U.S. patent application number 15/128970 was filed with the patent office on 2018-12-13 for steel plate with yield strength at 890mpa level and low welding crack sensitivity and manufacturing method therefor.
The applicant listed for this patent is Baoshan Iron & Steel Co.,Ltd.. Invention is credited to Pengjian WANG, Liandeng YAO, Sixin ZHAO.
Application Number | 20180355452 15/128970 |
Document ID | / |
Family ID | 50989971 |
Filed Date | 2018-12-13 |
United States Patent
Application |
20180355452 |
Kind Code |
A1 |
YAO; Liandeng ; et
al. |
December 13, 2018 |
STEEL PLATE WITH YIELD STRENGTH AT 890MPA LEVEL AND LOW WELDING
CRACK SENSITIVITY AND MANUFACTURING METHOD THEREFOR
Abstract
A steel plate with yield strength at an 890 MPa level and low
welding crack sensitivity and a manufacturing method therefor.
Weight percentages of components thereof are: C: 0.06-0.13 wt. %,
Si: 0.05-0.70 wt. %, Mn: 1.20-2.30 wt. %, Mo: 0-0.25 wt. %, Nb:
0.03-0.11 wt. %, Ti: 0.002-0.050 wt. %, Al: 0.02-0.15 wt. %, and B:
0-0.0020 wt. %, where 2Si+3Mn+4Mo.ltoreq.8.5, and others being Fe
and inevitable impurities. The use of controlled thermo-mechanical
rolling and cooling technologies is beneficial to improvement of
steel plate strength, plasticity and toughness. The yield strength
of the steel plate is greater than 890 MPa, the tensile strength is
greater than 950 MPa, the charpy impact work Akv (-20) is greater
than or equal to 120 J, and the welding crack sensitivity indicator
Pcm is less than or equal to 0.25%.
Inventors: |
YAO; Liandeng; (Shanghai,
CN) ; ZHAO; Sixin; (Shanghai, CN) ; WANG;
Pengjian; (Shanghai, CN) |
|
Applicant: |
Name |
City |
State |
Country |
Type |
Baoshan Iron & Steel Co.,Ltd. |
Shanghai |
|
CN |
|
|
Family ID: |
50989971 |
Appl. No.: |
15/128970 |
Filed: |
January 15, 2015 |
PCT Filed: |
January 15, 2015 |
PCT NO: |
PCT/CN2015/070729 |
371 Date: |
September 24, 2016 |
Current U.S.
Class: |
1/1 |
Current CPC
Class: |
C21D 9/46 20130101; C21D
8/0205 20130101; C22C 38/02 20130101; C22C 38/28 20130101; C22C
38/38 20130101; C22C 38/32 20130101; C21D 6/005 20130101; C22C
38/22 20130101; C21D 6/008 20130101; C21D 8/0226 20130101; C22C
38/04 20130101; C22C 38/26 20130101; C22C 38/06 20130101; C21D
8/0231 20130101; C22C 38/14 20130101; C21D 6/002 20130101 |
International
Class: |
C21D 9/46 20060101
C21D009/46; C21D 8/02 20060101 C21D008/02; C21D 6/00 20060101
C21D006/00; C22C 38/32 20060101 C22C038/32; C22C 38/28 20060101
C22C038/28; C22C 38/22 20060101 C22C038/22; C22C 38/38 20060101
C22C038/38; C22C 38/04 20060101 C22C038/04; C22C 38/02 20060101
C22C038/02; C22C 38/06 20060101 C22C038/06; C22C 38/26 20060101
C22C038/26 |
Foreign Application Data
Date |
Code |
Application Number |
Mar 25, 2014 |
CN |
201410114779.X |
Claims
1. A steel plate with a yield strength at an 890 Mpa level and low
welding crack sensitivity, wherein the steel plate has the
following components in weight percentage: C of 0.06-0.13 wt. %, Si
of 0.05-0.70 wt. %, Mn of 1.20-2.30 wt. %, Mo of 0-0.25 wt. %, Nb
of 0.03-0.11 wt. %, Ti of 0.002-0.050 wt. %, Al of 0.02-0.15 wt. %,
and B of 0-0.0020 wt. %, with 2Si+3Mn+4Mo.ltoreq.8.5, the balance
being Fe and inevitable impurities; and the steel plate meets the
welding crack sensitivity index Pcm.ltoreq.0.25%.
2. A method for manufacturing a steel plate with a yield strength
at 890 Mpa level and low welding crack sensitivity, comprising the
steps of: 1) smelting and casting the following chemical components
were smelt and casted to a continuous casting billet or steel ingot
of a thickness not less than 4 times of the thickness of the
finished steel plate; the continuous casting billet or steel ingot
has the following chemical components in weight percentage: C of
0.06-0.13 wt. %, Si of 0.05-0.70 wt. %, Mn of 1.20-2.30 wt. %, Mo
of 0-0.25 wt. %, Nb of 0.03-0.11 wt. %, Ti of 0.002-0.050 wt. %, Al
of 0.02-0.15 wt. %, and B of 0-0.0020 wt. %, with
2Si+3Mn+4Mo.ltoreq.8.5, the balance being Fe and inevitable
impurities; and the steel plate meets the welding crack sensitivity
index Pcm.ltoreq.0.25%; 2) heating and rolling the heating
temperature is 1050-1180.degree. C., and the holding time is 120 to
180 minutes; the rolling is divided into a first stage of rolling
and a second stage of rolling; during the first stage of rolling,
the start rolling temperature is 1050-1150.degree. C., and when the
thickness of the rolled piece reached 2-3 times of the thickness of
the finished steel plate, it is stayed in the roller bed until the
temperature reached 800-860.degree. C.; during the second stage of
rolling, the pass deformation rate is 10-28%, and the finish
rolling temperature is 780-840.degree. C.; 3) Cooling the steel
plate is cooled to 220-350.degree. C. at a speed of 15-30.degree.
C./S, and air cooled after being out of water.
3. The method for manufacturing the steel plate with a yield
strength at 890 Mpa level and low welding crack sensitivity of
claim 2, characterized in that, in step 3), the air cooling is
cooling in packed formation or in a cold bed.
Description
FIELD OF THE INVENTION
[0001] The present invention relates to a steel plate with a high
strength and low welding crack sensitivity, and in particular, the
present invention relates to a steel plate with a yield strength at
an 890 Mpa level and low welding crack sensitivity and a method for
manufacturing the same.
BACKGROUND
[0002] The steel for high strength mechanical equipment and
engineering construction requires a relatively high strength and an
excellent toughness, wherein contribution to the strength from
various factors can be expressed by the following formula:
.sigma.=.sigma..sub.f+.sigma..sub.p+.sigma..sub.sl+.sigma..sub.d
[0003] wherein .sigma..sub.f is fine grain strengthening,
.sigma..sub.p is precipitation strengthening, .sigma..sub.sl is
solid solution strengthening, and .sigma..sub.d is dislocation
strengthening. The thermo-mechanical treatment of the steel plate
is usually done by a controlled rolling and controlled cooling
process (TMCP). The refinement of the microstructures or the
formation of the high strength structures such as ultrafine bainite
can be realized by controlling the deformation rate and cooling
rate, thus improving the yield strength of the steel.
[0004] Currently, the composition of the low-carbon and
high-strength steel produced using TMCP is mainly
Mn--Ni--Nb--Mo--Ti and Si--Mn--Cr--Mo--Ni--Cu--Nb--Ti--Al--B
systems.
[0005] For example, the chemical composition of a low-alloy and
high-strength steel produced by the TMCP process in two temperature
stages disclosed in the international publication no. WO 99/05335
is as follows (wt. %): C: 0.05-0.10%, Mn: 1.7-2.1%, Ni: 0.2-1.0%,
Mo: 0.25-0.6 Mo %, Nb: 0.01-0.10%, Ti: 0.005-0.03%,
P.ltoreq.0.015%, and S.ltoreq.0.003%.
[0006] Also for example, the chemical composition of a superlow
carbon bainitic steel disclosed in the Chinese patent publication
no. 1521285 is as follows (wt. %): C: 0.01-0.05%, Si: 0.05-0.55%,
Mn: 1.0-2.2%, Ni: 0.0-1.0%, Mo: 0.0-0.5%, Cr: 0.0-0.7%, Cu:
0.0-1.8%, Nb: 0.015-0.070%, Ti: 0.005-0.03%, B: 0.0005-0.005%, and
Al: 0.015-0.07%.
[0007] The alloying element designs of the above two types of the
steels disclosed are an Mn--Ni--Nb--Mo--Ti and an
Si--Mn--Cr--Mo--Ni--Cu--Nb--Ti--Al--B system respectively; since Mo
and Ni are both precious metals, the production costs of such steel
plates are relatively high from the analysis of the type and the
total amount of the alloying elements added.
SUMMARY OF THE INVENTION
[0008] An object of the present invention is to provide a steel
plate with a yield strength at an 890 Mpa level and low welding
crack sensitivity and a method for manufacturing the same, using
the type of steel of an Si--Mn--Nb--Mo--V--Ti--Al--B system, by the
controlled thermo-mechanical rolling and cooling technologies,
without tempering, and the steel plate has a welding crack
sensitivity index Pcm.ltoreq.0.25%, a yield strength of greater
than 890 MPa, a tensile strength of greater than 950 MPa, a Charpy
impact energy Akv (-20.degree. C.).gtoreq.120 J, a plate thickness
of up to 60 mm, has a good low-temperature toughness and
weldability, and is a low-carbon superfine bainite lath steel plate
with low welding crack sensitivity.
[0009] To achieve the above-mentioned object, the technical
solution of the present invention consists in:
[0010] A steel plate with a yield strength at an 890 Mpa level and
low welding crack sensitivity, wherein the steel plate has the
following components in weight percentage: C of 0.06-0.13 wt. %, Si
of 0.05-0.70 wt. %, Mn of 1.20-2.30 wt. %, Mo of 0-0.25 wt. %, Nb
of 0.03-0.11 wt. %, Ti of 0.002-0.050 wt. %, Al of 0.02-0.15 wt. %,
and B of 0-0.0020 wt. %, with 2Si+3Mn+4Mo.ltoreq.8.5, the balance
being Fe and inevitable impurities; and the steel plate meets the
welding crack sensitivity index Pcm.ltoreq.0.25%.
[0011] In the composition design of the present invention:
[0012] C: C can enlarge an austenitic area, and carbon in a
supersaturated ferrite structure formed in the quenching can
increase the strength. C has an adverse impact on welding
performance. The higher the content of C, the poorer the welding
performance; for a bainitic steel produced using a TMCP process,
the lower the content of C, the better the toughness, and a lower
carbon content can results in a high toughness steel plate having a
greater thickness; therefore, the content of C in the present
invention is controlled at 0.06 to 0.13%.
[0013] Si: Si cannot be formed into a carbide in the steel, but
exists in the bainite ferrite or austenite in the form of a solid
solution. It can improve the strength of the bainite austenite or
ferrite in the steel. The solid solution strengthening effect of Si
is stronger than that of Mn, Nb, Cr, W, Mo and V. Si can reduce the
diffusion rate of carbon in the austenite, and makes the ferrite
CCT curve and pearlite C curve shift to the right, thus
facilitating the formation of a bainite structure in a continuous
cooling process. In the inventive steel, no more than 0.70% of Si
is added, which facilitates to improve the matching relationship
between the strength and toughness of the steel.
[0014] Mn: Mn and Fe can form a solid solution, which improves the
strength and hardness of the bainite ferrite and austenite in the
steel. Mn can enlarge the austenitic area in the iron-carbon
equilibrium phase diagram, so that the ability of the steel to form
a stable austenite structure is second only to that of Ni, which
strongly increases the hardenability of the steel. When the Mn
content is relatively high, it has the tendency to grain coarsening
of the steel. In the present invention, 1.20-2.30% of Mn is added,
and the speed of the ferrite and pearlite transform is slowed,
which is beneficial for the formation of the refined bainite
structure, and impart the steel with a certain strength.
[0015] Mo and Cr: Mo and Cr are ferritizing elements, which reduce
the austenitic area. Mo and Cr are in a solid solution in austenite
and ferrite to increase the strength, improve the hardenability of
the steel and prevent the temper brittleness. Mo is a very
expensive element, and the present invention does not require the
tempering treatment; in the present invention, only no more than
0.25% of Mo and no more than 0.20% of Cr are added to achieve the
purpose of reducing the cost.
[0016] Nb: in the present invention, a relatively high amount of Nb
is added in order to, on the one hand, achieve the purpose of
refining crystal grains and increasing the thickness of the steel
plate, and on the other hand to increase the non-recrystallization
temperature of the steel, which facilitates the use of a relatively
high finish rolling temperature in the rolling process, thus
accelerating the rolling speed and increasing the production
efficiency. In addition, since the grain refining effect is
strengthened, a thicker steel plate can be produced. In the present
invention, 0.03-0.10 wt. % of Nb is added to give consideration to
the solid solution strengthening effect and the fine grain
strengthening effect of Nb.
[0017] Ti: Ti is a ferritizing element, which reduces the
austenitic area significantly. The carbide of Ti, i.e. TiC, is
relatively stable, and can inhibit the growth of the crystal
grains. Ti, solid solved in austenite, is favourable to improve the
hardenability of the steel. Ti can reduce the first type of temper
brittleness at 250-400.degree. C.; however, the present invention
does not require the tempering, so the addition amount of Ti can be
reduced. In the present invention, an amount of 0-0.050 wt. % is
added, which forms fine carbonitride to precipitate out, thus
refining the Bainite laths.
[0018] Al: Al can increase the driving force of the phase change in
the transition from austenite to ferrite and is an element which
can intensively reduce the phase circle of the austenite. Al
interacts with N in the steel to form fine and diffusive AlN, which
precipitates out and can inhibit the growth of the crystal grains,
thus achieving the purpose of refining crystal grains and improving
the low temperature toughness of the steel. A too high content of
Al will have an adverse impact on the hardenability and welding
property of the steel. In the present invention, no more than 0.15%
of Al is added to refine the crystal grains, improving the
toughness and ensuring the welding property of the steel plate.
[0019] B: B can dramatically increase the hardenability of the
steel, in the present invention, 0-0.002% of B is added so that one
can relatively easily obtain a high strength bainite structure from
steel under certain cooling conditions.
[0020] The contents of the three elements, Si, Mn and Mo, should
comply with the following relationship: 2Si+3Mn+4Mo.ltoreq.8.5, to
meet that the steel plate of the present invention has a good
welding property. Specifically, it can be ensure that the steel
plate having a thickness of 60 mm or less has no cracks upon
welding at relatively low preheating temperature (normal
temperature to 50.degree. C.) conditions.
[0021] The steel plate having a maximum thickness of 60 mm is
produced using the chemical composition designed in the present
invention and by reasonably using the action of various alloying
elements.
[0022] The welding crack sensitivity index Pcm of the steel plate
with low welding crack sensitivity can be determined according to
the following formula:
Pcm=C+Si/30+Ni/60+(Mn+Cr+Cu)/20+Mo/15+V/10+5B
[0023] The welding crack sensitivity index Pcm is an indicator for
judging the weld cold cracking inclination of the steel, wherein
the smaller the Pcm, the better the weldability, and conversely,
the worse the weldability. Good weldability means that the
occurrence of weld cracking is not easy during welding; in
contrast, cracks easily occur in the steel having poor weldability;
in order to prevent cracking, steel is preheated before welding;
the better the weldability, the lower the preheating temperature
required, inversely, a higher preheating temperature is required.
According to the stipulations of the Chinese ferrous metallurgy
industry standards YB/T 4137-2005, a Pcm value for the type of
steel of trademark Q800CF should be lower than 0.28%. The superfine
bainite lath steel plate with a high strength and low welding crack
sensitivity involved in the present invention has a welding crack
sensitivity of lower than 0.20%, and has an excellent welding
property.
[0024] A method for manufacturing a steel plate with a yield
strength at 890 Mpa level and low welding crack sensitivity of the
present invention comprises the steps of:
[0025] 1) Smelting and Casting
[0026] The following chemical components were smelt and casted to a
continuous casting billet or steel ingot of a thickness not less
than 4 times of the thickness of the finished steel plate; wherein
the steel plate has the following components in weight percentage:
C of 0.06-0.13 wt. %, Si of 0.05-0.70 wt. %, Mn of 1.20-2.30 wt. %,
Mo of 0-0.25 wt. %, Nb of 0.03-0.11 wt. %, Ti of 0.002-0.050 wt. %,
Al of 0.02-0.15 wt. %, and B of 0-0.0020 wt. %, with
2Si+3Mn+4Mo.ltoreq.8.5, the balance being Fe and inevitable
impurities; and the steel plate meets the welding crack sensitivity
index Pcm.ltoreq.0.25%;
[0027] 2) Heating and Rolling
[0028] The heating temperature is 1050-1180.degree. C., and the
holding time is 120 to 180 minutes;
[0029] the rolling is divided into a first stage of rolling and a
second stage of rolling;
[0030] during the first stage of rolling, the start rolling
temperature is 1050-1150.degree. C., and when the thickness of the
rolled piece reached 2-3 times of the thickness of the finished
steel plate, it is stayed in the roller bed until the temperature
reached 800-860.degree. C.;
[0031] during the second stage of rolling, the pass deformation
rate is 10-28%, and the finish rolling temperature is
780-840.degree. C.;
[0032] 3) Cooling
[0033] The steel plate is cooled to 220-350.degree. C. at a speed
of 15-30.degree. C./S, and air cooled after being out of water.
[0034] further, in step 3), the air cooling is cooling in packed
formation or in a cold bed.
[0035] In the manufacturing method of the present invention:
[0036] (1) Rolling Process
[0037] When the thickness of the rolled piece reached 2-3 times of
the thickness of the finished steel plate, it is stayed in the
roller bed until the temperature reached 800-860.degree. C. For the
steel containing Nb, the non-recrystallizing temperature is about
950-1050.degree. C. It is firstly rolled at a relatively high
temperature, and there is a certain dislocation density in the
austenite. During the relaxation process of lowering the
temperature of the rolled billet to 800-860.degree. C., a recovery
and static recrystallization process inside the austenite crystal
grains occur, thus refining the austenite crystal grains. In the
relaxation process, individual precipitation and complex
precipitation of carbonitrides of Nb, V and Ti occur
simultaneously. The precipitated carbonitrides pin the dislocation
and subgrain boundary movements, reserves a lot of dislocation in
the austenite crystal grains, and provides a lot of nucleation
sites for the formation of bainite during the cooling process.
Rolling at 800-860.degree. C. greatly increases the dislocation
density in the austenite. The carbonitride precipitated at the
dislocation inhibits the coarsing of the deformed crystal grains.
Due to the precipitating effect induced by deformation, a
relatively large pass deformation rate will facilitate the
formation of finer and more diffusive precipitates. The
precipitates from high density dislocation and fine diffusion
provide high density of nucleation sites for bainite, and the
pining effect of the second phase particles on the bainite growth
interface inhibits the growth and coarsing of the bainite laths,
which is beneficial for both the strength and toughness of the
steel.
[0038] The finish rolling temperature is controlled in the low
temperature section of the non-recrystallization zone, and at the
same time, this temperature zone is close to the phase transmission
point Ar.sub.3, i.e. the finish rolling temperature is
780-840.degree. C., and finishing rolling within this temperature
range can increase the defects in the austenite by increasing the
deformation and inhibiting the recovery, thus providing higher
energy accumulation for the bainite phase change without bringing
about a too high load to the roller, being suitable for producing a
thick plate.
[0039] (2) Cooling Process
[0040] After the rolling is complete, the steel plate is sent to an
accelerated cooling device, and cooled to 450-550.degree. C. at a
rate of 15-30.degree. C./s. A faster cooling speed can avoid the
formation of ferrite and pearlite, and directly enters the bainite
transition area of the CCT curve. The phase change driving force of
the bainite can be expressed by:
.DELTA.G=.DELTA.G.sub.chem+.DELTA.G.sub.d
[0041] wherein .DELTA.G.sub.chem is a chemical driving force, and
.DELTA.G.sub.d is a strain storage energy resulting from defects. A
faster cooling speed results in the overcooling of the austenite,
increases the driving force of a chemical phase change, and
increase the driving force of the bainite nucleation when
considered by combining the strain storage energy .DELTA.G.sub.d
caused during the rolling process. Due to the high dislocation
density in the crystal grains, the nucleation sites of bainite
increase. Considered by combining both the thermodynamic and
dynamic factors, the bainite can nucleate at a very large speed. A
faster cooling speed enables the bainite transformation to be
completed quickly and inhibits the coarsing of the bainite ferrite
laths. Air cooling in packed formation at 450-550.degree. C. can
enable a more complete precipitation of the carbide of V in the
ferrite, thus enhancing the contribution of the precipitation
strengthening to the strength. Therefore, the matrix structure
composed mainly of the refined bainite can be obtained by the heat
treatment process of the present invention, so as to produce steel
plates having a higher strength and a good toughness.
[0042] The steel for high strength mechanical equipment and
engineering construction requires a relatively high strength and an
excellent toughness, wherein contribution to the strength from
various factors can be expressed by the following formula:
.sigma.=.sigma..sub.f+.sigma..sub.p+.sigma..sub.sl+.sigma..sub.d
[0043] wherein .sigma..sub.f is fine grain strengthening,
.sigma..sub.p is precipitation strengthening, .sigma..sub.sl is
solid solution strengthening, and .sigma..sub.d is dislocation
strengthening. The thermo-mechanical treatment of the steel plate
is usually done by a controlled rolling and controlled cooling
process (TMCP). The refinement of the microstructures or the
formation of the high strength structures such as ultrafine bainite
can be realized by controlling the deformation rate and cooling
rate, thus improving the yield strength of the steel. In the
composition of the present invention, a microalloy element Nb is
added, and during the heat treatment Nb may form a carbonitride,
which has a precipitation strengthening effect. Nb in a solid
solution in the matrix has a solid solution strengthening effect.
During the heat treatment, modified TMCP and Relaxation
Precipitation Controlling (RPC) technologies are used to form a
stable dislocation network, and diffusive and fine second phase
particles precipitate out at the dislocation and subgrain boundary;
the refinement of the bainite lath is achieved by promoting the
nucleation and inhibiting the growth, and a combined action of
dislocation strengthening, precipitation strengthening and fine
grain strengthening is formed, thus improving the strength and
roughness of the steel, its principle mechanism being as
follows:
[0044] the steel plate fully deforms in the recrystallization zone,
such that a high defect accumulation occurs in the deformed
austenite, thus greatly increasing the dislocation density in the
austenite. Recovery and recrystallization occurring during the
rolling refine the original austenite crystal grains. Dislocation
within the crystals will be re-arranged during the controlled
cooling relaxation after rolling and deforming. Since a hydrostatic
pressure field exits in the edge dislocation, interstitial atoms
such as B will enrich to the dislocation, grain boundary and
subgrain boundary, and reduce the dislocation mobility. The high
density dislocation resulting from the deformation will evolve
during the recovery to form a stable dislocation network. During
the relaxation, the microalloy elements such as Nb, V, Ti
precipitate out at the grain boundary, subgrain boundary and
dislocations in the form of carbonitrides of different
stoichiometric ratios such as (Nb,V,Ti).sub.x(C,N).sub.y. The
second phase particles, such as the precipitated carbonitrides, pin
the dislocations and subgrain boundary within the crystal grains
and stabilize the substructures, such as dislocation walls and the
like. After relaxation, the dislocation density of the steel is
further increased by rolling. After relaxation, when the deformed
austenite is accelerated cooled, the deformed austenite crystal
grains with dislocation and carbonitride precipitation
configuration at the beginning of the phase change is different
from the circumstance that after deformation, no relaxation occurs
and there are a lot of dislocations disorderly distributed.
Firstly, a subgrain boundary with a certain orientation difference
is a preferred nucleation site, and if a second phase, which has a
heterophasic interface with the matrix, precipitates out nearby,
this will be more advantageous for the new phase nucleation during
phase change. After relaxation, a lot of new phase crystal grains
will nucleate within the original austenite crystal grains.
Secondly, since after relaxation, a certain amount of dislocations
move to the subgrain boundary, the orientation difference between
the subgrains is increased to a certain extent. After the medium
temperature transformed product, such as bainite, nucleates at the
subgrain boundary, it is hindered by the front subgrain boundary
during the growth. When the bainite ferrite forms, its phase change
interface is daggled by the precipitated second phase carbonitride
particles, which inhibits its growth process. The TMCP RPC process
results in a high density dislocation network structure, and the
second phase precipitation particle points provide a lot of
potential nucleation sites for the nucleation of the bainite
ferrite The daggling effect of the second phase particles to the
moving interface and the evolved subgrain boundary have an
inhibiting effect on the growth of the bainite. The combined effect
of promoting the nucleation and inhibiting the growth in the
process refines the bainite ferrite laths of the final
structure.
[0045] With regard to high strength steels used in the mechanical
structure and engineering construction, no preheating or a little
preheating is required before welding, without the generation of
crack, which mainly solves the welding construction problem of
large steel structures. The only method to reduce Pcm is to reduce
the addition amounts of carbon and alloying elements; however, for
the high strength steel produced by a quenching tempering process,
reducing the addition amounts of carbon and alloying elements will
inevitably lead to the reduction of the steel strength, while the
use of the modified TMCP RPC process in the present invention can
remedy such a defect. The composition system used in the present
invention ensures that the steel plate has a high strength and a
low-temperature toughness, and at the same time a welding crack
sensitivity index Pcm.ltoreq.0.20%, and has an excellent welding
property.
[0046] Advantageous effects of the present invention lie in:
[0047] 1. The content of C is greatly reduced by reasonably
designing the chemical composition, replacing part of Mo with cheap
alloy elements such as Mn, replacing the precipitation
strengthening effect of Cu with the precipitation strengthening
effect of precipitated fine particles of carbonitride of Nb,
without adding noble elements such as Ni; the content of the alloy
element is low, the cost of the raw materials is relatively low,
the welding crack sensitivity is relatively low and no preheating
is required before welding.
[0048] 2. The steel plate of the present invention does not require
any additional thermal treatment, thus simplifying the
manufacturing procedure and reducing the manufacture cost of the
steel.
[0049] 3. Due to the reasonable composition and process design, the
process system is relatively loose in view of the implementing
effects and the steel plate can be produced stably in a medium, and
thick steel plate production line.
[0050] 4. The steel plate with low welding crack sensitivity of the
present invention has a yield strength of greater than 890 MPa, a
tensile strength of greater than 950 MPa, a Charpy impact energy
Akv (-20.degree. C.).gtoreq.100 J, and a plate thickness of up to
60 mm. The steel plate has a welding crack sensitivity index
Pcm.ltoreq.0.25%, and has an excellent welding property.
[0051] 5. A thick plate having a maximum thickness of 60 mm can be
produced by the present invention.
PREFERRED EMBODIMENTS OF THE INVENTION
[0052] The present invention is described by the following examples
in further detail. These examples are only intended to describe the
preferred embodiments of the invention, but not to limit the scope
of the invention in any way.
[0053] Table 1 is the chemical composition (wt. %) of the steel
plate of the examples of the present invention and the Pcm (%)
values. Table 2 is the mechanical property of the steel plate of
the examples of the present invention. Table 3 is the test (small
Tekken test) results of the welding property of the steel plate
with an 890 Mpa level and low welding crack sensitivity of Example
1 of the present invention.
EXAMPLE 1
[0054] The chemical components as shown in Table 2 are smelt in an
electric furnace or a converter and casted to a continuous casting
billet or steel ingot, which is then heated to 1110.degree. C. for
a holding time of 120 min and is subjected to a first stage of
rolling in a middle, and thick rolling mill, wherein the start
rolling temperature is 1050.degree. C., when the thickness of the
rolled piece is 60 mm, it is stayed in the roller bed until the
temperature reached 850.degree. C., and then a second stage of
rolling is performed, wherein the pass deformation rate in the
second stage of rolling is 15-28%, the finish rolling temperature
is 830.degree. C., and the thickness of the finished steel plate is
20 mm. After the rolling is complete, the steel plate is sent to an
accelerated cooling (ACC) device, and cooled to 300.degree. C. at a
rate of 30.degree. C./s, followed by cooling in packed formation or
in a cold bed after being out of water.
EXAMPLE 2
[0055] It is performed as in Example 1,wherein the heating
temperature is 1050.degree. C. and holding time is 240 min; the
start rolling temperature in the first stage of rolling is
1040.degree. C., and the thickness of the rolled piece is 90 mm;
the start rolling temperature in the second stage of rolling is
840.degree. C., the pass deformation rate is 15-20%, the finish
rolling temperature is 810.degree. C., and the thickness of the
finished steel plate is 30 mm; and the cooling rate of the steel
plate is 25.degree. C./S, and the final temperature is 350.degree.
C.
EXAMPLE 3
[0056] It is performed as in Example 1,wherein the heating
temperature is 1150.degree. C. and the holding time is 150 min; the
start rolling temperature in the first stage of rolling is
1080.degree. C., and the thickness of the rolled piece is 120 mm;
the start rolling temperature in the second stage of rolling is
830.degree. C., the pass deformation rate is 10-15%, the finish
rolling temperature is 820.degree. C., and the thickness of the
finished steel plate is 40 mm; and the cooling rate of the steel
plate is 20.degree. C./S, and the final temperature is 330.degree.
C.
EXAMPLE 4
[0057] It is performed as in Example 1,wherein the heating
temperature is 1120.degree. C. and the holding time is 180 min; the
start rolling temperature in the first stage of rolling is
1070.degree. C., and the thickness of the rolled piece is 150 mm;
the start rolling temperature in the second stage of rolling is
830.degree. C., the pass deformation rate is 10-20%, the finish
rolling temperature is 800.degree. C., and the thickness of the
finished steel plate is 50 mm; and the cooling rate of the steel
plate is 15.degree. C./S, and the final temperature is 285.degree.
C.
EXAMPLE 5
[0058] It is performed as in Example 1,wherein the heating
temperature is 1130.degree. C. and the holding time is 180 min; the
start rolling temperature in the first stage of rolling is
1080.degree. C., and the thickness of the rolled piece is 150 mm;
the start rolling temperature in the second stage of rolling is
840.degree. C., the pass deformation rate is 10-15%, the finish
rolling temperature is 810.degree. C., and the thickness of the
finished steel plate is 60 mm; and the cooling rate of the steel
plate is 15.degree. C./S, and the final temperature is 220.degree.
C.
EXAMPLE 6
[0059] It is performed as in Example 1,wherein the heating
temperature is 1120.degree. C. and the holding time is 180 min; the
start rolling temperature in the first stage of rolling is
1050.degree. C., and the thickness of the rolled piece is 120 mm;
the start rolling temperature in the second stage of rolling
820.degree. C., the pass deformation rate is 15-25%, the finish
rolling temperature is 780.degree. C., and the thickness of the
finished steel plate is 40 mm; and the cooling rate of the steel
plate is 20.degree. C./S, and the final temperature is 300.degree.
C.
TABLE-US-00001 TABLE 1 unit: weight percentage Examples C Si Mn Nb
Al Ti Cr Mo B Fe Pcm 1 0.09 0.35 1.80 0.070 0.02 0.015 0.16 0.25
0.0018 the 0.217 balance 2 0.06 0.70 2.25 0.045 0.06 0.020 0 0
0.0010 the 0.201 balance 3 0.08 0.40 2.06 0.085 0.04 0.050 0.20
0.10 0.0011 the 0.218 balance 4 0.13 0.55 1.20 0.110 0.15 0 0.16
0.25 0.0015 the 0.183 balance 5 0.06 0.05 1.45 0.065 0.07 0.020
0.12 0.20 0.0010 the 0.241 balance 6 0.10 0.15 1.90 0.095 0.09
0.008 0.15 0.22 0.0020 0.232
TABLE-US-00002 TABLE 2 -20.degree. C. Yield Tensile Longitudinal
Exam- strength strength Elongation impact energy ples MPa MPa % J 1
940 1050 16.0 189 216 204 965 1065 16.5 2 950 1060 15.9 208 190 209
975 1070 15.2 3 955 1058 15.0 195 202 212 960 1065 15.0 4 945 1154
16.1 191 208 206 940 1150 16.3 5 988 1169 15.0 207 201 224 995 1173
14.5 6 910 1067 17.3 202 210 210 915 1082 17.3
[0060] From tables 1 and 2, it can be seen that the Pcm of the
steel plate with a yield strength at an 890 Mpa level and low
welding crack sensitivity involved in the present invention is
.ltoreq.0.25%, the yield strength is larger than 890 MPa, the
tensile strength is larger than 950 MPa, the Charpy impact energy
Akv (-20.degree. C.) is .gtoreq.120 J, and the plate thickness can
be up to 60 mm, and the steel plate has an excellent
low-temperature toughness and weldability.
[0061] The steel plate of Example 1 of the present invention is
tested for the welding property (small Tekken test), under
conditions of room temperature and 50.degree. C., and no crack is
observed (see table 3), indicating that the type of steel of the
present invention has an excellent welding property, and generally
does not require preheating when welding.
TABLE-US-00003 TABLE 3 Surface Root section Environ- Test crack
crack crack mental temper- rate rate rate temper- Relative ature
Examples % % % ature humidity RT 1 0 0 0 22.degree. C. 60% 2 0 0 0
3 0 0 0 50.degree. C. 4 0 0 0 5 0 0 0
* * * * *