
[Table 1]

[0062]
[Table 2]

[0063]
[Table 3]

[0064]
【The invention's effect】
ADVANTAGE OF THE INVENTION According to this invention, while providing the toughness after strain aging, it is possible to provide the manufacturing method of 60 kg grade non-heat-treated steel which is excellent in weldability, and the effect is very large industrially.

Claims (6)

In mass% , C: 0.04 to 0.09%, Si: 0.1 to 0.5%, Mn: 1.2 to 1.8%, Nb: 0.01 to 0.05%, sol . Al: 0.002 to 0.07%, N: 0.001 to 0.004%, and satisfy Pcm ≦ 0.20%, Ceq (WES) ≦ 0.42% , remaining Fe and inevitable The steel made of impurities is heated to 1150 ° C. or lower, heated and then rolled at 1 d / hm ≧ 1.0 in the temperature range of 900 to 1000 ° C. for one or more passes, and then cumulatively reduced in the temperature range of Ar 3 to less than 900 ° C. 60kg class excellent in weldability and toughness after strain aging, characterized by cooling to a temperature range of 300-600 ° C at a cooling rate of 2 ° C / second or higher than Ar3 after rolling at a rate of 10-60% Production method of non-tempered high-tensile steel.
However, Pcm = C + Mn / 20 + Si / 30 + Cu / 20 + Ni / 60 + Cr / 20
+ Mo / 15 + V / 10 + 5B,
Ceq (WES) = C + Mn / 6 + Si / 24 + Ni / 40 + Cr / 5 + Mo / 4
+ V / 14.
Id: Projected contact arc length Id = (R · (hi−ho)) ^1/2
hm: Average plate thickness hm = (hi + 2ho) / 3
R: roll radius, hi: plate thickness before rolling, ho: plate thickness after rolling The method for producing a 60 kg grade non-tempered high strength steel excellent in weldability and toughness after strain aging according to claim 1, further comprising Cr: 0.1 to 0.5% by mass as a steel composition. The weldability and strain aging according to claim 1 or 2, further comprising, as a steel composition, one or two of Mo: 0.02-0.3% and Cu: 0.1-0.6% in mass%. Of 60kg grade non-tempered high strength steel with excellent toughness. The steel composition further contains Ni: 0.1 to 0.5% by mass% . 60 kg class non-heat treated high excellent in weldability and toughness after strain aging according to any one of claims 1 to 3. Tensile steel manufacturing method. The steel composition further contains V: 0.01 to 0.08% by mass% , 60 kg class non-refined high excellent in weldability and toughness after strain aging according to any one of claims 1 to 4. Tensile steel manufacturing method. The weldability according to any one of claims 1 to 5, wherein the steel composition further contains one or two of Ti: 0.005 to 0.02% and Ca: 0.001 to 0.004% in terms of mass% . And the manufacturing method of 60 kg class non-tempered high-tensile steel excellent in toughness after strain aging.