EP2576848B1

EP2576848B1 - Method for producing a hot-rolled steel product, and a hot-rolled steel

Info

Publication number: EP2576848B1
Application number: EP11743330.0A
Authority: EP
Inventors: Jarkko Vimpari; Tommi Liimatainen; Mikko HEMMILÄ; Ari Hirvi; Jaakko Savola; Juha Kuoppala; Sakari Tihinen; Päivi TAMMINEN; Teemu Peltonen
Original assignee: Rautaruukki Oyj
Current assignee: Rautaruukki Oyj
Priority date: 2010-06-07
Filing date: 2011-06-07
Publication date: 2016-12-07
Anticipated expiration: 2031-06-07
Also published as: EP2576848A1; CN103097556A; FI122313B; WO2011154831A4; CN103097556B; FI20100239A0; WO2011154831A1

Description

Background of the invention

The invention relates in particular to direct quenched martensitic sheet-like steels, on which temper annealing is conducted, i.e. quenched and tempered steels and their production.
The object of the invention is a method for producing a hot-rolled steel according to claim 1.
The object of the invention is also a hot-rolled steel according to claim 8.
From EP1860205A1 is known a martensitic hot-rolled steel with a tensile strength greater than 980MPa, which is very capable of being mechanically cut. The composition of the steel as percentages by weight is: 0.03 - 0.10% of carbon, C; 0.2 - 2.0% of silicon, Si; 0.5 - 2.5% of manganese, Mn; 0.02 - 0.10% of aluminium, Al; 0.20 - 1.5% of chromium, Cr; 0.1 - 0.5% of molybdenum, Mo; and to which can further be added 0.0005 - 0.005% of boron, B; 0.1 - 2.0% of nickel, Ni; and 0.0005- 0.0050% of calcium, Ca. The steel is produced by direct quenching at a temperature of less than 400°C, such as, for example, at a temperature of 250 - 300°C. Temper annealing is not conducted on the steel. The purpose of the publication is to achieve mechanical properties without precipitation hardening alloying elements, such as titanium Ti, niobium Nb or vanadium V, as well as by decreasing the carbon C and increasing the molybdenum Mo content. According to the teaching, the effect of molybdenum Mo ends at an upper limit of 0.5% Mo, after which alloying it pointlessly increases costs. Additionally, the publication teaches that nickel can be added 0.1 - 2.0%.
The disadvantage of this known steel composition and method is that the steel presented therein it is not suitable for use as a structural steel in application sites, because its elongation and impact toughness are not remarkably good. Elongation and impact toughness are difficult to improve in the steel in question, because it is not particularly temper-resistant. In addition, a disadvantage is that it is not well suited for steel products that during usage will have to be for long periods of time in the temperature range of 450 - 600°C, which is a dangerous temperature range due to the upper temper brittleness. Steel can be subjected to this temperature range in different usage situations, such as in heat treatments, or in a situation, in which steel structures are reworked hot (in shape corrections by heating) or during bell furnace annealing, in which annealing occurs a slow cooling in said temperature range. When subjected to a upper temper brittleness, the steel becomes fragile at room temperature and thus quite useless. Temper brittleness causes, among other things, atomic segregations forming at the grain boundaries, which weaken the structure.
In addition, generally known are conventional quenched and tempered steels, whose carbon content is high, such as on the level of C 0.12 - 0.18% and/or into these is alloyed more nickel Ni, copper Cu or niobium Nb than in a hot-rolled steel according to the invention. With quenched and tempered steels, particularly direct quenched tempered steels, all important properties, such as yield strength, elongation, impact toughness, flangeability and temper resistance, are difficult to achieve at good levels simultaneously in the same steel.
EP1764423 discloses a method for forming a high tensile strength steel plate which involves direct quenching from the Ar₃ transformation point to below 400°C.
JP8-143954 discloses a quenched and tempered steel plate.
EP1375694 discloses a hot rolled steel strip with a microstructure comprising at least 95% martensite and/or bainite.
Metallurgical transations, Vol 22A, Oct 1991, pp2359-2374 discloses the effects of vanadium and processing parameters on the structures and properties of a direct quenched low carbon Mo-B steel.

Brief description of the invention

The object of the invention is to eliminate the disadvantages related to known art and to achieve a high-strength hot-rolled steel that is very temper-resistant after the direct quenching process, wherein by tempering it is made further high-strength (R_p0.₂ ≥ 890MPa) combined at the same time with good impact toughness (Charpy V (-20°C) ≥ 37J/cm²) and flangeability as well as good weldability.
Another object of the invention is to provide a hot-rolled steel production method that is as easy as possible in relation to tempering treatment, i.e. a hot-rolled steel according to the invention must be as robust as possible in relation to tempering, or easily tempered, wherein it is preferable to implement a tempering treatment. Steel is not critical, for example, in relation to tempering temperature and the time used for tempering and its tendency for upper temper brittleness is low.

The invention consists among others in technical solutions according to the following items:
1. A method for producing a hot-rolled steel product, characterized in that it is arranged a steel billet, whose composition as a percentage by weight is

C	0.075 - 0.12%
Si	0.1 - 0.8%
Mn	0.8 - 1.7%
Al	0.015 - 0.08%
P	less than 0.012%
S	less than 0.005%
Cr	0.2 - 1.3%
Mo	0.15 - 0.80%
Ti	0.01 - 0.05%
B	0.0005 - 0.003%
V	0.02 - 0.10%
Nb	less than 0.3%
Ni	less than 1%
Cu	less than 0.5%

the rest being Fe and unavoidable impurities, in which method the steel billet having said composition is
heated (1) to the austenitizing temperature of 1200 - 1350 °C, and
hot-rolled (2, 3) to the desired thickness such that the rolling temperature of the billet at the last pass is 760 - 960 °C, and
direct quenched (4) after the last pass conducted using one-step cooling at a cooling rate of 30 - 150 °C /s to a temperature of 300 °C at the most, which direct quenching is conducted at the latest 15 s after the last hot-rolling pass,
characterized in that the hot-rolled steel product is a strip steel, which, after being direct quenched (4), is coiled (5) and subsequently temper annealed (6) at a temperature of 200 - 700 °C for 24 hours at the most, and
in that the austenite of the hot rolled steel billet is flattened before the direct quenching so that the average flattening ratio of the grain of the microstructure of the steel is greater than 2.

item

characterized

item

characterized

item

characterized

²

characterized

In order to implement the objects of the invention, the method according to the invention is characterized in that it is arranged a steel slab, whose composition as percentages by weight is

the rest being Fe and unavoidable impurities, in which method the steel slab is
heated to austenitizing temperature of 1200 - 1350 °C, and heat-rolled to the desired thickness such that the roller temperature of the slab at the last pass is 760 - 960 °C,
direct quenched after the last pass conducted using one-step cooling at a cooling rate of 30 - 150 °C /s to a temperature of 300 °C at the most, which direct quenching is conducted at the latest 15 s after the last hot roller pass,
temper annealed at a temperature of 200 - 700 °C for 24 hours at the most, and such that the austenite of the hot-rolled steel billet is flattened before the direct quenching so that average flattening ratio of the grain of the microstructure of the steel is greater than 2.

The preferred embodiments of the method according to the invention are presented in claims 2 - 7.
As a result of the method according to the invention, the microstructure of a hot-rolled steel is preferably tempering martensitic, i.e. in the steel has, as a result of direct quenching, formed an essentially martensitic microstructure, after which the steel is subjected to temper annealing, wherein the final result is a hot-rolled steel product, whose impact toughness and strength are of the desired level.
The greatest advantage of the method according to the invention is that temper annealing, which substantially improves the impact toughness and elongation of the steel product, is simple to implement on a hot-rolled steel according to the invention. The strength and impact toughness properties of the steel are not sensitive to changes in tempering temperature and time nor to the cooling rate of the sheet after tempering. Using direct quenching, it is also achieved a good flangeability for the steel, which is typically more difficult to achieve for the direct quenched tempering steel in comparison to traditionally furnace-quenched steel.
In order to achieve the objects of the invention, the composition of a hot-rolled steel according to the invention is in particular characterized in that the carbon C and manganese Mn contents are low, being in the range presented, and, additionally, into the steel is always alloyed the presented contents of boron B, vanadium V and titanium Ti, in order that the objects of the invention can be achieved. It is not absolutely necessary to alloy Niobium Nb and if it is alloyed, its content is limited. Additionally, the nickel Ni and copper Cu contents can be quite low, being even at the level of impurities. The significance and effects of the alloying elements are described further in the detailed section of the description.
The properties of a hot-rolled steel are presented in independent claim 8. The preferred embodiments of a hot-rolled steel are presented in claims 9 - 15.
The greatest advantages of a hot-rolled steel according to the invention are that, in addition to high strength, its impact toughness and flangeability can, using the method according to the invention, be simultaneously produced to a good level. Additionally, the hot-rolled steel according to the invention is exceptionally temperresistant, because its composition enables that high-strength martensitic steel can be tempered, for example, in a bell furnace and, additionally, at the same time efficiently limit the detrimental effects of upper temper brittleness. The impact toughness properties of the steel are, indeed, excellent also as measured from HAZ (heat affected zone) area of the welding seam, which is exceptionally important for building steel use. The steel is also quite suitable for use particularly in the welded boom structures of cranes. Additionally, the steel possesses excellent usability due to good weldability and flangeability.
In the invention, it is surprisingly found that, using said composition, a steel is achieved that can, after direct quenching, be temper annealed even in a temperature area (450 - 600 °C) of the upper temper brittleness that is typical for quenched and tempered steels and nonetheless achieve the objects of the invention in a structural steel.

Brief description of the figures

The invention will now be described in greater detail by means of examples, with reference also to the accompanying figures, in which

Fig. 1 shows the main steps of a preferred embodiment of the method according to the invention as a time-temperature curve, in which the reference numbers for the process steps of the method are: 1 = furnace heating, 2 = pre-rolling, 3 = strip rolling, 4 = direct quenching, 5 = coiling, 6 = temper annealing,
Fig. 2 shows a welding test arrangement, which shows the measuring point of the fusion line FL,
Fig. 3 shows the microstructure of a hot-rolled steel according to the invention in a tempering martensitic state, and
Fig. 4 shows the main steps of another preferred embodiment of the method according to the invention as a schematic time-temperature curve, in which the reference numbers for the process steps of the method are: 1 = furnace heating, 2 = pre-rolling, 3 = strip rolling, 4 = direct quenching, 5 = coiling, 6 = temper annealing.
Fig. 5 shows an example of the effect of high carbon content on the impact toughness of the base material (steel C)
Fig 6. shows an example of the effect of high Si content on the impact toughness of the base material (steel F)
Fig 7. shows an example of the effect of high Mn content on the impact toughness of the base material (steel G)
Fig. 8 shows an example of the uselessness of high nickel content for impact toughness of the base material (steel B)
Fig. 9 shows an example of the effect of a hot-rolled steel according to the invention on impact toughness, which is excellent both transversally and longitudinally (steel K of table 1)
Fig. 10 shows the correlation between TBI and impact toughness
Fig 11. Shows the TBI value of different test steels as a function of measuring temperature, as measured longitudinally in relation to the direction of rolling
Fig. 12 shows the TBI value of different test steels as a function of measuring temperature, as measured transversally in relation to the direction of rolling

Detailed description of the invention

The composition of a hot-rolled steel according to the invention as percentages by weight is:

the rest being iron, Fe, and unavoidable impurities.

In the following is described in greater detail the composition of a hot-rolled steel according to the invention and, by way of example, the properties achieved by each composition with the most important production parameters. Additionally, preferred embodiments and their advantages are presented. The contents are percentages by weight.

Table 1. The chemical composition of test steels

steel	C	Si	Mn	Cr	N	S	P	Cu	Mo	Ti	V	Al	Nb	Ni	B	CEV	*
A	0,14	0,2	1,12	0,38	0,004	0,002	0,01	0,02	0,49	0,012	0,041	0,04	0,026	0,37	0,0019	0,53	R
B	0,078	0,23	1,76	1,49	0,007	0,004	0,011	0,02	0,24	0,01	0,072	0,037	0,002	1,14	0,0003	0,81	R
C	0,149	0,2	1,56	0,98	0,005	0,003	0,009	0,02	0,24	0,023	0,068	0,026	0,004	1,15	0,0022	0,74	R
D	0,081	0,21	1,5	0,97	0,006	0,001	0,01	0,02	0,47	0,016	0,015	0,026	0,052	0,06	0,0019	0,63	R
E	0,151	0,17	1,48	1,02	0,006	0,003	0,009	0,02	0,23	0,014	0,07	0,025	0,004	1,19	0,0002	0,74	R
F	0,085	0,51	1,51	1,02	0,007	0,001	0,008	0,02	0,48	0,029	0,082	0,033	0,004	0,05	0,0023	0,66	I
G	0,081	0,2	1,45	0,99	0,006	0,001	0,01	0,03	0,5	0,027	0,076	0,034	0,002	0,05	0,0021	0,64	I
H	0,081	0,51	1,47	0,98	0,006	0,001	0,009	0,02	0,47	0,015	0,011	0,028	0,051	0,05	0,0019	0,62	I
I	0,076	0,22	1,04	0,7	0,006	0,001	0,009	0,02	0,65	0,028	0,082	0,036	0,004	0,06	0,0016	0,54	I-PREF
J	0,087	0,22	0,99	0,72	0,007	0,001	0,011	0,04	0,64	0,032	0,082	0,038	0,003	0,07	0,0016	0,55	I-PREF
K	0,103	0,2	0,99	1	0,006	0,001	0,01	0,03	0,65	0,028	0,079	0,034	0,003	0,06	0,0014	0,62	I-PREF
L	0,099	0,23	1,01	0,99	0,006	0	0,012	0,02	0,65	0,031	0,084	0,032	0,002	0,06	0,0015	0,62	I-PREF
M	0,084	0,18	1,04	1,33	0,006	0,004	0,009	0,03	0,14	0,027	0,046	0,028	0,003	0,06	0,0019	0,57	R

* I the composition of a hot-rolled steel according to the invention
* I-PREF the composition of a hot-rolled steel according to the preferred embodiment according to the invention
* R the composition outside the composition of a hot-rolled steel according to the invention

Table 2. Production parameters and mechanical properties of the tests, sheet thickness t approx. 6 mm.

Steel	Rolling Temperature *	Heat Treatment **	Rp0.2 (MPa)	Rm (MPa)	A5(%)	CV(J/cm², -40°C)
A	2	2	1101	1176	10.7	81
B	2	1	1059	1120	11.7	55
C	2	2	1073	1126	13.1	40
D	2	3	1083	1112	12.4	76
E	1	2	1039	1088	13.0	47
F	2	3	986	1019	13.4	81
G	2	3	1021	1053	13.4	83
H	1	4	1014	1035	11.6	108
I	2	3	1024	1060	13.3	142
K	1	3	1091	1138	11.1	114
L	1	4	1071	1130	11.5	104
M***	2	3	890	913	14.2	120

* = Temperature of the steel during the last rolling pass: 1 = below 900 °C, 2 = above 900 °C.
** = Heat treatment temperature: 1 ≤ 500 °C < 2 ≤ 550 °C < 3 ≤ 600 °C, 4 > 600 °C
*** = annealed in a conventional furnace, holding time 1 hour.

Mechanical properties are defined according to the testing instructions of standard ISO 10025-6.

All the steels of the tables are produced by the method according to the invention, i.e. by direct quenching to a low temperature, wherein the coiling temperature has been below 300 °C and by the subsequent tempering treatment, which is performed, for example, in a Bell-type type of furnace. Impact toughness tests are performed as Charpy V tests using a 6 mm thick test material.

Table 3. Steel flanging results

Steel	Direction *	R/t **	W (mm) ***
A	Longitudinal	2.5	75/100
A	Transversal	2.0	75/100
E	Longitudinal	2.5	100
E	Transversal	2.4	75/100
F	Longitudinal	2.7	100
	Transversal	2.5	75
H	Longitudinal	3.4	75/100
H	Transversal	5.1	75/100
I	Longitudinal	2.0	75/100
I	Transversal	1.4	75/100
K	Longitudinal	2.2	75
K	Transversal	3.1	75
L	Longitudinal	2.7	75/100
L	Transversal	3.3	75/100

* direction of flanging; longitudinal = edge lengthwise in relation to the direction of rolling, transversal = edge crosswise in relation to the direction of rolling
** R = bending radius, t = sheet thickness
*** W = the width of the opening (mm), to which the flanging is made

Flanging is implemented by a known method as a V-bending between an upper-lower tool. Free flanging is used as the manner of flanging.
The carbon equivalent of the steel C + Mn/6 + (Cr + Mo + V)/5 + (Ni + Cu)/15 is somewhat high, but despite this good weldability can be achieved using a low carbon content, as is observed in the following.
In the invention, it is found that for conventional and relatively high-carbon quenched and tempered steels combined with conventional, or somewhat high, Mn content, impact toughness values remained low, wherein it was found as important to limit the maximum carbon C and manganese Mn contents.
Limiting the maximum carbon and manganese contents is particularly important, when temper annealing occurs below a temperature of 600°C or the steel cools down slowly after tempering through the temperature range in question.
According to the objects of the invention, a high impact toughness is produced for both the base material and the welding HAZ area, particularly such that the Charpy V impact toughness of the base material is at least 37J/cm² as measured longitudinally in relation to the direction of rolling and at a temperature of -20°C. Preferably, the impact toughness of the base material is at least 33J/cm² as measured transversally in relation to the direction of rolling and at a temperature of at least -20°C. Most preferably, said impact toughness requirements are also achieved as measured at a temperature of -40°C.
The impact toughness of the steel was defined as a Charpy V test using three welding HAZ (heat affected zone) areas, forming a notch in the following sites:

1. In the fusion line, where impact toughness was measured from the site FL, where a segment placed in the direction of the sheet and mid-way in the thickness of the sheet cuts the fusion line formed while welding, Fig. 2.
2. In the area of the coarse grained zone (CGHAZ, coarse grained HAZ), where impact toughness was measured from a site located at a 1 mm distance from the measuring site of the fusion line FL towards the base material (FL+1)
3. Partially in the area of the austenitized zone (ICHAZ, intercritical HAZ), where impact toughness was measured from a site located at a 3 mm distance from the measuring site of the fusion line FL towards the base material (FL+3)

Due to the lower carbon content of the steel particularly in the partially austenitized zone (ICHAZ), in which the temperature is at the maximum 700 - 850 °C, impact toughness is retained in the steel while welding as better than in the higher-carbon quenched and tempered steel produced in a typical traditional manner. In this zone (ICHAZ), austenitizing occurs only there, where nucleation of the austenite has been easy, i.e. mainly there, where carbon content has been high. The high-carbonous austenitized part changes upon cooling to martensite and bainite. Upon cooling, the high-carbonous local austenite area can form as a hard MA island, weakening the impact toughness of the zone, wherein the lower carbon content of the developed steel is of advantage, because the formation of hard and more fragile microstructures is lesser in the area of the ICHAZ.
The composition of a hot-rolled steel according to the invention achieves exceptionally good impact toughness particularly in the area of the partially austenitized zone (ICHAZ), which is measured from the site FL+3. The hot-rolled steel according to the invention is thus quite weldable also without expensive alloying of nickel, when the steel is achieved alloyed with vanadium, wherein impact toughness in the HAZ zone is at least on the level of typical quenched and tempered steels or better. Table 4. Charpy V impact toughness values typical for welds, t = 6 mm.

Arc energy E1=0.6kJ/mm

-20 °C FL+1 >50J

FL+3 >46J

-40 °C FL+1 >25J

FL+3 >37J

Arc energy E2=0.8kJ/mm

-20 °C FL+1 >45J

FL+3 >50J

-40 °C FL+1 >20J

FL+3 >40J
Table 4 shows typical impact toughness values with different heat inputs for the composition K, which is presented in table 1. As the welding method is used MAG welding in the flat position without preheating and a 50 ° V-groove as the groove shape. Tests were conducted using two different heat inputs Q1=0.48 kJ/mm (arc energy E1=0.6 kJ/mm) and Q2=0.64 kJ/mm (arc energy E2=0.8 kJ/mm), wherein the calculated cooling time of the joints in the heat interval 800...500 °C (T8/5) was 7 s and 13 s.
Carbon content as a percentage by weight 0.075 - 0.12% is low in comparison to typical quenched and tempered steels, wherein impact toughness remains at a good level. If the carbon content of the steel as a percentage by weight is less than 0.075%, then it is difficult to get the steel strong and impact tough enough, because, in this case, enough martensite is not formed as a consequence of direct quenching. If the carbon content as a percentage by weight is above 0.12%, then impact toughness is weakened too much and the objects of the invention are not achieved.
Preferably, the carbon content of the steel as a percentage by weight is 0.08 - 0.11%, more preferably 0.09 - 0.11%, wherein, in welding, also the HAZ zone achieves adequate uniform strength with the base material while at the same time the impact toughness of the base material is adequate.

It is generally known that, in general, a low carbon equivalent value and carbon content are more advantageous for weldability than high. However, in the invention, it is surprisingly found that tensile tests pulled transversally over the welding seam were weaker with composition I than with composition K, the carbon equivalent (CEV) and carbon content of which composition K are greater than of composition I. As an example of this is enclosed comparison table 5. Steel K achieves also in the HAZ zone excellent impact toughness, because its carbon content is in the preferred carbon area of the invention 0.09 - 0.11%.

Table 5. Example of the mechanical properties of two example steels

	Arc Energy E1= O.6kJ/mm		Arc Energy E2= 0.8kJ/mm
Steel	I	K	I	K
R_p0.2 (MPa)	952	1034	921	1022
R_m (MPa)	1035	1116	1024	1111
HV10 (FL+1mm)	332	355	335	353
HV10 (FL+3mm)	346	352	359	321

As an example of the detrimental effect of high carbon content to the base material is steel C, whose composition is shown in table 1 as well as rolling and temper annealing parameters, and mechanical properties in table 2. From Fig. 5, it is observed that transversal impact toughness is poor, when carbon content is greater than the carbon content of a hot-rolled steel according to the invention.
Silicon content as a percentage by weight is 0.1 - 0.8%. Preferably, silicon content as a percentage by weight is 0.1 - 0.4%, more preferably 0.1 - 0.3%. However, in the invention, it is surprisingly found that too high a silicon content, such as a 0.5% content as a percentage by weight, can detrimentally effect on the impact toughness of the steel. This can be clearly seen in steel F from Fig. 6.
For said reason, silicon content as a percentage by weight is preferably at the most 0.4%. Silicon contents less than 0.1% are not recommended, because desulphurisation of the steel and form control of inclusions are easier, when the steel contains some silicon.
Additionally, silicon, Si, increases the strength of the steel without a rise in carbon equivalent, which is an advantage especially if carbon content is close to the carbon content upper limit 0.11 - 0.12% of a hot-rolled steel according to the invention.
Manganese content as a percentage by weight is 0.8 - 1.7%. Preferably, manganese content as a percentage by weight is 0.8 - 1.4%, more preferably 1.0 - 1.2%. In order to assure good hardenability, manganese content as a percentage by weight must be at least 0.8%, preferably at least 1%. On the other hand, unfavourable segregation of manganese is less, when manganese content as a percentage by weight is limited to at the most 1.4%, preferably at the most 1.2%.
As an example is shown in Fig. 7 the detriment of high manganese content to the base material steel G, whose content is shown in table 1 as well as rolling and temper annealing parameters, and mechanical properties in table 2.
Chromium content as a percentage by weight is 0.2 - 1.3%, more preferably 0.5 - 1.3%, in order that the high strength steel is achieved and hardenability is good.
More preferably, chromium content as a percentage by weight is 0.8 - 1.2%. There is preferably at least 0.8% of chromium, in order that a welding joint of adequately uniform strength is achieved at a low carbon content, on the other hand, there is preferably at the most 1.2% of chromium due to the excessive rise in carbon equivalent, which is of particular detriment, when carbon content is near the 0.11 - 0.12% carbon content upper limit of the invention.
Boron content as a percentage by weight is 0.0005 - 0.003%, because alloying with boron is a preferred means to assure the good hardenability of the steel. At contents above 0.003%, the hardenability-increasing effect of boron weakens and, additionally, too much boron weakens the weldability of the steel. Preferably, boron is alloyed 0.0008 - 0.002% as a percentage by weight both to retain good impact toughness of the weld and to assure adequate hardenability.
Nickel content must be limited to a content of less than 1% as a percentage by weight, because nickel can, under some circumstances, even decrease somewhat the impact toughness of the tempered steel or its effect is slight. Additionally, nickel is an expensive alloying element. Preferably, the content of nickel is to be limited to a content of less than 0.1% as a percentage by weight, more preferably less than 0.05%, wherein the alloying costs of the steel can be kept as low as possible. The composition of a nickel-alloyed steel B after tempering treatment is of modest impact toughness, transversal impact toughness results in particular are modest, which is observed from Fig. 8. Tempering treatment is performed in a Bell-type furnace for 24 hours at the most and at a temperature below 500 °C.
Molybdenum content as a percentage by weight is 0.15 - 0.80%. Preferably, molybdenum content as a percentage by weight is 0.30 - 0.80%, because, with a molybdenum content of less than 0.30% in a steel according to the invention, adequate strength is not achieved without the needing to alloy into the steel large contents of other alloying elements, such as carbon C, silicon Si, nickel Ni or manganese Mn, the detrimental effects of which are described earlier and also later in connection with the indexes TBI and UTBI presented in the description.
Molybdenum precipitates in temper annealing, which decreases the lowering of strength caused by tempering treatment and thus helps in achieving high strength. Additionally, molybdenum is used i.a. to prevent the upper temper brittleness of steel by slowing segregation of i.a. phosphorus, P, to the grain boundaries during temper annealing at the critical temperature range of 450 - 600 °C. Molybdenum also efficiently increases the hardenability of steel.
In order to assure temper resistance, according to one embodiment, molybdenum is alloyed 0.50 - 0.70% as a percentage by weight. Contents exceeding a 0.8% molybdenum content increase the carbon equivalent value and increase excessively alloying element costs, because molybdenum is an expensive alloying element. On the other hand, at a Mo content less than 0.15%, as in steel M, whose composition is shown in table 1 and test results in table 2, show that strength remains low in temper annealing of 500 - 600 °C already for a relatively short 1 hour temper annealing time. For this reason, i.e. to achieve adequate strength, molybdenum must be alloyed at least 0.15% as a percentage by weight, preferably at least 0.30% or even at least 0.50%.
Although niobium alloy is used in many conventionally produced, well-flangeable quenched and tempered steels, in the invention, it surprisingly was found that the flangeability of direct quenched steel is not achieved at a good level, neither in the hardened nor in the tempered state, if the steel contains large amounts of niobium, Nb. As an example of this is steel H in table 3. Thus, in the invention, it is surprisingly found that niobium can crucially weaken steel flangeability in a hot-rolled steel according to the invention, especially at large contents.
It is thus not absolutely necessary to alloy niobium, but if it is alloyed, its content is limited to a Nb content of less than 0.3% as a percentage by weight, wherein it can, in some situations, effect on strength. Preferably, niobium content is to be limited to 0.03% Nb at the most, because, at the 0.05% niobium content of steel H, it was observed a clear weakening of flangeability. More preferably, niobium content is limited to less than 0.005%, wherein the best possible flangeability properties for the steel are assured.
Vanadium content must be 0.02 - 0.1% as a percentage by weight. In order to assure strength, vanadium, V, is to be alloyed at least 0.02% as a percentage by weight. As vanadium content increases, weldability can weaken and, for this reason, the vanadium content maximum value as a percentage by weight is 0.1% at the most.
According to the preferred embodiment, the vanadium content must be 0.04 - 0.1% as a percentage by weight, when niobium, Nb, is not alloyed, i.e. when Nb is less than 0.005%. Vanadium is thus alloyed in particular without alloying of niobium, in order that flangeability would be as good as possible. In the invention, it is surprisingly found that alloying of vanadium is not detrimental to flangeability with the composition of the invention, as is observed from tables 2 and 3, although niobium, Nb, was found to have a flangeability-weakening effect, when steels are compared at the same strength and carbon levels.
According to one embodiment of the invention, vanadium contents and niobium contents are selected as follows: V 0.04 - 0.10% as a percentage by weight and Nb 0.008 - 0.03% as a percentage by weight, wherein it is achieved a good combination of impact toughness and strength while flangeability still remains reasonable.
According to one embodiment of the invention, vanadium contents and niobium contents are selected as follows: V 0.02 - 0.03% as a percentage by weight and Nb 0.008 - 0.03% as a percentage by weight, wherein it is achieved, above all, a combination of HAZ zone strength and impact toughness in the highest possible quality, particularly by severely limiting the content of vanadium, but by, however, still reasonably alloying niobium. Alloying of niobium is of advantage particularly in achieving adequate strength and impact toughness in the base material.
Copper content is limited to less than 0.5% as a percentage by weight. It is not absolutely necessary to alloy copper, but it can be used in a small amount as needed to increase strength or improve weather resistance of the steel. If copper, Cu, is alloyed more than 0.3%, nickel must be alloyed at least 0.33 * Cu content, in order that the surface quality of the steel remains good in hot-rolling.
Preferably, copper content as a percentage by weight is less than 0.05%, wherein its content is on the level of impurities, and adequate strength can be attained less expensively in terms of costs and properties without alloying copper.
Aluminium content as a percentage by weight is 0.015 - 0.08%. Aluminium, Al, is used to kill steel, i.e. to bind oxygen from the steel. Preferably, aluminium content is 0.02 - 0.06% as a percentage by weight.
Titanium content is 0.01 - 0.05% as a percentage by weight, because titanium is required for binding nitrogen, N, in the steel, in order that boron, B, functions efficiently as an improver of hardenability and does not form boron nitrides. Titanium is used, because it works more reliably with direct quenched steel than aluminium, Al. Preferably, there is 0.02 - 0.03% of titanium as a percentage by weight, because at lower contents, if nitrogen remains high for some reason, it is not possible to get all the nitrogen bound. On the other hand, higher contents increase amounts of the relatively large-sized TiN, which is detrimental in terms of impact toughness. Ti/N ratio is preferably 3-4.
Phosphorus content as a percentage by weight must be limited to P less than 0.012%, because phosphorus weakens impact toughness. Preferably, phosphorus content as a percentage by weight is limited to less than 0.008%.
Sulphur content is limited as an impurity to a level of less than 0.005% as a percentage by weight to assure good impact toughness and formability.
Next, Fig. 9 presents as an example (steel K of table 1) the excellent effect achieved by the composition of a hot-rolled steel according to the invention on the impact toughness of the steel, which is excellent both transversally and longitudinally.
A hot-rolled steel means a steel hot-rolled to be sheet-like, such as a hot-rolled heavy plate or hot-rolled strip steel. According to the most preferred embodiment, the hot-rolled steel is a hot-rolled strip steel, because it is most easily achieved as excellent in terms of production efficiency, costs, surface quality and measurement tolerances. The thickness of the strip steel can be 2 - 10 mm, however, preferably in the range of 4 - 8 mm.
A hot-rolled steel means in particular a direct quenched steel, whose microstructure is essentially martensitic. Most preferably, after direct quenching, tempering treatment is conducted on the hot-rolled steel, wherein it is a question of a direct quenched and tempered steel, whose microstructure is essentially tempering martensitic.
The microstructure of the steel before tempering treatment is preferably as perfectly composed as possible (above 90%) of martensite and self-tempered martensite. In any event, the majority of the microstructure must be like this, wherein bainite may appear in the structure to some degree. The content of ferrite and perlite before tempering must generally be in total less than 10%.
The austenite of a hot-rolled steel is flattened before direct quenching. The flattening ratio of the grain is the numeric ratio of average grain height (H) / width (W) defined from the microsection. Grain is measured from the section, the section surface of which is in the direction of rolling and in the direction of sheet thickness as well as at the inspection site of a depth about ¼ the thickness of the sheet.
The flattening ratio of the grain must be greater than 2.0, which is formed, when steel is direct quenched directly from the hot-rolling occurring in the austenite area and the steel does not have time to re-crystallize. In traditional furnace-tempered steels, the ratio is less than 2.0. Most preferably, the average flattening ratio of the grain structure of a hot-rolled steel according to the invention is greater than 4.0.
Fig. 3 shows a picture of the microstructure of a steel product produced by the method according to the invention, in which is shown the height (H) and width (W) of the grain. The figure shows thus the preferred embodiment of a hot-rolled steel according to the invention in a direct quenched and tempered state, i.e. as tempering martensitic, in which the flattening of the microstructure is still recognizable. In the example, the flattening ratio of the grain structure W1/H1 is approx. 16 and W2/H2 is approx. 28. The flattening of the grain structure is significantly affected by the rolling temperature used, which, in the method according to the invention, is at the last rolling pass in the range of 760 - 960 °C.
The yield strength of a hot-rolled steel according to the invention is 890 - 1200 MPa, most preferably 960 - 1100 MPa. This is achieved by immediate direct quenching after the rolling of hot-rolling, after which tempering treatment is conducted. Tempering treatment can be conducted either immediately or later. Elongation at break (A5) is at least 8%, most preferably more than 10%.
Yield ratio is typically somewhat high in structural steels and the yield ratio (yield strength/breaking strength) of a hot-rolled steel according to the invention is above 0.85.

The method according to the invention is characterized in that it is arranged a steel slab, whose composition as percentages by weight is

the rest being iron, Fe, and unavoidable impurities, in which method the steel slab is
heated to the austenitizing temperature of 1200 - 1350 °C, (reference number 1),
hot-rolled to the desired thickness such that the rolling temperature of the slab at the last pass is 760 - 960 °C, (reference numbers 2, 3),
direct quenched after the last pass conducted using one-step cooling at a cooling rate of 30 - 150 °C /s to a temperature of 300 °C at the most, which direct quenching is conducted at the latest 15 s after the last hot-rolling pass, (reference numbers 4, 5), and
temper annealed at a temperature of 200 - 700 °C for 24 hours at the most (reference number 6).

The preferred embodiments of the method according to the invention are presented in claims 2 - 7.

Fig. 1 shows the steps of the method according to the invention for producing a hot-rolled steel product. As the starting material is a steel slab, whose composition as percentages by weight is

the rest being iron and unavoidable impurities.

In step 1 of the method, the steel slab is heated to the austenitizing temperature of 1200 - 1350 °C. The thickness of the steel slab is, for example, 210 mm and it is heated to the austenitizing temperature of 1280 °C, where it is kept until it is of adequately even warmth and the alloying elements have adequately dissolved into the matrix, in practise for several hours. Naturally, the thickness of the steel slab can vary from that presented and the austenitizing temperature can be selected differently, but it is recommended that it is in the range of 1200 - 1350 °C. If the austenitizing temperature is below said lower limit, then there is a danger that not all microalloying elements dissolve into the austenite, i.e. the austenite is not made homogenous and, in precipitation, annealing strength may remain low. On the other hand, a higher temperature would lead to exceptionally large grain size of the austenite and increased oxidation of the slab surface. Annealing time can most suitably vary in the range of 2 - 4 hours, but, depending on the selected furnace technology and the thickness of the slab, it can also be significantly longer or shorter.
After heating, in the second step, hot-rolling 2 is conducted, which comprises pre-rolling step 2 and the subsequent strip rolling step 3. The temperature of hot-rolling at the last pass is 760 - 960 °C. Preferably, the end temperature at the last pass of the hot-rolling is 800 - 900 °C. The end temperature of hot-rolling is at least 800 °C, in order that rolling forces remain reasonable and at the most 900 °C, wherein i.a. excellent surface quality is assured.
After hot-rolling, the steel is direct quenched, i.e. cooled at an accelerated rate. Preferably, the speed of direct quenching 4 is at the most 120 °C /s, because, in this case, such a microstructure is achieved for the steel that gives the steel exceptionally good mechanical properties, including good impact toughness, combined with good flangeability. Quenching can be conducted, for example, with water.
Preferably, the end temperature of direct quenching 4 is at the most 130 °C, because, in this case, after quenching, a planar strip is achieved, the edges of which are also even and planar.
Preferably, direct quenching 4 of the steel strip is conducted directly at the coiling temperature and it is coiled 5.
The hot-rolled steel product is preferably a steel strip, which, after being direct quenched 4, is coiled and subsequently temper annealed 6.
Preferably, temper annealing treatment 6 is conducted on the steel in the temperature range of 450 - 599 °C, wherein the composition of a low-carbon steel according to the invention can be formed to be inexpensive in terms of both its total amounts of alloying elements and its cost.
Alternatively, tempering treatment 6 of the steel can be conducted in the temperature range of 200 - 449 °C or 600 - 650 °C.
Thus, after direct quenching, temper annealing treatment 6 of the method according to the invention can be implemented for the strip sheet cut from the coil or for a sheet continuously unwinding from the coil. On the other hand, after direct quenching, temper annealing treatment can alternatively be implemented also for a whole coil, for example, in a bell furnace, in which the temperature rises and falls slowly. Temperature variation between the midpoint and the surface specific to tempering of the coil is not a problem, because a hot-rolled steel according to the invention is exceptionally robust in terms of tempering. Robust means in this connection that for the steel homogenous mechanical properties are achieved in every part of the coil regardless of how the steel is tempered. Due to robustness, the method can very well be implemented also for sheet-rolled sheets of different thicknesses and strip sheets cut from the coil without the tempering furnace technology needing to be adjustable for exceptionally exact tempering temperature and time. This, in turn, enables the use of inexpensive and simple furnace technology and decreases the risk of rejection of material.
According to one embodiment of the method of the invention, the hot-rolled steel, on which direct quenching 4 is conducted, is cut as sheets, after which the sheets are straightened and only at the end is tempering treatment conducted. Thus is achieved temper annealing treatment 6 for straightened sheets, in the straightening of which could have formed detrimental stresses. The final result is an exceptionally even steel sheet of exceptionally even quality and 2 - 12 mm in thickness, in which elongation and impact toughness are somewhat better than with other embodiments.
If the steel is not disposed to upper temper brittleness, using a low annealing temperature, high strength may be achieved more easily than with a higher, alloying costs may be saved or use could even be made of simple and inexpensive, but efficient in production capacity, bell furnace annealing, in which cooling and heating occurs slowly.
Possible development of brittleness (or toughness) of the steel in tempering was examined by annealing test steels in different types of furnaces (bell furnace and conventional), using different tempering times (0.5 - 24 h), and temperatures (200 - 650 °C).
On the basis of tests, the indexes describing impact toughness (or temper brittleness) TBI (temper brittleness index) and UTBI (upper temper brittleness index) were defined for direct quenched strip steels made on a production scale (full scale test).
TBI describes a measured impact energy value in a Charpy V test, when the steel is annealed in the non-critical area for upper temper brittleness, i.e. above or below the temperature range of 450 - 599 °C (at temperature T below 450 °C or T above 599 °C). UTBI describes a measured impact energy value in a Charpy V test, when the steel is annealed in the critical area for upper temper brittleness T = 450 - 599 °C.
TBI is defined on the basis of the tensile strength of the steel, from the direction of the impact test bar in relation to the direction of rolling, the measuring temperature of the impact test and the composition of the alloying elements according to the following equation: $TBI (temper brittleness index) = 190 - 0.121 Rm (MPa) - 0.516 direction (°) + 0.944 Test temperature (°C) - 87.3 Si - 39.1 Mn + 3335 Nb + 2054 V - 16.0 Ni - 21618 Nb * V,$
in which

* Rm is the tensile strength of the sample (MPa)
* Direction is the measurement direction for impact toughness in relation to the direction of rolling:
- direction = 0, if the measurement direction is longitudinal (impact test specimen longitudinal to rolling direction)
- direction = 90, if the measurement direction is transversal (impact test specimen transversal to rolling direction)
* Test temperature is the testing temperature of the Charpy V test (°C)

UTBI is defined on the basis of the tensile strength of the steel, from the direction of the impact test bar in relation to the direction of rolling, the measuring temperature of the impact test and the composition of the alloying elements according to the following equation: $UTBI (upper temper brittleness index) = 458 - 0.427 direction (°) - 0.254 Rm (MPa) + 1.06 Test temperature (°C) - 37.9 Si - 77.1 Mn + 1749 Nb + 691 V - 27261 Nb * V$
in which

The higher the UTBI, the better the steel resists upper temper brittleness retaining good impact toughness, even if the steel were to be annealed or it were to cool down slowly in the temperature range of 450 - 599 °C.
The values of both TBI and UTBI are dependent on temperature such that, as the testing temperature rises, the index value also rises.
According to the TBI, which describes achievable impact toughness after tempering treatment (at temperature T below 450°C or T above 599 °C) detrimental alloying elements for tempering steel are Si, Mn and Ni, but surprisingly the effects of Nb and V are the opposite. Thus, to achieve the objects of the invention, the composition of a hot-rolled steel according to the invention is limited on the part of these alloying elements to the limits presented earlier.
According to the invention, preferably, the TBI index describing impact toughness is for the longitudinal impact test bar at least 120, as defined at a temperature of -40 °C.
According to the invention, preferably, the UTBI index describing impact toughness is for the longitudinal impact test bar at least 100, as defined at a temperature of -40 °C.
The behaviour of UTBI differs from TBI mainly in that the multipliers of the factors are different, but the alloying elements effect in the same direction, so according to the invention it is possible to optimise the steel such that the values of both indexes UTBI and TBI are high, wherein, in accordance with the invention, steel can be produced with such a composition that it retains its impact toughness in over a wide tempering temperature range as well as in a upper temper brittleness range. An example of this is in table 6.
The correlation between TBI and the impact toughness of the test results can be seen below in Fig. 10.
The following Figs. 11 and 12 show the TBI value of different test steels as a function of impact toughness measuring temperature, impact toughness as measured both longitudinally (Fig. 11) and transversally (Fig. 12) in relation to the direction of rolling. The four uppermost examples (steels I, L, F and H of table 1) are steels according to the invention. The very uppermost two examples (steels I and L of table 1) are steels according to the preferred embodiment of the invention.
From Figs. 11 and 12, it can be clearly observed, how the composition of a steel according to the invention (steels I, L, F and H of table 1), in particular the compositions of steels according to the preferred embodiments of the invention (steels I and L of table 1) achieve considerably better impact toughness properties, as measured in both longitudinal and transversal directions, than the comparison steels (steels B, C oftable 1).
The conventional furnace type of table 6 (conventional) describes a manner, in which the steel is tempered in the conventional manner one sheet at a time in a furnace, wherein the sheet cools down slowly. Furnace type (Bell-type) means a furnace, in which the steel is annealed as a coil, where the temperature falls slowly, particularly the core of the steel coil cools down slowly.
As an example of the robustness of a hot-rolled steel according to the invention in relation to tempering are example steels K and L (see table 1), which are exceptionally close to each other in composition, in the composition area of a steel according to the invention, and achieve exceptionally good mechanical values regardless of whether tempering is conducted using an conventional or a bell furnace type of tempering furnace. Additionally, the composition achieves uniform mechanical properties and good impact toughness regardless of at how high a temperature tempering treatment is conducted, of which in table 6 is example steel L in comparison to example steel K.
Additionally, from table 6 can be observed that steels B and C belonging outside the invention become significantly brittle in tempering treatment in a bell furnace.

From table 6 is then observed that a hot-rolled steel according to the invention can be successfully tempered by highly varied means. Tempering temperature and furnace type can be selected with surprising freedom and the final result is nonetheless surprisingly good. The steel is thus exceptionally easy to produce, i.e. robust on the part of production, which facilitates production in many ways.

Table 6. The effect of tempering treatment using a bell furnace and an conventional furnace type

Heat	Direction	Furnace	t(h)	T(°C)	Rp0.2 (MPa)	Rm (MPa)	A5 (%)	CV (-40°C, J/cm2)	TBI (-40 °C)	UTBI (-40°C)
B	Longitud.	Bell-type	24*	450	1031	1094	14.2	55	-	55
C	Longitud.	Bell-type	24*	550	1021	1098	14.4	40	-	67
K	Longitud.	Conventional	24	570	1084	1089	15.5	106	-	117
K	Longitud.	Bell-type	24*	570	1081	1135	13.6	114	-	117
L	Longitud.	Conventional	0.5	650	1081	1081	13.3	159	131	-

* = Holding time at a constant temperature (does not contain heating and cooling)

The invention is illustrated above by means of examples and preferred embodiments.
The invention in its details can be implemented in many ways within the scope of the accompanying claims.

Claims

A method for producing a hot-rolled steel product, characterized in that
it is arranged a steel billet, whose composition as a percentage by weight is C 0.075 - 0.12% Si 0.1 - 0.8% Mn 0.8 - 1.7% Al 0.015 - 0.08% P less than 0.012% S less than 0.005% Cr 0.2 - 1.3% Mo 0.15 - 0.80% Ti 0.01 - 0.05% B 0.0005 - 0.003% V 0.02 - 0.10% Nb less than 0.3% Ni less than 1% Cu less than 0.5%

the rest being Fe and unavoidable impurities,
in which method the steel billet having said composition is
heated (1) to the austenitizing temperature of 1200 - 1350 °C, and
hot-rolled (2, 3) to the desired thickness such that the rolling temperature of the billet at the last pass is 760 - 960 °C, and
direct quenched (4) after the last pass conducted using one-step cooling at a cooling rate of 30 - 150 °C /s to a temperature of 300 °C at the most, which direct quenching is conducted at the latest 15 s after the last hot-rolling pass,
and in that the hot-rolled steel product is a strip steel, which, after being direct quenched (4), is coiled (5) and subsequently temper annealed (6) at a temperature of 200 - 700 °C for 24 hours at the most, and
in that the austenite of the hot-rolled steel billet is flattened before the direct quenching (4) so that average flattening ratio of the grain of the microstructure of the steel is greater than 2.
The method according to claim 1 for producing a hot-rolled steel product, characterized in that the steel is temper annealed (6) at a temperature of 450 - 599 °C, of 200 - 449 °C, or of 600 - 650 °C.
The method according to claims 1 or 2, characterized in that, in the method, it is arranged a steel billet, whose V content as a percentage by weight is 0.04 - 0.10% and Nb content as a percentage by weight is 0.005%.
The method according to any one of claims 1 - 3, characterized in that, in the method, it is arranged a steel billet, whose V content as a percentage by weight is 0.04 - 0.10% and Nb content as a percentage by weight is 0.008 - 0.03%.
The method according to any one of claims 1 - 4, characterized in that, in the method, it is arranged a steel billet, whose V content as a percentage by weight is 0.02 - 0.03% and Nb content as a percentage by weight is 0.008 - 0.03%.
The method according to any one of claims 1 - 5, characterized in that, in the method, it is arranged a steel billet, whose Mo content as a percentage by weight is 0.30 - 0.80%.
The method according to any one of claims 1 - 6, characterized in that, in the method, it is arranged a steel billet, whose Ni content as a percentage by weight is less than 0.1%, more preferably less than 0.05%.
A hot-rolled steel whose composition as a percentage by weight is C 0.075 - 0.12% Si 0.1 - 0.8% Mn 0.8 - 1.7% Al 0.015 - 0.08% P less than 0.012% S less than 0.005% Cr 0.2 - 1.3% Mo 0.15 - 0.80% Ti 0.01 - 0.05% B 0.0005 - 0.003% V 0.02 - 0.10% Nb less than 0.3% Ni less than 1% Cu less than 0.5%

the rest being iron and unavoidable impurities, characterized
in that the hot-rolled steel is a hot-rolled strip steel having tempering martensitic microstructure and that the yield strength of the steel is at least 890MPa and Charpy V impact toughness as measured longitudinally in relation to the direction of rolling at a temperature of -20°C (preferably at a temperature of -40°C) is at least 37 J/cm²,
in that the average flattening ratio of the grain of the microstructure of the steel is greater than 2, and in that the hot-rolled strip steel has a thickness of 2 to 10 mm.
The hot-rolled steel according to claim 8, characterized in that the index TBI (temper brittleness index) describing the impact toughness of the steel as defined at a temperature of -40 °C for longitudinal impact test bar is at least 120 when calculated according to the following equation: $TBI = 190 - 0.121 Rm (MPa) - 0.516 direction (°) + 0.944 Test temperature (°C) - 87.3 Si - 39.1 Mn + 3335 Nb + 2054 V - 16.0 Ni - 21618 Nb * V,$
in which
* Rm is the tensile strength of the sample (MPa)

* Direction is the measurement direction for impact toughness in relation to the direction of rolling:
direction = 0, if the measurement direction is longitudinal (impact test specimen longitudinal to rolling direction)

direction = 90, if the measurement direction is transversal (impact test specimen transversal to rolling direction)

* Test temperature is the testing temperature of the Charpy V test (°C).
The hot-rolled steel according to claims 8 or 9, characterized in that the index UTBI (upper temper brittleness index) describing the impact toughness of the steel as defined at a temperature of -40 °C for longitudinal impact test bar is at least 100 when calculated according to the following equation: $UTBI = 458 - 0.427 direction (°) - 0.254 Rm (MPa) + 1.06 Test temperature (°C) - 37.9 Si - 77.1 Mn + 1749 Nb + 691 V - 27261 Nb * V$
in which
* Rm is the tensile strength of the sample (MPa)

* Direction is the measurement direction for impact toughness in relation to the direction of rolling:
direction = 0, if the measurement direction is longitudinal (impact test specimen longitudinal to rolling direction)

direction = 90, if the measurement direction is transversal (impact test specimen transversal to rolling direction)

* Test temperature is the testing temperature of the Charpy V test (°C).
The hot-rolled steel according to any one of claims 8 - 10, characterized in that the V content of the steel as a percentage by weight is 0.04 - 0.10% and Nb content as a percentage by weight is less than 0.005%.
The hot-rolled steel according to any one of claims 8 - 11, characterized in that the V content of the steel as a percentage by weight is 0.04 - 0.10% and Nb content as a percentage by weight is 0.008 - 0.03%.
The hot-rolled steel according to any one of claims 8 - 12, characterized in that the V content of the steel as a percentage by weight is 0.02 - 0.03% and Nb content as a percentage by weight is 0.008 - 0.03%.
The hot-rolled steel according to any one of claims 8 - 13, characterized in that the Mo content of the steel as a percentage by weight is 0.30 - 0.80%.
The hot-rolled steel according to any one of claims 8 - 14, characterized in that the B content of the steel as a percentage by weight is 0.0008 - 0.002%.