1. A high-hole expansion performance DH1180MPa cold-rolled continuous annealing steel plate is characterized in that the steel plate comprises the following components in percentage by weight: c: 0.14% -0.16%, Si: 0.8% -1.2%, Mn: 2.2-3.0%, P is less than or equal to 0.015%, S is less than or equal to 0.005%, Ti/Nb: 0.02-0.03%, and satisfies that C + Si/4 is less than or equal to 0.4%, and the balance is Fe and inevitable impurities.

2. The high hole expansibility DH1180MPa cold-rolled continuous annealing steel plate according to claim 1, wherein the steel plate microstructure is epitactic ferrite + tempered martensite + bainite + retained austenite; wherein the steel plate structure comprises the following components in percentage by volume: ferrite: less than or equal to 10 percent, tempered martensite 70 to 80 percent, bainite 3.5 to 11 percent and residual austenite 5 to 8.5 percent.

3. The cold-rolled continuous-annealed steel plate with high hole expansibility DH1180MPa according to claim 1, wherein the steel plate has yield strength of 850-1050 MPa, tensile strength of 1180MPa or more, elongation after fracture of 15% or more, and hole expansibility of 25% or more.

4. A method for preparing a DH1180MPa cold-rolled continuous-annealed steel plate with high hole expansibility according to any one of claims 1 to 3, comprising smelting, casting, hot rolling, pickling, cold rolling, and continuous annealing; the method is characterized in that:

(1) hot rolling: the heating temperature is 1200-1280 ℃, and the heat preservation time is more than 2 hours; the initial rolling temperature is 1050-1150 ℃, and the final rolling temperature is above 900 ℃; the coiling temperature is 550-600 ℃;

(2) cold rolling: the cold rolling reduction rate is 50-60%;

(3) and (3) continuous annealing:

heating: heating to A at 10 deg.C/s₃+ (10-20) DEG C, isothermal time is 100-220 s;

secondly, slow cooling: slowly cooling at 680-720 ℃, and controlling the slow cooling speed at 0.5-5 ℃/s;

③ quick cooling: after slow cooling, the steel plate is rapidly cooled to the quenching temperature T at the cooling speed of more than 30 ℃/s_Q-(10～20)℃；

Fourthly, overaging: then heating to Ms + (10-20) DEG C-420 ℃ at a heating speed of more than 10 ℃/s for overaging, wherein the overaging isothermal time is 120-400 s;

and fifthly, cooling to room temperature at a cooling speed of 1-3 ℃/s.

5. The method for preparing the DH1180MPa cold-rolled continuous-annealed steel plate with high hole expansion performance according to claim 4, wherein the method comprises the following steps: a is described₃：820～870℃，Ms：335～380℃，T_Q：205～260℃。

6. The method for preparing the DH1180MPa cold-rolled continuous-annealed steel plate with high hole expansion performance according to claim 4, wherein the method comprises the following steps: and (3) in the step (3) and the step (II), the volume percentage of ferrite in the microstructure of the steel plate after slow cooling is less than 10%.

7. The method for preparing the DH1180MPa cold-rolled continuous-annealed steel plate with high hole expansion performance according to claim 4, wherein the method comprises the following steps: in the step (3), the volume percentage of the retained austenite in the microstructure of the steel plate after rapid cooling is 5-8.5%.

8. The method for preparing the DH1180MPa cold-rolled continuous-annealed steel plate with high hole expansion performance according to claim 4, wherein the method comprises the following steps: in the steps (3) to (4), the volume percentage of bainite in the microscopic structure of the steel plate after overaging is 3.5 to 11 percent.

Background

In recent years, energy crisis and environmental problems continuously promote the automobile industry to develop towards the direction of energy conservation and emission reduction, and the process of automobile light weight is accelerated. The high strength of automobiles as main materials of automobile bodies is advanced and the aim of 'high strength and thinning' is continuously strived. However, the incompatibility of strength and plasticity has been an insurmountable "gap" in the field of materials. Taking dual-phase steel as an example, when the strength reaches above 980MPa and 1180MPa, the plasticity is reduced to about 10%, and a complicated or even more complicated vehicle body structural member is difficult to finish forming under a cold deformation condition, and a part scheme is often realized by means of roll forming. Based on the problem, related studies in europe propose the concept of DH steel, i.e., plasticity-enhanced dual-phase steel, in which a small amount of bainite and retained austenite are introduced to improve the plasticity of the steel sheet. However, although the introduction of the complex phase improves the forming performance of the steel plate, the addition of the soft phase and the uneven strain distribution under the complex phase lead to the reduction of the hole expanding performance of the DH steel, which leads to the failure of the DH steel to be qualified for automobile structural members with high flanging performance requirements, such as engine side rails, seat side rails, longitudinal beams and other structural members with high flanging requirements. Meanwhile, even the plasticity of the DH steel with strengthened plasticity reaches about 15% when the strength reaches 980MPa or 1180MPa, and the DH steel cannot be used as the drawing performance structural member for various vehicle heights. Therefore, based on better drawing performance of the DH steel relative to DP steel, the flanging performance of the DH steel is improved, namely the improvement of the hole expanding performance of the DH steel is of great importance to the application of the high-strength DH steel on the automobile body.

A dual-phase steel disclosed in patent document 'a cold-rolled dual-phase steel with 1200MPa of tensile strength and a production method thereof' (publication No.: CN 108642379B) comprises the following main chemical components: c: 0.09% -0.13%, Si: 0.1-0.4%, Mn: 2.0-2.6%, P is less than or equal to 0.01%, S is less than or equal to 0.01%, Al: 0.02% -0.06%, Cr: 0.3-0.6%, Mo: 0.1 to 0.3%, Nb: 0.01 to 0.04%, Ti: 0.01% -0.04%, B: 0.001 to 0.003%, and the balance of Fe and unavoidable impurities. The obtained steel plate has the yield strength of 820MPa, the tensile strength of 1200MPa and the elongation (A80) of more than or equal to 6 percent. The product is added with comparatively expensive alloys such as Cr, Mo, Nb and the like, so that the alloy cost of the steel is improved; in addition, the elongation of 6% is difficult to realize the cold stamping forming of parts, and the use difficulty is large.

Patent document "a cold-rolled dual-phase steel with 1200 MPa-grade tensile strength and a preparation method thereof" (publication number: CN102021483B) discloses dual-phase steel which comprises the following main chemical components: c: 0.19 to 0.21%, Si: 0.7-0.9%, Mn: 1.9% -2.1%, Cr: 0.01 to 0.02%, Nb: 0.02-0.04%, P is less than or equal to 0.005%, S is less than or equal to 0.003%, and the balance is Fe and inevitable impurities. The product is a cold-rolled annealed steel plate, the steel is 1200 MPa-grade steel, the elongation is 8% -10%, and the yield strength is 580-660 MPa. Lower yield strength is indicative of poor hole expansion performance. Meanwhile, the plasticity of the steel is maintained in a lower range, and the drawing requirement of a more complex structural part is difficult to realize.

The patent document '1200 MPa grade ultra-fast cold-rolled dual-phase steel plate and preparation method thereof' (publication number: CN109280857A) discloses dual-phase steel which comprises the following main chemical components: c: 0.12% -0.17%, Si: 0.3% -0.6%, Mn: 2.0-2.4%, P is less than or equal to 0.015%, S is less than or equal to 0.008%, Als: 0.03-0.06%, Ti: 0.03-0.06%, N is less than or equal to 0.005%, and the balance is Fe and inevitable impurities. The steel has the following properties: yield strength 820-950 MPa, tensile strength: 1200-1350 MPa and 5-10% of elongation. The steel preparation method involves ultra-fast cooling at a rate of 130-150 ℃/s, which is difficult to realize in the actual production process, and the product drawing performance is difficult to meet the requirement of complex part forming.

Disclosure of Invention

The invention aims to overcome the problems and the defects and provides a DH1180MPa cold-rolled continuous annealing steel plate with high hole expansion performance, wherein the yield strength is 850-1050 MPa, the tensile strength is more than 1180MPa, the elongation after fracture is more than 15%, and the hole expansion rate is more than 25%, and a preparation method thereof.

The purpose of the invention is realized as follows:

a high-hole expansion performance DH1180MPa cold-rolled continuous annealing steel plate comprises the following components in percentage by weight: c: 0.14% -0.16%, Si: 0.8% -1.2%, Mn: 2.2-3.0%, P is less than or equal to 0.015%, S is less than or equal to 0.005%, Ti/Nb: 0.02-0.03%, and satisfies that C + Si/4 is less than or equal to 0.4%, and the balance is Fe and inevitable impurities.

The microstructure of the steel plate is ferrite, tempered martensite, bainite and residual austenite; wherein the steel plate structure comprises the following components in percentage by volume: ferrite: less than or equal to 10 percent, tempered martensite 70 to 80 percent, bainite 3.5 to 11 percent and residual austenite 5 to 8.5 percent. The invention introduces partial residual austenite on the precipitation of DP steel to improve the plasticity of the steel plate.

The yield strength of the steel plate is 850-1050 MPa, the tensile strength is more than 1180MPa, the elongation after fracture is more than 15%, and the hole expansion rate is more than 25%.

The reason for the alloy design of the present invention is as follows:

c: c is one of important elements in the invention. The austenitizing starting temperature (A) of the steel is secured₁) And the complete austenitizing temperature (A)₃) Thereby ensuring the austenite stabilization behavior in the steel sheet boundary region. In addition, the diffusion behavior of C atoms in the heat preservation process in the alloying galvanization stage and the diffusion behavior in the adjustment stage before galvanization promote the stability enhancement of residual austenite at room temperature, so that the performance of the TRIP effect in the deformation process of the steel plate is ensured, and the effects of strengthening the steel plate and improving the plasticity are achieved. However, too high C addition will result in a reduction in the weldability of the steel sheet. Therefore, the content of the C element is controlled to be 0.14-0.16 percent in the invention.

Si: si is one of the important elements in the present invention. In the present invention, Si mainly acts to suppress the precipitation of cementite at the overaging stage. However, adding too much Si reduces the surface quality of the steel. Therefore, the content of the Si element is controlled to be 0.8-1.2%.

Mn: mn is one of the important elements in the present invention. The amount of Mn element affects the hardenability of the super-cooled austenite in the cooling stage, delays the pearlite transformation in the cooling process and ensures the effective martensite phase transformation under a certain cooling speed condition. Meanwhile, the change of the austenitizing temperature is also determined by the addition of Mn element, the increase of higher Mn content greatly reduces the full-austenitizing temperature point, which causes overhigh austenitizing temperature and coarse grains at the common continuous annealing temperature; while lowering the austenitizing temperature will not match the annealing temperature. Therefore, the Mn content is controlled to be 2.2-3.0%.

Ti: ti can capture free N atoms in the steel and plays a role in fixing N. Meanwhile, TiN can be precipitated in the solidification process to play a role in pinning a crystal boundary, and the Ti (C, N) is precipitated in the hot rolling stage to play a role in pinning a prior austenite crystal boundary and refining the prior austenite crystal grain. Therefore, the content of Ti element is controlled to be 0.02-0.03 percent.

Nb: when the Ti element is not added in the present invention, Nb will be an essential additive element. The addition of Nb is formed in the prior austenite grains in a mode of strain-induced precipitation in the hot rolling stage, and plays a role in pinning the prior austenite grain boundary and refining the prior austenite grains.

The second technical scheme of the invention is to provide a preparation method of a DH1180MPa cold-rolled continuous annealing steel plate with high hole expansion performance, which comprises smelting, casting, hot rolling, acid washing, cold rolling and continuous annealing;

(1) smelting:

and smelting by a converter to obtain the alloy components within the range.

(2) Hot rolling:

hot rolling: the heating temperature is 1200-1280 ℃, and the heat preservation time is more than 2 hours; the initial rolling temperature is 1050-1150 ℃, and the final rolling temperature is above 900 ℃; the coiling temperature is 550-600 ℃;

the heating temperature is 1200-1280 ℃, and the heat preservation time is more than 2 hours: the alloy elements in the steel are homogenized, and when the Ti element is added into the steel, the activity of Ti atoms is improved, free N in the steel is captured, and the effect of fixing N is achieved; meanwhile, Ti (C, N) is promoted to be precipitated, and the functions of pinning the original austenite grain boundary and refining the original austenite grain are achieved.

The initial rolling temperature is 1050-1150 ℃, and the final rolling temperature is above 900 ℃: ensuring the rolling pass of a recrystallization zone, promoting the dynamic recrystallization behavior of prior austenite grains in a hot rolling stage, and refining the grains; meanwhile, when Nb is added into the steel, the strain-induced precipitation behavior in a recrystallization rolling interval is enhanced, the precipitation of Nb (C, N) is promoted, and the original austenite grain boundary is pinned.

The coiling temperature is 550-600 ℃, so that the coil collapse phenomenon caused by overhigh coiling temperature is prevented, the bainite formation caused by overlow coiling temperature is prevented, and the cold rolling difficulty is increased. The thickness of the hot rolled coil is between 2.8 and 3.5 mm.

(3) Acid washing: and removing the scale generated on the surface of the hot rolling, and ensuring the surface quality of the cold-rolled steel plate.

(4) Cold rolling: the cold rolling reduction rate is 50% -60%, the rolling reduction of more than 50% of cold rolling is ensured, and the tissue fibrosis in the cold rolling configuration is promoted; meanwhile, the cold rolling reduction rate is prevented from being too high, so that the deformation resistance is too large, and the target thickness is difficult to roll.

(5) Continuous annealing:

heating: heating to A at 10 deg.C/s₃+ (10-20) DEG C, isothermal time is 100-220 s; a is described₃：820～870℃；

The continuous annealing isothermal temperature is A₃The temperature is +/-10-20 ℃, so that the full austenitizing behavior of steel with different alloy components in the heat preservation stage is ensured, and the phenomenon that the ferrite in the critical region is not dissolved back in the heating process is prevented; meanwhile, the phenomenon that the whole performance of the steel plate is influenced by coarsening of austenite grains caused by overhigh full-austenite transformation temperature is prevented. The isothermal time is selected to prevent the phenomenon that the yield strength of the steel plate is reduced due to residual critical zone ferrite caused by insufficient austenite nucleation caused by too short isothermal time; meanwhile, the method also prevents the austenite grains from coarsening due to overlong full-austenite transformation time.

Secondly, slow cooling: slowly cooling at 680-720 ℃, and controlling the slow cooling speed at 0.5-5 ℃/s; the volume percentage of the oriented periphytic ferrite in the microstructure of the steel plate after slow cooling is less than 10 percent.

The phenomenon that the precipitation of the oriented ferrite in austenite is large exists in the cooling process, the yield strength of the steel plate is obviously reduced due to the precipitation of the ferrite, so the slow cooling temperature is controlled to be more than 700 ℃ to prevent the formation of excessive oriented ferrite, and the content of the oriented ferrite is controlled to be less than 10%.

③ quick cooling: after slow cooling, the temperature is higher thanCooling to T at a cooling rate of 30 ℃/s_Q- (10-20) DEG C; the volume percentage content of austenite in the microstructure of the steel plate after rapid cooling is less than or equal to 5 percent; t is_Q：205～260℃； T_QThe optimal quenching temperature corresponding to the alloy composition;

f_α＝1-exp{β(M_s-T_Q) Formula 1

The Ms point formula is Ms-545-423C-30.4 Mn-7.5Si +30Al-60.5V_γ ^-1/3Equation 2

Wherein f is_αIs the martensite phase transformation amount; β is a constant, for carbon steel-0.011; t is_QIs the quenching temperature.

Wherein, T and R are absolute temperature and molar gas constant respectively; f. of_i ^γAndthe mole fraction and carbon content, respectively, of the unconverted austenite at the beginning of the partitioning process;andrespectively the mole fractions of austenite and martensite at the end of the partitioning process;andrespectively for the carbon contents of austenite and martensite at the end of the partitioning processThe carbon content of the alloy is the mole fraction.

T can be obtained by the above formulas 1 to 6_QAnd the optimal quenching temperature.

T_QThe optimum quenching temperature for the corresponding alloy composition at which the retained austenite content is maximized and stabilized is the goal of the present invention to obtain a higher strength matrix. Therefore, the quenching temperature is continuously reduced on the basis of obtaining the maximum retained austenite, and the austenite content is controlled to be about 5-8.5%.

Fourthly, overaging: then heating to Ms + (10-20) DEG C-420 ℃ at a heating speed of more than 10 ℃/s for overaging, wherein the overaging isothermal time is 120-400 s; the volume percentage of bainite in the microscopic structure of the steel plate after overaging is 3.5 to 11 percent; ms: at 335-380 ℃.

C diffusion in the tempering process is ensured, so that the continuous annealing overaging temperature is controlled to Ms + (10-20) -420 ℃; and simultaneously, the precipitation of tempered and softened carbides caused by excessively high overaging temperature is prevented, so that the temperature is not higher than 420 ℃. The overaging isothermal time is 120-400 s, so that the C diffusion effect is not sufficient due to too low isothermal time, and tempered martensite carbide precipitation due to too long isothermal time is prevented. It is worth noting that about 3.5% to 11% of bainite will be formed in the process.

And fifthly, cooling to room temperature at a cooling speed of 1-3 ℃/s.

The microstructure of the final steel plate is ferrite, tempered martensite, bainite and residual austenite; wherein the steel plate structure comprises the following components in percentage by volume: steel plate ferrite: less than or equal to 10 percent, tempered martensite 70 to 80 percent, bainite 3.5 to 11 percent and residual austenite 5 to 8.5 percent. . The invention introduces partial residual austenite on the precipitation of DP steel to improve the plasticity of the steel plate.

The product performance obtained by the method can meet the requirements of yield strength of 850-1050 MPa, tensile strength of more than 1180MPa, elongation after fracture of more than 15% and hole expansion rate of more than 25%. The invention has the beneficial effects that:

(1) the steel has the chemical components mainly comprising C, Si and Mn as main elements, is simple in alloy design and suitable for industrial production, and simultaneously controls the ratio of C + Si/4 to be less than or equal to 0.4, thereby being beneficial to laser welding and resistance spot welding in the production and application processes;

(2) the final structure of the invention is tempered martensite, oriented periphytic ferrite, residual austenite and a small amount of bainite, the process is adjusted, the traditional dual-phase steel and the ferrite in the middle critical area of DH steel are abandoned, the yield strength is improved by depending on the oriented periphytic ferrite obtained in the slow cooling and quick cooling processes, meanwhile, the residual austenite is introduced to improve the plasticity of the steel plate, and the comprehensive effects of 'yield improvement' and 'plasticity improvement' of 1180MPa grade high-strength steel are realized.

(3) In the field of 1180MPa automobile high-strength steel, the 15% plasticity and the 25% reaming value indicate that the high-strength steel can be used as a more complex automobile body structural part, and more possibility of 1180MPa high-strength steel application in an automobile body is provided.

Description of the drawings:

FIG. 1 is a gold phase diagram of a microstructure of example 1 of the present invention.

Detailed Description

The present invention is further illustrated by the following examples.

According to the embodiment of the invention, smelting, casting, hot rolling, acid washing, cold rolling and continuous annealing are carried out according to the component proportion of the technical scheme.

(1) Hot rolling: the heating temperature is 1200-1280 ℃, and the heat preservation time is more than 2 hours; the initial rolling temperature is 1050-1150 ℃, and the final rolling temperature is above 900 ℃; the coiling temperature is 550-600 ℃;

(2) cold rolling: the cold rolling reduction rate is 50-60%;

(3) continuous annealing:

heating: heating to A at 10 deg.C/s₃+(10～20)℃，The isothermal time is 100-220 s;

secondly, slow cooling: the slow cooling temperature is above 700 ℃, and the slow cooling speed is controlled to be 0.5-5 ℃/s;

③ quick cooling: after slow cooling, the mixture is rapidly cooled to T at a cooling speed of more than 30 ℃/s_Q-(10～20)℃；

Fourthly, overaging: then heating to Ms + (10-20) DEG C-420 ℃ at a heating speed of more than 10 ℃/s for overaging, wherein the overaging isothermal time is 120-400 s;

and fifthly, cooling to room temperature at a cooling speed of 1-3 ℃/s.

Further, the method comprises the following steps of; a is described₃：820～870℃，Ms：335～380℃，T_Q：205～260℃。

Further, the method comprises the following steps of; and (3) in the step (3), the volume percentage of the oriented epiferrite in the microstructure of the steel plate after slow cooling is less than 10%.

Further, the method comprises the following steps of; in the step (3), the volume percentage of austenite in the microstructure of the steel plate after rapid cooling is 5-8.5%.

Further, the method comprises the following steps of; in the steps (3) to (4), the volume percentage of bainite in the steel plate microstructure after overaging is about 3.5-11%.

The compositions of the steels of the examples of the invention are shown in table 1. The main process parameters of the hot rolling of the steel of the embodiment of the invention are shown in Table 2. The main process parameters of the continuous annealing of steel in the embodiment of the invention are shown in Table 3. The structure of the steel of the examples of the present invention is shown in Table 4. The mechanical properties of the steels of the examples of the invention are shown in Table 5.

TABLE 1 composition (wt%) of steels of examples of the present invention

	C	Mn	Si	Al	P	S	Ti	Nb	A3	Ms	TQ
												1	0.151	2.8	1.0	--	0.01	0.003	--	--	816	332	244
2	0.160	2.7	0.9	--	0.01	0.005	0.02	--	812	327	239
												3	0.147	2.8	1.2	--	0.009	0.003	-	0.02	824	351	262
4	0.155	2.5	0.8	0.3	0.01	0.005	0.01	--	841	339	266
												5	0.152	2.2	0.8	0.2	0.005	0.003	--	0.02	835	348	238
6	0.154	2.3	1.2	--	0.01	0.005	0.02	--	827	338	242
												7	0.152	2.2	1.1	--	0.01	0.005	--	0.01	823	335	253
8	0.158	2.5	1.1	--	0.01	0.005	0.02	-	848	358	245
												9	0.146	2.4	0.8		0.009	0.003	0.02	--	834	346	234
10	0.157	2.5	0.8	--	0.01	0.005	0.02	-	829	341	245
												11	0.143	2.4	0.7	--	0.005	0.003	-	0.02	833	346	236
12	0.152	2.2	0.7	0.5	0.009	0.003	--	-0.02	829	332	229
												13	0.156	2.2	1.1	--	0.008	0.005	0.02	--	835	336	248
14	0.151	2.2	1.0	0.4	0.01	0.003	0.01	--	848	352	231
												15	0.154	2.3	1.2	0.3	0.02	0.005	-	0.03	839	349	243

TABLE 2 Main Process parameters for hot rolling of steels according to examples of the invention

TABLE 3 Main Process parameters for continuous annealing of steel in the examples of the present invention

TABLE 4 Structure of inventive example steels

TABLE 5 mechanical Properties of steels of examples of the invention

Examples	Rp0.2/MPa	Rm/MPa	A50/％	λ/％
					1	858	1205	16.7	30.4
2	866	1214	15.8	29.5
					3	906	1209	15.2	25.6
4	872	1198	16.1	27.5
					5	876	1190	15.4	29.3
6	863	1205	15.9	28.7
					7	902	1208	16.3	30.2
8	881	1188	16.5	26.6
					9	856	1206	15.4	28.5
10	903	1211	16.2	26.6
					11	886	1196	16.1	27.9
12	865	1197	15.3	27.6
					13	866	1209	15.5	26.7
14	903	1204	15.8	30.1
					15	878	1189	15.5	25.9

In order to express the present invention, the above embodiments are properly and fully described by way of examples, and the above embodiments are only used for illustrating the present invention and not for limiting the present invention, and those skilled in the relevant art can make various changes and modifications without departing from the spirit and scope of the present invention, and any modifications, equivalent substitutions, improvements, etc. made by the persons skilled in the relevant art should be included in the protection scope of the present invention, and the protection scope of the present invention should be defined by the claims.