CN117512460B - Si-Mn-Cr-Mo-V-Ti-Nb multi-alloyed ultrahigh-strength wire rod and preparation method thereof - Google Patents
Si-Mn-Cr-Mo-V-Ti-Nb multi-alloyed ultrahigh-strength wire rod and preparation method thereof Download PDFInfo
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- 238000002360 preparation method Methods 0.000 title claims abstract description 19
- 229910000831 Steel Inorganic materials 0.000 claims abstract description 95
- 239000010959 steel Substances 0.000 claims abstract description 95
- 229910052720 vanadium Inorganic materials 0.000 claims abstract description 22
- 229910052758 niobium Inorganic materials 0.000 claims abstract description 14
- 229910052757 nitrogen Inorganic materials 0.000 claims abstract description 8
- 229910052748 manganese Inorganic materials 0.000 claims abstract description 7
- 229910052804 chromium Inorganic materials 0.000 claims abstract description 6
- 229910052750 molybdenum Inorganic materials 0.000 claims abstract description 6
- 229910052717 sulfur Inorganic materials 0.000 claims abstract description 5
- 229910052782 aluminium Inorganic materials 0.000 claims abstract description 4
- 239000012535 impurity Substances 0.000 claims abstract description 4
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- 238000010438 heat treatment Methods 0.000 claims description 71
- 238000001816 cooling Methods 0.000 claims description 52
- 238000005098 hot rolling Methods 0.000 claims description 42
- 229910001562 pearlite Inorganic materials 0.000 claims description 40
- 150000003839 salts Chemical class 0.000 claims description 37
- 229910001566 austenite Inorganic materials 0.000 claims description 25
- 229910001567 cementite Inorganic materials 0.000 claims description 23
- KSOKAHYVTMZFBJ-UHFFFAOYSA-N iron;methane Chemical compound C.[Fe].[Fe].[Fe] KSOKAHYVTMZFBJ-UHFFFAOYSA-N 0.000 claims description 23
- 229910045601 alloy Inorganic materials 0.000 claims description 21
- 239000000956 alloy Substances 0.000 claims description 21
- 238000005266 casting Methods 0.000 claims description 16
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- 238000004321 preservation Methods 0.000 claims description 11
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- 229910001182 Mo alloy Inorganic materials 0.000 claims description 4
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- 238000000034 method Methods 0.000 abstract description 41
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- 238000005246 galvanizing Methods 0.000 abstract description 14
- 229910052742 iron Inorganic materials 0.000 abstract description 8
- 230000009467 reduction Effects 0.000 abstract description 6
- 206010003549 asthenia Diseases 0.000 abstract description 3
- 230000035882 stress Effects 0.000 description 27
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- 238000010622 cold drawing Methods 0.000 description 13
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- 238000007670 refining Methods 0.000 description 13
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- 238000005728 strengthening Methods 0.000 description 9
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- 238000000265 homogenisation Methods 0.000 description 8
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- OKTJSMMVPCPJKN-UHFFFAOYSA-N Carbon Chemical compound [C] OKTJSMMVPCPJKN-UHFFFAOYSA-N 0.000 description 6
- SOIFLUNRINLCBN-UHFFFAOYSA-N ammonium thiocyanate Chemical compound [NH4+].[S-]C#N SOIFLUNRINLCBN-UHFFFAOYSA-N 0.000 description 6
- 230000002829 reductive effect Effects 0.000 description 6
- 238000003303 reheating Methods 0.000 description 6
- 239000000243 solution Substances 0.000 description 5
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- 238000011105 stabilization Methods 0.000 description 5
- UFHFLCQGNIYNRP-UHFFFAOYSA-N Hydrogen Chemical compound [H][H] UFHFLCQGNIYNRP-UHFFFAOYSA-N 0.000 description 4
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- 229910000734 martensite Inorganic materials 0.000 description 2
- 229910052759 nickel Inorganic materials 0.000 description 2
- 238000005554 pickling Methods 0.000 description 2
- 238000011084 recovery Methods 0.000 description 2
- 238000010079 rubber tapping Methods 0.000 description 2
- 238000005482 strain hardening Methods 0.000 description 2
- 229910000599 Cr alloy Inorganic materials 0.000 description 1
- 229910000677 High-carbon steel Inorganic materials 0.000 description 1
- 229910018643 Mn—Si Inorganic materials 0.000 description 1
- 229910000611 Zinc aluminium Inorganic materials 0.000 description 1
- 230000032683 aging Effects 0.000 description 1
- 238000005275 alloying Methods 0.000 description 1
- HXFVOUUOTHJFPX-UHFFFAOYSA-N alumane;zinc Chemical compound [AlH3].[Zn] HXFVOUUOTHJFPX-UHFFFAOYSA-N 0.000 description 1
- 239000000470 constituent Substances 0.000 description 1
- 238000010276 construction Methods 0.000 description 1
- 229910052802 copper Inorganic materials 0.000 description 1
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Classifications
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- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
-
- B—PERFORMING OPERATIONS; TRANSPORTING
- B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
- B21B—ROLLING OF METAL
- B21B1/00—Metal-rolling methods or mills for making semi-finished products of solid or profiled cross-section; Sequence of operations in milling trains; Layout of rolling-mill plant, e.g. grouping of stands; Succession of passes or of sectional pass alternations
- B21B1/16—Metal-rolling methods or mills for making semi-finished products of solid or profiled cross-section; Sequence of operations in milling trains; Layout of rolling-mill plant, e.g. grouping of stands; Succession of passes or of sectional pass alternations for rolling wire rods, bars, merchant bars, rounds wire or material of like small cross-section
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/0205—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips of ferrous alloys
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
- C21D8/0226—Hot rolling
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
- C21D8/0263—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment following hot rolling
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C33/00—Making ferrous alloys
- C22C33/04—Making ferrous alloys by melting
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
- C22C38/22—Ferrous alloys, e.g. steel alloys containing chromium with molybdenum or tungsten
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
- C22C38/24—Ferrous alloys, e.g. steel alloys containing chromium with vanadium
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
- C22C38/26—Ferrous alloys, e.g. steel alloys containing chromium with niobium or tantalum
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
- C22C38/28—Ferrous alloys, e.g. steel alloys containing chromium with titanium or zirconium
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D2211/00—Microstructure comprising significant phases
- C21D2211/003—Cementite
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D2211/00—Microstructure comprising significant phases
- C21D2211/009—Pearlite
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- Chemical & Material Sciences (AREA)
- Engineering & Computer Science (AREA)
- Mechanical Engineering (AREA)
- Materials Engineering (AREA)
- Metallurgy (AREA)
- Organic Chemistry (AREA)
- Physics & Mathematics (AREA)
- Thermal Sciences (AREA)
- Crystallography & Structural Chemistry (AREA)
- Heat Treatment Of Steel (AREA)
- Heat Treatment Of Strip Materials And Filament Materials (AREA)
Abstract
The invention relates to a Si-Mn-Cr-Mo-V-Ti-Nb multi-element alloyed ultrahigh-strength wire rod and a preparation method thereof, belongs to the technical field of prestressed steel wire rods, and solves the problems of insufficient tensile strength, uncontrollable loss of processing strength after drawing and insufficient stress corrosion resistance or torsion resistance in the working process of the wire rod mainly faced after the ultrahigh strength of prestressed steel in the prior art. The invention discloses a Si-Mn-Cr-Mo-V-Ti-Nb multi-element alloyed ultra-high strength wire rod, which comprises ：C：0.90%～1.00%,Si：0.70%～1.20%,Cr：0.20%～0.80%,Mn：0.10%～0.60%,Mo：0.05%～0.35%,V：0.07%～0.15%,Ti：0.010%～0.025%,Nb：≤0.03%,Al：≤0.04%,P：≤0.02%,S：≤0.005%,N：≤0.004%, mass percent of Fe and unavoidable impurities as the rest. The tensile strength of the wire rod reaches 1400-1680 MPa, the reduction of area is more than or equal to 28 percent, the strength loss of hot galvanizing or stabilizing treatment after drawing is small, the preparation process is simple, and the wire rod is suitable for preparing the downstream products such as the ultra-high strength prestressed steel wires, ropes and the like.
Description
Technical Field
The invention relates to the technical field of prestressed steel, in particular to a Si-Mn-Cr-Mo-V-Ti-Nb multi-element alloyed ultrahigh-strength wire rod and a preparation method thereof.
Background
The prestressed steel comprises prestressed steel strands, cables, steel wire ropes and the like which are widely applied to the fields related to national life such as buildings, bridges, engineering machinery, water conservancy and hydropower, energy and ocean engineering. The requirements of ultra-high bearing, light weight and easy construction are put forward for prestressed steel strands, cables and steel wire ropes, and the strength of the prestressed steel is upgraded to be necessarily selected.
The strength upgrade of prestressed steel brings three problems: (1) The basic strength of the wire rod is improved, and meanwhile, high plasticity is ensured, so that the processability and the work hardening strength of the cold drawn wire are ensured; (2) The loss amount of the work hardening strength of the cold drawn steel wire in the subsequent hot galvanizing or stabilizing treatment process is controllable, otherwise, the final strength of the prestressed wire, rope and rope is difficult to ensure; (3) After the prestress line, the rope and the rope are subjected to ultra-high strength, how to ensure the service performance, such as stress corrosion resistance, torsion resistance and the like.
The invention discloses a prestressed steel strand with 2100 MPa-level strength and a production process (202010216245.3) of a C-Si-Mn-Cr-Al-V component system, wherein the tensile strength of the C-Si-Mn-Cr-Al-V component system is 2000MPa, the tensile strength of a 2100 MPa-level prestressed steel strand is 1270-1291 MPa, and the production process is implemented by using a wire rod. The invention discloses a C-Si-Mn-Cr-Al-V component system tensile strength 2200 MPa-level prestressed steel strand and a production process thereof, wherein the strength of the wire rod is 1330-1343 MPa. The wire rods related to the two patents are produced by adopting a hot rolling process, and the special design of the alloy for controlling the strength loss, stress corrosion resistance, torsion and other service performances in the hot galvanizing or stabilizing process is not involved. The invention patent 'a 2300MPa prestress steel strand wire rod and a production method thereof', wherein the C-Si-Mn-Cr-Al basic components are adopted, V, cu and Ni are optional, other alloy designs except 0.01 to 0.05 percent of Cu and 0.01 to 0.05 percent of Ni are more conventional, the special design of the alloy for controlling strength loss, stress corrosion resistance, torsion and other service performances in the hot galvanizing or stabilizing process is not involved, offline salt bath treatment is adopted after hot rolling production of the wire rod, the tensile strength of the wire rod is 1450 to 1480MPa, and the tensile strength of the prestress steel strand is 2280 to 2370MPa.
In summary, the hot rolled wire rod or the ultra-high strength wire rod, which are key process products of the prestressed steel in the prior art, have one of the defects of insufficient tensile strength and uncontrollable strength loss in the post-drawing treatment, and the downstream prestressed steel product has insufficient stress corrosion resistance and torsion resistance despite the improvement of the tensile strength because the pre-design is not fully performed on the subsequent prestressed steel product at the stage of the process product.
Disclosure of Invention
In view of the above analysis, the embodiment of the invention aims to provide a Si-Mn-Cr-Mo-V-Ti-Nb multi-element alloyed ultra-high strength wire rod and a preparation method thereof, which are used for solving one of the defects of insufficient tensile strength and uncontrollable strength loss after drawing treatment of the existing wire rod, and are beneficial to improving the stress corrosion resistance and torsion resistance of a downstream prestressed steel product.
The invention discloses a Si-Mn-Cr-Mo-V-Ti-Nb multi-element alloyed ultra-high strength wire rod, which comprises ：C：0.90%～1.00%,Si：0.70%～1.20%,Cr：0.20%～0.80%,Mn：0.10%～0.60%,Mo：0.05%～0.35%,V：0.07%～0.15%, Ti：0.010%～0.025%,Nb：≤0.03%,Al：≤0.04%,P：≤0.02%,S：≤0.005%,N：≤0.004%, mass percent of Fe and unavoidable impurities as the rest.
Specifically, si, mn, cr, mo alloy components in the ultra-high strength wire rod satisfy the following mass percentages: si+Mn+Cr+Mo is 1.2-2.4 and Cr+Mo is 0.40-0.85.
Specifically, the diameter of the ultra-high strength wire rod is less than or equal to 15mm, the microstructure is complete lamellar pearlite, and no network cementite exists.
Specifically, the tensile strength of the ultra-high strength wire rod reaches 1400-1680 MPa, and the area shrinkage is more than or equal to 28%.
The invention also discloses a preparation method of the ultra-high strength wire rod, which comprises the following steps:
s1: batching according to alloy composition, and smelting and casting to obtain a casting blank;
s2: homogenizing the casting blank, namely heating to the austenite homogenizing temperature and preserving heat;
S3: continuously rolling the heat-preserving casting blank after descaling until reaching the target diameter to obtain a semi-finished product wire rod;
S4: and curling and post-processing the semi-finished wire rod to obtain the finished product of the ultra-high strength wire rod.
Specifically, the austenite homogenization temperature in the step S2 is 1180-1250 ℃, and the heat preservation time is 0.5-3 h.
Specifically, the heating temperature before hot rolling in the step S3 is 1180-1250 ℃, and the hot rolling finishing temperature is 890-950 ℃.
Specifically, the post-treatment mode in step S4 is one of air cooling, isothermal salt bath heat treatment or off-line isothermal heat treatment, air cooling, isothermal salt bath heat treatment.
Specifically, in the air cooling step, the air cooling speed is 1.5 ℃/s-3.5 ℃/s.
The invention also discloses Si-Mn-Cr-Mo-V-Ti-Nb multi-element alloyed ultra-high strength prestressed steel, which is prepared from the ultra-high strength wire rod or the ultra-high strength wire rod prepared by the preparation method.
Compared with the prior art, the invention has at least one of the following beneficial effects:
1. The ultra-high strength wire rod disclosed by the invention has high tensile strength and high reduction of area. In the wire rod, the thought of taking Mn or Mn-Si as a main alloy design is changed, the Si-Cr-Mn-Mo multi-element main alloy design is adopted, the pearlite transformation is not delayed obviously, the complete pearlite structure can be obtained at the air cooling speed of 1.5 ℃/s-3.5 ℃/s before 475 ℃ after hot rolling finish rolling, or the complete pearlite transformation can be completed within 5min through isothermal salt bath heat treatment at 475-550 ℃, and the lamellar spacing of pearlite is thinned obviously, so that the strength of the wire rod is improved; the V-Ti (-Nb) strong carbide forming element microalloy design is adopted, and the MX type second phase particles are separated out in a grading way to refine austenite grains, so that the size of pearlite clusters is further refined, the orientation diversity of lamellar layers is increased, the area shrinkage is improved, and the drawing performance is improved. The tensile strength of the ultra-high strength wire rod reaches 1400-1680 MPa, and the area shrinkage rate is more than or equal to 28%.
2. The ultra-high strength wire rod disclosed by the invention has small strength loss after being pulled out. Cr, mo, V, ti, nb in the wire rod alloy composition is used as a strong carbide forming element, and by inhibiting the recovery of a drawing deformation structure of the prestressed steel or precipitating MX type second phase reinforcement, the high-temperature stability of the structure is improved, and the strength loss of the prestressed steel wire caused by hot galvanizing and stabilizing treatment after the wire rod is drawn is reduced.
3. The stress corrosion resistance of the downstream prestressed steel product is improved through the pre-design of the wire rod alloy composition and the preparation process. The V, ti and Nb strong carbide forming elements have the following grading precipitation characteristics and are synergistically reinforced with the medium-strength carbide forming element Mo. Ti fixes N and is almost completely precipitated as TiN or Ti (C, N) refined austenite grains homogenized before hot rolling, and Ti (C, N) may contain a small amount of Nb or V. Nb is partially separated out in the hot rolling process, and NbC particles and solid-solution Nb act together to refine austenite grains, so that austenite deformation energy storage is improved, pearlite transformation driving force is improved, and meanwhile, pearlite clusters and lamellar spacing are refined; the incompletely precipitated Nb is continuously precipitated in the cooling process, so that the strength is further improved; nb precipitates almost completely during austenitizing reheating in an off-line isothermal heat treatment, refining austenite grains. V precipitates in a small amount in the hot rolling process, but precipitates VC alone or precipitates (Nb, V) C together with Nb when austenitizing and reheating in off-line isothermal heat treatment; in the air cooling or isothermal salt bath heat treatment process after hot rolling, V is separated out of VC alone or together with Nb (Nb, V) C to play a role in precipitation strengthening, cementite separation is participated in on the other hand, the interlayer spacing of pearlite sheets is thinned, and the rest V is in a solid solution state. After pulling, a small amount of MX-type second phase particles are continuously separated out in the galvanization or stabilization treatment process due to the dissolution of C atoms in cementite. The precipitation of Ti and Nb accounts for 95-100% of the total addition, and the precipitation of V accounts for 50-80% of the total addition. Mo participates in precipitation of MX type second phase particles such as TiN, ti (C, N), nbC, (Nb, V) C and VC, so that the total precipitation amount of the MX type second phase is increased, and coarsening of the MX type second phase particles is prevented. The above-classified MX-type second phase particles collectively act as stress corrosion resistance. The diameter of the prestressed steel subjected to cold drawing and stabilizing heat treatment is less than or equal to 7mm, the tensile strength is 2160-2460 MPa, and the stress corrosion resistance meets the following requirements: and carrying out a solution A (analyzing pure ammonium thiocyanate aqueous solution) according to the specification of GB/T21839, carrying out a stress corrosion test by loading stress with 80% of the actual maximum force, wherein the test group is not less than 5 groups, the minimum value of the stress corrosion breaking time of the prestressed steel is not less than 2 hours, and the median value is not less than 5 hours.
4. The raw materials are easy to obtain, the preparation process is relatively simple, and the method is suitable for large-scale production. The preparation method of the ultra-high strength plate wire/prestressed steel disclosed by the invention has the advantages that various raw materials and production equipment required by the preparation method of the ultra-high strength plate wire/prestressed steel are common materials/equipment on the market, the process flow is relatively simple, the process condition is easy to realize, the process flow is mature, and the preparation method is suitable for large-scale production and large-scale popularization.
In the invention, the technical schemes can be mutually combined to realize more preferable combination schemes. Additional features and advantages of the invention will be set forth in the description which follows, and in part will be obvious from the description, or may be learned by practice of the invention. The objectives and other advantages of the invention may be realized and attained by the structure particularly pointed out in the written description and drawings.
Drawings
The drawings are only for purposes of illustrating particular embodiments and are not to be construed as limiting the invention, like reference numerals being used to designate like parts throughout the drawings;
FIG. 1 is a flow chart of ultra-high strength wire rod preparation;
fig. 2 is a photograph of a microstructure of the ultra-high strength wire rod of example 1.
Detailed Description
The following detailed description of preferred embodiments of the application is made in connection with the accompanying drawings, which form a part hereof, and together with the description of the embodiments of the application, are used to explain the principles of the application and are not intended to limit the scope of the application.
The invention discloses a Si-Mn-Cr-Mo-V-Ti-Nb multi-element alloyed ultra-high strength wire rod, which comprises ：C：0.90%～1.00%,Si：0.70%～1.20%,Cr：0.20%～0.80%,Mn：0.10%～0.60%,Mo：0.05%～0.35%,V：0.07%～0.15%, Ti：0.010%～0.025%,Nb：≤0.03%,Al：≤0.04%,P：≤0.02%,S：≤0.005%,N：≤0.004%, mass percent of Fe and unavoidable impurities as the rest.
The following is a specific description of the action and the selection of the amounts of the components contained in the invention:
C: and the steel can effectively improve the tensile strength of the wire rod, and the second phase strengthening elements such as cementite and MX are added in the steel. In order to ensure that the tensile strength of the wire rod reaches more than 1400MPa, the C content is not less than 0.90 percent, and the wire rod is hypereutectoid steel, but the wire rod is subjected to proper air cooling or isothermal salt bath heat treatment after hot rolling to obtain complete pearlite without network cementite; however, the higher the C content is, the more difficult the control of C segregation and network cementite is increased, and too high C content can also lead to the reduction of wire rod plasticity and influence drawing. Therefore, the upper limit of the C content must be limited not to exceed 1.00%. The C content of the steel is 0.90-1.00%.
Si: the steel has excellent relaxation resistance by replacing solid solution strengthening elements stronger than Mn, cr and Mo, improves the eutectoid phase transition temperature, reduces the eutectoid carbon content, refines the lamellar spacing of pearlite formed by being lower than the eutectoid phase transition temperature, has more remarkable strengthening and refining effects, but has larger plastic damage when Si is excessively high, such as more than 1.2 percent, and aggravates decarburization in the off-line heat treatment process. Si is also an alloy element for improving the hardenability of steel, but with Si addition, the effect of deferring pearlite transformation of which is lower than Mn, cr and Mo and exceeds 0.70%, the steel wire rod of the invention can obtain complete pearlite structure under the condition of slower cooling (namely air cooling and air cooling) after hot rolling, and can also obtain complete pearlite structure under the condition of isothermal salt bath heat treatment in a short time (less than or equal to 5 min) without forming bainite and martensite structures. In addition, si also plays a role in deoxidization. By combining the above, the Si content of the steel of the invention is 0.70% -1.20%.
Cr: the medium-strength carbide forming element and one of the alloy cementite constituent elements can improve the eutectoid phase transition temperature, reduce the eutectoid carbon content, obviously refine the lamellar spacing of pearlite formed below the eutectoid phase transition temperature, and the addition of 0.1% Cr has obvious refining effect. Therefore, the lower limit of the Cr content of the steel of the present invention is controlled to be 0.2%. However, too high Cr content also significantly delays the pearlite transformation, which is detrimental to the achievement of a complete pearlite structure in the case of slow cooling (i.e. air cooling and air cooling) of high carbon wire rods after hot rolling and isothermal salt bath heat treatment conditions for a short time (5 min or less). When the content of Si, mn and Mo alloy in the steel is close to the lower limit, the highest adding amount of Cr is not more than 0.80 percent. By combining the above, the Cr content of the steel is 0.20-0.80%.
Mn: the austenite element is stabilized, the eutectoid phase transition temperature and the eutectoid carbon content are reduced, and the influence on the equilibrium phase transition pearlite lamellar spacing is small. Therefore, the Mn content is not higher than 0.6%. In addition, the Mn is more than or equal to 0.1 percent and plays a role in deoxidization. The invention utilizes the deoxidization effect of Mn and the effect of stabilizing austenite. By combining the above, the Mn content of the steel of the invention is 0.1% -0.60%.
Mo: the medium-strength carbonitride forming element, one of the alloying cementite and MX carbonitride forming elements, improves the eutectoid phase transition temperature, reduces the eutectoid carbon content, obviously refines the lamellar spacing of pearlite formed below the eutectoid phase transition temperature, has a stronger effect than Cr, and has obvious effect of refining the lamellar spacing of the pearlite when Mo is more than or equal to 0.05 percent. However, mo has a remarkable effect of improving hardenability and postpones pearlite transformation. When the Si, mn and Cr alloy contents of the steel are close to the lower limit, mo exceeding 0.35% can obviously prolong the pearlite transformation time of the wire rod, and is unfavorable for obtaining a complete pearlite structure by the manufacturing method of the steel. In addition, mo inhibits the recovery of the drawing deformation structure of the prestressed steel or participates in the strengthening of the precipitated MX type second phase, improves the high-temperature stability of the structure, and reduces the strength loss of the prestressed steel wire caused by hot galvanizing and stabilizing treatment after wire rod drawing. By combining the above, the Mo content of the steel is 0.05-0.35%.
V: the strong carbonitride forms elements, the solid solution quantity of the elements in high-carbon steel is much larger than that of Ti and Nb, and the elements can be combined with C to form finer MX-type VC at a lower phase transition temperature, so that the elements play a remarkable role in resisting stress corrosion, and the action mechanism mainly comprises two aspects: on one hand, atoms in the VC lattice are in a shortage position to provide a hydrogen trap, and on the other hand, VC and an iron matrix form a semi-coherent interface to provide the hydrogen trap. In addition, V can also enter cementite to replace part of Fe atoms, namely participate in pearlite transformation, and play a role in refining the lamellar spacing of pearlite. Therefore, the addition amount of V in the steel of the invention is controlled relatively high, namely 0.07% -0.15%. In the high-temperature homogenization stage before hot rolling, V may be involved in Ti (C, N) precipitation in a very small amount. Because the hot rolling finishing temperature is higher, V is precipitated in a small amount in the hot rolling stage, and VC is mainly precipitated singly or (Nb, V) C is precipitated together with Nb during austenitizing and reheating of offline isothermal heat treatment; in the air cooling or isothermal salt bath heat treatment process after hot rolling, V is separated out of VC alone or together with Nb (Nb, V) C to play a role in precipitation strengthening, cementite separation is participated in on the other hand, the interlayer spacing of pearlite sheets is thinned, and the rest V is in a solid solution state. After pulling, the C atoms in the cementite are dissolved back, and a small amount of MX-type second phase particles are continuously separated out from V in the galvanization or stabilization treatment process. An excessively low V content is detrimental to the stress corrosion resistance of ultra-high strength pre-stressed steels of the order of 2160MPa and above (i.e. downstream products of the present wire rod), and when a higher V content (e.g. exceeding 0.15%) is added, V will precipitate at a higher temperature, although the total amount of precipitation will increase, the precipitated particles will also increase, which is detrimental to significantly improving the stress corrosion resistance, while also increasing the alloy cost. The V precipitation in the steel accounts for 50-80% of the total addition, wherein the hot rolling process can precipitate 50-65%, and the isothermal salt bath heat treatment process can precipitate 65-80%.
Ti: the strong carbonitride forming element mainly forms TiN or TiC. The solid solution amount of TiN in the steel is small, and the solid solution amount of TiC in the steel is also small. The steel of the invention utilizes TiN to fix N, but in order to avoid the precipitation of TiN in liquid, the lower Ti addition amount is controlled on the basis of controlling the N content to be not more than 0.004%, the atomic ratio of TiN is 3.42, the lower limit of Ti addition is about 3.42 times of the N content and is set to be 0.01%, and the upper limit is not more than 0.025%. The trace Ti after N fixation also precipitates TiC in a lower off-line austenitizing reheating process or an air cooling process or an isothermal salt bath heat treatment process. After pulling, the C atoms in cementite are dissolved back, and Ti atoms which are not fully separated out by hot rolling can be continuously separated out a small amount of MX type second phase particles in the stabilizing treatment process. The precipitation of Ti in the steel accounts for 95-100% of the total addition. The semi-coherent interface formed by TiC and the iron matrix can be a hydrogen trap to play a role in resisting stress corrosion.
Nb: the strong carbonitride forming element, nb combines with N to form NbN with a solid solubility that is an order of magnitude greater than TiN, so Ti in the steel of the present invention fixes N at a high temperature stage, while Nb mainly combines with C to precipitate NbC. Because the content of C in the steel is higher, the solid solubility of NbC is more than 1 order of magnitude less than that of VC, and therefore, the Nb content of the steel is controlled to be 0.01-0.03%. NbC is completely precipitated during the hot rolling process or the off-line austenitizing reheating process. Nb is partially separated out in the hot rolling process, and NbC particles and solid-solution Nb act together to refine austenite grains, so that austenite deformation energy storage is improved, pearlite transformation driving force is improved, and meanwhile, pearlite block mass and lamellar spacing are refined; the incompletely precipitated Nb is continuously precipitated in the cooling process, so that the strength is further improved; nb is almost completely precipitated during austenitizing reheating in off-line isothermal heat treatment, refining austenite grains, and further refining the pearlite block size. After pulling, a small amount of MX-type second phase particles can be continuously precipitated in the stabilizing treatment process of insufficient precipitation of hot rolling due to dissolution back of C atoms in cementite. The precipitation of Nb accounts for 95-100% of the total addition, and the stress corrosion resistance is achieved by taking the atomic defect positions in the NbC lattice and the semi-coherent interface of NbC and the iron matrix as hydrogen traps.
Al: the steel of the invention takes the strong deoxidizing element as an optional deoxidizing element, the control content is less than or equal to 0.04 percent, and the excessive Al content easily causes the coarse size of partial inclusion.
P: the toughness and plasticity of the steel are reduced, and the P content of the steel is controlled to be less than or equal to 0.02 percent.
S: the toughness of the steel is reduced, and the S content of the steel is controlled to be less than or equal to 0.005%.
N: an effective interstitial solid solution strengthening element and a second phase strengthening element in the steel. The free N atoms cause the strength of the wire rod to be increased, the toughness to be reduced, and phenomena such as aging and blue embrittlement to be easily caused, so that the total N content and the free N content are as low as possible. Ti is adopted to fix N, so that TiN is formed. In order to avoid the formation of coarse, large amounts of TiN and thus the deterioration of wire rod drawing processability, the total N content needs to be controlled as low as possible. In connection with the experimental or production equipment technical level, the N content of the steel according to the invention does not exceed 0.004%.
Specifically, si, mn, cr, mo alloy components in the ultra-high strength wire rod satisfy the following mass percentages: si+Mn+Cr+Mo is 1.2-2.4 and Cr+Mo is 0.40-0.85. The alloy composition ratio can further effectively refine the pearlite colony spacing of the ultra-high strength wire rod.
Preferably, in order to exert the strengthening effect of Nb microalloying in the hot rolling process and match with V microalloying, the ultra-high strength wire rod manufactured in a mode of hot rolling and air cooling or hot rolling and air cooling and on-line isothermal heat treatment (isothermal salt bath heat treatment) is adopted, and the ratio of the V, nb and Ti microalloying is as follows: v:0.07 to 0.15 percent of Ti:0.01% -0.025%, nb:0.01 to 0.03 percent.
Preferably, in order to exert the effect of refining austenite grains by Nb microalloying in the off-line heat treatment process and thereby refining the size of pearlite block, and in cooperation with the V microalloying refining effect, the ultra-high strength wire rod manufactured by the method of "hot rolling + off-line isothermal heat treatment + air cooling + on-line isothermal heat treatment (isothermal salt bath heat treatment)" is adopted, and the ratio of V, nb, ti microalloys follows the following two cases: (1) V:0.07 to 0.11 percent of Ti:0.01% -0.025%, nb:0.01% -0.03%; (2) V:0.12 to 0.15 percent of Ti:0.01% -0.025%, nb: less than 0.01%.
Specifically, the diameter of the ultra-high strength wire rod is less than or equal to 15mm, the microstructure is complete lamellar pearlite, and no network cementite exists.
The diameter specification of the wire rod used for preparing the prestressed steel wire, wire and rope is generally not more than 15mm, and the diameter specification is usually 10-15 mm. The diameter of the wire rod for preparing the prestress steel strand is commonly used to be 12.5-13 mm, and the wire rod can be expanded to 14mm when the strength and the plasticity are excellent; the diameter of the wire rod for preparing the prestressed steel wire is commonly used to be 10-13 mm; the diameter of the wire rod for preparing the prestressed wire rope is usually 12.5-15 mm.
The lamellar pearly-luster structure with refined mass size and lamellar spacing, diversified lamellar orientation has high strength, good deformability and excellent drawing performance. However, if some network cementite is formed, the proeutectoid cementite at these austenite grain boundaries is coarser in size, more brittle, not only is the strength impaired, but also the drawing performance is deteriorated.
Specifically, the tensile strength of the ultra-high strength wire rod reaches 1400-1680 MPa, and the area shrinkage is more than or equal to 28%. The ultra-high strength wire rod has excellent tensile strength and high area shrinkage, and is beneficial to preparing the downstream products such as prestressed steel/steel stranded wires with high strength and high toughness.
The invention also discloses a preparation method of the ultra-high strength wire rod, which comprises the following steps:
s1: batching according to alloy composition, and smelting and casting to obtain a casting blank;
s2: homogenizing the casting blank, namely heating to the austenite homogenizing temperature and preserving heat;
S3: continuously rolling the heat-preserving casting blank after descaling until reaching the target diameter to obtain a semi-finished product wire rod;
S4: and curling and post-processing the semi-finished wire rod to obtain the finished product of the ultra-high strength wire rod.
Specifically, in the step S1, a casting blank can be obtained by adopting vacuum induction smelting, die casting and forging modes.
Preferably, in the step S1, excessive raw materials of each alloy element can be weighed according to alloy composition, and added into a high-temperature converter, and casting blanks are formed through converter smelting, LF refining, RH or VD degassing, electromagnetic stirring and continuous casting. The method comprises the steps of controlling relevant parameters of converter smelting, LF refining, RH or VD degassing, carrying out real-time sampling analysis on molten steel in the converter, pouring out the molten steel when each element component reaches a preset value/range, and carrying out magnetic stirring and continuous casting. The process and the parameters are controlled to be a mature alloy preparation process in the prior art, and the parameters can be adjusted in the implementation process according to practical experience and specific alloy composition.
Specifically, the austenite homogenization temperature in the step S2 is 1180-1250 ℃, and the heat preservation time is 0.5-3 h. When the homogenization temperature is too high, the TiN part in the casting blank can be completely dissolved in a solid state, so that austenite grains are coarsened, and the energy cost is increased; too low homogenization may result in insufficient solid solution of strong carbonitride forming elements such as Nb, V, etc. in the steel, thereby affecting the subsequent finished product properties. The homogenization heat preservation time is too long, austenite grains coarsen, the energy cost is increased, and the production efficiency is not facilitated; the homogenization heat preservation is too short, and the thickness and temperature uniformity of the casting blank are difficult to ensure. Therefore, the heat preservation time is controlled to be 0.5-3 h.
Preferably, when the Nb content in the alloy is 0.01% -0.03%, the austenite homogenization temperature is 1200-1250 ℃, so that more than 80% of Nb is in a solid solution state.
Illustratively, the descaling operation in step S3 may employ a high-pressure water descaling technique or other conventional descaling process. The descaling operation has the function of removing the iron scales so as to prevent the iron scales from being pressed into the surface of the alloy/wire rod to generate defects, thereby improving the surface quality of the product.
Specifically, the heating temperature before hot rolling in the step S3 is 1180-1250 ℃, and the hot rolling finishing temperature is 890-950 ℃. The final rolling temperature is too high, the hot-rolled austenite grains are coarser, and the effect of refining the hot-rolled austenite grains by Nb microalloying is not obvious; too low a finish rolling temperature will result in more V precipitation and a significant increase in rolling force.
Specifically, the post-treatment mode in step S4 is one of air cooling, (air cooling+isothermal salt bath heat treatment) or (off-line isothermal heat treatment+air cooling+isothermal salt bath heat treatment).
Specifically, when the post-treatment mode is air cooling, the air cooling speed is 1.5 ℃/s-3.5 ℃/s after the hot rolling and before 475 ℃. The air cooling speed is more than or equal to 1.5 ℃/s, and the formation of proeutectoid cementite (network cementite) can be avoided; when the air cooling speed is more than or equal to 3.5 ℃, a large amount of bainite tissues can be formed, and even part of superhard martensite tissues are formed in the subsequent cooling process, so that the drawing is not facilitated, and the excessive air cooling speed is avoided.
Specifically, when the post-treatment mode is 'air cooling and isothermal salt bath heat treatment', the air cooling speed after hot rolling and before isothermal salt bath heat treatment is more than or equal to 1.5 ℃/s, the formation of proeutectoid cementite (network cementite) is avoided, the isothermal salt bath heat treatment temperature is controlled to 475-550 ℃, and the isothermal time is controlled to 1-5 min. Depending on the chemical composition, the temperature and time match to obtain a complete pearlitic structure. For the chemical composition of the invention, the higher the temperature, the greater the lamellar spacing of the pearlite is, the invention aims at ultra-high strength wire rods, so the temperature does not exceed 550 ℃, and the upper temperature does not exceed the fastest isothermal transformation temperature of the pearlite. On the other hand, the temperature is not lower than 475 ℃, and bainite transformation does not occur, as the lower limit temperature. The temperature is lower than the fastest pearlite isothermal transformation temperature, so that the lower the temperature is, the longer the pearlite isothermal transformation time is, the pearlite isothermal transformation time of the steel is less than or equal to 5min at the lower limit temperature of 475 ℃, the pearlite isothermal transformation time at the upper limit temperature of 550 ℃ is tens of seconds, and the shortest isothermal transformation time is controlled to be 1min, so that the pearlite isothermal transformation is fully completed.
Specifically, when the post-treatment mode is 'off-line isothermal heat treatment + air cooling + isothermal salt bath heat treatment', the off-line austenitizing temperature is controlled to be 900-980 ℃, the heat preservation time is controlled to be 10-1 min, the air cooling speed before entering the salt bath is more than or equal to 1.5 ℃/s, the formation of proeutectoid cementite (network cementite) is avoided, the isothermal salt bath heat treatment temperature T5 is controlled to be 475-550 ℃, and the isothermal time is controlled to be 1-5 min.
The offline austenitizing temperature control principle is as follows: the lower limit is about 30 ℃ above the equilibrium transformation temperature for complete austenitization of the steel of the invention so that complete austenitization can occur at temperature when the wire rod is heat treated offline; the upper limit must not be too high, on the one hand, based on controlling the austenite grains to be fine and, on the other hand, on controlling the energy consumption. The off-line heat preservation time is matched with the off-line austenitizing temperature, the higher the temperature is, the shorter the time is, the austenitizing heat preservation is carried out for 10s at 980 ℃, the austenitizing heat preservation is carried out for 1min at 900 ℃, the fineness of austenite grains and the energy consumption can be controlled at a high temperature in a short time, and meanwhile, the austenitic phase transformation is sufficiently carried out and the components are homogenized.
The invention also discloses Si-Mn-Cr-Mo-V-Ti-Nb multi-element alloyed ultra-high strength prestressed steel, which is prepared from the ultra-high strength wire rod or the ultra-high strength wire rod prepared by the preparation method.
Specifically, the ultra-high strength wire rod is subjected to cold drawing and stabilizing heat treatment to prepare the high strength prestressed steel, and the specific operation and parameters are as follows:
cold drawing: acid washing is carried out before cold drawing, the acid washing time is 15-25 min, a surface oxide layer is cleaned, and the surface quality of the wire rod before drawing is improved; because the wire rod has high strength, in order to establish better drawing lubrication conditions, the wire rod after pickling is subjected to phosphating treatment for 5-10 min, so that a compact phosphating film is formed on the surface of the wire rod; drawing times after phosphating are 7-15 times, and drawing speed is 2-4 m/s.
Stabilizing heat treatment: the cold drawn monofilament has a diameter of 5-7 mm, and the steel has high heat stability, high ductility, post-stranding stabilizing heat treatment temperature of 400-480 deg.c and linear speed of 1.2-2.0 m/s. The stabilization heat treatment process is fully separated out and matched with V, ti, nb, mo microalloy, the temperature is matched with the linear velocity, the linear velocity is high when the temperature is high, the linear velocity is low when the temperature is low, and the separation of V, ti, nb, mo nanometer MX second phase particles is facilitated, and the strong plasticity is improved.
Specifically, the diameter of the prestressed steel is less than or equal to 7mm, the tensile strength is 2160-2460 MPa, and the stress corrosion resistance meets the following requirements: and carrying out a solution A (analyzing pure ammonium thiocyanate aqueous solution) according to the specification of GB/T21839, carrying out a stress corrosion test by loading stress with 80% of the actual maximum force, wherein the test group is not less than 5 groups, the minimum value of the stress corrosion breaking time of the prestressed steel is not less than 2 hours, and the median value is not less than 5 hours. Solution a (analytically pure aqueous ammonium thiocyanate) was: 200g of NH 4 SCN was dissolved in 800mL of distilled or demineralised water to prepare an ammonium thiocyanate solution. Ammonium thiocyanate is analytically pure, wherein the content of NH 4 SCN is at least 98.5%, cl - is less than 0.005%,＜0.005%，S2-<0.001%。
Specifically, the ultra-high strength wire rod is subjected to cold drawing, hot galvanizing and stabilizing heat treatment to prepare the high strength prestressed steel, and specific operations and parameters are as follows:
cold drawing: acid washing is carried out before cold drawing, the acid washing time is 15-25 min, a surface oxide layer is cleaned, and the surface quality of the wire rod before drawing is improved; because the wire rod has high strength, in order to establish better drawing lubrication conditions, the wire rod after pickling is subjected to phosphating treatment for 5-10 min, so that a compact phosphating film is formed on the surface of the wire rod; drawing times after phosphating are 7-15 times, and drawing speed is 2-4 m/s.
Hot galvanizing: the diameter of the cold drawn monofilament is 3-7 mm, and the temperature of hot galvanizing (zinc aluminum, zinc aluminum magnesium) is 420-500 ℃ and the time is 20-40 s.
Stabilizing heat treatment: the hot galvanizing time is short, the heating processes experienced by the surface and the interior of the cold drawn steel wire are inconsistent, so that the original relatively uniform residual stress can become unevenly distributed, and macroscopic appearance is that the torsion performance of the steel wire is sharply reduced. The steel of the invention has good heat stability, and the post-stranding stabilization heat treatment temperature is 400-480 ℃ and the linear velocity is 0.5-2.0 m/s for obtaining higher ductility and torsion performance. The stabilization heat treatment process is fully separated out and matched with V, ti, nb, mo microalloy, the temperature is matched with the linear velocity, the linear velocity is high when the temperature is high, the linear velocity is low when the temperature is low, and the separation of V, ti, nb, mo nanometer MX second phase particles is facilitated, and the strong plasticity is improved.
The diameter of the prestressed steel subjected to cold drawing, hot galvanizing and stabilizing heat treatment of the ultra-high strength wire rod is less than or equal to 7mm, the tensile strength is 2160-2260 MPa, and the torsion performance reaches: the torsion is more than or equal to 18 times.
The advantages of the precise control of the composition and process parameters of the steel according to the invention will be demonstrated in the following in the specific examples and comparative examples. The chemical composition of the steels of examples 1 to 9 and comparative example 1 are shown in Table 1, the specific rolling and cooling process parameters are shown in Table 2, the heat treatment process parameters are shown in Table 3, and the mechanical properties are shown in Table 4.
Comparative examples 1 and 3, which selected conventional 82B hot rolled wire rods as example components, processes and properties, comparative examples 2,4 and 5 have higher carbon content with little or no microalloy elements added.
The specific steps are as follows:
Examples 1 to 2 were vacuum induction melted, die cast, and forged into a 160mm thick 160mm wide billet, 160mm thick 160mm wide 1000mm long 1 block was cut, welded to a conventional 82b 160mm thick 160mm wide 11000mm long billet, heated to 1220 c, heat-preserved for 1 hour, hot rolled into 13mm wire rods after tapping, and cooled to 475 c at a finish rolling temperature of 910 c at an air cooling rate of 1.8 to 2.8 c/s.
Examples 3-6 were smelted in a converter, LF refined, RH degassed, electromagnetic stirred, continuous cast to 160mm (thickness). Times.160 mm (width) square billets, cut to 12000mm billets, rolled to 1180-1250 ℃, respectively kept at temperature for 0.5-3 h, hot rolled to 10-15 mm diameter wire rods after tapping, and final rolling temperature 890-950 ℃. The air cooling speed is 1.5-3.3 ℃/s before cooling to 475 ℃ after hot rolling. Wherein, examples 3, 4 and 6 adopt on-line salt bath heat treatment, after hot rolling, the hot rolling is cooled to the air cooling speed of 1.5-2.2 ℃/s of on-line isothermal salt bath heat treatment, the on-line isothermal salt bath temperature is 475-550 ℃, and the isothermal time is 1-5 min; example 5 was air cooled on-line at a rate of 3.3 c/s before cooling to 475 c.
Examples 7 to 9 were smelted by a converter, LF refined, VD degassed, electromagnetic stirred, continuous cast into square billets 160mm (thick) ×160mm (wide), cut into billets 12000mm long, rolled and heated to 1180-1210 ℃, respectively kept at the temperature for 2-3 hours, hot rolled into wire rods 13-15 mm in diameter after being discharged, and cooled to 475 ℃ after final rolling at 890-950 ℃ at an air cooling rate of 1.6-2.3 ℃/s. Wherein, examples 7 and 9 adopt off-line salt bath heat treatment, the off-line austenitizing temperature is 900-980 ℃, the heat preservation is carried out for 10 s-1 min, the cooling is carried out to the air cooling speed of 1.5-2.0 ℃/s of the off-line salt bath heat treatment after austenitizing, the off-line isothermal salt bath temperature is 475-550 ℃, and the isothermal time is 1-5 min; example 8 was air cooled on-line at a rate of 1.9 c/s before cooling to 475 c.
Table 1 chemical composition wt% of examples and comparative examples
Table 2 specific hot rolling cooling process parameters of examples and comparative examples
Table 3 example on-line or off-line salt bath heat treatment process parameters
Table 4 diameter specification and mechanical properties of the example and comparative wire rods
As can be seen from tables 2, 3 and 4, examples 1,2, 5 and 8 are hot rolled + air cooled wire rods, satisfying the tensile strength of 1400MPa to 1550MPa (for example, 1430 to 1575 MPa), the reduction of area is more than or equal to 28% (for example, 28.5% to 31.0%), and the diameter specification is 13 to 15mm. And comparative examples 1 and 3 are also ordinary 82B hot rolled wire rods, and the wire rod tensile strength level of 13-14 mm diameter is only 1200MPa grade.
Examples 3, 4 and 6 are hot rolling, air cooling and on-line isothermal salt bath heat treatment wire rods, which meet the tensile strength of 1480-1680 MPa (for example, 1505-1685 MPa), the area reduction rate of more than or equal to 28% (for example, 26.5-29.5%), and the diameter specification of 10-15 mm. Comparative example 4 is a hot rolled + air cooled + on-line isothermal salt bath heat treated wire rod, the wire rod tensile strength of 15mm diameter is only 1380MPa.
Examples 7 and 9 are hot rolling, off-line isothermal heat treatment (austenitizing treatment), air cooling and isothermal salt bath heat treatment wire rods, and satisfy the tensile strength of 1450-1600 MPa (e.g. 1470-1620 MPa), the area reduction rate is more than or equal to 28% (e.g. 27.0% -32.0%), and the diameter specification is 14-15 mm.
Comparative examples 2 and 5 are hot rolling + off-line isothermal heat treatment (austenitizing treatment) +air cooling + isothermal salt bath heat treatment wire rods, and wire rods with diameters of 14-15 mm have tensile strengths of only 1400-1480 MPa.
The temperature of the stabilizing heat treatment is 400-480 ℃ and the linear speed is 0.5-2.0 m/s.
Table 5 example wire rod corresponding prestressed steel hot galvanizing and stabilizing heat treatment process and performance
Because the steel is matched and added with the heat-strength element and the strong carbide element, the steel can obtain higher basic strength before cold drawing, and can also adopt higher temperature and longer time for hot galvanizing and stabilizing heat treatment after cold drawing, thereby improving plasticity and residual stress distribution and having smaller strength loss. The temperature and time of the cold drawing post treatment of the comparative example are relatively low, the strength loss is still large, and the strong plastic matching and the service performance are far inferior to those of the invention example.
As shown in Table 5, the prestressed steel with the diameter of 5-7 mm, which is subjected to cold drawing and stabilizing heat treatment by the ultra-high strength wire rods of examples 1,2, 4, 6 and 9, has the tensile strength reaching 2160-2460 MPa, and the stress corrosion resistance meets the following requirements: and carrying out a solution A (analyzing pure ammonium thiocyanate aqueous solution) according to the specification of GB/T21839, carrying out a stress corrosion test by loading stress with 80% of the actual maximum force, wherein the test group is not less than 5 groups, the minimum value of the stress corrosion breaking time of the prestressed steel is not less than 2 hours, and the median value is not less than 5 hours. The high-strength wire rods of comparative examples 1, 3 and 4 were cold drawn, stabilized heat treated, prestressed steels having diameters of 5 to 7mm, and the minimum stress corrosion fracture time was 2.3 to 2.6 hours, and the median value was only 2.8 to 3.0 hours.
The prestressed steel with the diameter of 3-7 mm is prepared by carrying out cold drawing, hot galvanizing and stabilizing heat treatment on the ultrahigh-strength wire rods of examples 3, 4, 6 and 8, wherein the tensile strength reaches 2160-2300 MPa, and the torsion performance reaches: the torsion is more than or equal to 18 times. The high-strength wire rods of comparative examples 2 and 5 were cold drawn, hot dip galvanized, and stabilized heat treated to 5mm diameter prestressed steels having tensile strengths of 2060 to 2160MPa, but the number of pass of torsion was only 8 to 10.
The present invention is not limited to the above-mentioned embodiments, and any changes or substitutions that can be easily understood by those skilled in the art within the technical scope of the present invention are intended to be included in the scope of the present invention.
Claims (5)
1. A Si-Mn-Cr-Mo-V-Ti-Nb multi-element alloyed ultra-high strength wire rod is characterized in that: the ultra-high strength wire rod alloy comprises ：C：0.90%～1.00%,Si：0.70%～1.20%,Cr：0.20%～0.80%,Mn：0.10%～0.60%,Mo：0.05%～0.35%,V：0.07%～0.15%, Ti：0.010%～0.025%,Nb：≤0.03%,Al：≤0.04%,P：≤0.02%,S：≤0.005%,N：≤0.004%, parts by mass of Fe and unavoidable impurities as the rest;
The Si, mn, cr, mo alloy components in the ultra-high strength wire rod are as follows in percentage by mass: si+Mn+Cr+Mo is more than or equal to 1.2 and less than or equal to 2.4, and Cr+Mo is more than or equal to 0.40 and less than or equal to 0.85;
the ultra-high strength wire rod manufactured by adopting a mode of hot rolling, air cooling or hot rolling, air cooling and isothermal salt bath heat treatment is characterized in that the ratio of the V, nb and Ti microalloy is as follows: v:0.07 to 0.15 percent of Ti:0.01% -0.025%, nb:0.01% -0.03%;
The ultra-high strength wire rod manufactured by adopting a mode of hot rolling, off-line isothermal heat treatment, air cooling and isothermal salt bath heat treatment is characterized in that the ratio of the V, nb and Ti microalloys is as follows: (1) V:0.07 to 0.11 percent of Ti:0.01% -0.025%, nb:0.01% -0.03%; or (2) V:0.12 to 0.15 percent of Ti:0.01% -0.025%, nb: less than 0.01%;
the diameter of the ultra-high strength wire rod is less than or equal to 15mm, the microstructure is complete lamellar pearlite, and no network cementite exists;
The tensile strength of the ultra-high strength wire rod reaches 1400-1680 MPa, and the area shrinkage rate is more than or equal to 28%.
2. A method of making the ultra-high strength wire rod of claim 1, comprising the steps of:
s1: batching according to alloy composition, and smelting and casting to obtain a casting blank;
S2: homogenizing the casting blank, namely heating to the austenite homogenizing temperature and preserving heat;
s3: continuously rolling the heat-preserving casting blank after descaling until reaching the target diameter to obtain a semi-finished product wire rod;
s4: curling and post-processing the semi-finished wire rod to obtain a finished product of the ultra-high strength wire rod;
the post-treatment mode in the step S4 is one of air cooling, isothermal salt bath heat treatment or off-line isothermal heat treatment, air cooling and isothermal salt bath heat treatment;
Wherein the air cooling speed before 475 ℃ after hot rolling and final rolling is 1.5 ℃/s-3.5 ℃/s;
Wherein the isothermal salt bath heat treatment temperature is 475-550 ℃, and the isothermal time is controlled to be 1-5 min.
3. The preparation method according to claim 2, characterized in that: the austenite homogenizing temperature in the step S2 is 1180-1250 ℃, and the heat preservation time is 0.5-3 h.
4. The preparation method according to claim 2, characterized in that: the heating temperature before hot rolling in the step S3 is 1180-1250 ℃, and the hot rolling finishing temperature is 890-950 ℃.
5. A Si-Mn-Cr-Mo-V-Ti-Nb multi-alloyed ultra-high strength prestressed steel is characterized in that: the pre-stressed steel is made of the ultra-high strength wire rod of claim 1 or the ultra-high strength wire rod made by the manufacturing method of any one of claims 2 to 4.
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| JP2002194496A (en) * | 2000-12-28 | 2002-07-10 | Sumitomo Electric Ind Ltd | Steel wire for spring, spring and its production method |
| CN105671443A (en) * | 2016-02-25 | 2016-06-15 | 邢台钢铁有限责任公司 | Hot-rolled wire rod for 1,960MPa-level cable rope galvanized steel wire and production method |
| CN114015946A (en) * | 2021-11-10 | 2022-02-08 | 湖南三泰新材料股份有限公司 | High-strength corrosion-resistant stainless steel coated steel wire for bridge cable and preparation method thereof |
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| CN103805861B (en) * | 2014-02-11 | 2016-06-01 | 江苏省沙钢钢铁研究院有限公司 | A kind of carbon steel wire rod with high and its preparation method |
| JP6481770B2 (en) * | 2015-10-23 | 2019-03-13 | 新日鐵住金株式会社 | Steel wire rod for wire drawing |
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| JP2002194496A (en) * | 2000-12-28 | 2002-07-10 | Sumitomo Electric Ind Ltd | Steel wire for spring, spring and its production method |
| CN105671443A (en) * | 2016-02-25 | 2016-06-15 | 邢台钢铁有限责任公司 | Hot-rolled wire rod for 1,960MPa-level cable rope galvanized steel wire and production method |
| CN114015946A (en) * | 2021-11-10 | 2022-02-08 | 湖南三泰新材料股份有限公司 | High-strength corrosion-resistant stainless steel coated steel wire for bridge cable and preparation method thereof |
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