CN112921150A

CN112921150A - Manufacturing method of aluminum-free low-alloy steel plate suitable for large-line weldable

Info

Publication number: CN112921150A
Application number: CN202110126372.9A
Authority: CN
Inventors: ***; 赵和明; 王川; 曲之国; 杨海峰; 李廷刚; 刘振华
Original assignee: Minmetals Yingkou Medium Plate Co ltd
Current assignee: Minmetals Yingkou Medium Plate Co ltd
Priority date: 2021-01-29
Filing date: 2021-01-29
Publication date: 2021-06-08

Abstract

A smelting method of aluminum-free low alloy steel suitable for large-wire welding comprises the following steps of refining in a converter, LF (ladle furnace) and RH (relative humidity) treatment, wherein the molten steel is obtained, by weight, with the free [ O ] less than or equal to 15PPm, P less than or equal to 0.008%, Cu less than or equal to 0.30%, Ni less than or equal to 0.30%, Mo less than or equal to 0.06%, V: not less than 4X [ N ] to 0.06%, S: not more than 0.003 percent, B not more than 0.0008 percent, Cr not more than 0.30 percent, Alt not more than 0.007 percent, C: 0.03 to 0.07%, Mn: 1.30-1.60%, Si: 0.15-0.30%, Ti: 0.011-0.022%, [ N ] 35-60 ppm, Nb:0.015 to 0.025%, and the balance Fe and inevitable impurity elements in the steel. The molten steel is used as a raw material, and the produced casting blank can be used for rolling a large-wire weldable thick steel plate by a TMCP (thermal mechanical control processing) process, so that the process is simple, the production continuity is good, the component control is accurate, and the welding performance of the finished large-wire weldable steel plate is reliable.

Description

Manufacturing method of aluminum-free low-alloy steel plate suitable for large-line weldable

Technical Field

The invention relates to a smelting method of low alloy steel which does not contain aluminum and is suitable for large-scale wire welding, belonging to the field of low alloy steel manufacturing.

Background

Welding is the most prominent way of joining steel structures. During the welding process of the steel plate, the structure and the performance of a welding heat affected zone can be changed due to the action of welding heat circulation, and particularly, the low-temperature toughness can be greatly reduced. In order to ensure the quality of a steel structure welding joint, the traditional structural steel generally needs to strictly limit welding line energy, execute welding bead layout and control interlayer temperature, and sometimes needs to be preheated before welding. Although the reliability of the welded structure can be effectively guaranteed through strict standardization work, the welding work efficiency is greatly limited, and therefore the manufacturing efficiency of the steel structure is limited. Under the circumstances that the economic development is fast, the construction of various steel structures is increasingly demanded, and the welding work amount of thick steel plates is increasing because the steel structures are also becoming larger, engineers have proposed an idea of developing steel materials for large-line weldability, and have proposed various research schemes.

Summarizing the basic design idea, the method comprises three aspects: one is that the welding heat affected zone is inhibited from being excessively hardened and the M-A phase is rarely generated by controlling alloy elements and carbon equivalent; secondly, the influence of the segregation of impurity elements on the low-temperature toughness of the HAZ zone is reduced through high-cleanness control and organization and multi-interface of the HAZ zone; and thirdly, the steel matrix and the corresponding HAZ zone phase composition need to have good toughness. The second point is the most important point, because under the condition of large welding line energy, the degradation zone of the HAZ for long-time high-temperature retention is wider, and in the fusion line and the adjacent zone, because the temperature is higher, the dwell time is longer, the coarsening degree of the reheated austenite is more serious, the structure converted in the subsequent cooling process is coarser, and the reduction of the toughness is more obvious. Since the total amount of impurity elements can be controlled by clean metallurgy, the key issue becomes how to improve the texture of the HAZ deteriorated region. Therefore, the method is started from three aspects, one is to control the reheated austenite not to be excessively coarsened (so that the grain size after phase transformation is in accordance with the expectation); secondly, coarse grain boundary ferrite (net) does not appear in the reheated austenite grain boundary; and thirdly, controlling the reheated austenite to generate multi-point nucleation mass transfer in the crystal during the temperature reduction process to induce phase transformation, so that the side plate ferrite nucleated by the crystal boundary is shortened and is divided by the synchronous phase transformation structure in the crystal to form a more complex polycrystalline fit structure.

Based on the design idea, engineers introduce a method that TiN particles precipitated in a dispersion way and the like are not completely dissolved in a reheating process, and residual nitrides reach a critical dimension effective nail rolling austenite crystal boundary, so as to inhibit austenite coarsening.

Such as document 1 "CN 1962916A" is based on this assumption. However, this document only considers suppression of austenite coarsening in the HAZ region by undissolved TiN particles, does not consider austenite intragranular induced nucleation, and has limited improvement in HAZ toughness when large-wire-weldable.

Document 2 "CN 103343284A" also intends to suppress austenite coarsening in the HAZ region by using undissolved TiN, and does not consider the reduction of austenite grain-oriented nucleation phase, and the Nb and Ti contents used are high, which causes problems in slab quality and is not economical. In addition, too high Ti and N contents have made it possible to precipitate TiN in the molten steel in a liquid state, and such TiN particles are coarse, greatly affecting the homogeneity of the steel sheet base material, and are not reliable.

In the document 3, "CN 104498827A", the base material performance is improved by controlling the low carbon equivalent and the low carbon equivalent P, S, so as to ensure that the HAZ zone also has good toughness, and the toughness of the HAZ zone under the condition of large-line welding cannot be completely ensured.

Document 4 "CN 102459656A" intends to prevent the merging of austenite grain boundaries by forming an oxide of Ti or a composite oxide of Ti and Mg as a core of a superheated austenite intragrain induced phase transformation core and binding excess N and B to BN. Therefore, austenite coarsening and multi-phase synchronous phase transformation in original austenite crystal are simultaneously controlled, and a complex polycrystalline conjunction structure is formed. On one hand, BN is easy to dissolve, and the capability of preventing the combination of grain boundaries is limited when large wires can be welded; on the other hand, there is no indication of the process route for controlling the oxide to exert an induced nucleation effect, i.e. how to achieve a fine and diffuse presence of beneficial inclusions by metallurgical means. In the actual smelting process, oxygen reacts with various active elements in the range given in document 4, and thereafter the remaining portion is bonded with Ti. These oxide potentials must polymerize with the oxides of Ti to form large particles that emerge from the steel, making the amount and distribution of oxides in the steel uncontrollable. And large-particle impurities left in the steel can destroy the continuity of the steel matrix, and simultaneously, the function of intragranular induced mass transfer phase change can be lost due to overlarge size, so that the toughness is not improved.

Document 5 "CN 106574316B" proposes improvement of toughness by controlled rolling with a low carbon equivalent content. In the aspect of component control, the toughness of the matrix is guaranteed by controlling low carbon, low silicon, adding Ni and Cu and controlling carbon equivalent; controlling the sulfur content and the calcium content, wherein the mass transfer induced phase change nucleation points in the austenite crystal are obtained by forming CaS; controlling the content of titanium and nitrogen and the content of B, and inhibiting the coarsening of overheated austenite grains by forming TiN or using TiN as a nucleation core for inducing ferrite phase transition; meanwhile, nitrides and oxides forming elements such as Al \ Nb \ V \ Zr, Mg, REM and the like are also added into the steel. It is well known that the design of low Si, low carbon equivalent and addition of Ni and Cu is beneficial to improve the toughness of steel, and the control of the weight percentage of calcium remaining after the completion of the combination with oxygen is 1.25 times lower than that of sulfur, which is the first feature of this document, the concept of adding Ti \ B and controlling the N content is similar to documents 1 and 4, and the concept of adding Zr, Mg, REM, etc. is similar to document 4. There is a problem how various easily oxidizable elements do not form oxides in a liquid phase, and they are polymerized and floated out, and at the same time, the formed oxides can exist in a fine dispersion state in the steel. As described herein, the preferable range of Ti is 0.01 to 0.035%, the preferable range of Al is 0.01 to 0.10%, and the preferable range of B is 0.0008 to 0.0025%, in which case N in the steel should preferentially combine with Ti and Al, and the large amount of B content sharply increases the hardenability of the steel, i.e., B tends to be easily segregated to grain boundaries during reheating to inhibit diffusion of carbon, and when strong carbide-forming elements such as Mo, V, Cr, and Nb are present in the steel, side lath bainite or more MA phase appears after transformation, and the toughness of the HAZ region may be sharply impaired.

Document 6 "CN 103114241 a" proposes measures to control the ultra-low carbon content, to add Ni and Cu, to control low P and low S, to make the steel have high N and Ti/N values of 1 to 2, and to add B and to control Al to 0.01% or less, with the intention of suppressing the growth of reheated austenite by forming TiN with a high remelting temperature. However, the content of N is too high, coarse TiN is easily precipitated in a liquid state, the TiN distribution is not uniform, the effect of inhibiting the coarsening of the overheated austenite is not achieved, meanwhile, excessive N exists, the heat affected zone is easy to embrittle after welding on the basis of insufficient nitrogen fixation elements, the literature does not provide measures for controlling the uniform dispersion distribution of TiN, and the control of the size and the distribution of a second phase for inhibiting the grain boundary migration is undoubtedly one of the technical keys of the large-wire weldable steel.

The document 7 "CN 107287508A" is actually a low-carbon microalloyed steel, the components of which are very common, and the base metal can have high impact toughness by adopting the TMCP process, but the performance of the HAZ region which can be welded on a large line is uncertain, and the document does not actually provide a substantial technology suitable for the welding on the large line.

Document 8 "CN 106756543 a" is actually a low-carbon microalloyed steel, which substantially limits the carbon equivalent, reduces P \ S, and produces a steel sheet by a trace Ca treatment. Compared with the existing common high-performance steel, the method does not show a substantial technology aiming at large-wire welding.

Document 9 "CN 106906413 a" also proposes a composition scheme intended to realize a nucleation core with dispersed fine oxides as the reheat austenite intracrystalline mass transfer induced synergistic phase transformation in the HAZ region by composition control. The composition design emphasizes that the S content is less than 0.007 percent but not less than 0.0015 percent, limits the range of Al, Ti, Mg, N and the like, and proposes that the Mg/Ti value is not less than 0.017, and the S content after subtracting one-to-one molar ratio of S to Ca, REM, Zr, Mg and the like is 0.0003 to 0.003 percent and the like. However, it is difficult to obtain the desired effect regardless of the element added and the size and distribution of the inclusions which are not controlled in time and formation order in the metallurgical process, because it is determined whether the transformation is induced intragranular and the transformation is induced in the HAZ region, and the transformation is performed by the superheated austenite to have a suitable polycrystalline fit structure, and it is the size and distribution of the inclusions which are formed to ensure sufficient toughness in the superheated region under the excessively high weld line energy.

Document 10 "CN 108677088A" proposes a composition scheme and a smelting scheme in which the contents of boron and nitrogen in steel are increased based on low-carbon low-silicon nickel-added manganese-containing steel, and after one-to-one boron with the nitrogen content as a molar ratio is subtracted, the remaining boron is 0.001% to 0.0020%. On one hand, the proposal is more expensive and is not economical, on the other hand, BN is used as a core of an induced phase deformation core in HAZ overheated austenite crystal, which is actually possibly unfavorable because free boron and nitrogen are elements easy to segregate in a crystal boundary, and non-redissolved BN is too large and rare, and on the third hand, too high N is used, after excessive B is added, a coarse BN precipitated phase (inclusion grade) is formed in the crystal boundary in the solidification process, the continuity of a final steel plate matrix is damaged, on the fourth hand, the redissolved temperature of fine BN is too low and is difficult to play a role in nailing and rolling overheated austenite crystal boundary at too high temperature, and the coarse BN has no crystal lattice coordination relationship with the austenite phase and plays a role in nailing and rolling overheated austenite crystal boundary.

Document 11 "CN 109097685 a" discloses a method for producing a large-wire weldable steel plate in which coarsening of weld reheating austenite is suppressed by TiN, which is characterized in that based on a low-phosphorus low-sulfur low-carbon manganese-containing low-alloy steel, AL deoxidation is performed, acid-soluble aluminum is retained by 0.02% or more, microtitanium treatment is performed, 0.0008 to 0.0015% of B element is added, 0.04 to 0.10% of V is added, and N is controlled to be 0.007 to 0.009% and not more than 0.003+0.28Ti +0.03V +1.27B, and a cast slab is heated and then rolled by three-stage controlled rolling to obtain a steel plate. A toughness measurement of the HAZ zone of the steel sheet after welding at a weld heat input of 100KJ/cm is given. The method mainly depends on TiN to inhibit reheating austenite from coarsening to ensure HAZ toughness, the component control range of N is narrow, and the applicable welding heat input is low.

Document 11 "CN 109161671 a" discloses a method for manufacturing steel for large-wire welding based on low-carbon niobium-titanium microalloyed low alloy steel, which is characterized in that a converter and a refining electric furnace are used for smelting steel with the final component range of C: 0.06% -0.18%, Si: 0.15 to 0.50 percent of steel, 1.10 to 1.60 percent of Mn, less than or equal to 0.012 percent of P, less than or equal to 0.003 percent of S, 0.10 to 0.40 percent of Ni, 0.010 to 0.030 percent of Nb, less than or equal to 0.010 percent of Al, 0.010 to 0.030 percent of Ti and 0.002 to 0.010 percent of Ca, wherein Al deoxidation is not adopted in refining after smelting in a converter, Al element is controlled to be less than or equal to 0.0075 percent, free oxygen in the steel is in the range of 10 to 100PPm, Ti is added for microalloying, small Ti oxides are formed in the steel, and calcium treatment is carried out after LF refining and RH treatment. Although the manufacturing method can obtain the steel plate which is suitable for large-line welding and has good welding HAZ superheat zone performance, the Ti oxide formed after Ti deoxidation can be polymerized after a certain time, Ca line treatment is further carried out after RH treatment, CaO can be formed due to the presence of introduced oxygen element and other oxides in the steel, the Ti-Ca series oxide can be coarsened, the number of oxide particles in unit volume is too low, the large-line welding performance among batches is inconsistent, meanwhile, the nitrogen content is higher due to the fact that nitrogen is too little in the steel, and the HAZ zone outside the coarse crystal zone can be subjected to aging embrittlement caused by nitrogen during welding.

Document 12 "CN 109321847 a" discloses a thick steel plate for EH 420-grade ocean engineering capable of large heat input welding and a preparation method thereof, which is characterized in that the composition of inclusions with the diameter of 0.5-5 μm in the steel is analyzed, the steel is weakly deoxidized after a converter, free oxygen level is controlled to be 10-100 PPm for Ti deoxidation, magnesium wires and aluminum wires are fed when the oxygen level reaches 5-50 PPm, then RH treatment is carried out, and Ca wires are fed. The formation of Al-Mg-Ti-Ca-Mn-O-S composite inclusions is intended. The inclusion is formed in the LF refining stage and then undergoes RH treatment time, so that the inclusion is coarsened and floated, the finally obtained molten steel has uncertain quality, and the large-wire welding reliability is not high.

Document 13 "CN 109321815 a" discloses a method for manufacturing a high heat input welding resistant high strength thick steel plate, which is characterized in that the steel is alloyed, then secondary oxygen blowing is performed in an RH process, after oxygen is determined to 10 to 100PPm, Ti alloying is performed, then one of Ca, Mg, Ca, Zr and REM is added, argon blowing and stirring are performed, and continuous casting is performed, so as to obtain Ti-M-O composite inclusions. Since blowing oxygen at RH causes preferential oxidation of C, Si, Mn, Al, etc., resulting in loss of components, and the secondary oxidation product formed in this way has uncontrollable particle size, and the resulting inclusion particle size and distribution are unreliable, it is not a perfect solution.

Document 14 "CN 109321846 a" discloses a method for producing a steel for high heat input welding having a yield strength of 355MPa level, which is similar to document 12, but differs from the deoxidation process in that Si, Mn-based weak deoxidizers are used to deoxidize in the first half of the tapping process, Ti iron is added to deoxidize in the second half, and thereafter, white slag is formed in LF. Because the free oxygen after the smelting of the converter is up to over 700PPm, the turbulence of the molten steel is violent in the tapping process, so that Ti oxides in the steel are not easy to form uniform and stable distribution, and most of the Ti oxides float out. During the subsequent further refining, the oxides of Al, Ca and the like can continuously polymerize the original oxides and float out of the slag, so the granularity and the bulk density of the oxides in the finally obtained steel are uncertain, and the reliable welding performance of the final large-wire energy welding is not easy to guarantee.

Document 15 "CN 109321816 a" discloses a method for producing a steel for high heat input welding having a yield strength of 460MPa, which is obtained by deoxidizing with a Si or Mn-based weak deoxidizer in the first half of the tapping process of a converter and adding Ti iron in the second half of the tapping process to deoxidize so that the free oxygen content in molten iron reaches 100 ppm. This free oxygen content, Ti in the steel is present substantially in the form of large particle oxides and is difficult to retain in the steel during subsequent refining. After the RH high vacuum treatment, the free oxygen in the steel can not reach 50ppm, and after the calcium wire is fed, the oxygen content of the fixed oxygen in the steel can not reach 50ppm, and the rare earth wire is not fed.

Although these production methods can provide steel sheets suitable for large-wire weldability, it is difficult to stably obtain effective inclusions, which are so-called Ti-Ca-Mn-O-S and contain RE oxides and are advantageous for large-wire weldability, by the above-mentioned production methods of the prior art. And the prepared steel plate has too large fluctuation of welding performance, unstable performance and unreliable performance.

Disclosure of Invention

Based on the analysis of the prior art, aiming at the defects of the prior art, the invention aims at improving the performance of the coarse crystal area of the large linear weldable heat affected zone, realizes the structural ordering and the size optimization of the inclusions by controlling the formation process of the fine inclusions in the liquid steel, and ensures that the inclusions are uniformly distributed in the steel, so that the steel plate generates multi-directional acicular ferrite phase transformation induced by the inclusions in the coarse crystal area of the large linear weldable heat affected zone, thereby effectively ensuring the impact toughness of the welded heat affected zone.

The inclusion formed by this method is TiO_X-MgO-CaO-MnS-TiN, etc. During the welding temperature rise, nitrides and sulfides in the inclusions are decomposed. During cooling, nitrides and sulfides are deposited on the oxide matrix. Resulting in uneven distribution of elements around the inclusions and formation of acicular ferrite phase transformation conditions, and mostFinally, the reheated austenite is transformed to form a room-temperature structure with higher grain boundary density, the density of impurity elements distributed on the grain boundary is reduced, the capability of the grain boundary for inhibiting crack propagation at low temperature is improved, and the performance degradation degree of a coarse grain region of a welding seam is effectively reduced. Compared with the prior art, the beneficial nitrogen oxide particles formed in the steel have reasonable and ordered structure and uniform distribution, the specific surface area for inducing phase change is large, and the obtained material is more suitable for large-linear energy welding.

The specific scheme of the invention is as follows:

a manufacturing method of a low-alloy steel plate suitable for large-wire welding adopts low-carbon low-aluminum clean steel which is smelted in a converter or an electric furnace and subjected to vacuum treatment as a raw material, FeOx cored wires are fed into a molten steel tank in the vacuum treatment to increase oxygen, the oxygen increase amount in the molten steel is controlled to be 8-20 ppm, and the low-alloy steel plate is stirred for 2-5 min by weak argon blowing; then adding Mg and Ca into the steel in a wire feeding mode, and controlling the Mg increasing amount to be 5-10 PPm and the Ca increasing amount to be 10-20 PPm in the molten steel; then, argon is blown into the molten steel through an argon hole at the bottom of the tank for stirring, the argon blowing flow is controlled to enable the surface of the molten steel to roll weakly, the diameter of the exposed surface of the molten steel is less than or equal to 80mm, and the stirring time is 3-10 min; pouring on a machine within 10min after stirring. Continuous casting needs to control the superheat degree of the molten steel in the tundish to be 15-35 ℃.

Before the FeOx cored wire is fed for oxygenation treatment, the adopted raw materials of molten steel comprise, by weight, less than or equal to 15PPm of free [ O ], less than or equal to 0.008% of P, less than or equal to 0.025% of Nb, less than or equal to 0.30% of Cu, less than or equal to 0.30% of Ni, less than or equal to 0.06% of Mo, and V: not less than 4X (N) to 0.06 percent, not more than 0.003 percent of S, not more than 0.0008 percent of B, not more than 0.30 percent of Cr, not more than 0.007 percent of Alt, C: 0.03 to 0.07%, Mn: 1.30-1.60%, Si: 0.15-0.30%, Ti: 0.011 to 0.022%, 30 to 60ppm of N, and the balance Fe and inevitable impurity elements in the steel.

According to the manufacturing method of the low-alloy steel plate suitable for large-wire welding, the finally manufactured molten steel is cast into a plate blank on a continuous casting machine and then rolled into the steel plate through a medium-thickness steel plate rolling line by a TMCP (thermal mechanical control processing) process, and the steel plate is prepared from the following components, by weight, of less than or equal to 40PPm of total oxygen, less than or equal to 0.008% of P, 0.015-0.0.021% of Nb, 0.12-0.30% of Cu, and Ni: 0.05-0.30%, Mo is less than or equal to 0.06%, V is less than or equal to 0.020-0.050%, S is less than or equal to 0.003%, B is less than or equal to 0.0008%, Cr is less than or equal to 0.30%, Als: less than or equal to 0.005 percent, C: 0.04-0.065%, Mn: 1.30-1.60%, Si: 0.15-0.30%, Ti: 0.008-0.016%, 30-50 PPm of N, 3PPm or more of Mg, 10PPm or more of Ca, and the balance of Fe and inevitable impurity elements in the steel.

The production process of the continuously cast steel billet on the rolling line of the medium plate steel plate basically comprises the steps of heating the steel billet evenly at the temperature of 1100-1220 ℃, removing the iron oxide scale on the surface of the steel billet by high-pressure water spraying before rolling, and controlling the rolling to the thickness of the finished product by two stages.

The accumulated reduction rate of the first stage of rolling is more than or equal to 50 percent, and the rolling temperature is more than 930 ℃.

The accumulated reduction rate of the second stage of rolling is more than or equal to 50%, and the rolling temperature range is 780-840 ℃.

And (3) watering and cooling after rolling, wherein the cooling starting temperature is more than or equal to 720 ℃, the maximum temperature of the surface temperature return of the steel plate after cooling is within the range of 450-630 ℃, and the maximum temperature is reduced along with the increase of the thickness of the steel plate.

The average cooling speed of watering and cooling needs to be more than or equal to 8 ℃/S.

The steel plate can be tempered, and the tempering and heat preservation temperature range is 400-630 ℃.

The steel plate produced by the method has the structure of non-equiaxial ferrite and acicular ferrite, and the HAZ coarse crystal area forms the structure form mainly comprising the acicular ferrite after welding. Even if the welding line energy is increased to 400KJ/cm, the peak temperature of the coarse crystal region reaches more than 1350 ℃, and the HAZ coarse crystal region can also obtain a higher low-temperature (-40 ℃) impact toughness value.

The reason why the structure morphology mainly comprising acicular ferrite is formed in the HAZ coarse crystal region is that a sufficient number of beneficial inclusions having a diameter of 0.5 to 5 μm are formed in the steel by the method of the present invention. The structure of these inclusions is mainly TiO_X-MgO-CaO-MnS-TiN, etc. Wherein the oxide portion is formed before the molten steel solidifiesMost of the solder does not dissolve during the solder reheating process, and the position does not migrate. However, the nitrides and sulfides attached to the surface of the matrix are dissolved back, and in the subsequent cooling process, the nitrides and sulfides are re-attached to the oxides and gradually separated out, and the nucleation and growth of acicular ferrite at the interface of the inclusions and the matrix are promoted based on the mass transfer process of the separated out nitrides and sulfides.

The reason why the V content in the steel is controlled to 4X N-0.06% is to prevent the steel itself and the steel from suffering aging embrittlement due to the presence of excessive amount of solid-solution nitrogen after welding.

The reason why the upper limit of the element such as Mo, V, Nb, etc. is controlled is that these elements have a strong affinity with carbon and tend to promote the formation of a structure such as M-A which is unfavorable for low-temperature toughness.

The method has the beneficial effects that the oxide is enabled to have a strong function of promoting acicular ferrite phase change in the HAZ coarse crystal region by controlling the ordered generation of the oxide after vacuum refining, so that the steel has stable large-line weldability.

Drawings

FIG. 1 is the morphology and structure of beneficial inclusions (EDS photographs) about 5 microns in diameter for the steel sheet of example 1;

FIG. 2 is a particle size distribution of nitrogen oxides in a certain visual field of a metallographic sample of a steel sheet according to example 1.

Detailed Description

Reference will now be made in detail to embodiments of the present invention, examples of which are illustrated in the accompanying drawings, wherein like or similar reference numerals refer to the same or similar modules or modules having the same or similar functionality throughout. The embodiments described below with reference to the accompanying drawings are illustrative only for the purpose of explaining the present invention, and are not to be construed as limiting the present invention.

The scheme of the invention is as follows:

a method of manufacturing a low alloy steel sheet suitable for large wire weldable, containing no aluminium, comprising the steps of: feeding molten steel serving as a raw material into an FeOx cored wire, stirring by weak argon blowing, adding Mg and Ca, stirring by argon blowing, pouring, continuously casting to obtain a casting blank, rolling, and cooling by water to obtain the low-alloy steel plate which does not contain aluminum and is suitable for large-wire welding.

The manufacturing method of the low-alloy steel plate suitable for large-wire welding adopts low-carbon low-aluminum clean steel which is smelted in a converter or an electric furnace and subjected to vacuum treatment as a raw material, FeOx cored wires are fed into a molten steel tank in the vacuum treatment to increase oxygen, the oxygen increase amount in the molten steel is controlled to be 8-20 ppm, and the low-alloy steel plate is stirred for 2-5 min by weak argon blowing; then adding Mg and Ca into the steel in a wire feeding mode, and controlling the Mg increasing amount to be 5-10 PPm and the Ca increasing amount to be 10-20 PPm in the molten steel; then, argon is blown into the molten steel through an argon hole at the bottom of the tank for stirring, the argon blowing flow is controlled to enable the surface of the molten steel to roll weakly, the diameter of the exposed surface of the molten steel is less than or equal to 80mm, and the stirring time is 3-10 min; pouring on a machine within 10min after stirring. Continuous casting needs to control the superheat degree of the molten steel in the tundish to be 15-35 ℃.

The production process of the continuously cast blank on the rolling line of the medium-thickness steel plate basically comprises the steps of heating uniformly in the range of 1100-1220 ℃, removing the iron oxide scale on the surface of the billet by high-pressure water spraying before rolling, and controlling the rolling to the thickness of the finished product by two stages.

The steel plate can be further tempered after being cooled by water, and the tempering and heat preservation temperature range is 400-630 ℃.

The reason why the structure morphology mainly comprising acicular ferrite is formed in the HAZ coarse crystal region is that a sufficient number of beneficial inclusions with a diameter of 0.5 to 5um are formed in the steel by the method of the present invention. The structure of the inclusions is mainly a composite structure of TiOx-Mgo-CaO-MnS-TiN and the like. The oxide part is formed before the molten steel is solidified, most of the oxide part cannot be dissolved in the welding reheating process, and the oxide part cannot migrate. However, the nitrides and sulfides attached to the surface of the matrix are dissolved back, and in the subsequent cooling process, the nitrides and sulfides are re-attached to the oxides and gradually separated out, and the nucleation and growth of acicular ferrite at the interface of the inclusions and the matrix are promoted based on the mass transfer process of the separated out nitrides and sulfides.

The reason for controlling the V content in the steel to 4 xN-0.06% is to prevent the steel itself and the aging embrittlement after welding due to the presence of excessive solid-solution nitrogen in the steel.

The reason why the upper limit of Mo, V, Nb and the like is controlled is that these elements have a strong affinity for carbon and tend to promote the formation of a structure such as M-A which is unfavorable for low-temperature toughness.

DETAILED DESCRIPTION OF EMBODIMENT (S) OF INVENTION

In the embodiment of the invention, the billet is smelted by adopting a production mode of 120-ton converter-LF furnace refining-RH furnace vacuum treatment-250 mm section casting machine. The method comprises the following steps of smelting molten iron in a converter, wherein P is less than or equal to 0.110%, S is less than or equal to 0.010%, the temperature of the molten iron is more than or equal to 1300 ℃, deoxidizing is carried out by forming elements by weak oxides in the tapping process after smelting, and the tapping feeding sequence is as follows: active lime-manganese alloy-silicon alloy. The adding amount of the active lime is 350-450kg, and 30kg of calcium carbide is added during tapping. And (4) tapping is started until the tapping 3/4 is kept at argon full opening, then the flow is switched to a small flow, and the small flow is opened for 200L/min multiplied by 3-5 minutes after tapping. Target composition after furnace: less than or equal to 0.04 percent of C, 1.0 to 1.3 percent of Mn, less than or equal to 0.008 percent of P, less than or equal to 0.012 percent of S and less than or equal to 0.20 percent of Si. The temperature after the furnace is more than or equal to 1540 ℃. In the LF furnace treatment process, the white slag holding time is required to be more than or equal to 15min, and the components are adjusted to target components. During heating, the submerged arc operation is maintained, the micro-positive pressure state and the reducing atmosphere in the furnace are ensured, aluminum particles can be added for deoxidation, but the aluminum particles are gradually added for many times, the total Al content in the steel is controlled not to exceed 0.007%, and the FeO and MnO in the furnace slag after refining is controlled not to exceed 1.5%. And (5) performing Ti alloying treatment after LF (ladle furnace) is out of the station, and adjusting the titanium to a target range. RH processing vacuum pressure maintaining (less than or equal to 60Pa) time is more than or equal to 20min, fine adjustment is carried out on alloy components after the air is broken, and the components meet the requirements before FeOx cored wire feeding and oxygenation are ensured.

The compositions of examples 1 to 4 of the present invention before the oxygen increasing operation are shown in Table 1.

TABLE 1 ingredients before oxygen enrichment

In the embodiment 1-4, after the RH is out of the station, 280 plus 340 meters of FeOx cored wires are fed, the oxygen increasing amount is 8-19PPm respectively, and after the wires are fed and the oxygen is increased, weak argon blowing is carried out for stirring for 2-5 min, so that the increased oxygen is uniformly distributed in the steel. Then feeding a Mg and Ca composite alloy wire with the feeding amount of 300-550 m, controlling the Mg increasing amount to be 5-10 PPm and the Ca increasing amount to be 12-17 PPm. And then, weakly blowing argon gas at the bottom of the tank for stirring, controlling the flow of the argon gas according to the condition that the diameter (average diameter length) of the exposed surface of the molten steel surface at the gas outlet point is less than or equal to 80mm, stirring for 5-8 min, and loading the molten steel into a machine for pouring within 8 min.

The main relevant control operation parameters of the embodiment 1-4 of the invention from oxygenation to continuous casting on the machine are shown in the table 2.

TABLE 2 oxygenation to main relevant control operating parameters for continuous casting on a lead

The final melt composition of examples 1-4 is shown in Table 3.

TABLE 3 Final product test composition for each example

The test steel sheets were rolled on a medium plate rolling line to obtain steel sheets having a gauge of 60mm × 2800mm × Lmm. Actually heating to 1200 ℃, wherein the furnace time is more than 300 minutes, removing the iron scale on the surface of the billet by high-pressure water spraying before rolling, and controlling the rolling to the thickness of the finished product by two stages. The accumulated reduction rate of the first stage of rolling is more than or equal to 50 percent, the rough rolling is carried out for 8 times, and the reduction rate of the second time is more than or equal to 15 percent. The accumulated reduction rate of the second stage of rolling is 50%, and the rolling temperature interval is 790-800 ℃. And (3) watering and cooling after rolling, wherein the cooling starting temperature is more than or equal to 770 ℃, and the maximum temperature of the surface temperature return of the steel plate after cooling is 520-535 ℃.

The rolling process and properties of the steel sheets of the samples of Experimental examples 1-4 of the present invention are shown in tables 4 and 5.

TABLE 4 Process setting parameters for rolling of 60 mm. times.2800 mm. times.10000 mm steel sheets for each example

TABLE 5 Final Properties of the examples

The steel plates with the thickness of 60mm rolled by the components in the samples of the examples 1-4 of the invention are butt-welded by a double-wire electrogas welding machine, the adopted welding parameters and welding wires are shown in the table 6, and the impact toughness results of the welding joints are shown in the table 7.

Table 6 welding process of each ingredient number

TABLE 7 weld joint impact toughness for each ingredient number

Fig. 1 shows an EDS photograph of the structure of the inside of the steel sheet of example 1. The sample steel forms a sufficient amount of beneficial inclusions with a diameter of 0.5-5 mu m. The structure of the inclusions is mainly a composite structure of TiOx-MgO-CaO-MnS-TiN and the like; FIG. 2 is the grain size distribution of nitrogen oxides in a certain visual field of a metallographic sample of the steel of example 1, wherein the nitrogen oxides are intensively distributed in an interval of 1-2 μm.

Tables 3-7 also show that even if the energy of the welding line of the steel plate produced by the invention is improved to 400KJ/cm, the peak temperature of the coarse crystal zone reaches more than 1350 ℃, the HAZ coarse crystal zone can obtain higher low-temperature (-40 ℃) impact toughness value, and the minimum fusion line can reach more than 65J.

In the description herein, references to the description of "one embodiment," "another embodiment," etc., mean that a particular feature, structure, material, or characteristic described in connection with the embodiment is included in at least one embodiment of the invention. In this specification, the schematic representations of the terms used above are not necessarily intended to refer to the same embodiment or example. Furthermore, the particular features, structures, materials, or characteristics described may be combined in any suitable manner in any one or more embodiments or examples. Furthermore, various embodiments or examples and features of different embodiments or examples described in this specification can be combined and combined by one skilled in the art without contradiction.

Although embodiments of the present invention have been shown and described above, it is understood that the above embodiments are exemplary and should not be construed as limiting the present invention, and that variations, modifications, substitutions and alterations can be made to the above embodiments by those of ordinary skill in the art within the scope of the present invention.

Claims

1. A method of manufacturing a low alloy steel sheet suitable for large wire weldable, containing no aluminium, comprising the steps of: feeding molten steel serving as a raw material into an FeOx cored wire, stirring by weak argon blowing, adding Mg and Ca, stirring by argon blowing, pouring, continuously casting to obtain a casting blank, rolling, and cooling by water to obtain the low-alloy steel plate which does not contain aluminum and is suitable for large-wire welding.

2. The method of claim 1, wherein the raw material used before the FeOx cored wire aeration treatment is composed of, by weight, a molten steel having a composition of [ O ] 15PPm or less, P0.008% or less, Nb 0.025% or less, Cu 0.30% or less, Ni 0.30% or less, Mo 0.06% or less, V: not less than 4X (N) to 0.06 percent, not more than 0.003 percent of S, not more than 0.0008 percent of B, not more than 0.30 percent of Cr, not more than 0.007 percent of Alt, C: 0.03 to 0.07%, Mn: 1.30-1.60%, Si: 0.15-0.30%, Ti: 0.011 to 0.022%, 30 to 60ppm of N, and the balance Fe and inevitable impurity elements in the steel.

3. The method for manufacturing the aluminum-free low alloy steel plate suitable for large-wire welding according to claim 2, wherein FeOx cored wires are fed into a molten steel tank through vacuum treatment to increase oxygen, the oxygen increase amount in the molten steel is controlled to be 8-20 ppm, and the molten steel is stirred for 2-5 min by weak argon blowing; then adding Mg and Ca into the steel in a wire feeding mode, and controlling the Mg increasing amount to be 5-10 PPm and the Ca increasing amount to be 10-20 PPm in the molten steel; then, argon is blown into the molten steel through an argon hole at the bottom of the tank for stirring, the argon blowing flow is controlled to enable the surface of the molten steel to roll weakly, the diameter of the exposed surface of the molten steel is less than or equal to 80mm, and the stirring time is 3-10 min; and (3) pouring on a machine within 10min after stirring is finished, wherein the superheat degree of the tundish molten steel is controlled to be 15-35 ℃ in continuous casting.

4. The method for manufacturing the aluminum-free low-alloy steel plate suitable for large-line welding according to any one of claims 1 to 3, wherein the molten steel is cast into a slab on a continuous casting machine and then rolled into a steel plate through a medium steel plate rolling line by a TMCP process, and the steel plate is preferably prepared from the following components in parts by weight of less than or equal to 40PPm of total oxygen, less than or equal to 0.008 of P, 0.015 to 0.0.021 of Nb, 0.12 to 0.30 of Cu, Ni: 0.05-0.30%, Mo is less than or equal to 0.06%, V is less than or equal to 0.020-0.050%, S is less than or equal to 0.003%, B is less than or equal to 0.0008%, Cr is less than or equal to 0.30%, Als: less than or equal to 0.005 percent, C: 0.04-0.065%, Mn: 1.30-1.60%, Si: 0.15-0.30%, Ti: 0.008-0.016%, 30-50 PPm of N, 3PPm or more of Mg, 10PPm or more of Ca, and the balance of Fe and inevitable impurity elements in the steel.

5. The method of claim 1, wherein the billet produced by continuous casting is produced by a process on a pass line of a medium gauge steel plate, wherein the process comprises: heating evenly in the range of 1100-1220 ℃, removing the iron scale on the surface of the billet by high-pressure water spraying before rolling, and controlling the rolling to the thickness of the finished product by two stages.

6. The method for manufacturing a low alloy steel sheet suitable for large wire weldability according to claim 5, characterized in that the cumulative reduction at the first stage of rolling should be equal to or greater than 50%, and the rolling temperature is 930 ℃ or higher.

7. The method for manufacturing the aluminum-free low-alloy steel plate suitable for large-wire weldable according to claim 5 or 6, wherein the cumulative reduction rate of the second stage of rolling is equal to or more than 50%, and the rolling temperature range is 780-840 ℃.

8. The method for manufacturing a low alloy steel sheet free of aluminum suitable for large-line weldability according to claim 7, wherein water cooling is performed after rolling, the cooling start temperature is not less than 720 ℃, and the maximum temperature of the surface temperature return of the steel sheet after cooling is in the range of 450 to 630 ℃ and decreases as the thickness of the steel sheet increases.

9. The method for manufacturing a low alloy steel sheet suitable for large wire welding containing no aluminum as claimed in claim 8, wherein the average cooling rate of the water cooling is not less than 8 ℃/S.

10. The method for manufacturing a low alloy steel sheet suitable for large line welding free from aluminum according to any one of claims 1 to 9, wherein the steel sheet is optionally tempered after water cooling, and the tempering holding temperature is in the range of 400 to 630 ℃.