GB1573162A

GB1573162A - High tensile steel products

Info

Publication number: GB1573162A
Application number: GB1520577A
Authority: GB
Original assignee: Nippon Steel Corp
Current assignee: Nippon Steel Corp
Priority date: 1976-04-12
Filing date: 1977-04-12
Publication date: 1980-08-13
Also published as: CA1084310A; SE7704046L; DE2716081A1; IT1077777B; DE2716081C2

Description

(54) IMPROVEMENTS IN OR RELATING TO HIGH TENSILE STEEL PRODUCTS (71) We, NIPPON STEEL CORPORATION, a Japanese Company, of No. 6-3, 2-chome, Ote-machi, Chiyoda-ku, Tokyo, Japan, do hereby declare the invention, for which we pray that a patent may be granted to us, and the method by which it is to be performed, to be particularly described in and by the following statement: The present invention relates to a process for producing high tensile steel products, such as plates, sheets and strips (all of which are referred to herein as sheets) having an excellent low-temperature toughness and with a yield point typically of 40 kg/mm2 or higher, and also to such steel sheets themselves.

The steel products according to the present invention are useful as hot rolled or as heated to a temperature ranging from 300 to 750"C after the hot rolling.

Conventionally, steel sheets such as for making pipe lines for use in cold regions and which are required to have high strength and toughness in the "as rolled" condition have been produced by a method called "controlled rolling" (hereinafter abridged as CR), and mainly Nb-containing steels have been used for this purpose.

In general, the CR method can be broadly classified into two steps; the first step is a heating step and the second step is a rolling step (cooling). The following considerations must be taken into account respectively in these steps.

(1) In the heating step, it is required to dissolve elements such as Nb and V sufficiently for refinement of the structure and precipitation hardening, and it is required to maintain the austenite grains during the heating (heated y grains) as fine as possible.

(2) In the rolling step, it is necessary to recrystallize the heated y grains repeatedly by the rolling to obtain refined rolled austenite grains (rolled y grains), and it is necessary to elongate the rolled y grains and reduce their thickness by rolling in their non-recrystallization zone so as to obtain refinement of the rolled structure.

However, in case of Nb-containing steels as commonly used, Nb(CN) is stable at high temperatures and it is difficult to redissolve Nb(CN) consistently and satisfactorily even by heating for a long time, if the heating temperature is not higher than 1150"C.

If the heating temperature is raised, it is possible to attain satisfactory solid solution of Nb(CN), but on the other hand, the heated y grains grow excessively, thus resulting in considerable deterioration of the toughness of the rolled steel.

Therefore, in the CR method, it is necessary to lower the heating temperature and maintain the heated 7 grains smaller when severe low-temperature toughness requirement is to be satisfied. On the other hand, when the heating temperature is lowered, the amount of Nb in solid solution increases or decreases depending on a slight change in the heating temperature and time in case of the commonly used Nb-containing steels. However, even under the same rolling condition, the resultant strength fluctuates in a wide range depending on the change in the amount of the solid solution Nb, and a high strength, if obtained, is accompanied with deterioration of toughness. Thus, it is difficult to obtain a stable balance between strength and toughness.

The above difficulties can be attributed to the facts that the toughness lowers in proportion to the increase in strength, and the increase in strength corresponds to an increase in the amount of Nb(CN) in solid solution and the coarsening of the heated 7 grains, so that the steel structure will be of coarse grains and mixed grains.

However, in the conventional CR method, proper consideration has not been given to the fact that the heated y grains coarsen when enough Nb(CN) is dissolved in solid solution during the heating step, and thus the toughness is deteriorated.

As described above, it is necessary to prevent the growth of the heated y grains by means of precipitation in order to maintain fine heated 7 grains and improve toughness.

For this purpose, it is required to lower the heating temperature and keep the precipitates such as Nb(CN) from dissolving into solid solution during the heating. On the other hand, in order to maintain the strength, it is necessary to dissolve Nb(CN) into solid solution as much as possible during the heating so as to precipitate it during the cooling after the rolling to strengthen the steel. For this purpose, it is desirable to maintain the heating temperature as high as possible.

Therefore, it is a principal aim of this invention at least to reduce the completely contradictory problems mentioned above, and to a process for making a steel sheet having significantly smaller heated 7 grains than those of conventional steels, in spite of the requirement of Nb(CN) in solid solution for strength. By the appropriate application of rolling conditions, an aim is to give an excellent balance between strength and toughness.

The features of the present invention may be summarized as below.

(1) In order to attain satisfactory and stable Nb(CN) in solid solution, the carbon content is lowered to an extreme degree as understood from the solubility product relation.

(2) The growth of the heated y grains due to Nb(CN) in solid solution is prevented by TiN which is stronger than Nb(CN) for prevention of the growth of the heated y grains, and (3) Optimum rolling conditions are selected.

By combining the above features it is possible to utilize Nb(CN) and TiN separately for different purposes; the former for strengthening the steel and the latter for preventing the growth of the heated 7 grains, and thus the problem in the heating step can be solved.

Starting from the fine heated y grains, a rolled structure having still finer grains can be obtained by rolling under proper conditions and remarkable strength and toughness can be obtained through the decrease in the pearlite proportion attained by the lowered carbon content as well through grain refinement.

Regarding improvement of steel toughness by refinement of the heated fly grains, the present inventors have discovered a method therefor and they have conducted further various extensive studies on the production of a high tension steel having excellent toughness at low temperatures, and have found that the toughness can be stabilized and improved remarkably according to the production process of the present invention.

The production process according to the present invention is characterized in that a steel ingot or slab containing not less than 0.004% of TiN not larger than 0.02,u is heated to a temperature not higher than 1150"C and rolled, and the growth of the 7 grains during this heating and rolling step is prevented by TiN thereby to improving the toughness.

The present invention will be described in more detail referring to the attached drawings.

Figure 1 shows the relation between the heated y grain size and the percentage content of TiN not larger than 0.02,u in a steel, when heated to 1 1500C and held at the temperature for 60 minutes.

Figure 2 shows the relation between the heating temperatures and the heated y grain size when the steel No. 2 in Table 1 according to the present invention is heated to various temperatures and held at the various temperatures for 60 minutes.

Figure 3 shows the relation between the ratio of NaS to TiN/N (marked by 0) and the heating temperature when the steel No. 1 in Table 1 according to the present invention is heated to various temperatures and held at the various temperatures for 120 minutes and rapidly cooled in water, and the relation between the percentage content of TiN not larger than 0.2 (marked by A) when the same steel then is heated to 11500C and held at the temperature for 120 minutes.

Figure 4 shows the relation between the mean cooling rate and the content (%) of TiN not larger than 0.02,u when the steel No. 1 in Table 1 according to the present invention is cast at various solidification rates.

Figure 5 shows the relation between the amount of solid solution Nb and the carbon content when steels with different carbon contents are heated to various temperatures and held at the various temperatures for 30 minutes.

Figure 6 shows the relation between the heating temperature and the product of (solid solution Nb%) x (solid solution C%) when the steel according to the present invention is heated to various temperatures and held at the various temperatures for 60 minutes.

Figure 7 shows the relation between the reduction amount at temperatures not higher than 930"C and the yield point as well as vTrs in the steel No. 2 in Table 1 according to the present invention.

Figure 8 shows the relation between the finishing temperature and the yield point as well as vTrs in the steel No. 2 in Table 1 according to the present invention.

The term "TiN not larger than 0.02ju," is intended also to include Ti and N which are present in solid solution in the steel and TiN which is present in the form of precipitate and has a size not larger than 0.02. Ti and N which are present in solid solution in the steel precipitate as TiN not larger than 0.02cm during the subsequent heating and effectively prevent the coarsening of the heated fly grains. In this case, according to the studies made by the present inventors, there is a correlation between the heated 7 grain size and the heating rate, and when the heating rate starting from 800"C to a predetermined temperature is excessively high, Ti and N do not precipitate fully, thus failing to obtain a satisfactory refinement of the heated 7 grains. Therefore, in order to refine the heated 7 grains, it is necessary to decrease the heating rate to some degree, and it is preferable to control the heating rate starting from 800"C to a predetermined temperature to a rate not larger than 6 C/min.

Figure 1 shows the relation between the amount of TiN not larger than 0.02,u and the heated y grain size, when the steel is heated to 1 1500C and held at the temperature for 60 minutes. It is clear that unless TiN not larger than 0.02it. is present in an amount of 0.004%or more, no satisfactory refinement of the heated 7 grains can be expected. Therefore, it is necessary that TiN not larger than 0.02,u is present in an amount not less than 0.004% in the steel before the heating. However, even when this condition is satisfied, prevention effect on the coarsening of the heated 7 grains by TiN becomes unstable if the heating temperature is excessively high.

As will be understood from Figure 2, showing the relation between the heating temperature and the heated y grain size, it is necessary to maintain the heating temperature at 11500C or lower, preferably in the range from 900 to 1 1500C in order to obtain fine heated y grains (not lower than No. 3 of ASTM).

As described above, satisfactory refinement of the heated y grains can be obtained under appropriate conditions when TiN not larger than 0.02cm is present in an amount not less than 0.004%.

Hereinbelow descriptions will be made on the method of introducing into the steel not less than 0.004% of TiN not larger than 0.02,u in connection with the ingot-making method and the continuous casting method respectively.

In the ingot-making method, the coarse TiN which has precipitated during the solidification step of the molten metal is dissolved in solid solution in an amount not less than 0.004% during the ingot heating (soaking) step, and part of the solid solution TiN is precipitated during the break down rolling step and the cooling step to maintain not less than 0.004%of TiN, not larger than 0.02,u in the steel slab before the heating. In this case, if the Ti content is excessive, it is difficult to maintain not less than 0.004% of TiN in solid solution during the ordinary ingot heating step, because TiN precipitates in an excessively coarse form during the solidification step. Even in this case, the solid dissolution of TiN depends on the heating temperature and the holding time, but if the heating temperature is too high, there is caused a burning phenomenon and thus there is a certain maximum heating temperature inherent to the steel. Therefore, on the basis of the present steel making techniques it is necessary that the content of Ti is maintained not greater than 0.03% and the amount of Ti required for the minimum amount of 0.004% for TiN not larger than 0.02cm is 0.004% on a commercial production taking into consideration the amount of Ti consumed for formation of oxides etc.

Therefore, the content of Ti should be in a range from 0.004 to 0.03%.

Next, detailed explanations will be made hereinbelow on the limitations of the heating temperature for the solid dissolution of TiN which has precipitated during the solidification of the molten steel as well as the limitations of the N and TiN contents.

For economical and stable solid dissolution of TiN during the heating step on a commercial level, it is effective to limit not only the Ti content, but also the N content. The reason for setting the lower limit of the total N content at 0.001%, is that it is the minimum amount required for the lower limit of 0.004 % of TiN which must be dissolved in solid solution during the heating step. Further, in order to prevent TiN precipitating and to maintain an sufficient amount of TiN in solid solution during the heating step, the upper limit of the total N content should not exceed that which corresponds to the upper limit of the Ti content having regard to the formation of TiN in solid solution. Therefore the upper limit of the total N content is set at 0.009% which corresponds to 0.03% of Ti. On the other hand, if the TiN content exceeds 0.04%, the toughness of the steel sheet deteriorates and thus it is necessary to set the upper limit of the TiN content at 0.04it, but as long as the Ti and total N contents are within the ranges defined above, the TiN content does not exceed 0.04%.

When the Ti and N contents are within the ranges defined in the present invention, the lower limit of the heating temperature for dissolving not less than 0.004% of TiN into solid solution may be 12500C as shown in Figure 3 and confirmed by experiments, while the upper limit is set at 14000C as a practically feasible temperature inspite of partial burning of the iron oxide on the steel surface.

In the continuous casting method, where the steel slabs are made directly from the molten steel, if the Ti and N contents are excessive, coarse TiN precipitates form during the solidification so that it is impossible to maintain not less than 0.004% of TiN not larger than 0.02E.L, therefore, just as in the ingot-making method, it is necessary to limit the Ti and N contents respectively to 0.004 to 0.03%Ti and 0.001 to 0.009%N. Even when the Ti and N contents are within these ranges, it is impossible to obtain not 13ss than 0.004% of TiN not larger than 0.02cm, if the solidification coolingrate is too slow. Therefore, itis desirable for the mean cooling rate of the centre portion of the steel slabe to be not less than 8"C/min. from the molten steel temperature at the time of casting to 1100 C. When the cooling rate is below 8"C/min., it is difficult to attain not less than 0.004% of TiN not larger than 0.02,u in the steel slab as shown in Figure 4 and no effective prevention of the coarsening of the heated y grains can be assured.

The basic features of the present invention have been described hereinabove.

It has been further found by the present inventors that when the hot rolled steel material obtained by the above production process is reheated to a temperature ranging from 300 to 750"C, part of the fine carbides or the solid solution carbon coagulates into carbides of favourable size so that toughness is improved due to the relief of stress by the precipitation hardening of the matrix, and the arrest property as represented by BDWTT (Battele Drop Wear Tear Test), as well as the yield strength are still remarkably improved.

Explanations will be made on the limitation on the steel compositions defined in the present invention.

The base steel composition applicable to the present invention comprises 0.01 to 0.13 % C, 0.01 to 1.0% Si, 0.7 to 2.0% Mn, not more than 0.10% total Al, 0.004 to 0.03% Ti, 0.001 to 0.009% total N, 0.01 to 0.10% Nb, one or more of 0.01 to 0.15% V and 0.05 to 0.4% Mo and satisfying the condition of NbSo) x (C%) < 5 x 10-3 Now the lower limit of 0.01% for the carbon content is set because it is a minimum amount for assuring the grain refinement of the steel material and strength of weld joints as well as full development of effects of carbide forming elements such as Nb and V. On the other hand, when the carbon content is excessive, the amount of Nb in solid solution readily increases or decreases depending on even slight changes in the heating conditions as mentioned hereinbefore, and thus the strength-toughness balance becomes unstable. Therefore, it is effective to define an upper limit for the carbon content for assuring a stable solid solution of Nb(CN) in the steel slab to maintain desired strength and toughness even in cases where the heating temperature is below 1150"C.

In Figure 5 which shows the relation between the amount of the solid solution Nb and the heating temperature in connection with various carbon contents, it will clearly be seen that when the carbon content is lowered, the amount of the solid solution Nb (at a constant total Nb content of 0.05 So) increases, and when the carbon content is not higher than 0.13 So, Nb is completely dissolved in solid solution at 11500 C.

The reason for defining the total Nb content of 0.05%, is that this amount is enough for obtaining desired strength and toughness in case of 0.13% C. Thus the upper limit of the carbon content is set at 0.13%. In cases where the Nb content is large or the heating temperature is below 1150"C, it is necessary further to lower the carbon content in order to ensure that there is a sufficient amount of stable Nb(CN) in solid solution. But for this purpose it is favourable to limit not only the carbon content by itself, but also the carbon content in relation with the Nb content.

Figure 6 shows the experimental results concerning the relation between the heating temperature and (solid solution NbSo) x (solid solution C%), and it will be seen that Nb(CNa can be stably dissolved in solid solution when (C%) x (Nb%) S (solid solution NbSo) x (solid solution Coo).

Within the preferable heating temperature range from 1050 to 1150 according to the present invention, it is preferable to adjust (C%) and (NbSo) to satisfy the following formula desDite some fluctuation in the data.

(CSo) x (NbSo) < 5.0 x 10-3 For the reasons set above, the upper limit of the carbon content is set at 0.13% and the carbon content is further limited in relation with the Nb content in accordance with the following formula: (CSo) x (Nb 3to) < 5 x 10-3 Silicon is an element which comes into the steel unavoidably during the deoxidation step, but less than 0.1% silicon causes deterioration of the toughness. Therefore, the lower limit of the silicon content is set at 0.1 So. On the other hand, when the silicon content is excessive it damages the cleanness of the steel. Thus, the upper limit of the silicon content is set at 1.0%.

Manganese is an important element for assuring the desired strength and toughness of the low-carbon steel applicable to the present invention, and with manganese contents less than 0.7%the strength and toughness are low. Thus the lower limit of the manganese content is set at 0.7 %. On the other hand, when the manganese content is excessive, the toughness of HAZ (heat affected zone) deteriorates. Thus the upper limit is set at 2.0%.

Aluminium is contained in a killed steel unavoidably from the deoxidation step. However, when the total Al content exceeds 0.1% not only the toughness of HAZ but also the toughness of the weld metal are remarkably deteriorated. Thus, the upper limit of the total Al content is set at 0.1%.

Regarding the Ti and total N contents, they are limited to 0.004 to 0.03% Ti and 0.001 to 0.009% total N respectively as mentioned hereinbefore. So far as Ti and N are within these ranges the TiN content does not exceed 0.04%.

Niobium is added for improving the toughness of the steel material and expanding the feasible range of the plate thickness as well as for assuring the joint strength of the welded portion. The lower limit of the Nb content is set at 0.01 %for the reason that with Nb contents less than 0.01%, the desired refinement of the structure and the precipitation strengthening by Nb cannot be attained, and thus it is difficult to obtain the desired strength and toughness.

However, Nb addition beyond 0.10% causes difficulty in attaining stable and enough solid solution Nb with a heating temperature not higher than 1150"C, and causes HAZ toughness deterioration.

Vanadium, similar as niobium, may be contained up to 0.15%.

Molybdenum, similar as niobium and vanadium, increases hardening of HAZ and lowers HAZ toughness and cracking resistance, if present in an excessive amount. Therefore, the upper limit of the molybdenum is set at 0.40%. The lower limits of V and Mo are set at 0.01% and 0.05% respectively, because these amounts are minimum amounts for development of full effectiveness of these elements.

The steel applicable to the present invention contains phosphorus and sulphur as impurities. The phosphorus content is usually not more than 0.03% and a lower phosphorus assures improvement of toughness. The sulphur content is usually not more than 0.02%, and it is possible to lower the sulphur content to about 0.0005% by the present level of the techniques and thereby the toughness of the steel sheet is improved. In the present invention, neither phosphorus nor sulphur is added intentionally.

According to one modification of the present invention, one or both of 0.001 to 0.03% REM (mainly Ce, La, Pr) and 0.0005 to 0.03% preferably 0.0005 to 0.003% Ca is added under the condition of REM/S = 1.0 to 6.0. With this modification, the toughness of the steel product obtained by the present invention is still further improved as shown in Table 2.

REM contents less than 0.001% produce no practical improvement of toughness, while REM contents exceeding 0.03% cause increase not only in size but also in amount of REM-oxysulphides, so that large inclusions are formed, which damage remarkably the toughness as well as the cleanness of the steel product.

Therefore, the REM content is limited to the range from 0.001 to 0.03 7a Meanwhile REM is effective to improve and stabilise the toughness of the steel sheet in correlation with the sulphur content, and the optimum range for this purpose is 1.0 to 6.0 of REM/S. Calcium has similar effects as REM is limited to the range from 0.0005 to 0.003%.

According to another modification of the present invention, one or more of not more than 0.6% Cr, not more than 1.0% Cu and not more than 4.0% Ni is added under the condition of (Cu + Ni)/5 + Cr + Mo < 0.90% The main object of addition of these elements is to improve the strength and toughness of the steel product and to expand the feasible plate thickness range. Naturally, the addition of these elements has limitation in their amounts, but in the low-carbon steel applicable to the present invention, their upper limits may be higher than those in an ordinary carbon steel.

Regarding chromium, an excessive chromium content increases the hardenability of HAZ and lowers the toughness and cracking resistance. Therefore, the upper limit of the chromium content is 0.6%.

Nickel is effective to improve the strength and toughness of the steel product without adverse effect on the hardenability and toughness of HAZ, but nickel contents exceeding 4.0% are not favourable on the hardenability and toughness of HAZ even in case of a low-carbon steel as used in the present invention. Therefore, the upper limit of the nickel content is set at 4.0%.

Copper has almost similar effects as nickel and further improves the hydrogen-induced cracking resistance, but copper contents beyond 1.0% cause copper-cracking during the rolling. Therefore, the upper limit of the copper content is set at 1.0%.

Further, the above addition elements are not added independently within their respective ranges but they are are added under the condition of (Cu + Ni)/5 + Cr + Mo s 0.90% Otherwise the hardness of HAZ is remarkably higher so that HAZ is susceptible to cracking during a small heat-input welding, and thus the steel can not be used for welding.

In a still further modification of the present invention where the steel product after hot rolling is reheated in a temperature range from 300 to 750"C, the basic steel composition should be limited.

First of all, when the carbon content is more than 0.10%, the amount of Nb, V or Mo which is dissolved in solid solution during the slab heating decreases so that the amount of the fine carbide precipitates of Nb, V or Mo during the reheating which is favourable for the strength, particularly the tensile strength, decreases.

Further, in the reheating step after the hot rolling, the fine carbides are coagulated into suitable size so as to improve the toughness. For this purpose, carbon contents less than 0.08 % are remarkably effective without formation of excessively large coagulated carbides.

Regarding the aluminium content, deoxidation of the molten steel by aluminium is particularly necessary for assuring enough precipitates of fine carbides of Nb, V or Mo during the reheating, which are required for the desired strength. Therefore, aluminium is present in an amount of 0.005% at least.

The sulphur content should be limited 0.010%or lower so as fully to develop the toughness improvement by the reheating.

Descriptions have been made on the limitations of the various elements of the steel composition used in the present invention. It has been further found that it is difficult to produce a steel sheet having an excellent low-temperature toughness and high strength of not lower than 40 kg/mm2 yield point by rolling the steel of the defined composition within the defined range merely in an ordinary way. Therefore, in the present invention the final rolling conditions are limited.

As the basic feature of the present invention, the rolling condition has been defined as below.

The total reduction amount in the temperature range not higher than 930"tis not less than 50%and the finishing temperature is not higher than 830"C. Under this rolling condition, the strength and toughness of the steel product are improved remarkably.

Explanations will be given for the limitations of the rolling condition.

When the total reduction amount of 930"C or lower is not less than 50%, the yield point and toughness are remarkably improved as shown in Figure 7, but if the total reduction amount in the temperature range is less than 50%, it is impossible to obtain a hot rolling step, a plate rolling mill is desirable, but the present invention is not limited thereto and is applicable also to the production of hot steel strips and steel wire.

The basic rolling condition in the present invention has been described above, but this basic rolling condition should be further limited as below when the reheating is added according to a modification of the present invention. First, the total reduction amount should be limited as below. Thus, the total reduction amount at 9000C or lower should be 60% or more. At amounts less than 60%, the amount of the fine precipitates Nb, V or Mo which are required for remarkably increasing the strength and the toughness after the reheating is not enough, and thus the resultant strength and toughness are not satisfactory. On the other hand, when the total reduction amount at 9000 C or lower is more than 95 So, Nb, V or Mo precipitates are coarse so that it is difficult to obtain the desired fine carbides, and it is difficult to maintain the desired strength, particularly the desired strength after the reheating.

Regarding the finishing rolling temperature, it should be further limited to 8000C or lower.

Otherwise the amount of the fine precipitates is not enough and the resultang strength and toughness are not satisfactory. On the other hand, when the finishing temperature is below 500"C, it causes deterioration of toughness due to intermittent workings and excessive precipitation of the fine carbides of Nb, V or Mo which coagulate into coarse form during the reheating step so that satisfactory strength can not be maintained.

When the finishing temperature is low, the rolling is done in a ferrite-predominant zone so that there is an excess of precipitates of fine carbides of Nb, V or Mo are formed in the worked ferrite matrix. This is unfavourable for the strength-toughness balance. Therefore the finishing rolling temperature should be preferably not lower than 700"C. On the other hand, when particularly excellent toughness is to be obtained, coarse precipitates of carbides of Nb, V or Mo are promoted by excessive working in the austenite zone of higher temperatures, and the coarse precipitates coagulate excessively in the reheating step and produce adverse effects on the toughness. Thus, in this case, it is preferable to maintain the finishing temperature not higher than 780"C. Therefore, in respect of both the strength and the toughness, the most preferable finishing temperature range is from 700 to 7800C.

Regarding the heating step after the hot rolling step, this step is required for uniformly and appropriately coarsening the fine carbides of Nb, V and Mo, thus relieving the stress of the matrix due to the precipitation hardening and improving the toughness. For this purpose, a minimum temperature of 300"C is sufficient. On the other hand, when the reheating temperature is higher than 750"C, the above fine carbides become excessively coarse, thus lowering the strength considerably. The most preferable reheating temperature range for both the strength and the toughness is from 500 to 7000 C. Meanwhile, regarding the holding time in the reheating step, it should be at least one minute for uniformly and appropriately coarsening the fine carbides, thus relieving the stress of the matrix due to the precipitation hardening and improving the toughness.

On the other hand, if the holding time is longer than 10 hours, the fine carbides become excessively coarse, thus lowering the strength considerably. The most preferable holding time range is from 10 minutes to 2 hours for both the strength and the toughness.

The reheating step as defined above may be performed before the hot rolled steel sheet cools down near the ordinary temperature. In this case, the reheating has also effect of hydrogen removal.

The limitations of the production conditions in case where the reheating step is added have been explained before. The steel products obtained by this modification have been found to have also an excellent resistance against the hydrogen-induced cracking.

Although it has not been fully clarified, the hydrogen-induced cracking resistance may be attributed to the facts that the carbon content is low with less segregation, that formation of coarse carbides is prevented by the formation of fine carbides of Nb, V or Mo, and that the stress of the matrix is relieved by the uniform coarsening of the fine carbides during the reheating step.

The present invention will be more clearly understood from the examples shown in the tables.

Tables 1 to 3 show examples according to the basic process of the present invention.

Table 4 shows examples according to the modification of the present invention. In these examples, various steel compositions as shown (G: electric furnace steel; C1, C2, C3: refined in converter and with special phosphorus treatment) were made into slabs (L, M: continuous casting) and hot rolled. The conditions of slab making the hot rolling are shown in Table 4.

The thickness of the products and tensile strength (API test piece in the direction at right angles to the rolling direction, the 2 mmV Charpy impact property, the B, DWTT 85 % SATT property, and the 2mmV Charpy impact values of 50% bond portion of sub-merged arc welding joints welded with 30 KJ/cm input are shown in Table 4.

Table 4 also shows the number of cross sectional crackings (per 5 mm thickness) of the test pieces (ground 1 mm on both sides) after immersion in 100%H2S saturated aqueous solution (25"C) for 96 hours.

As is clearly shown in Table 4, the steels Al, Bl, Cl, M and N according to the present invention show excellent tensile strength property and toughness, particularly DWTT property, as well as excellent weld toughness and hydrogen-induced cracking resistance.

The steels A2, A3, B2, B3, C2, and C3 having the steel composition within the range defined in the present invention but outside the scope of the present invention in respect of the rolling condition and the reheating condition show inferior properties.

As is clearly understood from the examples, the steel product according to the present invention has excellent strength toughness and additionally excellent weldability and hydrogen-induced cracking resistance.

The steel product according to the present invention is most suitable for production of steel pipes and also is useful for fittings tank structural components, ship-building materials frame members of various machine and apparatus for cold regions, etc. where the arrest property is required.

Table 1 Chemical Composition Steel No.

C Si Mn Nb V Mo Al Ti N

00rz N -r(00 \O n o0\0o0 00--0- 0--000 00'0 0000 c; cj c; ci - cj c; cj I N 0 o o:: I I N oo o o o oo os ~ oy 1 0.03 0.28 1.46 0.05 0.07 N (s a 2 0.05 ooo a 3 1 1 1 o 4 o o Zoo a 5 cs N l ~ H o II I, 7 0.05 m xD m a 0 8 1 1 1I oo 0. I, 0. 0. 0. II 10 0.12 0.21 1.46 0.05 0.06 o 0.023 oo 0.0081 Ii m t 0.03 \5 t ur 12 0.08 ' ' o ' o ' ' o o o oo o o ooo t N t t vM oo u) t m v v H H ~ H ~ ~ H oo (s (s (s H t N v N v 's oo o o ooo e m m vE (s e oo o o o o - ~ o o v oo o o ooo ~ 's e t m F oo o H (s üolausAUI XUasoIa uosuedm D Table 1 (cont'ed)

Production Conditions Ingot or Slab Production Step Thick plate Production Step No.

Continuous TiN Final Casting Ingot-Making Method #0.02 Heating Heated Reduc- Finish- Thick Temp. γ grain tion at ing ness Average Cooling Soaking Cooling (%) ( C) Size 930 C Temp. (mm) Rate( C/min)*1 Temp. ( C) Rate( C/min.) *2 (ASTM No.) or below ( C) 1 20.0 - - 0.0093 1050 7.0 80 760 20 2 - 1350 1.0 0.0062 1150 4.5 75 730 20 3 - " " " " " 75 800 20 4 - " " " " " 55 740 20 5 - 1320 60 0.0052 1150 5.0 65 740 25 6 - " " " " " 80 690 25 7 - 1350 1.0 0.0062 1150 5.0 35 730 20 8 - " " " " " 55 860 20 9 - " " " 1150 0.5 75 740 20 10 - 1350 1.0 - 1150 1.0 75 720 16 11 - 1350 1.0 - 1150 0.5 80 760 20 12 20.0 - - 0.0026 1150 1.0 75 720 20 *1) Average cooling rate at the center portion of the slab between the molten steel temperature and 1100 C *2) Contents in the steel before the heating for the final rolling Table 1 (cont'ed)

Properties*3 No.

Yield Point Tensile Strength Elongation vE-60 C*4 vTrs pT100 *5 (kg/mm) (kg/mm) (%) (kg-m) ( C) ( C) 1 35.0 59.8 44 24.1 -138 -80 2 55.6 65.7 41 22.8 -132 -74 3 49.1 63.6 42 24.2 -123 -66 4 53.2 63.1 41 25.3 -114 -58 5 47.8 65.9 48 25.9 -142 -78 6 53.9 70.3 44 21.2 -128 -69 7 50.3 60.9 43 21.9 - 76 -34 8 38.8 55.3 44 23.2 - 82 -32 9 56.2 65.8 41 20.1 - 92 -46 10 53.2 59.9 41 16.8 - 91 -26 11 53.8 63.3 43 20.2 - 98 -48 12 52.8 61.9 42 17.2 - 82 -32 *3) All properties are values at right angle to the final rolling direction.

*4) Charpy test pieces were taken from the center of the plate thickness.

*5) 100% ductility transition temperature by 2mmV press notch Charpy test.

Table 2

mO, O q g 00 g I 8 " o= 8 c;c; c; Steel ooo: o a0 1 0.32 a . . . . . . o \o It C tI tt I, It oo: c;c; o oE N Z f o. o: N. u o o o S oo t A tHt t (s oo H X oo o t HNe = > uo!lu3AUI X S XU3S31d 0E Table 2 (cont'ed)

Production Conditions No. Ingot or Slab Production Step Thick Plate Production Step Continuous Ingot-Making Method TiN Heating Heated Reduc- Finish- Final Casting #0.02 Temp. γ grain tion at ing Thick Average Cooling Soaking Cooling (%) *1 ( C) Size 930 C Temp. ness Rate ( C.min) Temp.( C) Rate ( C/min.) (ASTM No.)or below ( C) (mm) 1 - 1350 1.0 0.0062 1150 4.5 75 740 25 2 - 1350 60 0.0058 1150 5.0 75 770 20 3 - " " " 1050 6.5 70 760 20 4 - 1320 1.0 - 1150 0.5 75 720 20 *1) Contents in the steel before the heating for the final rolling Table 2 (cont'ed)

Properties No.

Yield Point Tensile Strength Elongation vE-60 C vTrs pT100 (kg/mm) (kg/mm) (%) (kg-m) ( C) ( C) 1 53.6 63.7 45 26.3 -140 - 87 2 52.5 62.6 42 28.9 -148 - 81 3 50.3 60.9 43 29.2 -152 -103 4 52.1 60.6 42 20.6 -103 - 59 Table 3

ccrc0o0 t vlrte I oO 0-000- 0- 0 0-000 0 0 c; c;cj o o c; Composition rlooccr oO Steel ooo o Mn Nb o d | 'lo 0.03 - 0.21 - 2.30 - 0.024 0.016 0.0073 2 tt tt II I tt tt - tI a 3 0.04 0.32 1.46 0.05 0.04 0.18 - 1.30 0.32 0.021 0.012 0.0053 o ono 0.38 1.28 o oo 5 0.05 me ns nv t s > ' Ho ' Nv a a "0 7 0.P?6 0;2t2 0;8t2 0;03 - 0.21 - 2.30 - 0.024 0.016 0.0073 - tt II ~~ I II 8 e u 10 It II It II tt It tI tI moN It o ooo o o *e m > I 100 1 lto meW n S o. o o, o. o, o. o ooo o o > vo oo H N as A o ' o cs N oo oe (s H o ooo o Hnetm W ooo oE o t , H X uol}U3AUI a Xuasold uosuedtU D .

Table 3 (cont'ed)

Continuous TiN Casting Ingot-Making Method #0.02 Heating Heated Reduc- Finish- Final Temp. γ grain tion at ing Thick Average Cooling Soaking Cooling (%) *2 ( C) Size 930 C Temp. ness Rate ( C/min.) *1 Temp.( C) Rate( C/min.) (ASTM No.) or below ( C) (mm) 1 - 1350 60 0.0069 1150 5.0 80 720 35 2 - " " " 1050 7.0 55 720 35 3 - 1350 1.0 0.0054 1050 6.5 75 740 35 4 - 1350 1.0 0.0058 1150 4.5 65 720 25 5 20.0 - - 0.0082 1150 5.0 75 740 35 6 " - - " " " 65 810 35 7 - 1350 60 0.0069 1150 5.0 40 720 35 8 - " " " 1050 7.0 55 830 35 9 20.0 - - - 1150 1.5 75 760 35 10 " - - - 1050 2.5 75 760 35 *1) Average cooling rate at the center portion of the slab between the molten steel temperature and 1100 C *2) Contents in the steel before the heating for the final rolling Table 3 (cont'ed)

Properties No. Yield Point Tensile Strength Elongation vE-60 vTrs pT100 (kg/mm) (kg/mm) (%) (kg-m) ( C) ( C) 1 52.3 67.3 49 23.2 -128 -71 2 53.6 68.2 50 25.8 -134 -79 3 46.8 62.4 49 24.3 -133 -80 4 50.4 65.0 46 22.5 -126 -72 5 53.8 68.7 50 23.8 -135 -79 6 49.2 62.9 51 26.2 -121 -70 7 51.9 68.8 50 17.4 - 96 -48 8 39.5 54.6 51 20.1 - 89 -39 9 47.4 63.5 47 15.2 - 97 -38 10 48.6 64.2 48 19.2 -102 -43 Table 4

Chemical Composition (wt.%) Steel No. C Si Mn P S Al O N Ti Nb V Mo Others O A1 0.04 0.26 1.52 0.016 0.004 0.036 0.006 0.0050 0.014 0.037 0.061 0.24 A2 " " " " " " " " " " " " A3 " " " " " " " " " " " " O B1 0.08 0.20 1.38 0.012 0.003 0.028 0.005 0.0049 0.012 0.030 - 0.28 Ni 1.60 B2 " " " " " " " " " " - " " B3 " " " " " " " " " " - " " O C1 0.03 0.25 1.90 0.008 0.004 0.035 0.004 0.0051 0.007 0.041 - 0.31 Rare earth element 0.02 C2 " " " " " " " " " " - " " C3 " " " " " " " " " " - " " I 0.13 0.18 1.26 0.020 0.003 0.050 0.006 0.0051 0.015 0.036 0.071 0.30 Ni 0.60 0.08 0.19 1.52 0.013 0.006 0.041 0.006 0.0075 - 0.051 0.060 - Rare earth J element 0.02 0.06 0.26 1.48 0.019 0.005 0.036 0.005 0.0061 0.046 - - - K 0.08 0.25 1.36 0.025 0.014 0.030 0.006 0.0050 0.010 0.041 - 0.30 L O 0.02 0.11 1.65 0.020 0.003 0.026 0.005 0.0055 0.013 0.011 0.056 0.27 Ni 0.20; M Cu 0.26; Rare earth element 0.009 O N 0.08 0.15 1.38 0.018 0.003 0.030 0.006 0.0070 0.012 0.060 0.080 - Ni 0.76; Ca 0.008 O indicates steels according to the present invention.

Table 4 (cont'ed) Ingot or Slab Production Step TiN #0.02 Continuous Ingot-Making (%) Casting Method Average Cooling Soaking Cooling *2 Rate ( C/min.) *1 Temp. ( C) Rate ( C/min.) - 1350 1.0 0.0068 ~ " " z 0.0060 - " " 0.0062 - 1320 1.0 0.0059 - " " 0.0057 - " " 0.0056 - 1350 60 0.0060 - " " 0.0059 - " " 0.0061 - 1320 1.0 0.0048 1350 60 1350 1.0 0.0030 19.0 - - 0.0076 18.5 - - 0.0088 - 1350 60 0.0069 *1) Average cooling rate at the center portion of the slab between the molten steel temperature and 1100 C *2) Contents in the steel before the heating for the final rolling Table 4 (cont'ed)

Hot Rolling Conditions Reheating Step Slab Heated Total Finishing Holding Holding Heating γ Grain Reduction Rolling Temperature Time Temperature Size at 900 C Temperature ( C) (min.) ( C) (ASTM No.) or below (%) ( C) 1150 6.5 70 720 630 30 " 6.0 " " - " 6.5 45 800 630 30 1150 6.0 70 720 600 20 " 5.5 " " - " 5.5 55 770 600 20 1150 6.5 70 700 660 40 1250 6.5 55 750 " " " 6.5 60 820 " " 1150 4.0 70 720 630 30 1150 0.5 70 730 640 30 1250 1.5 70 700 650 20 1150 6.5 65 690 630 30 1050 7.0 70 720 530 5 1150 6.5 65 700 590 20 Table 4 (cont'ed)

Tensile Properties Toughness Properties Charpy Hydrogen Impact Induced Product Yield Tensile ImmV Charpy B.DWTT Absorbed Cracking Point Strength Impact Test 85% Energy at Resistance Thickness (Kg/mm) (Kg/mm) Absorbed Transition SATT Welded Number of (mm) Impact Temperature ( C) Portion Cross Energy at vTrs ( C) -40 C 2mmV Sectional VE-40 C (Kg-m) Crackings (Kg-m) (mm) 32 54 27 -120 -55 18 0 " 50 62 16 - 80 -35 16 4 " 48 61 14 - 85 -30 17 1 26 57 68 30 -140 -80 16 1 " 54 63 17 - 90 -50 15 8 " 55 66 19 -100 -40 16 5 41 65 75 13 -110 -45 8 1 " 59 72 7 - 95 -20 5 7 " 58 71 4 - 90 -15 7 6 26 50 63 8 - 45 -25 12 7 26 52 64 4 - 40 -15 6 3 32 51 63 7 - 60 -30 7 5 26 50 62 6 - 40 -15 2 9 32 53 64 25 -120 -55 24 0 26 56 66 18 -105 -50 12 1

Claims

WHAT WE CLAIM IS:

1. A process for producing high tensile steel sheet (as defined herein) which comprises heating a steel ingot or slab to a temperature not higher than 1150 C, and rolling the steel ingot or slab thus heated with a total reduction amount of not less than 50% in a temperature range not higher than 930 C and a finishing temperature of not higher than 830 C, said steel ingot or slab containing (by weight) 0.01 to 0.13% C, 0.1 to 1.0% Si, 0.7 to 2.0% Mn, not more than 0.1% total Al, 0.004 to 0.03% Ti, 0.001 to 0.009% total N, 0.01 to 0.10% Nb, one or both of 0.01 to 0.15% V, 0.05 to 0.40% Mo, with the balance being unavoidable impurities and Fe, in which slab (Nb%) x (C%) # 5 x 10-3 and the amount of TiN not larger than 0.02 is not less than 0.004%.

2. A process for producing high tensile steel sheet according to claim 1, in which the steel ingot or slab additionally contains one or both of 0.001 to 0.03% REM and 0.0005 to 0.03% Ca, with the limitation that if REM is added, the ratio of REM/S is from 1.0 to 6.0.

3. A process for producing high tensile steel sheet according to claim 1 or claim 2, in which the steel ingot or slab additionally contains one or more of not more than 0.6% Cr, not more than 1.0% Cu, and not more than 4.0%Ni with the limitation that (Cu + Ni)/5 + Cr + Mo 0.90%.

4. A process for producing high tensile steel sheet according to any of claims 1 to 3, in which the steel ingot or slab contains not more than 0.10% C, not less than 0.005 Al, and not more than 0.010% S.

5. A process for producing a high tensile steel sheet as claimed in claim 4, in which the hot rolling total reduction amount is from 60 to 95% in a temperature range not higher than 900 C and a finishing temperature from 800 to 500 C, and the hot rolled ingot or slab is reheated in a temperature range from 300 to 750 C.

6. A process for producing high tensile steel sheet (as defined herein) according to claim 1 and substantially as hereinbefore described.

7. High tensile steel sheet whenever produced in accordance with a process a cording to any of claims 1 to 6.