High Tension Steel Plate with Small Acoustic Anisotropy and with Excellent Weldability and Method of Production of Same

Abstract

The present invention provides a high tension steel plate of a tensile strength of 570 MPa or more with a small acoustic anisotropy and with excellent weldability tensile strength and a method of production of steel plate predicated on the high productivity accelerated cooling stop process, which steel sheet is a high tension steel plate with a value of 0.045%≦Nb+2Ti≦0.105%, A=(Nb+2Ti)×(C+N×12/14) satisfying 0.0025 to 0.0055 and with a steel composition containing bainite in an amount of 30 vol % or more and pearlite+martensite-austenite constituent (MA) in an amount of less than 5 vol %. The steel is heated to 1200° C. or more, rough rolled at 1020° C. or more, then rolled at a cumulative reduction rate from over 920° C. to less than 1020° C. of 15% or less and a cumulative reduction rate from 860 to 920° C. of 20 to 50%, then starting accelerated cooling giving a cooling rate of 2 to 30° C./sec at 800° C. or more, stopping the accelerated cooling at 600 to 700° C., then cooling by 0.4° C./sec or less.

Description

TECHNICAL FIELD

The present invention relates to a method of production of high tension steel plate with a small acoustic anisotropy and with excellent weldability having a tensile strength of the 570 MPa or more with a high productivity without requiring off-line heat treatment. The steel of the present invention is used for thick plate, steel pipe, or steel shapes as structural members of bridges, ships, building structures, marine structures, pressure vessels, penstocks, line pipes, and other welded structures.

BACKGROUND ART

High tension steel plate of a tensile strength of the 570 MPa or more used for structural members of bridges, ships, building structures, marine structures, pressure vessels, penstocks, line pipes, etc. is required to feature not only strength, but also toughness and weldability and in recent years is often particularly required to feature weldability with large heat input. Numerous studies have been made in the past to improve these characteristics. The composition and production conditions of such steel plate are disclosed in for example Japanese Patent Publication (A) No. 53-119219, Japanese Patent Publication (A) No. 01-149923, etc. These relate to methods of production comprising rolling steel plate, then reheat quenching it off-line and further reheating it for tempering heat treatment. Further, for example, Japanese Patent Publication (A) No. 52-081014, Japanese Patent Publication (A) No. 63-033521, Japanese Patent Publication (A) No. 02-205627, etc. disclose technology relating to production by quenching steel plate on-line after rolling, that is, so-called direct quenching. These require off-line tempering heat treatment both with reheat quenching and direct quenching. However, the requirement of a heat treatment process off-line inevitably ends up obstructing the productivity, so to raise the productivity, a so-called “as-rolled” method of production eliminating tempering heat treatment and not requiring off-line heat treatment is desirable.

Technology relating to the as-rolled method of production has been disclosed in several publications, for example, Japanese Patent Publication (A) No. 54-021917, Japanese Patent Publication (A) No. 54-071714, Japanese Patent Publication (A) No. 2001-064723, Japanese Patent Publication (A) No. 2001-064728, etc. These relate to the accelerated cooling stop process which stops the accelerated cooling after rolling the steel plate. Thus uses accelerated cooling to quench the steel down to the transformation temperature or less to obtain a hardened structure and stops the water cooling in a state of a relatively high temperature after transformation to shift to the gradual cooling process to obtain the tempering effect by this gradual cooling process and thereby eliminate the reheat tempering. However, these production technologies all require controlled rolling under a relatively low temperature for obtaining toughness and strength. The temperature where the rolling is finished is around 800° C., so time is required for waiting for the temperature and the productivity cannot be said to be high. On the other hand, in particular in bridges, buildings, and other applications, a small acoustic anisotropy is required since anisotropy affects the precision of the angle beam ultrasonic testing of the weld zone. In controlled rolling finishing the rolling at a 800° C. or so temperature, texture is formed, so the steel plate becomes larger in acoustic anisotropy and therefore will not necessarily match these applications.

Further, the above Japanese Patent Publication (A) No. 2001-064728 discloses technology for production of high tension steel plate having a tensile strength of 570 MPa or more by the accelerated cooling stop process. However, in this patent, V is considered to contribute to precipitation strengthening even at the gradual cooling stage after stopping the water-cooling midway, but the inventors' studies found that, as explained later, V has a slower precipitation rate in the gradual cooling stage after stopping the water-cooling midway compared with Nb and Ti and is not that effective for strengthening. With this composition of ingredients, a stable strength probably cannot necessarily be stably obtained.

Further, Japanese Patent Publication (A) No. 2002-053912 discloses an as-rolled process not performing water cooling after rolling. Since this does not involve controlled rolling at a low temperature, the acoustic anisotropy does not become large, but instead, to obtain strength, the amounts of addition of alloys of Cu, Ni, Mn, etc. become greater and therefore there is a problem with economy.

DISCLOSURE OF INVENTION

Therefore, the present invention has as its object to obtain high tension steel plate of a tensile strength of 570 MPa or more with a small acoustic anisotropy and with excellent weldability by an economical composition of ingredients with little amounts of addition of alloys and a high productivity method of production predicated on an accelerated cooling stop process. The thickness of the steel plate covered is up to 100 mm.

There are several means for strengthening high tension steel, but the method of utilizing of precipitation strengthening by carbides or nitrides of Nb, V, Ti, Mo, and Cr etc. enables strengthening with relatively small amounts of alloy ingredients. At that time, to obtain a large amount of precipitation strengthening, it becomes important to form a precipitate with coherence with the base material.

In the accelerated cooling stop process, at the state of rolling, the steel structure is austenite. The accelerated cooling after the end of the rolling causes this to transform to a bainite or ferrite or other ferrite base material structure. The precipitate formed in the austenite during rolling loses its coherence with the ferrite base material and becomes smaller in strengthening effect after transformation. Further, the precipitate formed at an early stage of the rolling becomes coarser and becomes a factor lowering the toughness. Therefore, it is important to suppress precipitation during the rolling and cause precipitation as much as possible in the bainite or ferrite structure at the stage of gradual cooling after stopping the water-cooling. If a process involving water cooling, then reheating for tempering heat treatment, the temperature and time for precipitation can be sufficiently secured, so large precipitation strengthening can be easily obtained. As opposed to this, in the case of the accelerated cooling stop process without reheat tempering, precipitation can be expected during the gradual cooling after the water-cooling is stopped, but to obtain a hardened structure, the water-cooling stop temperature has to be made a low temperature to a certain extent, so both the temperature and time for precipitation are restricted. This is generally disadvantageous for precipitation strengthening. From this, as explained above, the as-rolled process has a high productivity, but requires a larger amount of alloy elements for obtaining the same strength as the thermal refinement process or requires controlled rolling at a low temperature.

Therefore, the inventors engaged in in-depth studies on a method for utilizing precipitation strengthening to the maximum extent so as to obtain a high strength predicated on the high productivity accelerated cooling stop process and without adding a large amount of alloy elements or controlled rolling at a low temperature.

First, to clarify the precipitation behavior in the process of gradual cooling after stopping the water-cooling, the inventors studied in depth the relationships between the rates of precipitation of carbides, nitrides, and carbonitrides of the various alloy elements in a bainite or ferrite structure and the amount of precipitation strengthening and the temperature and holding time. As a result, in a bainite or ferrite structure or their mixed structures, the rates of precipitation of Nb carbonitrides and Ti carbides are faster than those of V and other elements and these form precipitates with coherence with the base material, so the amount of strengthening is large. In particular, the rate of precipitation is fast and the amount of strengthening is large in the 600° C. to 700° C. temperature range. Further, when making joint use of Nb and Ti or Nb, Ti, and Mo for composite precipitation, due to the synergistic effect, even with short time holding, precipitates coherent with the base material finely disperse and large precipitation strengthening can be obtained.

However, if the amounts of addition of Nb and Ti are too large, the precipitates formed tend to become coarse and the number of precipitates conversely becomes smaller, so the amount of precipitation strengthening falls. The rates of precipitation and the form of precipitation of carbides, nitrides, and carbonitrides of Nb and Ti in austenite and ferrite are greatly affected by the amounts of addition of Nb and Ti and by the amounts of C and N. The inventors engaged in various experiments and analyses and obtained the discovery that the rates of precipitation and the form of precipitation of carbides, nitrides, and carbonitrides of Nb and Ti can be organized well by the parameter A=([Nb]+2×[Ti])×([C]+[N]×12/14) and that keeping this value within a certain range enables sufficient precipitation during the gradual cooling after stopping water-cooling while suppressing precipitation during rolling. That is, the greater the amounts of addition of Nb and Ti, the smaller the amounts of addition of C and N must be made. If the value of A is too small, the rate of precipitation in ferrite becomes slower and sufficient precipitation strengthening cannot be obtained. Conversely, if the value of A is too large, the rates of precipitation of carbides, nitrides, and carbonitrides in the austenite become too fast, the precipitates become coarser, and the amounts of coherent precipitation during the gradual cooling after stopping the accelerated cooling become insufficient, so the amount of precipitation strengthening falls. Further, Si also has an effect on the rate of formation of carbides, so a certain range of amount of addition is necessary.

These precipitation strengthening effects are greatly influenced by the structure. With a bainite structure, as compared with ferrite, it is easy to maintain the dislocation density and other worked structures. To promote fine coherent precipitation, the sufficient presence of dislocation, deformation zones, or other precipitation sites in the worked structure acts extremely effectively. According to the studies of the inventors, to obtain sufficient strengthening, a single bainite phase or a mixed structure of bainite and ferrite with a volume fraction of bainite of 30% or more is necessary. Further, pearlite, martensite-austenite constituent (MA), etc. precipitate at the phase interfaces resulting in a smaller strengthening effect and a drop in toughness etc., so the sum of the volume fractions of pearlite and martensite-austenite constituent (MA) has to be suppressed to 3% or less.

The inventors further studied the specific production conditions for obtaining the maximum extent of precipitation strengthening and obtained the following discoveries.

The precipitation of Nb and Ti at the rolling stage is promoted by rolling strain, so the rolling conditions in the high temperature region of austenite, the so-called rough rolling conditions, have a large effect on the final precipitation strengthening effect. Specifically, finishing the rough rolling in the temperature range of 1020° C. or more and avoiding rolling as much as possible in the temperature range of 1020° C. to 920° C. are requirements for suppressing precipitation during rolling. However, if finishing all rolling in the temperature range of 1020° C. or more, almost no worked structure will remain after the accelerated cooling is stopped due to recovery and recrystallization, so dislocation, deformation zones, and other precipitation sites will not be sufficiently present and sufficient precipitation strengthening cannot be obtained. Therefore, it is essential to perform the necessary and sufficient rolling in the not yet recrystallization temperature range, then quickly perform the accelerated cooling after rolling. Specifically, relatively light rolling with a cumulative reduction rate of 20% to 50% is performed in the range of 920° C. to 860° C. Under these conditions, it is possible to suppress unnecessary precipitation of Nb and Ti and leave suitable precipitation sites even after stopping the water-cooling stop. Further, under these conditions, no strong texture is formed, so the acoustic anisotropy also does not become great.

The temperature of stopping the water-cooling in the accelerated cooling stop process is made a temperature of the center of the sheet thickness of 600° C. to 700° C. advantageous for the precipitation of Nb and Ti. To obtain a steel structure with a volume fraction of bainite of 30% or more even with t his stopping temperature, it is necessary to limit the composition of ingredients of the steel to the later explained specific range and ensure at cooling rate of 2° C./sec to 30° C./sec in accelerated cooling. Further, it is essential to heat the slab at a high temperature for dissolution of the Nb and Ti. A heating temperature of 1200° C. or more is necessary.

The discovery obtained here is the new idea of controlling on-line the precipitation of carbides or carbonitrides of Nb and Ti during the rolling including the high temperature range, during the accelerated cooling, and until the gradual cooling process after stopping the cooling. A precipitation strengthening equal to or greater than that of the conventional thermal refining process can be realized by an accelerated cooling stop process not requiring off-line heat treatment.

Further, according to this production process, it is possible to keep down the weld crack sensitivity index Pcm (Pcm=[C]+[Si]/30+[Mn]/20+[Cu]/20+[Ni]/60+[Cr]/20+[Mo]/15+[V]/10+5[B]: where [C], [Si], [Mn], [Cu], [Ni], [Cr], [Mo], [V], and [B] mean the mass % of C, Si, Mn, Cu, Ni, Cr, Mo, V, and B) of the composition of the steel material and possible to provide a steel material with a high welding heat affected zone toughness even with a large heat input and with excellent weldability.

The gist of the present invention is as follows:

(1) A high tension steel plate of a tensile strength of 570 MPa or more with a small acoustic anisotropy and with excellent weldability, characterized by having a composition of steel containing, by mass %,

• • C: 0.03% to 0.07%,
  • Si: 0.1 to 0.6%,
  • Mn: 0.8 to 2.0%,
  • Al: 0.003% to 0.1%,
  • Nb: 0.025 to 0.1%,
  • Ti: 0.005 to 0.1%,
  • [Nb]+2×[Ti]: 0.045 to 0.105%
  • N: over 0.0025% to 0.008%,
    containing Nb, Ti, C, and N in a range where the value A expressed by the A of the following equation (1) is 0.0022 to 0.0055, and containing a balance of Fe and unavoidable impurities, having a steel structure with a volume fraction of bainite of 30% or more and a sum of volume fractions of pearlite and martensite-austenite constituent (MA) of less than 5%:

A=([Nb]+2×[Ti])×([C]+[N]×12/14)  (1)

• • where [Nb], [Ti], [C], and [N] mean the mass$ of Nb, Ti, C, and N

(2) A high tension steel plate of a tensile strength of 570 MPa or more with a small acoustic anisotropy and with excellent weldability as set forth in (1), characterized in that said steel sheet further contains, by mass %, one or more of

• • Mo: 0.05% to 0.3%,
  • Cu: 0.1% to 0.8%,
  • Ni: 0.1% to 1%,
  • Cr: 0.1% to 0.8%,
  • V: 0.01% to 0.03%,
  • W: 0.1% to 3%,
  • B: 0.0005% to 0.005%,
  • Mg: 0.0005% to 0.01%,
  • Ca: 0.0005% to 0.01%.

(3) A high tension steel plate of a tensile strength of 570 MPa or more with a small acoustic anisotropy and with excellent weldability characterized by heating a slab having a composition of ingredients as set forth in (1) or (2) to 1200° C. to 1300° C., rough rolling it at a temperature range of 1020° C. or more, then hot rolling it by a cumulative reduction rate of 15% or less in the temperature range of less than 1020° C. to over 920° C. and by a cumulative reduction rate of 20% or more in the temperature range of 920° C. to 860° C., then starting accelerated cooling by a cooling rate of 2° C./sec to 30° C./sec from a temperature of 800° C. or more, stopping said accelerated cooling at a temperature of the center of plate thickness of 700° C. to 600° C., then cooling by a cooling rate of 0.4° C./sec or less.

BEST MODE FOR WORKING THE INVENTION

Below, the reasons for limitation of the ingredients and method of production of the present invention will be explained.

C is an important element forming carbides and carbonitrides with Nb and Ti and becoming a principal factor in the strengthening mechanism of the steel of the present invention. If the amount of C is insufficient, the amount of precipitation during the gradual cooling after the accelerated cooling is stopped becomes insufficient and strength cannot be obtained. Conversely, even if in excess, the precipitation rate in the austenite region during rolling becomes faster. As a result, the amount of coherent precipitation during the gradual cooling after the accelerated cooling is stopped becomes insufficient and strength cannot be obtained. Therefore, the amount of C is limited to 0.03% to 0.07% in range.

Si is an element necessary as a deoxidizing element in steel making and also has an effect on the rate of precipitation of carbides. Suitable addition of Si has the effect of suppressing the precipitation of carbides in the austenite region during rolling. To achieve this object, Si is added in an amount of 0.1% or more, preferably 0.3% or more. However, if added over 0.6%, the precipitation rate becomes too slow. Further, the welding heat affected zone is made to drop in toughness, so 0.6% is made the upper limit.

Mn is an element necessary for improving the hardenability and obtaining a single bainite phase or obtaining a mixed structure of bainite and ferrite with a bainite percentage of 30% or more. To obtain this object, 0.8% or more is necessary, but if added over 2.0%, sometimes the matrix material is made to drop in toughness, so the upper limit is made 2.0%.

Al is added in the usual range as a usual deoxidizing element, that is, 0.003% to 0.1%.

Nb and Ti are important elements forming NbC, Nb(CN), TiC, TiN, Ti(CN), or their composite precipitates and composite precipitates of these with Mo and thereby forming main elements of the strengthening mechanism of the steel of the present invention. In the accelerated cooling stop process, to obtain a sufficient composite precipitate, suitable-addition considering the precipitation rate is necessary. That is, Nb is 0.025% or more, preferably 0.035% or more, and Ti is 0.005% or more, conditional on 0.045%≦([Nb]+2×[Ti])≦0.105% and, when A=([Nb]+2×[Ti])×([C]+[N]×12/14), the value of A being 0.0022 to 0.0055 (where, [Nb], [Ti], [C], and [N] respectively mean the mass % of Nb, Ti, C, and N). Note that the upper limit values of Nb and Ti are respectively made 0.1%.

Mo improves the hardenability, forms a complex precipitate with Nb and Ti, and greatly contributes to strengthening, so 0.05% or more is added. However, if excessively added, the welding heat affected zone toughness is inhibited, so the amount added is made 0.3% or less.

N bonds with Ti to form TiN. When TiN finely disperses, its pinning effect suppresses the coarsening of the structure of the welding heat affected zone and improves the welding heat affected zone toughness, but if the N is insufficient, the TiN becomes coarser and a pinning effect cannot be obtained. To make the TiN finely disperse, N is added in an amount over 0.0025%, preferably over 0.004%. Further, if excessively containing N, conversely sometimes it causes the matrix material to drop in toughness, so the upper limit is made 0.008%.

Cu has to be added in at least 0.1% to exhibit its effect when added as a strengthening element, but even if added over 0.8%, the effect does not increase in proportion to the amount of addition and if excessively added, sometimes impairs the welding heat affected zone toughness, so the amount is made 0.8% or more.

Ni has to be added in at least 0.1% when added to raise the matrix material toughness, but if excessively added, sometimes inhibits the weldability. It is also an expensive element, so the upper limit of addition is made 1%.

Cr has the effect of raising the hardenability in the same way as Mn and of facilitating obtaining a bainite structure. For this purpose, 0.1% or more is added, but if excessively added, impairs the welding heat affected zone toughness, so the upper limit is made 0.8%.

V has the effect of increasing the precipitation strengthening and hardenability a certain extent though having less of a strengthening effect compared with Nb and Ti. To obtain this effect, addition of 0.01% or more is necessary. If excessively added, a drop in the welding heat affected zone toughness is caused, so even when added, the amount is made less than 0.03%.

B increases the hardenability. When added to obtain strength, addition of 0.0005% or more is necessary, but even if added over 0.005%, the effect does not change, so the amount of addition is made 0.0005% to 0.005%. By adding one or both of Mg and Ca, it is possible to form sulfides or oxides and improve the matrix material toughness and welding heat affected zone toughness. To obtain this effect, Mg or Ca has to be added in amounts of 0.0005% or more. However, if excessively added over 0.01%, coarse sulfides or oxides are formed, so conversely the toughness is lowered in some cases. Therefore, the amounts of addition are 0.0005% or more and 0.01% or less, respectively.

As unavoidable impurities other than the above ingredients, P and S are harmful elements causing the matrix material to drop in toughness, so the amounts are preferably small. Preferably, P is made 0.02% or less, while S is made 0.02% or less.

Next, the method of production will be explained.

To cause sufficient solid solution of Nb and Ti, the heating temperature of the slab at the time of rolling must be made 1200° C. or more. However, even if the heating temperature exceeds 1300° C., the effect of dissolution does not change that much and the energy cost becomes higher, so the heating temperature of the slab at the time of rolling is made 1200° C. to 1300°.

In the rolling, to suppress precipitation of Nb and Ti during rolling as much as possible, after the rough rolling at a suitable reduction rate at a temperature range of 1020° C. or more, rolling in the range of less than 1020° C. to over 920° C. is performed by a cumulative reduction rate of 15% or less. Further, to obtain a worked structure necessary and sufficient as a precipitation site, the rolling is performed in a range of 920° C. to 860° C. by a cumulative reduction rate of 20% to 50%. With these rolling conditions, formation of texture is suppressed, so the acoustic anisotropy does not become large.

To suppress recovery of the worked structure and precipitation after work, soon after the finish of rolling, accelerated cooling is performed. The cooling is performed by water cooling from 800° C. or more under conditions of a cooling rate at the center part in the plate thickness of 2° C./sec to 30° C./sec. To make the volume fraction of bainite 30% or more, a 2° C./sec or higher cooling rate is necessary. To make the sum of the volume fractions of pearlite and martensite-austenite constituent (MA) less than 3%, the cooling rate is made 30° C./sec or less. The water cooling is stopped midway so that the temperature at the center of the plate thickness becomes 700° C. to 600° C., then cooling rate is made 0.4° C./sec or less by natural cooling. The object is to secure sufficient temperature and time for precipitation of Nb and Ti and composite precipitation of these and composite precipitation with Mo. If the water-cooling stop temperature is too high, a bainite structure becomes hard to obtain, the precipitation becomes slower at a low temperature, and sufficient strengthening cannot be obtained.

The steel of the present invention is used in the form of thick plate, steel pipe, or steel shapes as structural members of bridges, ships, building structures, marine structures, pressure vessels, penstocks, line pipes, and other welded structures.

EXAMPLES

Slabs obtained by producing steel of the compositions of ingredients shown in Table 1 were processed under the production conditions shown in Table 2 and Table 3 to 20 to 100 mm thick steel plates. Among these, 1-A to 14-N are steels of the present invention, while 15-O to 43-A are comparative examples. In the tables, underlined numerical values indicate ingredients or production conditions outside the scope of the patent or characteristics not satisfying the following target values.

The results of measurement of the tensile strength, welding heat affected zone toughness, and acoustic anisotropy for these steel plates are shown in Table 2. The tensile strength was measured by obtaining a No. 10 rod test piece defined in JIS Z2201 and testing it by the method defined in JIS Z2241. The matrix material toughness was evaluated by obtaining an impact test piece defined in JIS Z2202 from the center of thickness of the plate in the direction perpendicular to the rolling direction and finding the fracture appearance transition temperature (vTrs) by the method defined in JIS Z2242. The welding heat affected zone toughness was evaluated by the absorption energy at −20° C. (vE₋₂₀) of an impact test piece defined in JIS Z2202 given a thermal cycle corresponding to submerged arc welding with an amount of input heat of 20 kJ/mm. For a steel material with a thickness of 32 mm or less, the original thickness steel material was prepared. For a steel material with a thickness of over 32 mm, a steel plate reduced in thickness to 32 mm was prepared. The V-shaped abutted part was welded by large input heat submerged arc welding with an amount of input heat of 20 kJ/mm. An impact test piece defined in JIS Z2202 was taken so that the notch bottom ran along the fusion line and evaluated by the absorption energy (vE₋₂₀) at −20° C. The acoustic anisotropy was evaluated in accordance with the Japanese Society for Non-Destructive Inspection standard NDIS2413-86. A sound speed contrast of 1.02 or less was evaluated as a small acoustic anisotropy. The target values of the characteristics are a yield strength of 450 MPa, a tensile strength of 570 MPa or more, a vTrs of −20° C. or less, a vE₋₂₀ of 70 J or more, and a sound speed contrast of 1.02 or less.

Examples 1-A to 14-N all have yield strengths over 450 MPa, tensile strengths over 570 MPa, welding heat affected zone toughnesses vE₋₂₀ over 200 J, and sound speed contrasts of 1.02 or less or small acoustic anisotropies.

As opposed to this, Comparative Example 15-O has a low C, Comparative Example 16-P has a high C, Comparative Example 17-Q has a low Si, Comparative Example 19-S has a low Mn, Comparative Example 21-U has a low Mo, Comparative Example 23-W has a low Nb, Comparative Example 25-Y has a low Ti, Comparative Example 27-AA has a value of the parameter A (A=([Nb]+2×[Ti])×([C]+[N]×12/14)) of less than 0.0025, Comparative Example 37-A has a low heating temperature, Comparative Example 40-A has a high cumulative reduction rate in the range from 920° C. to 860° C., Comparative Example 41-A has a small plate thickness center cooling rate, Comparative Example 42-A has a high accelerated cooling stop temperature, and Comparative Example 43-A has a low accelerated cooling stop temperature, so their tensile strengths were less than 570 MPa.

Comparative Example 18-R has a high Si, Comparative Example 22-V has a high Mo, Comparative Example 24-X has a high Nb and an Nb+2Ti over 0.105%, Comparative Example 26-Z has a high Ti and an Nb+2Ti over 0.105%, Comparative Example 2 g-AC has a low N, Comparative Example 31-AE has a high V, Comparative Example 32-AF has a high Cu, Comparative Example 33-AG has a high Ni, Comparative Example 34-AH has a high Cr, Comparative Example 35-AI has a high Mg, and Comparative Example 36-AJ has a high Ca, so their welding heat affected zone toughnesses are low.

Comparative Example 20-T has a high Mn, Comparative Example 28-AB has a value of the parameter A of over 0.005, and Comparative Example 30-AD has a high N, so their matrix material toughnesses are low.

Comparative Example 38-A has a high cumulative reduction rate in the range from less than 1020° C. to over 920° C. and Comparative Example 3 g-A has a low cumulative reduction rate in the range from 920° C. to 860° C., so have low tensile strengths and have low welding heat affected zone toughnesses.

Comparative Example 3 g-A has a high cumulative reduction rate in the range from 920° C. to 860° C., so has a low tensile strength and a large acoustic anisotropy.

	TABLE 1
	Steel	Chemical composition (mass %)
	mat..	C	Si	Mn	P	S	Cu	Ni	Cr	Mo	Al	Nb	Ti	Nb + 2Ti	A**	V	B	Mg	Ca	N	Pcm
Inv.	A	0.04	0.35	1.55	0.008	0.003				0.07	0.03	0.051	0.014	0.079	0.0036					0.0053	0.134
steel	B	0.04	0.31	1.28	0.008	0.005				0.27	0.03	0.038	0.015	0.068	0.0032					0.0056	0.132
	C	0.03	0.37	1.52	0.006	0.002				0.13	0.04	0.049	0.022	0.093	0.0036					0.0071	0.127
	D	0.05	0.33	1.75	0.015	0.005				0.10	0.02	0.078	0.006	0.090	0.0049					0.0042	0.155
	E	0.06	0.51	1.56	0.009	0.002				0.24	0.05	0.051	0.014	0.079	0.0052					0.0055	0.171
	E	0.05	0.41	1.49	0.012	0.002				0.09	0.05	0.050	0.016	0.082	0.0045	0.020				0.0042	0.146
	G	0.04	0.35	1.54	0.008	0.003				0.15	0.01	0.028	0.032	0.092	0.0045		0.0009			0.0077	0.143
	H	0.04	0.22	1.42	0.005	0.004		0.41		0.12	0.01	0.054	0.015	0.084	0.0038					0.0041	0.133
	I	0.07	0.33	0.98	0.006	0.005			0.67	0.12	0.05	0.043	0.011	0.065	0.0049					0.0048	0.172
	J	0.05	0.54	1.37	0.007	0.011	0.32			0.19	0.04	0.041	0.021	0.083	0.0048					0.0064	0.165
	K	0.04	0.41	0.88	0.012	0.007				0.24	0.02	0.046	0.008	0.062	0.0027	0.015		0.009		0.0029	0.115
	L	0.03	0.37	1.32	0.008	0.004	0.31	0.24		0.08	0.07	0.051	0.016	0.083	0.0030		0.0032			0.0055	0.149
	H	0.04	0.33	1.54	0.010	0.004				0.12	0.006	0.056	0.012	0.080	0.0037			0.0034		0.0050	0.136
	N	0.04	0.30	1.38	0.007	0.003				0.24	0.02	0.057	0.009	0.075	0.0034				0.0019	0.0049	0.135
Comp.	O	0.01	0.41	1.61	0.006	0.004				0.21	0.03	0.065	0.018	0.101	0.0017					0.0058	0.118
steel	P	0.09	0.25	1.32	0.005	0.004				0.22	0.03	0.036	0.019	0.074	0.0071					0.0051	0.179
	Q	0.06	0.11	1.58	0.005	0.003				0.16	0.05	0.037	0.019	0.075	0.0049					0.0050	0.153
	R	0.06	1.22	1.37	0.002	0.002				0.18	0.03	0.048	0.015	0.078	0.0051					0.0049	0.181
	S	0.07	0.38	0.52	0.006	0.003				0.21	0.04	0.043	0.013	0.069	0.0053					0.0054	0.123
	T	0.05	0.25	2.22	0.005	0.006				0.19	0.03	0.047	0.011	0.069	0.0038					0.0048	0.182
	U	0.05	0.44	1.51	0.003	0.005				0.02	0.03	0.048	0.021	0.090	0.0050					0.0049	0.142
	V	0.06	0.51	1.39	0.060	0.003				0.51	0.05	0.047	0.011	0.069	0.0046					0.0055	0.181
	W	0.06	0.55	1.55	0.010	0.005				0.24	0.03	0.015	0.022	0.059	0.0039					0.0055	0.172
	X	0.04	0.29	1.35	0.005	0.003				0.19	0.02	0.102	0.009	0.120	0.0054					0.0041	0.130
	Y	0.06	0.30	1.52	0.004	0.004				0.25	0.03	0.061	0.002	0.065	0.0042					0.0046	0.163
	Z	0.04	0.29	1.61	0.009	0.003				0.18	0.03	0.026	0.041	0.108	0.0050					0.0051	0.142
	AA	0.03	0.31	1.49	0.006	0.003				0.21	0.04	0.036	0.011	0.058	0.0020					0.0042	0.129
	AB	0.06	0.33	1.56	0.008	0.002				0.19	0.04	0.052	0.023	0.098	0.0065					0.0055	0.162
	AC	0.04	0.33	1.59	0.003	0.003				0.18	0.02	0.048	0.024	0.096	0.0041					0.0022	0.143
	AD	0.04	0.34	1.57	0.004	0.003				0.19	0.03	0.049	0.016	0.081	0.0044					0.0021	0.143
	AE	0.05	0.35	1.45	0.005	0.005				0.23	0.02	0.039	0.015	0.069	0.0039	0.06				0.0052	0.156
	AF	0.04	0.36	1.05	0.006	0.004	1.55			0.11	0.04	0.036	0.031	0.098	0.0045					0.0049	0.189
	AG	0.06	0.35	1.32	0.050	0.030		1.81		0.12	0.04	0.045	0.017	0.079	0.0052					0.0048	0.176
	AH	0.05	0.32	0.96	0.004	0.003			1.22	0.21	0.03	0.051	0.021	0.093	0.0053					0.0061	0.184
	AI	0.05	0.29	1.52	0.005	0.005				0.22	0.03	0.056	0.018	0.092	0.0051			0.015		0.0044	0.150
	AJ	0.04	0.37	1.41	0.003	0.002				0.15	0.04	0.051	0.016	0.083	0.0039				0.013	0.0055	0.133
*Pcm = C + Si/30 + Mn/20 + Cu/20 + Ni/60 + Cr/20 + Mo/15 + V/10 + 5B
**A = (Nb + 2Ti) × (C + N × 12/14)

	TABLE 2
				Cumulative	Cumulative	Cooling	Water					Welded
	Pro-			reduction	reduction	rate at	cooling				Matrix	heat	Acoustic
	duction		Heating	rate	rate	plate	stop	Plate			material	affected	anisotropy
	con-		temp. at	at 1020° C.	at 920° C. to	thickness	temper-	thick-	Yield	Tensile	toughness	zone	(sound
	dition	Steel	rolling	920° C.	860° C.	center	ature	ness	strength	strength	vTrs	toughness	speed
	No	mat.	(° C.)	(%)	(%)	(° C./sec)	(° C.)	(mm)	(Mpa)	(Mpa)	(° C.)	vE-20 (J)	contrast)
Inv.	1	A	1200	0	46	13	660	32	520	617	−63	215	1.01
ex.	2	B	1220	0	45	30	620	20	532	635	−72	204	1.01
	3	C	1230	0	50	9	590	50	515	610	−61	213	1.01
	4	D	1230	0	40	5	610	75	502	605	−51	229	1.02
	5	E	1220	0	35	3.5	570	100	522	633	−44	208	1.01
	6	F	1230	0	40	25	650	20	546	650	−69	232	1.01
	7	G	1200	10	46	16	630	32	489	589	−70	202	1.01
	8	H	1200	0	22	10	650	50	533	625	−50	212	1.00
	9	I	1220	0	29	8	580	75	509	607	−47	204	1.01
	10	J	1250	0	36	7	600	75	518	625	−57	216	1.01
	11	K	1220	0	45	13	640	32	555	660	−60	231	1.02
	12	L	1220	0	42	11	590	50	512	613	−57	222	1.01
	13	M	1230	0	40	14	670	32	542	628	−72	233	1.02
	14	N	1220	0	40	24	680	25	526	629	−68	218	1.01

	TABLE 3
				Cumulative	Cumulative	Cooling	Water					Welded
	Pro-			reduction	reduction	rate at	cooling				Matrix	heat	Acoustic
	duction		Heating	rate	rate	plate	stop	Plate			material	affected	anisotropy
	con-		temp. at	at 1020° C.	at 920° C. to	thickness	temper-	thick-	Yield	Tensile	toughness	zone	(sound
	dition	Steel	rolling	920° C.	860° C.	center	ature	ness	strength	strength	vTrs	toughness	speed
	No	mat.	(° C.)	(%)	(%)	(° C./sec)	(° C.)	(mm)	(Mpa)	(Mpa)	(° C.)	vE-20 (J)	contrast)
Comp.	15	0	1220	0	40	13	660	32	427	543	−72	215	1.02
ex.	16	P	1220	0	42	13	630	32	444	563	−25	110	1.02
	17	Q	1220	0	38	18	640	32	426	551	−30	128	1.01
	18	R	1250	0	35	22	640	32	468	602	−39	43	1.01
	19	S	1220	0	46	10	600	40	433	555	−66	222	1.02
	20	T	1220	0	33	21	630	32	537	677	−5	125	1.01
	21	U	1220	0	40	12	590	40	444	546	−79	204	1.01
	22	V	1220	0	45	24	660	32	498	667	−24	37	1.01
	23	W	1220	0	28	10	670	40	419	547	−66	215	1.02
	24	x	1220	0	33	12	650	32	457	575	−35	24	1.01
	25	Y	1250	0	36	13	650	32	435	550	−80	220	1.02
	26	z	1220	0	40	13	660	32	468	602	−40	21	1.02
	27	AA	1220	0	40	14	670	32	422	537	−52	210	1.01
	28	AB	1220	0	40	13	650	32	439	561	−24	18	1.01
	29	AC	1220	0	40	14	650	32	514	612	−21	51	1.01
	30	AD	1220	0	38	13	640	32	527	630	−10	56	1.01
	31	AE	1230	0	32	11	680	40	518	617	−35	31	1.02
	32	AF	1220	0	40	13	600	32	531	635	−44	48	1.01
	33	AG	1220	0	33	15	630	32	514	622	−80	46	1.02
	34	AH	1220	0	33	14	590	32	510	630	−26	19	1.02
	35	AI	1220	0	33	16	660	32	511	635	−8	27	1.01
	36	AJ	1220	0	40	15	670	32	520	625	−10	50	1.02
	37	A	1150	0	33	18	670	32	441	546	−65	232	1.01
	38	A	1220	33	45	16	660	32	418	525	−40	52	1.02
	39	A	1220	0	10	17	670	32	438	557	−22	22	1.02
	40	A	1220	0	66	15	630	32	443	551	−38	95	1.04
	41	A	1220	0	36	1	620	32	405	522	−45	202	1.02
	42	A	1220	0	33	22	740	32	444	541	−35	110	1.02
	43	A	1220	0	40	19	480	32	446	563	−30	100	1.02

INDUSTRIAL APPLICABILITY

According to the present invention, it is possible to obtain high tension steel plate with a small acoustic anisotropy, with excellent weldability, and with a tensile strength of 570 MPa or more up to a plate thickness of 100 mm by economical ingredients with little amounts of addition of alloys and by a high productivity as-rolled method of production. The effect on the industry is extremely great.

Claims

1. A high tension steel plate of a tensile strength of 570 MPa or more with a small acoustic anisotropy and with excellent weldability, characterized by having a composition of steel containing, by mass %, containing Nb, Ti, C, and N in a range where the value A expressed by the A of the following equation (1) is 0.0022 to 0.0055, and containing a balance of Fe and unavoidable impurities, having a steel structure with a volume fraction of bainite of 30% or more and a sum of volume fractions of pearlite and martensite-austenite constituent (MA) of less than 5%:

C: 0.03% to 0.07%,

Si: 0.1 to 0.6%,

Mn: 0.8 to 2.0%,

Al: 0.003% to 0.1%,

Nb: 0.025 to 0.1%,

Ti: 0.005 to 0.1%,

[Nb]+2×[Ti]: 0.045 to 0.105%

N: over 0.0025% to 0.008%,

A=([Nb]+2×[Ti])×([C]+[N]x×12/14)  (1)

where [Nb], [Ti], [C], and [N] mean the mass$ of Nb, Ti, C, and N

2. A high tension steel plate of a tensile strength of 570 MPa or more with a small acoustic anisotropy and with excellent weldability as set forth in claim 1, characterized in that said steel sheet further contains, by mass %, one or more of

Mo: 0.05% to 0.3%,

Cu: 0.1% to 0.8%,

Ni: 0.1% to 1%,

Cr: 0.1% to 0.8%,

V: 0.01% to 0.03%,

W: 0.1% to 3%,

B: 0.0005% to 0.005%,

Mg: 0.0005% to 0.01%,

Ca: 0.0005% to 0.01%.

3. A high tension steel plate of a tensile strength of 570 MPa or more with a small acoustic anisotropy and with excellent weldability characterized by heating a slab having a composition of ingredients as set forth in claim 1 to 1200° C. to 1300° C., rough rolling it at a temperature range of 1020° C. or more, then hot rolling it by a cumulative reduction rate of 15% or less in the temperature range of less than 1020° C. to over 920° C. and by a cumulative reduction rate of 20% or more in the temperature range of 920° C. to 860° C., then starting accelerated cooling by a cooling rate of 2° C./sec to 30° C./sec from a temperature of 800° C. or more, stopping said accelerated cooling at a temperature of the center of plate thickness of 700° C. to 600° C., then cooling by a cooling rate of 0.4° C./sec or less.