The present disclosure relates to steel plates and manufacturing methods therefor, and particularly to hot-rolled high-strength steel and manufacturing methods therefor.

With the development of automotive lightweighting technology, high-strength steel plates are playing an increasingly important role in automotive structural members. Currently, many car models use grade 80 kg steel plates to produce automobile chassis components. Automobile chassis components, such as control arms, undergo forming processes including stamping, flanging, hole expansion, and the like, requiring extremely high hole expansion performance.

Chinese patent application CN104513930A, published on April 15, 2015, entitled "ULTRA-HIGH STRENGTH HOT-ROLLED COMPLEX PHASE STEEL PLATES AND STRIP STEEL WITH GOOD BENDING AND HOLE EXPANSION PERFORMANCE AND MANUFACTURING METHOD THEREFOR", discloses a steel plate with a microstructure of at least 80% bainite, ferrite, martensite, and retained austenite, having a hole-expansion ratio of over 50%, and a 180° cold bending of 0a.

The Chinese patent literature with publication number CN105154769A, published on December 16, 2015, entitled "GRADE 780MPA HOT-ROLLED HIGH-STRENGTH AND HIGH-EXPANSION STEEL AND MANUFACTURING METHOD THEREFOR", discloses that the hot-rolled coiling temperature of the steel plate is 600-700°C, and there is sufficient time for precipitation within the temperature range of 600-700°C. By controlling the cooling rate after coiling to be less than or equal to 20°C/h, the microstructure is ferrite.

Chinese patent literature with publication number CN112575267A, published on March 30, 2021, entitled "HIGH HOLE-EXPANSION COMPLEX PHASE STEEL AND MANUFACTURING METHOD THEREFOR", discloses a transverse tensile strength of ≥ 780 MPa, a yield strength of ≥ 700 MPa, an A50 elongation of ≥ 15%, and a punching hole-expansion ratio of ≥ 50%. The microstructure consists of bainite and ferrite.

It can be seen that the above patent literatures are silent on the effect of grain boundary density on hole-expansion ratio and bending performance.

One of the objectives of the present disclosure is to provide a hot-rolled high-strength steel having excellent hole-expansion and bending performance, and a manufacturing method therefor. The hot-rolled high-strength steel possesses high punching quality, excellent hole-expansion ratio and bending performance, and can be used as automotive body structural members and automotive chassis components.

For the above purposes, the present disclosure provides a hot-rolled high-strength steel having excellent hole-expansion and bending performance, comprising Fe and inevitable impurities, and further comprising the following chemical elements in percentage by mass:

• C: 0.030-0.080%, Si: 0.01-1.20%, Mn: 1.20-1.80%, S: 0.0005-0.0080%, Al: 0.020-1.000%, and B≤0.0035%;
• at least one of Mg: 0.0002-0.0100% and Ca: 0.0002-0.0100%; and
• at least one of 0<Ti≤0.13%, 0<Nb≤0.06% and 0<V≤0.20%;
• wherein the grain boundary density of a microstructure of the hot-rolled high-strength steel is 0.6-2.0µm^-1.

Further, the hot-rolled high-strength steel of the present disclosure has the following percentages by mass of the chemical elements:
C: 0.030-0.080%, Si: 0.01-1.20%, Mn: 1.20-1.80%, S: 0.0005-0.0080%, Al: 0.020-1.000%, and B≤0.0035%; at least one of Mg: 0.0002-0.0100% and Ca: 0.0002-0.0100%; at least one of 0<Ti≤0.13%, 0<Nb≤0.06% and 0<V≤0.20%; with the balance being Fe and inevitable impurities.

In some embodiments, the hot-rolled high-strength steel of the present disclosure has a grain boundary density of 1.0-2.0µm^-1.

The following describes in detail the design principles of each chemical element of the hot-rolled high-strength steel according to the present disclosure.

C: In the hot-rolled high-strength steel according to the present disclosure, the carbon content largely determines the tensile strength grade of a steel plate. Elemental carbon can be used for solid solution strengthening and it can form sufficient precipitation strengthening phases with elements such as titanium to ensure the strength of the steel. However, an excessively high mass percentage of carbon will result in coarse carbide particles and tend to form excessive carbides (pearlite, cementite), which is detrimental to the hole-expansion performance. In order to achieve high hole expandability for the steel grade having such high strength, as well as good formability and weldability, in the hot-rolled high-strength steel according to the present disclosure, the mass percentage of C is controlled to be between 0.030% and 0.080%.

Si: In the hot-rolled high-strength steel according to the present disclosure, elemental silicon can function to solid-solution strengthening to improve the strength of the steel plate. Meanwhile, the addition of silicon inhibits the formation of harmful carbides, increases the ferrite fraction, and helps improve the elongation of the steel plate. However, an excessively high content of silicon in the steel tends to cause surface defects of 2FeO-SiO₂ scale on the steel plate, which adversely affects the surface quality. Based on this, in the hot-rolled high-strength steel according to the present disclosure, the mass percentage of silicon is controlled to be between 0.01% and 1.20%.

Mn: Manganese is a solid-solution strengthening element in the hot-rolled high-strength steel according to the present disclosure. Manganese can delay the pearlite transformation, improve the hardenability of steel, and lower the bainite transformation temperature, thereby refining the substructure of the steel and ensuring the formation of lath substructure. This allows the product to possess excellent formability while maintaining the tensile strength. An excessively low mass percentage of manganese results in insufficient strength. However, an excessively high mass percentage of manganese reduces the plasticity of the steel plate and causes centerline segregation, which impairs formability. Furthermore, an excessively high Mn content tends to form MnS with S, which promotes cracking during punching or cutting of the steel strip and further degrades formability. Based on this, in the hot-rolled high-strength steel according to the present disclosure, the mass percentage of Mn is controlled to be between 1.20% and 1.80%.

S: In the hot-rolled high-strength steel according to the present disclosure, elemental sulfur (S) is a beneficial element rather than an impurity. It can form sulfides with magnesium (Mg) or calcium (Ca), act as nucleus for (Ti,Nb)N or TiN, and facilitate the refinement of (Ti,Nb)N or TiN, thereby improving the punching and trimming quality, and enhancing the hole-expansion performance. Nevertheless, an excessive S content leads to formation of MnS by reacting with Mn, which conversely degrades the hole expansion performance. Accordingly, in the hot-rolled high-strength steel according to the present disclosure, the mass percentage of S is controlled to be between 0.0005% and 0.0080%.

Al: In the hot-rolled high-strength steel according to the present disclosure, aluminum (Al) serves as a deoxidizing element for the steel. It can reduce oxide inclusions in the steel and thus purify the steel, which is beneficial to improving the formability of the steel plate. Similar to silicon, aluminum inhibits the formation of harmful carbides, increases the ferrite fraction, and helps improve the elongation of the steel plate. However, an excessively high mass percentage of aluminum causes oxidation, which further affects continuous casting. Based on this, in the hot-rolled high-strength steel according to the present disclosure, the mass percentage of Al is controlled to be between 0.020% and 1.000%.

B: In the hot-rolled high-strength steel according to the present disclosure, boron (B) is beneficial for expanding the bainite phase region and ensuring that a bainite structure can be obtained in the steel plate during cooling after rolling, which significantly improves the strength and hardness of the steel. However, an excessive amount of B element leads to formation of an excessive amount of massive martensite in the steel plate, resulting in a decrease in both the hole-expansion ratio and elongation of the steel. Based on this, in the hot-rolled high-strength steel according to the present disclosure, the mass percentage of B is controlled as B ≤ 0.0035%.

Mg: magnesium (Mg) can form oxides and sulfides with oxygen and sulfur, respectively, in the hot-rolled high-strength steel according to the present disclosure. Compared with steel without Mg, the formed Mg-based oxides and Mg-based sulfides reduce the size of TiN and (Ti,Nb)N precipitates and make them uniformly dispersed, which is beneficial to the improvement of hole expansion performance. Nevertheless, when the Mg content is less than 0.0002%, the effect is insufficient; when it exceeds 0.01%, excessive oxides and sulfides are formed, which conversely degrades the hole expansion performance. Based on this, in the hot-rolled high-strength steel according to the present disclosure, the mass percentage of Mg is controlled to be between 0.0002% and 0.0100%.

Ca: In the hot-rolled high-strength steel according to the present disclosure, calcium (Ca) has functions similar to magnesium. It can improve the morphology of sulfides such as MnS, transforming elongated sulfides like MnS into spherical MnS, which is beneficial to improving the morphology of inclusions and thus reducing the adverse effects of elongated sulfides on the hole expansion and forming performance. However, excessive addition of calcium will increase the amount of calcium oxide, which is detrimental to hole expansion performance. Based on this, in the hot-rolled high-strength steel according to the present disclosure, the mass percentage of Ca is controlled to be between 0.0002% and 0.0100%.

Ti: In the hot-rolled high-strength steel according to the present disclosure, titanium (Ti) is one of important fine-grain strengthening and precipitation strengthening elements. Ti can increase the recrystallization temperature and refine the grain size during hot rolling. Meanwhile, the combination of Ti and C provides excellent strengthening effect. However, an excessive mass percentage of Ti is unfavorable, as it tends to form TiN having a relatively large size, which is detrimental to the impact toughness of the steel. Accordingly, in the hot-rolled high-strength steel according to the present disclosure, the upper limit of the mass percentage of Ti can be controlled at 0.13%. In some preferred embodiments, the amount of Ti can be controlled in a range of from 0.05% to 0.13%.

Nb: In the hot-rolled high-strength steel according to the present disclosure, niobium (Nb) is one of important precipitation strengthening and fine-grain strengthening elements. However, when the mass percentage of Nb exceeds 0.06%, the strengthening effect of Nb becomes nearly saturated, and the production cost is relatively high. Therefore, in order to exert the beneficial effect of Nb while controlling the production cost, in the hot-rolled high-strength steel according to the present disclosure, the upper limit of the mass percentage of Nb can be controlled at 0.06%. In some preferred embodiments, the mass percentage of Nb can be controlled in a range of from 0.01% to 0.05%.

V: In the hot-rolled high-strength steel according to the present disclosure, vanadium (V) is one of important precipitation strengthening and fine-grain strengthening elements. Cooperating with niobium, V exhibits favorable effect in refining austenite grains and precipitation strengthening. It exists as fine precipitates during cooling after rolling or after coiling, thereby increasing strength via precipitation strengthening. Accordingly, in the hot-rolled high-strength steel according to the present disclosure, the upper limit of the mass percentage of vanadium can be controlled at 0.20%. In some preferred embodiments, the amount of V can be controlled in a range of from 0.05% to 0.20%.

Further, the chemical elements of the hot-rolled high-strength steel of the present disclosure satisfy at least one of the following formulas: $$\mathrm{Mg}/24+\mathrm{Ca}/40\ge \mathrm{O}/16\times 0.83$$ $$\mathrm{S}/32\le \mathrm{Mg}/24+\mathrm{Ca}/40-\mathrm{O}/16\times 0.83$$ and $$\mathrm{S}/32\times \mathrm{Mn}/55\le 8.0\times {10}^{-6}$$

In the formulas, each chemical element is substituted with the numerical value before the percent sign of its mass percentage content.

In the present disclosure, the oxygen content that can effectively combine with Mg and Ca accounts for 83% of the total oxygen content. Therefore, the contents of Mg and Ca preferably satisfy the formula Mg/24+Ca/40≥O/16×0.83. In some embodiments, Mg/24+Ca/40 ranges from 0.00010 to 0.00045.In some embodiments, O/16×0.83 ranges from 0.0003 to 0.00025.

In the present disclosure, a high sulfur content in steel tends to form MnS, which is detrimental to the hole expansion performance and bending performance of the material. Therefore, the addition of elemental Mg and Ca shall not only fix oxygen but also further fix sulfur to reduce the formation of MnS. Accordingly, the contents of Mg and Ca preferably simultaneously satisfy the formula: S/32≤Mg/24+Ca/40-O/16×0.83. In some embodiments, S/32 ranges from 0.00002 to 0.00035. In some embodiments, Mg/24+Ca/40-O/16×0.83 ranges from 0.00005 to 0.00040.

In the present disclosure, the high-strength steel contains a relatively high Mn content. When the S content is also high, MnS precipitates at elevated temperatures, inhibiting the formation of MgS and CaS and degrading hole-expansion and bending performance. Therefore, the contents of Mn and S preferably simultaneously satisfy the formula: S/32×Mn/55≤8.0×10^-6. In some embodiments, S/32×Mn/55≤5.0×10^-6. In some embodiments, S/32×Mn/55≤3.0×10^-6.In some embodiments, S/32×Mn/55 ranges from 4.0×10^-7 to 5.0×10^-6.

Furthermore, in the hot-rolled high-strength steel according to the present disclosure, when the corresponding elements are contained, the mass percentages further satisfy:

• Ti: 0.05-0.13%;
• Nb: 0.01-0.05%; and
• V: 0.05-0.20%.

Further, in the hot-rolled high-strength steel according to the present disclosure, the mass percentages of chemical elements satisfy: 0.5≤4×C/(3.3Nb+3.4V+Ti)≤2.25 (formula 4), where each chemical element is substituted with the numerical value before the percent sign of its mass percentage content.

During the slab heating process of the present disclosure, precipitates of Nb, Ti and V prevent the growth of original austenite grains. During hot rolling, (Nb,Ti,V)C contributes to increased recrystallization temperature and further refines austenite grains. Precipitated (Nb,Ti,V)(C,N) or (Nb,Ti,V)(Cr,Mo)(C,N) facilitates the refinement of transformed bainite and a small amount of martensite grains. During laminar cooling, nanoscale precipitation of (Nb,Ti,V)(C,N) or (Nb,Ti,V)(Cr,Mo)(C,N) exerts a strong precipitation-strengthening effect. Precipitation of them on bainitic and ferritic matrices enhances strength, especially that of the ferrite matrix, further reduces the strength difference between ferrite and bainite/martensite phases, improves the hole-expansion ratio, and refines the grain sizes of bainite, quasi-polygonal ferrite, polygonal ferrite, martensite, etc., in the microstructure, thereby enhancing hole-expansion and bending performance. Therefore, the contents of Nb, V and Ti satisfy: 0.5≤4×C/(3.3Nb+3.4V+Ti) ≤ 2.25, such that a better precipitation-strengthening effect is achieved.

Furthermore, the hot-rolled high-strength steel according to the present disclosure further contains at least one of the following chemical elements: 0<Cr≤0.7% or 0<Mo≤0.25%.

In a preferred embodiment of the present disclosure, at least one of Cr and Mo may be added, wherein:
Cr: In the hot-rolled high-strength steel according to the present disclosure, chromium (Cr) can inhibits pearlite formation and promotes the formation of bainitic microstructure, ultimately improving strength and increasing the hole-expansion ratio. An excessively low Cr content has an insignificant effect on the phase transformation curve, while an excessively high mass percentage of Cr not only increases cost but also tends to generate a relatively large amount of martensitic microstructure. Accordingly, the upper limit of the mass percentage of Cr in the hot-rolled high-strength steel of the present disclosure may be controlled at 0.7%.

Mo: In the hot-rolled high-strength steel according to the present disclosure, elemental molybdenum (Mo) not only inhibits pearlite formation but also facilitates the formation of bainitic microstructure and a small amount of martensite-austenite islands. In addition, elemental Mo can promote bainitic microstructure transformation at relatively high temperatures, allowing coiling of the steel at relatively high temperatures. Such a high coiling temperature provides sufficient precipitation kinetics, thereby inducing significant precipitation strengthening. In the present disclosure, elemental Mo also plays a role in precipitation associated with Nb and Ti, and reduces the possibility of coarsening of precipitated particles. However, it is noted that the Mo content in steel should not be too high. An excessively high Mo content in steel not only increases in alloy cost but also tends to form much martensite and austenite, which is detrimental to the performance of the steel. Accordingly, the upper limit of the mass percentage of elemental Mo in the hot-rolled high-strength steel of the present disclosure may be controlled at 0.25%, in order to exert the beneficial effects of Mo.

Further, in the hot-rolled high-strength steel of the present disclosure,

• when Cr is < 0.40%, the Mo content is 0.05%-0.25%; and
• When Cr ranges from 0.40% to 0.70%, Mo is not added.

In the present disclosure, Cr and Mo play important roles in the formation of bainite. Addition of Cr and/or Mo at an appropriate amount is intended to obtain bainitic structure and fine martensite and pearlite during hot rolling and coiling, so as to increase the grain boundary density of the microstructure, suppress crack propagation during hole expansion, and further improve hole-expansion and bending performance.

Further, among the inevitable impurities of the hot-rolled high-strength steel according to the present disclosure: P is ≤ 0.020%, N is ≤ 0.0050%, and O is ≤ 0.0040%.

The inevitable impurities in the hot-rolled high-strength steel according to the present disclosure mainly include P, N and O. Under permissible technical conditions, the impurity content in steel shall be reduced as much as possible to obtain steel with better performance and higher quality, wherein:
P: An excessive P content degrades weldability, workability and toughness. Accordingly, in some embodiments, the mass percentage of P may be controlled to P≤0.02%.

N: A lower N content is preferred. Nevertheless, nitrogen is an inevitable element during steelmaking. Although N is at a low content, it combines with strong carbide-forming elements such as Ti. TiN and (Ti,Nb)N exhibit a square shape with sharp corners causing great stress concentration between the sharp corners and the matrix, which readily induces cracks and significantly affects fracture toughness and hole expansion performance. Accordingly, in some embodiments, the mass percentage of N may be controlled to N≤0.005%.

O: Oxygen is an inevitable element during steelmaking. For the present disclosure, a certain amount of oxygen remains in steel after deoxidation and tends to form oxides. Inclusions themselves do not exert obvious adverse effects on the performance of the steel plate. However, Al₂O₃ readily acts as a nucleation site for TiN and promotes TiN growth. Based on this, in the hot-rolled high-strength steel according to the present disclosure, the mass percentage of O is controlled to 0 ≤ 0.0040%.

Therefore, in some embodiments, the mass percentages of chemical elements in the hot-rolled high-strength steel according to the present disclosure are as follows: C: 0.030-0.080%; Si: 0.01-1.20%; Mn: 1.20-1.80%; S: 0.0005-0.0080%; Al: 0.020-1.000%; B≤0.0035%; at least one of Mg: 0.0002-0.0100% and Ca: 0.0002-0.0100%; at least one of 0<Ti≤0.13%, 0<Nb≤0.06% and 0<V≤0.20%; and at least one of 0<Cr≤0.7% and 0<Mo≤0.25%; with the balance being Fe and inevitable impurities, wherein among the inevitable impurities, P is ≤ 0.020%, N is ≤ 0.0050%, O is ≤ 0.0040%.

Further, the hot-rolled high-strength steel according to the present disclosure has a microstructure (microstructure) comprising bainite, quasi-polygonal ferrite, polygonal ferrite and martensite.

Further, in the hot-rolled high-strength steel according to the present disclosure, the area fraction of bainite + quasi-polygonal ferrite is 70-95%, the area fraction of polygonal ferrite is 3-30%, and the area fraction of martensite is ≤5%.

Further, in the hot-rolled high-strength steel according to the present disclosure, the grain size of bainite and quasi-polygonal ferrite is ≤4.0 µm, the grain size of polygonal ferrite is 55.0 µm, and the grain size of martensite is ≤3.0 µm.

In some embodiments, the grain size of bainite and quasi-polygonal ferrite ranges from 2.0 µm to 4.0 µm. In some embodiments, the grain size of polygonal ferrite ranges from 3.5 µm to 5.0 µm. In some embodiments, the grain size of martensite, if present, ranges from 1.0 µm to 3.0 µm.

In such embodiments, refining the grain sizes of bainite, quasi-polygonal ferrite, polygonal ferrite and martensite in the microstructure further improves hole expansion and bending performance.

Further, in the hot-rolled high-strength steel of the present disclosure, the precipitates have a size of ≤8 µm.

Further, in the hot-rolled high-strength steel according to the present disclosure, precipitates with a size of ≤3.0 µm accounts for ≤50% of all precipitates. In some embodiments, precipitates with a size of ≤3.0 µm accounts for ≥60% of all precipitates. In some embodiments, precipitates with a size of ≤3.0 µm accounts for 60% to 75% of all precipitates.

Further, in the hot-rolled high-strength steel according to the present disclosure, the precipitates include MgO, CaO, MgS, CaS, TiN and (Nb,Ti)N particles.

In the present disclosure, addition of Mg and Ca forms Mg- and Ca-based oxides and sulfides, which functions as nucleation sites for TiN and (Nb,Ti)N precipitates. Mg- and Ca-based oxides and sulfides precipitated at low temperatures suppress the growth of as well as MnS by competing with TiN and (Ti, Nb)N for precipitation, thereby reducing the formation of fine and uniform voids on the punched section, alleviating stress concentration during hole-expanding processing, and ultimately improving hole-expansion and bending performance.

Further, the hot-rolled high-strength steel according to the present disclosure has a yield strength of ≥660 MPa, a tensile strength of ≥780 MPa, an elongation A50 of ≥15.0%, a punched hole-expansion ratio of ≥75%, and a VDA bending angle of ≥120°.

In some embodiments, the hot-rolled high-strength steel of the present disclosure has a yield strength of ≥680 MPa. In some embodiments, the hot-rolled high-strength steel of the present disclosure has a yield strength of ≥700 MPa. In some embodiments, the hot-rolled high-strength steel of the present disclosure has a yield strength of between 660 MPa and 760 MPa.

In some embodiments, the hot-rolled high-strength steel of the present disclosure has a tensile strength of ≥800 MPa. In some embodiments, the hot-rolled high-strength steel of the present disclosure has a tensile strength of ≥820 MPa. In some embodiments, the hot-rolled high-strength steel has a tensile strength between 780 MPa and 890 MPa. In some embodiments, the hot-rolled high-strength steel has a tensile strength of between 820 MPa and 890 MPa.

In some embodiments, the hot-rolled high-strength steel of the present disclosure has an elongation A50 of ≥17.0%. In some embodiments, the hot-rolled high-strength steel of the present disclosure has an elongation A50 of between 15.0% and 25.0%.

In some embodiments, the hot-rolled high-strength steel of the present disclosure has a hole-expansion ratio of ≥ 85%. In some embodiments, the hot-rolled high-strength steel of the present disclosure has a hole-expansion ratio of ≥ 90%. In some embodiments, the hot-rolled high-strength steel of the present disclosure has a hole-expansion ratio of between 75% and 115%.

In some embodiments, the hot-rolled high-strength steel of the present disclosure has a VDA bending angle of 2 125°. In some embodiments, the hot-rolled high-strength steel of the present disclosure has a VDA bending angle of between 120° and 165°.

In addition, another objective of the present disclosure is to provide a method of manufacturing the hot-rolled high-strength steel. The hot-rolled high-strength steel plate manufactured with the present method possesses high punching quality, excellent hole-expansion ratio and bending performance.

For the above purpose, the present application provides a method of manufacturing hot-rolled high-strength steel having excellent hole-expansion and bending performance, the method comprising steps of:

• smelting and casting, wherein superheat is controlled at 15 to 55°C;
• heating, wherein the slab is heated to a temperature ranging from 1200°C to 1300°C and kept for 1 to 3 h;
• rolling, wherein the rough rolling exit temperature is controlled in a range from 1000°C to 1080°C and the final rolling temperature is controlled in a range from 840°C to 950°C, wherein laminar cooling to a coiling temperature of 300°C to 620°C at an average cooling rate of ≥30°C/s is performed after rolling, and cooling to room temperature at a rate of ≤15°C/s is performed after coiling; and
• acid pickling.

Further, the manufacturing method according to the present disclosure further comprises a galvanizing step after acid pickling, thereby obtaining products with various coatings such as an electro-galvanized coating, a hot-dip galvanized coating, a Zn-Al-Mg coating, and the like.

During smelting and casting of the manufacturing method according to the present disclosure, the levels of central segregation and inclusions in the continuous casting slab can be controlled by controlling the superheat and the secondary cooling water, and using proper soft reduction, with the inclusion grade controlled to be less than 1.5.

In the manufacturing method of the present disclosure, the superheat is controlled to be in a range from 15°C to 55°C in order to properly control the size and precipitation density of MgO, CaO, MgS, CaS, (Nb,Ti)N and TiN. Large, brittle (Nb,Ti)N and TiN with sharp edges and corners act as potential crack sources and drastically degrade the hole-expansion and bending performance of such steel grade. Moreover, the superheat in steelmaking has a significant influence: the greater the superheat, the more beneficial it is to control inclusions. However, higher superheat also promotes TiN growth. Therefore, with the addition of Ca and Mg during steelmaking, the present disclosure can control the superheat at 15°C to 55°C to obtain precipitates including MgO, CaO, MgS, CaS, TiN and (Nb,Ti)N in the steel plate with a size of ≤8 µm, wherein further the proportion of precipitates with a size of ≤3.0 µm reaches ≥50%, which further improves hole-expansion and bending performance.

For micro alloy steels containing Nb, Ti, V or the like, the heating temperature for slabs is particularly critical for performance. A large amount of (Nb,Ti,V)(C,N) having a large size precipitates form during continuous casting. In the process of heating the slab, alloy elements such as Nb, V and Ti need to be dissolved as much as possible to ensure subsequent nanoscale precipitation of micro alloys such as Ti during hot rolling and coiling. Accordingly, the heating temperature is set to ≥1200°C. When the temperature exceeds 1300°C, grain coarsening tends to occur, which is detrimental to the toughness of the steel plate. Additionally, the relatively thick iron oxide scale impairs descaling. Therefore, the heating temperature in the manufacturing method according to the present disclosure is set at 1200°C to 1300°C.

In the present disclosure, control over the rough rolling temperature during hot rolling exerts a great influence on micro alloys such as Nb, V and Ti. At a relatively low rough rolling temperature and during finish rolling, for the micro alloys such as Nb, V and Ti, carbides and carbonitrides of Ti precipitate with large sizes, which is unfavorable for improving the final strength. Accordingly, the rough rolling exit temperature is controlled at 1000°C to 1080°C. In addition, the final rolling temperature of the present disclosure is controlled at 840°C to 950°C to enable rolling in the non-recrystallization zone for grain refinement.

In the present disclosure, to obtain a microstructure of bainite + quasi-polygonal ferrite + polygonal ferrite + martensite, especially a microstructure with a grain boundary density of 0.6µm^-1 to 2.0µm^-1, a combined effect of composition design and manufacturing process is required. With respect to the composition, one or both of Cr and Mo may be selectively added. Addition of Cr and/or Mo effectively suppresses pearlite formation, which is beneficial for increasing grain boundary density, and further improving hole-expansion and bending performance. However, excessive addition tends to form blocky secondary martensite. Therefore, addition of Cr and/or Mo must be coordinated with laminar cooling for hot rolling.

Laminar cooling after rolling in the present disclosure significantly affects the grain boundary density of the microstructure and the volume fractions of bainite, quasi-polygonal ferrite, granular ferrite and martensite phase transformations. In the present disclosure, a reasonable laminar cooling process and hot rolling coiling temperature must be strictly controlled in order to obtain an appropriate microstructure proportion. An excessively high coiling temperature leads to secondary martensite and polygonal ferrite having large sizes, which is unfavorable for the improvement of hole-expansion and bending performance; whereas an excessively low coiling temperature may produce primary martensite, resulting in low hole-expansion ratio and elongation. Therefore, the coiling temperature may be controlled between 300°C and 620°C. Further, the coiling temperature may be controlled between 350°C and 600°C.

After coiling, cooling to room temperature at a rate ≤15°C/s facilitates further bainite transformation and reduces the proportion of martensite, thereby improving hole-expansion and bending performance.

Further, in the method for manufacturing hot-rolled high-strength steel according to the present disclosure, the rolling speed is controlled at 7.0 m/s to 13.0 m/s during the hot rolling step.

Further, in the method for manufacturing hot-rolled high-strength steel according to the present disclosure, during the hot rolling step, controlling is performed such that the total reduction rate is ≥80%, the total reduction rate for finish rolling is ≥50%, and the reduction rate for single-pass in final rolling is ≤15%.

Further, in the method for manufacturing hot-rolled high-strength steel according to the present disclosure, a stage cooling process is adopted for laminar cooling: in the first stage, the steel is cooled to an intermediate point temperature of 610 to 750°C at an average cooling rate of ≥100°C/s (e.g., 100 to 200°C/s), followed by air cooling for 4.0 to 10.0 s; then in the second stage, the steel is cooled to the coiling temperature at an average cooling rate of (e.g., 30 to 140°C/s).

In some embodiments, the tension leveling elongation during pickling is controlled at 0.2% to 1.8% and the pickling speed is 60 to 150 m/min. The temperature of the last pickling tank in the pickling process may be 80 to 90°C, and the iron ion concentration may be 30 to 40 g/L.

In the present disclosure, the laminar cooling process exerts important influence on both the microstructure proportion and precipitation of micro alloys of Nb, V and Ti. The most intensive precipitation temperature range of Nb, V and Ti is from 610°C to 750°C. In practice, the actual coiling temperature is lower than this temperature in order to better exert the nanoscale precipitation strengthening effect of (Nb, Ti, V)(C, N) or (Nb, Ti, V)(Cr, Mo)(C, N). Controlling the temperature during air cooling within a range from 610°C to 750°C also facilitates the formation of quasi-polygonal ferrite. A two-stage cooling is further adopted in laminar cooling, where the intermediate point temperature is 610 to 750°C, the air cooling duration is 4.0 to 10.0 s, and the average cooling rate for the second stage is ≥30°C/s. Meanwhile, precipitation of (Nb,Ti,V)(C,N) or (Nb,Ti,V)(Cr,Mo)(C,N) increases the grain boundary density of the microstructure and improves hole-expansion and bending performance.

In some preferred embodiments, the coiling temperature may be controlled between 350°C and 600°C to further regulate the transformation of bainite and quasi-polygonal ferrite, precipitation of micro alloys, and thus control the strength, hole-expansion and bending performance of the steel plate.

The hot-rolled high-strength steel having excellent hole-expansion and bending performance, and manufacturing method therefor of the present disclosure have the following advantages and beneficial effects:
The hot-rolled high-strength steel having excellent hole-expansion and bending performance according to the present disclosure adopts precise addition of Mg and Ca combined with a tailored manufacturing process to control the grain boundary density, the type and proportion of the microstructure, as well as the size and quantity of precipitates. Thus, the hot-rolled high-strength steel plates can be produced on a conventional hot continuous rolling line or a low-carbon short-process line.

The hot-rolled high-strength steel having excellent hole-expansion and bending performance according to the present disclosure has a yield strength of ≥660 MPa, a tensile strength of ≥780 MPa, an elongation A50 of ≥15%, a punched hole-expansion ratio of ≥75%, and a VDA bending angle of ≥120°. It can be used as automotive body structural parts and automotive chassis parts, and it can be used in other fields requiring high strength and weight reduction.

• Figure 1 shows an image showing the grain boundary density, according to Example 1 of the present disclosure.
• Figure 2 shows an image showing the grain boundary density, according to Comparative Example 3.
• Figure 3 shows the effect of grain boundary density on the hole-expansion ratio.
• Figure 4 shows the effect of grain boundary density on the bending performance.
• Figure 5 shows hole-expansion cracking caused by coarse MnS.
• Figure 6 shows (Nb,Ti)(C,N) precipitate particles with a relatively large quantity in Comparative Example 10.

The hot-rolled high-strength steel having excellent hole-expansion and bending performance according to the present disclosure will be further explained and illustrated below with reference to the accompanying drawings and specific examples. Such explanation and illustration shall not improperly limit the technical solution of the present disclosure.

The hot-rolled high-strength steels having excellent hole-expansion and bending performance in the examples of the present disclosure were produced by the following steps:

1. (1) Smelting and casting to obtain a slab:
   In some embodiments, conventional methods may be used for smelting. For example, converter steelmaking was adopted, and the liquid steel was subjected to RH vacuum degassing and LF furnace desulfurization. The superheat during steelmaking was controlled at 15°C to 55°C. Tables 1-1, 1-2 and 1-3 list the chemical element contents and synergistic relationships of Comparative Examples and Examples of the present disclosure.
2. (2) Heating the slab to a temperature ranging from 1200°C to 1300°C and holding for 1 to 3 h.
3. (3) Hot rolling: laminar cooling was performed at an average cooling rate of 2 30°C/s to a coiling temperature ranging from 300°C to 620°C, followed by coiling; after coiling, the steel was cooled to room temperature at a rate of ≤ 15°C/s.

In some embodiments, a stage cooling process was adopted for laminar cooling, where in the first stage, the steel was cooled to an intermediate point temperature of 610 to 750°C at an average cooling rate of ≥100°C/s, followed by air cooling for 4 s to 10.0 s; then in the second stage, the steel was cooled to the coiling temperature at an average cooling rate of ≥30°C/s.

In some embodiments, the rough rolling exit temperature was controlled at 1000°C to 1080°C; the final rolling temperature was controlled in a range from 840°C to 950°C, and the rolling speed was controlled at 7.0 to 13.0m/s.

In some embodiments, controlling was performed such that the total reduction rate was ≥80%, the total reduction rate for finish rolling was ≥50%, and the reduction rate for single-pass in final rolling was ≤15%.

The tension leveling elongation during pickling was controlled at 0.2% to 1.8% and the pickling speed was 60 to 150 m/min. The temperature of the last pickling tank in the pickling process may be 80 to 90°C, and the iron ion concentration may be 30 to 40 g/L.

The hot-rolled high-strength steel plate obtained by the present disclosure may be further subjected to a galvanizing step after pickling, thereby obtaining products with various coatings such as an electro-galvanized coating, a hot-dip galvanized coating, a Zn-Al-Mg coating, and the like. Steel Grade C Si Mn S Al B Mg Ca P N O A 0.072 0.3 1.2 0.005 0.038 0.0005 0.0002 0.01 0.011 0.0045 0.0005 B 0.08 1 1.45 0.0011 0.02 0.0002 0.006 0.0008 0.013 0.0042 0.004 C 0.064 0.65 1.45 0.005 0.072 0.0004 0.01 0.0002 0.003 0.004 0.001 D 0.053 1.2 1.5 0.0031 1.0 0.0010 0.006 0 0.014 0.0023 0.0026 E 0.042 0.01 1.8 0.0023 0.1 0.0035 0 0.008 0.011 0.005 0.0005 F 0.03 0.1 1.6 0.0005 0.35 0 0.009 0 0.018 0.0025 0.0012 G 0.06 0.05 1.35 0.0013 0.068 0.0008 0.0003 0.0035 0.02 0.0025 0.001 H 0.045 0.8 1.5 0.0009 0.25 0 0.005 0 0.011 0.0045 0.0026 I 0.064 0.65 1.45 0.01 0.072 0.0004 0.01 0.0002 0.003 0.004 0.001 J 0.064 0.65 1.45 0.005 0.072 0.0004 0.001 0.0001 0.003 0.004 0.001 K 0.064 0.65 1.45 0.005 0.072 0.0004 0 0.0002 0.003 0.004 0.005 L 0.09 0.65 1.45 0.005 0.072 0.0004 0.001 0.0002 0.003 0.004 0.001 M 0.064 0.65 1.1 0.005 0.072 0.0004 0 0.0002 0.003 0.004 0.001 N 0.02 0.65 1.45 0.005 0.072 0.0004 0.001 0.0002 0.003 0.004 0.001 Steel Grade Ti Nb V Cr Mo A 0.13 0 0 0.7 0 B 0 0 0.2 0.38 0.05 C 0.07 0.05 0 0.4 0 D 0.05 0 0.05 0.05 0.05 E 0 0.06 0 0 0.08 F 0.1 0 0 0 0.25 G 0.1 0 0.06 0 0.13 H 0 0.01 0.1 0.55 0 I 0.07 0.03 0 0.4 0 J 0.07 0.03 0 0.4 0 K 0.07 0.03 0 0.4 0 L 0.14 0 0 0.4 0 M 0.07 0.03 0 0.4 0 N 0.07 0.03 0 0.4 0 Table 1-3 (wt%, balance being Fe and other inevitable impurities excluding P, O and N) Steel Grade Mg/24+Ca/40 O/16×0.83 S/32 Mg/24+Ca/40-O/16×0.83 S/32×Mn/55 4×C/(3.3Nb+3.4V+Ti) A 0.00026 0.00003 0.00016 0.00023 3.41×10^-6 2.22 B 0.00027 0.00021 0.00003 0.00006 9.06×10^-7 0.53 C 0.00042 0.00005 0.00016 0.00037 4.12×10^-6 1.09 D 0.00025 0.00013 0.00010 0.00012 2.64×10^-6 1.06 E 0.00020 0.00003 0.00007 0.00017 2.35×10^-6 0.85 F 0.00038 0.00006 0.00002 0.00031 4.55×10^-7 1.20 G 0.00010 0.00005 0.00004 0.00005 9.97×10⁷ 0.86 H 0.00021 0.00013 0.00003 0.00007 7.67×10^-7 0.54 I 0.00042 0.00005 0.00031 0.00037 8.24×10^-6 1.51 J 0.00004 0.00005 0.00016 -0.00001 4.12×10^-6 1.51 K 0.00001 0.00026 0.00016 -0.00025 4.12×10^-6 1.51 L 0.00005 0.00005 0.00016 -0.00001 4.12×10^-6 2.57 M 0.00001 0.00005 0.00016 -0.00005 3.13×10^-6 1.51 N 0.00005 0.00005 0.00016 -0.00001 4.12×10^-6 0.47

Note: Steel grades A-H in Table 1-1, Table 1-2 and Table 1-3 are those used in Examples of the present disclosure, and steel grades I-N are those used in Comparative Examples. In the formulas shown in Table 1-3, each chemical element is substituted with the numerical value before the percent sign of the mass percentage content.

Tables 2-1 and 2-2 list the specific process parameters in the above process steps for Examples and Comparative Examples of the present disclosure. No. Steel Grade Steel making Hot Rolling Super Heat (°C) Heating Temperature (°C) and holding Time(h) Rough Rolling Exit Temperature (°C) Final Rolling Temperature (°C) Total Reduc tion (%) Total Reduction For Finish Rolling (%) Reduction For Single-pass In Final Rolling (%) Ex.1 A 35 1235×2.0 1055 915 93 86 10 Ex.2 A 15 1220×2.0 1080 915 93 86 10 Ex.3 A 35 1280×2.0 1070 915 93 86 15 Ex.4 A 55 1255×3.0 1055 880 93 80 5 Ex.5 B 35 1260×1.0 1055 840 93 86 2 Ex.6 B 20 1255×2.5 1030 950 88 75 8 Ex.7 B 35 1255×1.5 1055 915 93 86 6 Ex.8 C 30 1220×2.2 1060 950 93 86 8 Ex.9 C 30 1255×2.0 1055 950 80 50 2 Ex.10 D 25 1300×1.8 1080 915 85 60 10 Ex.11 E 42 1250×2.0 1000 860 92 80 7 Ex.12 F 50 1200×3.0 1010 915 98 86 10 Ex.13 G 36 1250×1.3 1055 930 93 85 5 Ex.14 H 30 1250×2.8 1050 915 95 86 5 Comp. Ex. 1 A 35 1255×2.0 1055 915 93 86 10 Comp. Ex.2 A 35 1255×2.0 1055 915 93 86 10 Comp. Ex.3 A 35 1255×2.0 1055 915 93 86 10 Comp. Ex.4 B 35 1255×2.0 1055 915 93 86 10 Comp. Ex.5 B 35 1255×2.0 930 780 93 86 10 Comp. Ex.6 B 35 1100×2.0 920 915 93 86 10 Comp. Ex.7 C 10 1255×2.0 1055 915 93 86 10 Comp. Ex.8 C 80 1255×2.0 1055 915 93 86 10 Comp. Ex.9 I 30 1250×2.0 1050 910 93 86 10 Comp. Ex.10 J 30 1255×2.0 1055 915 93 86 10 Comp. Ex.11 K 30 1255×2.0 1055 915 93 86 10 Comp. Ex.12 L 30 1255×2.0 1055 915 93 86 10 Comp. Ex.13 M 30 1255×2.0 1055 915 93 86 10 Comp. Ex.14 N 30 1255×2.0 1055 915 93 86 10 No. Steel Grade Laminar cooling Rolling Speed (m/s) Coiling Tempera ture (°C) Cooling Rate After Coiling (°C/s) First Stage Cooling Rate (°C/s) Intermediate Point Temperature (°C) Air Cooling Time After Interme diate Point (s) Second Stage Cooling Rate (°C/s) Ex.1 A 140 685 8.0 80 8.2 480 5 Ex.2 A 100 750 5.0 30 6.7 380 5 Ex.3 A 180 655 10.0 120 7.8 480 5 Ex.4 A 200 610 9.8 40 7.0 580 8 Ex.5 B 176 685 5.1 140 13.0 300 5 Ex.6 B 140 685 9.0 90 7.9 520 10 Ex.7 B 140 650 8.0 90 9.0 550 5 Ex.8 C 140 685 8.0 80 8.2 430 5 Ex.9 C direct single-stage cooling, cooling rate100 10 430 5 Ex.10 D 140 660 8.0 80 8.0 450 15 Ex.11 E 140 685 8.0 80 8.4 480 5 Ex.12 F 140 630 8.0 80 8.4 600 5 Ex.13 G 140 685 8.0 80 7.8 620 7 Ex.14 H 140 685 8.0 80 8.2 350 5 Comp.Ex.1 A 140 685 8.0 80 8.2 480 20 Comp.Ex.2 A 140 685 8.0 80 7.3 280 5 Comp.Ex.3 A 140 685 8.0 80 9.5 650 5 Comp.Ex.4 B 185 780 12.0 70 6.7 400 5 Comp.Ex.5 B 95 685 8.0 80 8.4 520 5 Comp.Ex.6 B 140 685 8.0 80 8.2 520 5 Comp.Ex.7 C 140 685 8.0 80 8.2 430 5 Comp.Ex.8 C 140 685 8.0 80 8.2 430 5 Comp.Ex.9 I 140 685 8.0 80 8.2 480 5 Comp.Ex.10 J 140 685 8.0 80 8.2 480 5 Comp.Ex.11 K 140 685 8.0 80 8.2 480 5 Comp.Ex.12 L 140 685 8.0 80 8.2 580 5 Comp.Ex.13 M 140 685 8.0 80 8.2 450 5 Comp.Ex.14 N 140 685 8.0 80 8.2 600 5

Samples were taken from the hot-rolled high-strength steels having excellent hole-expansion and bending performance of Examples (Ex.) 1-14 and Comparative Examples (Comp. Ex.) 1-14, respectively. The grain boundary density of the materials was analyzed with EBSD and TEM. The EBSD test had the following parameters: accelerating voltage: 20 kV, electron beam current: 18 µA, working distance: 16 mm, step size: 0.1 µm, and tilt angle: 70°.After the EBSD test, the crystallographic data obtained were processed using HKL Channel 5 software to analyze the grain boundary density distribution of the samples. The test results are listed in Table 3.

In addition, the precipitated phase in the sample were electrolyzed out or taken out through extraction replica, and then were taken a photograph and subjected to single-point energy-dispersive spectroscopy analysis with transmission electron microscopy. Specifically, for the physicochemical phase analysis, a mixed aqueous solution of 5% potassium chloride and 1% citric acid was used, with the current density controlled at 0.05 A/cm².The morphology, size, composition and other characteristics of precipitated phase in the specimens prepared with different processes were characterized and analyzed by transmission electron microscopy to determine the type and quantity of precipitates. The test results are listed in Table 3. No. Steel Grade Microstructure Grain Boundary Density (µm^-1) Bainite + Quasi-polygonal Ferrite Polygonal Ferrite Martensite Precipitated Particles such as MgS, CaS and TiN Area Fraction (%) Grain Size (µm) Area Fraction (%) Grain Size (µm) Area Fraction (%) Grain Size (µm) Particle Size (µm) Propor tion of Particles ≤ 3.0 µm (%) Ex.1 A 1.1 90 2.8 8 3.5 2 1.5 ≤8 70 Ex.2 A 0.7 94 3.5 3 4.7 3 1.5 ≤8 70 Ex.3 A 1.1 90 2.5 8 4.5 2 1.0 ≤8 70 Ex.4 A 1.8 75 2.0 25 4.9 - - ≤8 70 Ex.5 B 0.6 95 2.8 3 4.5 2 2.0 ≤3 75 Ex.6 B 1.4 90 3.2 9 4.7 1 1.5 ≤3 75 Ex.7 B 1.6 88 4.0 10 4.9 2 2.5 ≤3 75 Ex.8 C 0.9 92 3.8 5 4.5 3 1.5 ≤6 70 Ex.9 C 0.9 92 3.8 3 4.5 5 1.5 ≤6 70 Ex.10 D 1.0 91 3.5 6 4.5 3 1.5 ≤7 60 Ex.11 E 1.1 90 2.5 8 4.5 2 2.5 ≤8 65 Ex.12 F 1.9 72 4.0 28 4.9 - - ≤8 60 Ex.13 G 2.0 70 3.5 30 5.0 - - ≤6 70 Ex.14 H 0.8 95 3.5 3 4.6 2 1.5 ≤7 65 Comp. Ex.1 A 1.1 84 3.5 3 4.5 13 3.0 ≤8 70 Comp. Ex.2 A 0.4 28 3.0 - - 72 6.0 ≤8 70 Comp. Ex.3 A 2.1 52 5.5 48 6.8 - - ≤8 70 Comp. Ex.4 B 1.4 92 3.0 2 4.7 4 1.5 ≤8 75 Comp. Ex.5 B 2.1 72 3.0 35 4.0 3 1.5 ≤8 75 Comp. Ex.6 B 1.4 92 2.8 5 3.8 3 1.5 ≤8 40 Comp. Ex.7 C 0.9 92 3.8 5 4.5 3 1.5 12 40 Comp. Ex.8 C 0.9 92 3.8 5 4.5 3 1.5 9 45 Comp. Ex.9 I 1.1 90 2.5 8 4.5 2 1.5 12 30 Comp. Ex.10 J 1.1 90 2.5 8 4.5 2 1.5 9 40 Comp. Ex.11 K 1.1 90 2.5 8 4.5 2 1.5 11 35 Comp. Ex.12 L 1.8 75 2.0 25 4.9 - - 14 45 Comp. Ex.13 M 1.0 91 3.5 6 4.5 3 1.5 ≤8 70 Comp. Ex.14 N 1.9 72 4.0 28 4.9 - - ≤8 70

In addition, mechanical properties of the specimens in each Example and Comparative Example were tested in accordance with GB/T 228.1-2010. Yield strength and tensile strength were measured using JIS 5# tensile specimens taken along the longitudinal direction. The hole-expansion ratio was tested according to the method specified in GB/T 24524-2021: the specimen with a central hole was pressed into a die by a punch to expand the central hole of the specimen until the edge of the hole exhibited necking or through cracks. Bending performance was tested in accordance with VDA 238-100, and evaluated by the bending angle. No. Steel Grade Mechanical Properties Yield Strength (MPa) Tensile Strength (MPa) A50 (%) Hole-expansion ratio (%) VDA Bending Angle (°) Ex.1 A 702 832 18.5 91 135 Ex.2 A 743 828 16.0 75 128 Ex.3 A 713 822 17.5 105 138 Ex.4 A 685 846 21.5 85 152 Ex.5 B 755 829 15.0 79 135 Ex.6 B 709 843 19.5 88 138 Ex.7 B 682 826 20.5 102 150 Ex.8 C 718 868 18.0 95 141 Ex.9 C 738 878 16.0 85 128 Ex.10 D 741 892 16.5 75 129 Ex.11 E 728 879 18.5 85 131 Ex.12 F 693 836 22.5 103 163 Ex.13 G 708 837 24.5 112 165 Ex.14 H 742 838 16.0 75 125 Comp.Ex.1 A 758 889 15.5 58 110 Comp.Ex.2 A 809 932 13.5 45 115 Comp.Ex.3 A 658 753 25.5 68 112 Comp.Ex.4 B 691 815 14.5 75 128 Comp.Ex.5 B 645 753 24.0 125 160 Comp.Ex.6 B 657 756 22.5 98 154 Comp.Ex.7 C 695 843 17.5 62 119 Comp.Ex.8 C 705 851 17.5 55 119 Comp.Ex.9 I 719 836 16.5 45 101 Comp.Ex.10 J 720 843 16.5 48 116 Comp.Ex.11 K 709 832 15.5 43 106 Comp.Ex.12 L 785 936 18.5 38 106 Comp.Ex.13 M 643 752 18.5 76 129 Comp.Ex.14 N 623 736 22.5 103 163

As shown in Tables 3 and 4, through reasonable chemical composition design combined with optimized process parameters, Examples 1-14 of the present disclosure achieved ideal microstructural characteristics, leading to steel plates with excellent performance. The precipitates in the steel plates of Examples were MgO, CaO, MgS, CaS, TiN and (Nb,Ti)N. All the Examples of the present disclosure exhibited a yield strength of greater than 660 MPa, a tensile strength of greater than 780 MPa, an elongation A50 of ≥ 15%, a punched hole-expansion ratio of ≥ 75%, and a VDA bending angle of greater than 120°. The precipitates had a particle size of ≤8 µm. The proportion of precipitates with particle size of ≤3.0 µm was ≥50%.

In addition, as shown in Table 3, the grain boundary density of each Example of the present disclosure is in the range of 0.6µm ^-1 to 2.0 µm^-1. In addition, Figure 1 shows an image showing the grain boundary density of Example 1 of the present disclosure. Figure 2 shows an image showing the grain boundary density of Comparative Example 3.

The grain boundary density of the hot-rolled high-strength steel plate according to the present disclosure has a direct influence on the punched hole expansion performance and bending performance.

Figure 3 shows the effect of grain boundary density on the hole-expansion ratio. As shown in Figure 3, when the grain boundary density is in the range of 0.6 µm^-1 to 2.0 µm^-1, the hole-expansion ratio increases with increasing grain boundary density.

Figure 4 shows the effect of grain boundary density on the bending performance. As shown in Figure 4, when the grain boundary density is in the range of 0.6 µm^-1 to 2.0 µm^-1, the bending angle increases with increasing grain boundary density.

In contrast to the present disclosure, each Comparative Example failed to achieve the technical effects of the present disclosure due to the composition or process parameters unsatisfying the design requirements of the present disclosure, as detailed below:
In Comparative Example 1, the cooling rate after hot rolling and coiling was excessively high, leading to an excessive proportion of martensitic transformation from supercooled austenite after coiling. The excessively high martensite grain fraction resulted in low hole-expansion ratio and poor bending performance.

Comparative Example 2 adopted a relatively low coiling temperature, leading to an excessively high primary martensite content and excessively large martensite grain size in the microstructure, as well as excessively high grain boundary density in the microstructure, which ultimately resulted in low hole-expansion ratio, poor bending performance and low elongation.

Comparative Example 3 adopted a relatively high coiling temperature, leading to a high ferrite content and coarse grains in the microstructure. The area fraction of polygonal ferrite reached 48% with a grain size of 6.8 µm, and the bainite had a grain size of 5.5 µm, which was relatively large. As shown in Figure 2, the grain boundary density of the microstructure was only 0.5 µm^-1, resulting in insufficient strength, hole-expansion ratio and bending performance.

In Comparative Example 4, the intermediate point temperature was relatively high, leading to insufficient supercooling in the air-cooling stage during laminar cooling and absence of ferrite transformation, finally resulting in relatively low elongation of 14.5%.

Comparative Example 5 adopted relatively low rough rolling exit temperature and finishing delivery temperature, leading to polygonal ferrite transformation with a large size and high area fraction during rolling, and excessively low grain boundary density, finally resulting in low yield strength and tensile strength.

In Comparative Example 6, the heating temperature was low, leading to insufficient solid solution of V. Coarse V(C,N) particles formed during continuous casting were not fully dissolved, contributing little to strength and resulting in insufficient tensile strength of the steel plate.

In Comparative Examples 7 and 8, the superheat was not controlled within the specified range. Low superheat was unfavorable for precipitated particles such as MgS, CaS and TiN to float, and thus detrimental to inclusion removal. Excessively high superheat promoted the growth of precipitated particles such as MgS, CaS and TiN, finally leading to large particle sizes of MgS, CaS, TiN and other precipitates in the steel and relatively high proportions of particles ≤ 3.0 µm. The particle sizes of Comparative Examples 7 and 8 reached 12 µm and 9 µm, respectively, and the proportions of particles ≤ 3.0 µm were only 40% and 45% respectively, thus exhibiting low hole-expansion ratio and poor bending performance.

In Comparative Example 9, high S content resulted in coarse MnS with a size up to 12 µm in the final microstructure, and the proportion of particles ≤ 3.0 µm was only 40%, which significantly affected the final hole-expansion ratio and bending performance, resulting in insufficient hole-expansion ratio and bending performance. Figure 5 showed hole expansion cracking caused by coarse MnS, and MnS also exhibited brittle fracture during hole expansion.

In Comparative Example 10, low contents of Mg and Ca failed to refine precipitates, resulting in (Nb,Ti)(C,N) or (Nb,Ti,V)(Cr)(C,N) particles having sizes up to 9 µm in the final precipitates. The sizes were relatively high. The proportion of particles ≤ 3.0 µm was only 40%, which greatly impaired the final hole-expansion ratio and bending performance, which were only 48% and 116°, respectively. Figure 6 showed (Nb,Ti)(C,N) precipitate particles with a relatively large quantity in Comparative Example 10.

In Comparative Example 11, the excessively high O content led to excessive CaO in the final precipitates. CaO had a large size and could not function to refine precipitates. In addition, the high O content resulted in the formation of Al₂O₃ which readily acted as nucleation sites for TiN and promoted TiN growth. Consequently, the size of (Nb,Ti)(C,N) or (Nb,Ti,V)(Cr)(C,N) particles in the final precipitates reached 11 µm, and the proportion of particles ≤ 3.0 µm was only 35%, seriously affecting the final hole-expansion ratio and bending performance and leading to insufficient hole-expansion ratio and bending performance.

In Comparative Example 12, addition of C and Ti at high amounts led to large-sized and high-quantity of (Nb,Ti)(C,N) or (Nb,Ti,V)(Cr)(C,N) on one hand, and relatively high strength on the other hand, both of which significantly affected the final hole-expansion ratio and bending performance, resulting in insufficient hole-expansion ratio and bending performance.

In Comparative Example 13, addition of Mn at a low amount resulted in low yield strength and tensile strength.

In Comparative Example 14, addition of C at a low amount resulted in low yield strength and tensile strength.

It is noted that the combinations of technical features in the present application are not limited to those described in the claims or the specific examples. All technical features described in the present application can be freely combined or integrated in any manner unless they are mutually exclusive.

It is also noted that the examples listed above are merely specific embodiments of the present disclosure. Obviously, the present disclosure is not limited to the above examples. Similar changes or modifications derived directly or easily associated by those skilled in the art from the disclosure of the present disclosure shall all fall within the protection scope of the present disclosure.