Patent Application
20080175743
This disclosure relates to an alloyed hot-dip galvanized steel sheet having high strength, excellent surface appearance, and excellent anti-secondary work embrittlement and a method of producing the same. The alloyed hot-dip galvanized steel sheet is suitably applied to automobile steel sheets.
In recent years, from the viewpoint of conservation of the global environment, there have been advanced thinning of steel sheets by weight reduction of car bodies for improving mileages and increases in high strength steel sheets for improving safety. However, increases in strength of steel sheets decrease ductility and toughness, and thus there have been desired steel sheets having both high strength and high formability and excellent toughness after forming (anti-secondary work embrittlement).
For such a requirement, there have been developed various steel sheets such as ferrite-martensite dual phase steel (so-called Dual-Phase steel) and steel using transformation-induced plasticity of residual austenite (so-called TRIP steel).
In some cases, the surfaces of these steel sheets are galvanized for improving rust prevention in practical use. As such galvanized steel sheets, alloyed hot-dip galvanized steel sheets subjected to heat treatment for diffusing Fe of the steel sheets into plating layers after hot-dip galvanization are widely used from the viewpoint of securing press property, spot weldability, coating adhesion.
For example, in Japanese Unexamined Patent Application Publication No. 11-279691, there has been proposed an alloyed hot-dip galvanized steel sheet in which residual austenite (also referred to as “residual γ”) is secured by adding a large amount of Si to the steel sheet, thereby achieving high ductility. However, Si decreases the adhesion of zinc plating, and thus a complicated process of Ni pre-plating, applying a special chemical, reducing an oxide layer on the surface of a steel sheet, and/or appropriately controlling the thickness of an oxide layer is required for adhesion of zinc plating to such high-Si steel.
In Japanese Unexamined Patent Application Publication No. 2002-030403, there has been proposed an alloyed hot-dip galvanized steel sheet in which instead of Si, Al with a small adverse effect on zinc plating adhesion is added to the steel sheet, thereby securing excellent ductility and improving wettability and anti-powdering qualities of zinc plating. However, in high-Al steel, N and Al in the steel forms AlN which precipitates in large amounts at austenite grain boundaries during continuous casting, thereby embrittling the grain boundaries. Since bending correction from the vertical direction to the horizontal direction is performed in usual continuous casting, the embrittlement of the grain boundaries easily causes cracking in a slab at a corrected portion. When a slab having a crack is rolled, the crack remains in a final product to significantly deteriorate the surface appearance. In this case, a process of removing the crack of the slab with a grinder is required, thereby significantly increasing cost.
In U.S. Pat. No. 3,596,316, there has been proposed a method of avoiding slab cracking in which Ti is added to a steel sheet to fix N as TiN in order to avoid the slab cracking. However, precipitation of AlN actually starts at a temperature higher than the temperature where N is completely fixed as TiN, and thus it is difficult to completely avoid the slab cracking. Further, Al is a strong ferrite stabilizing element and increases the A3 transformation point, and thus ferrite is easily produced, due to an increase in the transformation point, at a slab corner where a temperature drop easily occurs until the slab width is reduced after the slab is discharged from a heating furnace during hot rolling. As a result, local distortion is concentrated at a corner during width reduction, thereby easily causing surface defects such as scabs.
Any one of the proposed techniques has the problem of high surface sliding resistance and low anti-secondary work embrittlement in comparison to cold-rolled steel sheets because the high-tension steel sheets are hot-dip galvanized in a subsequent process.
On the other hand, in Japanese Unexamined Patent Application Publication No. 2004-211140, there has been proposed a hot-dip galvanized steel sheet in which in order to improve the anti-secondary work embrittlement of a galvanized steel sheet, the amount of B added to a Dual-Phase steel sheet used as a steel sheet is controlled to decrease the grain size of ferrite grains, improving grain boundary strength. However, a steel sheet utilizing a martensite phase cannot utilize an improvement in ductility (TRIP effect) due to strain-induced transformation of residual austenite and thus does not have sufficient ductility.
It could therefore be advantageous to provide an alloyed hot-dip galvanized steel sheet having high strength and excellent surface appearance and excellent anti-secondary work embrittlement and a method of producing the same.
We provide an alloyed hot-dip galvanized steel sheet having a composition containing about 0.05 to about 0.25 wt % of C, about 0.5 wt % or less of Si, about 1 to about 3 wt % of Mn, about 0.1 wt % or less of P, about 0.01 wt % or less of S, about 0.1 to about 2 wt % of Al, and less than about 0.005 wt % of N, and satisfying the relations, Si+Al≧0.6 wt %, (0.0006×Al) wt %≦N≦(0.0058−0.0026×Al) wt %, and Al≦(1.25×C0.5−0.57×Si+0.625×Mn) wt %, the balance including Fe and inevitable impurities.
The alloyed hot-dip galvanized steel sheet has the composition which may further contain at least one element selected from the group consisting of about 1 wt % or less of Cr, about 1 wt % or less of V, and about 1 wt % or less of Mo.
Each of the alloyed hot-dip galvanized steel sheet described above has the composition which may further contain at least one element selected from the group consisting of about 0.1 wt % or less of Ti, about 0.1 wt % or less of Nb, about 0.005 wt % or less of B, and about 1 wt % or less of Ni.
Each of the alloyed hot-dip galvanized steel sheet described above has the composition which may further contain at least one element selected from the group consisting of Ca and REM in a total of about 0.01 wt % or less.
Each of the alloyed hot-dip galvanized steel sheet described above has a metal structure preferably containing a residual austenite phase at a volume ratio of about 3 to about 20%.
Further, we provide a method of producing an alloyed hot-dip galvanized steel sheet, the method including casting, hot-rolling, and cold-rolling steel for a steel sheet having any one of the above-described compositions, holding the steel at about 730° C. to about 900° C. for about 60 to about 300 seconds, cooling the steel at about 3 to about 100° C./s, further holding the steel at about 350° C. to about 600° C. for about 30 to about 250 seconds, hot-dip galvanizing the steel, and then alloying the steel at about 470° C. to about 600° C.
We found that precipitation and coarsening of AlN influence not only the surface quality of a final product due to slab cracking, but also the anti-secondary work embrittlement of a final product and the anti-secondary work embrittlement is improved when an N amount is N≦(0.0058−0.0026×Al) wt %. Although details of reasons are not entirely understood, a conceivable reason is that breakage starts from coarse AlN due to embrittlement at a low temperature and, when the Al amount is increased, precipitation of AlN starts from a high temperature to easily coarsen the precipitate, thereby deteriorating the anti-secondary work embrittlement.
On the other hand, fine AlN precipitating at a high temperature suppresses coarsening of austenite grains, and thus ferrite grains are made fine, thereby improving the anti-secondary work embrittlement. Therefore, we found that there is a required minimum N amount, and a suitable N amount satisfies (0.0006×Al) wt %≦N for improving the anti-secondary work embrittlement.
We also found that the occurrence of scabs in hot rolling is suppressed by controlling the Al amount to Al≦(1.25×C0.5−0.57×Si+0.625×Mn) wt %. This is adapted for balancing an increase in the Ar3 transformation point due to addition of Al and Si and a decrease in the transformation point due to addition of C and Mn. The reason for this is that when the components are controlled in the above range, ferrite formation at a corner of a slab before width reduction is suppressed.
Therefore, slab cracking in continuous casting, scabs in hot rolling, and non-plating after hot-dip galvanization are suppressed to improve the surface appearance of products. In addition, the anti-secondary work embrittlement is improved. Further, an alloyed hot-dip galvanized steel sheet with high strength can be obtained without passing through a complicated process.
Our steel sheets will be described in detail below.
First, the reasons for specifying the composition of the alloyed hot-dip galvanized steel sheet will be described. Hereinafter, “%” represents “% by mass.”
C: about 0.05 to about 0.25%
C is an element for stabilizing austenite and a necessary element for securing a martensite amount and causing austenite to remain at room temperature. When the C amount is less than about 0.05%, it is difficult to simultaneously secure the strength of the steel sheet and the amount of residual austenite to achieve high ductility. On the other hand, when the C amount exceeds about 0.25%, welds and heat-affected zone are significantly hardened, thereby deteriorating weldability. Therefore, the C amount is in the range of about 0.05 to about 0.25%.
Si: about 0.5% or Less
Si is an element effective in strengthening steel. Si is also a ferrite forming element which promotes the concentration of C in austenite and suppresses the formation of carbides and thus has the function of promoting the formation of residual austenite. However, when the Si amount exceeds about 0.5%, galvanizing is deteriorated, and plating is difficult in a usual hot-dip galvanization process. Therefore, the Si amount is about 0.5% or less. Usually, the Si amount is preferably about 0.01 to about 0.5% and more preferably about 0.01 to about 0.4%.
Mn: about 1 to about 3%
Mn is an element effective in strengthening steel. Mn is also an element for stabilizing austenite and an element necessary for increasing residual austenite. However, when the Mn amount is less than about 1%, these effects cannot be easily obtained. On the other hand, when the Mn amount exceeds about 3%, a second phase fraction is excessively increased, and the amount of solid-solution hardening is increased, thereby significantly increasing strength and greatly decreasing ductility. Therefore, such a Mn amount is unsuitable for application to automobile steel sheets. Thus, the Mn amount is in the range of about 1 to about 3%.
P: about 0.1% or Less
P is an element effective in strengthening steel. However, when the P amount exceeds about 0.1%, plating defects or non-plating occurs, and P segregates at grain boundaries to deteriorate the anti-secondary work embrittlement. Therefore, the P amount is about 0.1% or less. From the viewpoint of anti-secondary work embrittlement, the P amount is preferably about 0.05% or less. Usually, the P amount is preferably about 0.001 to about 0.05%.
S: about 0.01% or Less
S forms inclusions such as MnS and causes deterioration in impact resistance and cracking along metal flow of welds. Therefore, the S amount is preferably as small as possible. However, from the viewpoint of production cost, the S amount is about 0.01% or less. Usually, the S amount is preferably about 0.0001 to about 0.01%.
Al: about 0.1 to about 2%
Si+Al≦0.6%
Al and Si are ferrite forming elements which promote the concentration of C in austenite and suppress the formation of a carbide and thus have the function of promoting the formation of residual austenite. When a total of Al and Si added is less than about 0.6%, sufficient ferrite or residual γ cannot be obtained, thereby significantly decreasing ductility. Even when Al is added in an amount of less than about 0.1% and Si is added up to an upper limit, the total of Al+Si is less than about 0.6%. On the other hand, when the Al amount exceeds about 2%, the amount of the inclusion in the steel sheet is increased, deteriorating ductility. Therefore, the Al amount is in the range of about 0.1 to about 2%, and Si+Al≧0.6% is satisfied.
Al≦(1.25×C0.5−0.57×Si+0.625×Mn) %
When the Al content exceeds (1.25×C0.5−0.57×Si+0.625×Mn) %, scabs easily occur in hot rolling. Therefore, the Al amount satisfies Al≦(1.25×C0.5−0.57×Si+0.625×Mn) %.
N: Less than about 0.005%
(0.0006×Al) %≦N≦(0.0058−0.0026×Al) %
N is an important element and causes deterioration in the anti-secondary work embrittlement when the amount of AlN precipitate is increased with an increase in the N amount. To avoid such deterioration in the anti-secondary work embrittlement, the N amount is limited to less than about 0.005%, and the relational expression, (0.0006×Al)≦N≦(0.0058−0.0026×Al) %, is satisfied. When the N amount is about 0.005% or more, AlN excessively precipitates, causing cracking in a slab. Therefore, the N amount is less than about 0.005%. Usually, the N amount is preferably about 0.001 to less than about 0.005%. Cr, V, Mo: each about 1% or less
Cr, V, and Mo have the function of suppressing the formation of pearlite in cooling from an annealing temperature and thus can be added according to demand. When the amount of each of these elements is about 1% or less, appropriate steel sheet strength is obtained, and ductility and adhesion of zinc plating are little adversely affected. Therefore, when Cr, V, and Mo are added, the amount of each of the elements is preferably about 1% or less. In general, the amount is preferably about 0.01 to 1% and more preferably about 0.01 to about 0.5%.
Ti, Nb: Each about 0.1% or Less
Ti and Nb are effective in precipitation strengthening of steel and can thus be added according to demand. However, when the amount of each of Ti and Nb is about 0.1% or less, formability and shape fixability can be easily maintained. Therefore, when Ti and Nb are added, the amount of each of the elements is preferably about 0.1% or less. In general, the amount is preferably about 0.01 to about 0.1% and more preferably about 0.01 to about 0.05%.
B: about 0.005% or Less
B is effective in strengthening steel and can thus be added according to demand. When the B amount is about 0.005% or less, an increase in strength can be easily controlled, thereby achieving excellent formability. Therefore, when B is added, the amount is preferably about 0.005% or less. In general, the amount is preferably about 0.0001 to about 0.005% and more preferably about 0.0001 to about 0.003%.
Ni: about 1% or Less
Ni is an austenite stabilizing element which causes austenite to remain and is effective in increasing strength, and thus can be added according to demand. When the Ni amount is about 1% or less, ductility of a steel sheet can be easily increased. Therefore, when Ni is added, the amount is preferably about 1% or less. In general, the amount is preferably about 0.01 to about 1%.
Ca and REM: at Least One in Total of about 0.01% or Less
“REM” represents at least one of the rare earth elements. Ca and REM have the function of controlling the form of sulfide inclusions and thus have the effect of improving elongation and stretch flange formability of a steel sheet, and thus can be added according to demand. When the total of these elements exceeds about 0.01%, these effects are saturated. Therefore, when Ca and REM are added, the total of at least one of the elements is preferably about 0.01% or less. In general, the total is preferably about 0.001 to about 0.01% and more preferably about 0.001 to about 0.005%.
Besides the above-described elements and Fe in the balance, various impurities in the production process and trace amounts of essential elements in the production process are inevitably mixed. However, these inevitable impurities are permissible as long as they have no particular influence on the disclosed advantage.
From the viewpoint of anti-secondary work embrittlement, it is preferable to satisfy the relational expression, (−10×C+5×Si−Mn+6×Al) %≧−0.5. Although details of reasons for this are not completely understood, a conceivable reason is that to suppress coarsening of austenite grains at a high temperature due to a decrease in the Ac3 point, a decrease in Ac3 point due to the addition of C and Mn and an increase in Ac3 point due to the addition of Si and Al are balanced by satisfying the above range, thereby improving the anti-secondary work embrittlement.
Next, the metal structure of the steel sheet will be described.
Residual Austenite Phase: Volume Ratio of about 3 to about 20%
The strain-induced transformation of a residual austenite phase is effectively utilized so that excellent surface appearance and anti-secondary work embrittlement and not only high strength but also high ductility can be imparted to the alloyed hot-dip galvanized steel sheet as a final product. Therefore, it is very important to control the volume ratio of the residual austenite. From the viewpoint of securing high ductility, the ratio of the residual austenite phase is preferably about 3% or more. On the other hand, when the ratio of the residual austenite phase is about 20% or less, the formation of martensite after forming is suppressed, and thus such a ratio is suitable in view of brittleness. Therefore, the ratio of the residual austenite phase is preferably about 20% or less. The metal structure of a steel sheet used as the alloyed hot-dip galvanized steel sheet includes a ferrite main phase and a residual austenite phase as a second phase. However, the volume ratio of the ferrite phase is preferably about 40 to about 90% from the viewpoint of securing high ductility. Examples of a metal structure other than the residual austenite phase in the second phase include a bainite phase, a martensite phase and/or a pearlite phase. The total volume ratio of these phases is preferably about 7 to about 50%. The average grain size of the ferrite main phase is preferably about 15 μm or less because the anti-secondary work embrittlement can be easily improved.
Next, a preferred method of producing the alloyed hot-dip galvanized steel sheet will be described.
Steel satisfying the conditions of the above composition is continuously cast to form a cast slab, and then the cast slab is hot-rolled and cold-rolled to prepare a steel sheet. However, the conditions for these processes are not particularly limited. Then, in a continuous galvanizing line, the steel sheet is annealed by holding in a temperature range of about 730° C. to about 900° C. for about 60 to about 300 seconds, cooled at about 3 to about 100° C./s, held in a temperature range of about 350° C. to about 600° C. for about 30 to about 250 seconds, hot-dip galvanized, and then alloyed at about 470° C. to about 600° C.
In the production method, excellent surface appearance and anti-secondary work embrittlement and not only high strength but also high ductility can be imparted to the alloyed hot-dip galvanized steel sheet as the final product.
Each of the production conditions will be described in further detail below.
Annealing Temperature: about 730 to about 900° C.
Holding Time: about 60 to about 300 Seconds
Annealing is performed in the austenite region or intercritical region including an austenite phase and a ferrite phase. When the annealing temperature is about 730° C. or more and the holding time is about 60 seconds or more, it is easy to dissolve a carbide in the steel sheet, completely recrystallize ferrite, and impart ductility. On the other hand, when the annealing temperature is about 900° C. or less, coarsening of austenite grains is suppressed, and the number of ferrite nucleation sites formed from the second phase by subsequent cooling is easily increased. Further, when the holding time is about 300 seconds or less, coarsening of AlN is suppressed, and the anti-secondary work embrittlement is easily improved. Therefore, the annealing temperature is about 730 to about 900° C., and the holding time is about 60 to about 300 seconds.
Cooling Rate: about 3 to about 100° C./s
When the cooling rate is about 3° C./s or more, precipitation of pearlite is suppressed, dissolved Carbon content in untransformed austenite tends to increase, and thus the intended metal structure (residual austenite phase) can be easily obtained. When the cooling rate is about 100° C./s or less, growth of ferrite is promoted to increase the volume ratio of ferrite, and thus sufficient ductility can be easily secured. Therefore, the cooling rate is preferably about 3 to about 100° C./s. Although the scope of our disclosure includes a case in which the cooling rate changes during cooling, the average cooling rate is preferably about 10° C./s or more and more preferably over about 20° C./s from the viewpoint of productivity.
Holding Temperature Range: about 350° C. to about 600° C.
When the holding temperature is about 600° C. or less, precipitation of a carbide in untransformed austenite is suppressed. When the holding temperature is about 350° C. or more, precipitation of a carbide in bainitic ferrite due to lower bainite transformation is suppressed to easily form stable residual austenite. Therefore, the holding temperature is preferably about 350° C. to about 600° C. In order to stably produce residual austenite, the holding temperature is preferably about 500° C. or less.
Holding Time: about 30 to about 250 Seconds
The holding time pays a very important role for controlling residual austenite. Namely, when the holding time is about 30 seconds or more, stabilization of untransformed austenite proceeds, and thus the amount of residual austenite is easily secured, easily obtaining desired properties. On the other hand, when the holding time is about 250 seconds or less, line speed need not be extremely decreased even in a CGL line where austempering cannot be performed for a long time, thereby causing an advantage of productivity. Therefore, the holding time is about 30 to about 250 seconds. The holding time is preferably about 70 seconds or more in order to stably secure residual austenite and preferably about 200 seconds or less from the viewpoint of productivity.
Alloying Temperature: about 470° C. to about 600° C.
The alloying temperature after hot-dip galvanization must be higher than the plating bath temperature, and the lower limit is about 470° C. When the alloying temperature is about 600° C. or less, like in the case where the holding temperature is about 600° C. or less, precipitation of a carbide in untransformed austenite is suppressed, and thus stable residual austenite can be easily obtained. Therefore, the alloying temperature is about 470° C. to about 600° C.
In the production method, the specified annealing temperature, holding temperature, and alloying temperature need not be constant as long as they are in the above-respective ranges. The plating conditions may be in a usual operation range, i.e., coating weight may be about 20 to about 70 g/m2, and the amount of Fe in a plating layer may be about 6 to about 15%.
Molten steel having each of the compositions shown in Table 1 was prepared by a converter and continuously cast to form a cast slab. The resulting slab was heated to 1250° C. and then hot-rolled at a finish rolling temperature of 900° C. to prepare a hot-rolled steel sheet having a thickness of 3.0 mm. The hot-rolled steel sheet produced as described above was visually observed for the occurrence of scabs. The presence of scabs is shown in Table 2.
After hot-rolling, the hot-rolled steel sheet was pickled and further cold-rolled to prepare a cold-rolled steel sheet having a thickness of 1.2 mm. Then, in a continuous galvanizing line, each cold-rolled steel sheet was annealed at 820° C., cooled at a rate of 10° C./s, galvanized with coating weight of 50/50 g/m2 by a zinc plating bath at 460° C., and then alloyed at 520° C. to prepare an alloyed hot-dip galvanized steel sheet.
Since the surface appearance of an alloyed hot-dip galvanized steel sheet are significantly inhibited by scabs, cracking in a slab, and plating defects, the surface appearance of the alloyed hot-dip galvanized steel sheets are also shown in Table 2. Good surface appearance represent a state in which a final product has a uniform surface with beauty without defects, and poor surface appearance represent a state in which surface defects such as scabs or non-plating occur.
Further, each of the resulting alloyed hot-dip galvanized steel sheets was temper-rolled of 0.5%, and mechanical properties were examined. As the mechanical properties, tensile strength (TS) and elongation (El) were measured using a JIS No. 5 tensile specimen obtained from each steel sheet in a direction perpendicular to the rolling direction. The measured values and values of TS×El are also shown in Table 2.
The anti-secondary work embrittlement was evaluated by the following method:
The above-mentioned results indicate that an alloyed hot-dip galvanized steel sheet obtained from steel satisfying our conditions has excellent surface appearance, high strength, and excellent anti-secondary work embrittlement.
Molten steel having each of the compositions shown in Table 3 was prepared by a converter and a hot-rolled steel sheet having a thickness of 3.0 mm was prepared by the same method as in Example 1. The hot-rolled steel sheet produced as described above was visually observed for the occurrence of scabs. The presence of scabs is shown in Tables 4-1 and 4-2.
After hot-rolling, the hot-rolled steel sheet was pickled and further cold-rolled to prepare a cold-rolled steel sheet having a thickness of 1.2 mm. Then, in a continuous galvanizing line, each cold-rolled steel sheet was heat-treated under the conditions shown in Tables 4-1 and 4-2, plated at 50/50 g/m2, and then alloyed to prepare an alloyed hot-dip galvanized steel sheet.
Like in Example 1, the surface appearances of the resulting alloyed hot-dip galvanized steel sheets are also shown in Tables 4-1 and 4-2.
Further, each of the resulting alloyed hot-dip galvanized steel sheets was measured with respect to tensile strength (TS) and elongation (El) by the same method as in Example 1. The measured values and values of TS×El are also shown in Tables 4-1 and 4-2.
With respect to the anti-secondary work embrittlement, a critical temperature (longitudinal crack transition temperature) where a brittle crack occurred in the side wall of a cylindrical cup was determined by the same method as in Example 1. The critical temperature was used as an index for the anti-secondary work embrittlement, for evaluating steel sheets of Invention Examples 2-1 to 2-28 and Comparative Examples 2-29 to 2-38. The results are also shown in Tables 4-1 and 4-2.
Tables 3, 4-1, and 4-2 indicate that in the alloyed hot-dip galvanized steel sheets obtained from steel Nos. 2-S, 2-U to 2-W, and 2-Y not satisfying any one of the N amount of less than 0.005%, Al≦(1.25×C0.5−0.57×Si+0.625×Mn) %, and the Si amount of 0.5% or less, the surface appearance were deteriorated. In addition, in the alloyed hot-dip galvanized steel sheets obtained from steel Nos. 2-Q to 2-U not satisfying (0.0006×Al) %≦N≦(0.0058−0.0026×Al) %, the anti-secondary work embrittlement was deteriorated.
Further, Tables 4-1 and 4-2 indicate that in the alloyed hot-dip galvanized steel sheets of Invention Examples 2-2, 2-5, 2-8, 2-12, 2-14, 2-19, 2-20, 2-25, and 2-28 containing small amounts of residual austenite, both the elongation and the values of TS×El are low, but these steel sheets have excellent surface appearance and anti-secondary work embrittlement and high tensile strength (TS). On the other hand, in the alloyed hot-dip galvanized steel sheets of Comparative Examples 2-29 to 2-38, any one of the surface appearance, anti-secondary work embrittlement, and tensile strength (TS) is deteriorated.
Since the alloyed hot-dip galvanized steel sheets of Invention Examples 2-1, 2-3, 2-4, 2-6, 2-7, 2-9 to 2-11, 2-13, 2-15 to 2-18, 2-21 to 2-24, 2-26, and 2-27 were produced under the preferred production conditions and satisfied the condition of the amount of residual austenite, these steel sheets have not only excellent surface appearance and anti-secondary work embrittlement and high tensile strength but also high elongation (El) and high values of TS×El.
Molten steel having each of the compositions shown in Table 5 was prepared by a converter and continuously cast to form a cast slab. The occurrence of cracking in the slab is shown in Tables 6-1 and 6-2. The occurrence of cracking was determined by visual observation as well as color check after the slab was cooled to room temperature.
The resulting slab was heated to 1250° C. and then hot-rolled at a finish rolling temperature of 900° C. to prepare a hot-rolled steel sheet having a thickness of 3.0 mm. The hot-rolled steel sheet produced as described above was visually observed for the occurrence of scabs. The presence of scabs is shown in Tables 6-1 and 6-2.
After hot-rolling, the hot-rolled steel sheet was pickled and further cold-rolled to prepare a cold-rolled steel sheet having a thickness of 1.2 mm. Then, in a continuous galvanizing line, each cold-rolled steel sheet was heat-treated under the conditions shown in Tables 6-1 and 6-2, plated at 50/50 g/m2, and then alloyed so that the Fe amount in the plating layer was 9%.
Further, each of the resulting alloyed hot-dip galvanized steel sheets was measured with respect to tensile strength (TS) and elongation (El) by the same method as in Example 1. The measured values and values of TS×El are also shown in Tables 6-1 and 6-2.
Table 6-1 indicates that the alloyed hot-dip galvanized steel sheets of Invention Examples 3-5 and 3-8 to 3-19 satisfying our steel sheet compositions cause no cracking in the slabs and no scabs in the hot-rolled steel sheets and have excellent surface appearance. Also, the anti-secondary work embrittlement and tensile strength are excellent.
On the other hand, in the alloyed hot-dip galvanized steel sheets Comparative Examples 3-1 to 3-4, 3-6, 3-7, and 3-20 to 3-30 not satisfying any one of the steel sheet components, i.e., the N amount, (0.0006×Al) %≦N≦(0.0058-0.0026×Al) %, and Al≦(1.25×C0.5−0.57×Si+0.625×Mn) %, a problem occurred in at least one of deterioration in the surface appearance due to cracking in the slabs, scabs in the hot-rolled steel sheets, or non-plating, and the anti-secondary work embrittlement. Comparative Example 3-30 also had low tensile strength (TS).
In particular, as shown in Comparative Example 3-31 in Table 6-2, steel 3-W containing a large amount of Mn exhibits a significant increase in strength but very low elongation. Further, as shown in Comparative Example 3-32 in Table 6-2, steel 3-X having a small total of Al+Si exhibits very low elongation for strength and a low value of TS×El.
Since the alloyed hot-dip galvanized steel sheets of Invention Examples 3-5, 3-8, 3-11, 3-12, and 3-15 to 3-18 satisfy the preferred production conditions, these steel sheets have appropriate amounts of residual austenite and have not only excellent surface appearance and anti-secondary work embrittlement and high strength but also high elongation (El) and high values of TS×El.
0.0027
0.0037
0.0010
0.0005
0.0037
1.47
0.0054
0.0032
0.0056
0.0035
0.0051
0.0024
1.19
0.03
0.4
0.80
0.0029
0.5
0.63
X
X
X
X
Present
X
X
X
X
X
Prsent
X
X
X
X
Present
X
0.0011
0.0040
0.0057
0.0024
0.0027
0.0019
1.31
1.16
0.90
0.02
0.4
0.0160
0.50
0.4
Present
Present
Present
X
X
X
X
X
0.0022
0.0027
0.0037
0.0011
0.0055
0.0040
0.60
0.0053
0.0027
0.0011
1.64
0.0016
1.30
0.0029
0.91
0.0019
0.04
0.0022
1.26
0.70
0.80
3.50
Present
Present
Present
Present
Present
Present
Preesnt
Present
Preesnt
Preesnt
Present
Present
Preesnt
X
X
X
X
X
X
X
X
X
X
X
An alloyed hot-dip galvanized steel sheet causing no cracking in a slab in continuous casting, no scab in hot rolling, and no plating defect after hot-dip galvanizing and thus having excellent surface appearance can be obtained without passing through a complicated process. In addition, the alloyed hot-dip galvanized steel sheet has both excellent anti-secondary work embrittlement and high strength, and is thus suitable as an automobile steel sheet and can widely contribute to the industrial field.
| Number | Date | Country | Kind |
|---|---|---|---|
| 2005-103833 | Mar 2005 | JP | national |
| 2006-058460 | Mar 2006 | JP | national |
This is a §371 of International Application No. PCT/JP2006/307396, with an international filing date of Mar. 31, 2006 (WO 2006/104275 A1, published Oct. 5, 2006), which is based on Japanese Patent Application Nos. 2005-103833, filed Mar. 31, 2005, and 2006-058460, filed Mar. 3, 2006.
| Filing Document | Filing Date | Country | Kind | 371c Date |
|---|---|---|---|---|
| PCT/JP06/07396 | 3/31/2006 | WO | 00 | 9/6/2007 |
