Method for manufacturing and utilizing ferritic-austenitic stainless steel with high formability

Abstract

The invention relates to a method for manufacturing a ferritic-austenitic stainless steel having good formability and high elongation. The stainless steel is heat treated so that the microstructure of the stainless steel contains 45-75% austenite in the heat treated condition, the remaining microstructure being ferrite, and the measured M.sub.d30 temperature of the stainless steel is adjusted between 0 and 50° C. in order to utilize the transformation induced plasticity (TRIP) for improving the formability of the stainless steel.

Claims

1. A method for selecting a ferritic-austenitic stainless steel, comprising: applying a heat treatment to one or more stainless steels within a temperature range of 900-1200° C., followed by air cooling or water cooling so that the one or more stainless steels have 45-75% austenite phase with a remaining phase being ferrite; and straining to 0.30 true strain within the temperature range 0-50° C., yielding 50% transformation of the austenite phase to martensite phase in order to utilize the transformation induced plasticity (TRIP) and hence a measured M.sub.d30 of 0-50° C.; and selecting a stainless steel with a measured M.sub.d30 of 0-50° C. and an elongation value (A.sub.50) from 39% to 47%, wherein the stainless steel does not include vanadium.

2. The method according to claim 1, wherein the heat treatment is carried out as solution annealing.

3. The method according to claim 1, wherein the heat treatment is carried out as high-frequency induction annealing.

4. The method according to claim 1, wherein the heat treatment is carried out as local annealing.

5. The method according to claim 1, wherein the heat treatment is carried out at a temperature range of 1000-1150° C.

6. The method according to claim 1, wherein the measured M.sub.d30 is between 10° C. and 45° C.

7. The method according to claim 1, wherein the stainless steel contains in weight % less than 0.05% C, 0.2-0.7% Si, 2-5% Mn, 19-20.5% Cr, 0.8-1.35% Ni, less than 0.6% Mo, less than 1% Cu, 0.16-0.24% N, the balance Fe and inevitable impurities.

8. The method according to claim 1, wherein the stainless steel further contains one or more of 0-0.5% W, 0-0.2% Nb, 0-0.1% Ti, 0-0.2% V, 0-0.5% Co, 0-50 ppm B, and 0-0.04% Al.

9. The method according to claim 7, wherein the stainless steel contains 0-50 ppm O, 0-50 ppm S, and 0-0.4% P as inevitable impurities.

10. The method according to claim 7, wherein the stainless steel contains in weight % 0.01-0.04% C.

11. The method according to claim 7, wherein the stainless steel contains in weight % 1.0-1.35% Ni.

12. The method according to claim 7, wherein the stainless steel contains in weight % 0.18-0.22% N.

13. A method for selecting a ferritic-austenitic stainless steel, comprising heat treating ferritic-austenitic stainless steels based on a calculated M.sub.d30 temperature and austenite fraction in order to tune the transformation induced plasticity (TRIP) effect, including determining a Nohara M.sub.d30 temperature of each stainless steel using the following expression: M.sub.d30=551−462(C+N)−9.2Si−8.1Mn−13.7Cr−29(Ni+Cu)−18.5Mo−68Nb and selecting a stainless steel with a Nohara M.sub.d30 temperature within the range of −24-37° C. and an elongation value from 39% to 47%, wherein the stainless steel does not include vanadium.

14. The method according to claim 13, wherein the heat treatment is carried out as solution annealing.

15. The method according to claim 13, wherein the heat treatment is carried out as high-frequency induction annealing.

16. The method according to claim 13, wherein the heat treatment is carried out as local annealing.

17. The method according to claim 2, wherein the solution annealing is carried out at 1100° C. and the measured M.sub.d30 is between 20° C. and 35° C.

18. The method according to the claim 1, further including a step of determining a Nohara M.sub.d30 temperature of each stainless steel using the following expression:
M.sub.d30=551−462(C+N)−9.2Si−8.1Mn−13.7Cr−29(Ni+Cu)−18.5Mo−68Nb.

19. The method according to the claim 18, further including selecting from the one or more stainless steels, a stainless steel with a Nohara M.sub.d30 temperature within 20-35° C.

20. The method according to claim 1, wherein the austenitic phase includes, in weight %, from 0.05% C to 0.09% C, 0.28% N to 0.42% N, 0.25% Si to 0.31% Si, 2.25% Mn to 5.37% Mn, 18.67% Cr to 19.64% Cr, 0.79% Ni to 1.52% Ni, and 0.46% Cu to 0.63% Cu.

21. The method according to claim 20, wherein the austenitic phase further includes, in weight %, from 0.01% Mo to 0.4% Mo.

Description

(1) The present invention is described in more details referring to the drawings, where

(2) FIG. 1 is a diagram showing results of the M.sub.d30 temperature measurement using Satmagan equipment,

(3) FIG. 2 shows the influence of the M.sub.d30 temperature and the martensite content on strain-hardening and uniform elongation of the steels of the invention annealed at 1050° C.,

(4) FIG. 3a shows the influence of the measured M.sub.d30 temperature on elongation,

(5) FIG. 3b shows the influence of the calculated M.sub.d30 temperature on elongation,

(6) FIG. 4 shows the effect of the austenite content on elongation,

(7) FIG. 5 shows the microstructure of the alloy A of the invention using electron backscatter diffraction (EBSD) evaluation when annealed at 1050° C.,

(8) FIG. 6 shows the microstructures of the alloy B of the invention, when annealed at 1050° C., and

(9) FIG. 7 is a schematical illustration of the toolbox model.

(10) Detailed studies of the martensite formation were performed for some lean duplex alloys. Particular attention was paid on the effect of martensite formation and M.sub.d30 temperature on mechanical properties. This knowledge, crucial in designing a steel grade of optimum properties, is lacking from the prior art patents. Tests were done for some selected alloys according to Table 1.

(11) TABLE-US-00001 TABLE 1 Chemical composition of tested alloys Alloy C % N % Si % Mn % Cr % Ni % Cu % Mo % A 0.039 0.219 0.30 4.98 19.81 1.09 0.44 0.00 B 0.040 0.218 0.30 3.06 20.35 1.25 0.50 0.49 C 0.046 0.194 0.30 2.08 20.26 1.02 0.39 0.38 D 0.063 0.230 0.31 4.80 20.10 0.70 0.50 0.01 LDX 2101 0.025 0.226 0.70 5.23 21.35 1.52 0.31 0.30

(12) The alloys A, B and C are examples of the present invention. The alloy D is according to US patent application 2007/0163679, while LDX 2101 is a commercially manufactured example of SE 517449, a lean duplex steel with an austenite phase that has good stability to deformation martensite formation.

(13) The steels were manufactured in a vacuum induction furnace in 60 kg scale to small slabs that were hot rolled and cold rolled down to 1.5 mm thickness. The alloy 2101 was commercially produced in 100 ton scale, hot rolled and cold rolled in coil form. A heat treatment using solution annealing was done at different temperatures from 1000 to 1150° C., followed by rapid air cooling or water quenching.

(14) The chemical composition of the austenite phase was measured using scanning electron microscope (SEM) with energy dispersive and wavelength dispersive spectroscopy analysis and the contents are listed in Table 2. The proportion of the austenite phase (% γ) was measured on etched samples using image analysis in light optical microscope.

(15) TABLE-US-00002 TABLE 2 Composition of the austenite phase of the alloys after different treatments Alloy/treatment C % N % Si % Mn % Cr % Ni % Cu % Mo % C + N % % γ A (1000° C.) 0.05 0.28 0.28 5.37 18.94 1.30 0.59 0.00 0.33 73 A (1050° C.) 0.05 0.32 0.30 5.32 18.89 1.27 0.55 0.00 0.37 73 A (1100° C.) 0.06 0.35 0.28 5.29 18.67 1.32 0.54 0.00 0.41 68 B (1000° C.) 0.05 0.37 0.27 3.22 19.17 1.47 0.63 0.39 0.42 62 B (1050° C.) 0.06 0.37 0.27 3.17 19.17 1.52 0.57 0.40 0.43 62 B (1100° C.) 0.06 0.38 0.26 3.24 19.38 1.46 0.54 0.38 0.44 59 C (1050° C.) 0.07 0.40 0.26 2.25 19.41 1.32 0.51 0.27 0.47 53 C (1100° C.) 0.08 0.41 0.28 2.26 19.40 1.26 0.48 0.28 0.49 49 C (1150° C.) 0.09 0.42 0.25 2.27 19.23 1.27 0.46 0.29 0.51 47 D (1050° C.) 0.08 0.34 0.31 4.91 19.64 0.80 0.60 0.01 0.42 73 D (1100° C.) 0.09 0.35 0.31 5.00 19.51 0.79 0.52 0.01 0.44 72 LDX 2101 0.04 0.39 0.64 5.30 20.5 1.84 0.29 0.26 0.43 54 (1050° C.)

(16) The actual M.sub.d30 temperatures (M.sub.d30 test temp) were established by straining the tensile samples to 0.30 true strain at different temperatures and by measuring the fraction of the transformed martensite (Martensite %) with Satmagan equipment. Satmagan is a magnetic balance in which the fraction of ferromagnetic phase is determined by placing a sample in a saturating magnetic field and by comparing the magnetic and gravitational forces induced by the sample. The measured martensite contents and the resulting actual M.sub.d30 temperatures (M.sub.d30 measured) along with the predicted temperatures using Nohara expression M.sub.d30=551−462(C+N)−9.2Si−8.1Mn−13.7Cr−29(Ni+Cu)−18.5Mo−68Nb (M.sub.d30 Nohara) for the austenite composition are listed in Table 3. The measured proportion of austenite transformed to martensite at true stain 0.3 versus testing temperature is illustrated in FIG. 1.

(17) TABLE-US-00003 TABLE 3 Details of M.sub.d30 measurements Mart M.sub.d30 Mar- %/ Alloy/ Initial test tensite Initial M.sub.d30 ° C. M.sub.d30 ° C. treatment % γ temp % % γ measured (Nohara) A (1000° C.) 73 23° C. 44 61 29 37 40° C. 23 31 A (1050° C.) 73 23° C. 36 50 23 22 40° C. 17 23 60° C. 4 5 A (1100° C.) 68 23° C. 37 55 26 8.5 40° C. 15 22 B (1000° C.) 62 23° C. 35 57 27 −4 40° C. 17 27 B (1050° C.) 62 23° C. 28 45 17 −6 40° C. 13 27 60° C. 4 6 B (1100° C.) 59 23° C. 30 51 23.5 −13 40° C. 13 23 C (1050° C.) 53 23° C. 44 82 44 −12 40° C. 28 51 C (1100° C.) 49 23° C. 44 89 45 −18 40° C. 29 58 C (1150° C.) 47 23° C. 35 74 40 −24 40° C. 23 49 D (1050° C.) 73  0° C. 38 53 5 3 23° C. 23 32 D (1100° C.) 72  0° C. 37 52 3 −2 23° C. 19 26 LDX 2101 54 −40° C.   22 40 −52 −38 (1050° C.)  0° C. 7 14 LDX 2101 52 −40° C.   18 34 −59 −48 (1100° C.)  0° C. 8 15

(18) Measurements of the ferrite and austenite contents were made using light optical image analysis after etching in Beraha's etchant and the results are reported in Table 4. The microstructures were also assessed regarding the structure fineness expressed as austenite width (γ-width) and austenite spacing (γ-spacing). These data are included in Table 4 as well as the uniform elongation (Ag) and elongation to fracture (A.sub.50/A.sub.80) results in longitudinal (long) and transversal (trans) directions.

(19) TABLE-US-00004 TABLE 4 Micro-structural parameters, M.sub.d30 temperatures and ductility data γ- γ-width spacing M.sub.d30 ° C. *A.sub.50 % *A.sub.50 % Ag (%) Ag (%) Alloy/treatment % γ (μm) (μm) measured (long) (trans) (long) (trans) A (1000° C.) 73 5.0 2.5 29 44.7 41 A (1050° C.) 73 4.2 2.2 23 47.5 46.4 43 42 A (1100° C.) 68 5.6 3.5 26 46.4 42 B (1000° C.) 62 2.8 2.2 27 43.8 38 B (1050° C.) 62 4.2 3.0 17 45.2 44.6 40 40 B (1100° C.) 59 4.7 4.1 23.5 46.4 41 C (1050° C.) 53 3.3 3.4 44 41.1 40.3 38 37 C (1100° C.) 49 4.5 4.7 45 40.8 37 C (1150° C.) 47 5.5 5.9 40 41.0 37 D (1050° C.) 73 4.9 2.4 5 38 39 D (1100° C.) 72 6.4 2.8 3 40 39 LDX 2101 54 2.9 3.3 −52 36 30.0 24 21 (1050° C.) LDX 2101 52 3.3 4.2 −59 (1100° C.) *Tensile tests performed according to standard EN10002-1

(20) Examples of the resulting microstructures are shown in FIGS. 5 and 6. The results from tensile testing (standard strain rate 0.001 s.sup.−1/0.008 s.sup.−1) are presented in Table 5.

(21) TABLE-US-00005 TABLE 5 Full tensile test data Rp0.2 Rp1.0 Rm Ag A.sub.50 Alloy/treatment Direction (MPa) (MPa) (MPa) (%) (%) A (1000° C.) Trans 480 553 825 45 A (1050° C.) Trans 490 538 787 46 A (1050° C.) Long 494 542 819 43 48 A (1100° C.) Trans 465 529 772 46 B (1000° C.) Trans 492 565 800 44 B (1050° C.) Trans 494 544 757 45 B (1050° C.) Long 498 544 787 40 45 B (1100° C.) Trans 478 541 750 46 C (1050° C.) Trans 465 516 778 40 C (1050° C.) Long 474 526 847 38 41 C (1100° C.) Trans 454 520 784 41 C (1150° C.) Trans 460 525 755 41 D (1050° C.) Trans.sup.1) 548 587 809 45.sup.2) D (1050° C.) Long.sup.1) 552 590 835 38 44.sup.2) D (1100° C.) Trans.sup.1) 513 556 780 46.sup.2) D (1100° C.) Long.sup.1) 515 560 812 40 47.sup.2) LDX 2101 Trans 602 632 797 21 30 (1050° C.) LDX 2101 Long 578 611 790 24 36 (1050° C.) .sup.1)Strain rate 0.00075 s.sup.−1/0.005 s.sup.−1 .sup.2)A80

(22) To investigate the resistance to corrosion, the pitting potentials of the alloys were measured on samples, which were wet-ground to 320 mesh surface finish, in 1M NaCl solution at 25° C. using Standard Calomel electrode with a voltage scan of 10 mV/min. Three individual measurements were made for each grade. The results are shown in Table 6.

(23) TABLE-US-00006 TABLE 6 Pitting corrosion tests Result 1 Result 2 Result 3 Average Std dev Max Min Alloy mV mV mV mV mV mV mV A 341 320 311 324 15 17 13 B 380 400 390 14 10 10 C 328 326 276 310 29 18 34 304L 373 306 307 329 38 44 23

(24) Table 2 reveals that the phase balance and composition of the austenite phase vary with the solution annealing temperature. The austenite content decreases with increasing temperature. The compositional change in substitutive elements is small while the interstitial elements carbon and nitrogen show greater variation. As the carbon and nitrogen elements according to available formulas have a strong effect on the austenite stability against martensite formation, it appears to be crucial to control their level in the austenite. As shown in Table 3, the calculated M.sub.d30 temperatures are clearly lower for the heat treatments at higher temperature, indicating a greater stability. However, the measured M.sub.d30 temperatures do not display such dependence. For the alloys A, B and C the M.sub.d30 temperature is slightly reduced with just 3-4° C. when increasing the solution temperature with 100° C. This difference can be attributed to several effects. For example, the higher annealing temperature results in a coarser microstructure, which is known to affect the martensite formation. The tested examples have an austenite width or an austenite spacing in the order of about 2 to 6 μm. The products with the coarser microstructure show different stability and deviating description. The results show that the prediction of the martensite formation using current established expressions is not functional, even if advanced metallographic methods are employed.

(25) In FIG. 1 the results from Table 3 are plotted and the curves show that the influence of temperature on the martensite formation is similar for the tested alloys. Such dependence is an important part of the empirical descriptions for designed formability, as in industrial forming processes temperature can vary considerably.

(26) FIG. 2 illustrates the strong influence of the M.sub.d30-temperature of the austenite (measured) and the amount of the transformed strain-induced martensite (c.sub.α′) on the mechanical properties. In FIG. 2, the true stress-strain curves of the tested steels are shown with thin lines. The thick lines correspond to the strain-hardening rate of the steels, obtained by differentiating the stress-strain curves. According to Considére's criterion, the onset of necking, corresponding to uniform elongation, occurs at the intersection of the stress-strain curve and the strain-hardening curves, after which the strain-hardening cannot compensate the reduction of the load bearing capacity of the material caused by thinning.

(27) The M.sub.d30-temperatures and the martensite contents at uniform elongation of the tested steels are also shown in FIG. 2. It is obvious that the strain-hardening rate of the steel is essentially dependent on the extent of martensite formation. The more martensite is formed, the higher strain-hardening rate is reached. Thus, by carefully adjusting the M.sub.d30-temperature, the mechanical properties, namely the combination of tensile strength and uniform elongation can be optimized.

(28) Apparently, based on the present experimental results, the range of optimum M.sub.d30-temperature is substantially narrower than indicated by the prior art patents. A too high M.sub.d30-temperature causes rapid peaking of the strain-hardening rate. After peaking the strain-hardening rate drops rapidly, resulting in early onset of necking and low uniform elongation. According to the experimental results, the M.sub.d30-temperature of the steel C appears to be close to the upper limit. If the M.sub.d30-temperature was much higher, the uniform elongation would be substantially decreased.

(29) On the other hand, if the M.sub.d30-temperature is too low, not enough martensite is formed during deformation. Therefore, the strain-hardening rate remains low, and consequently, the onset of necking occurs at a low strain level. In FIG. 2, LDX 2101 represents typical behaviour of a stable duplex steel grade with low uniform elongation. The M.sub.d30-temperature of the steel B was 17° C., which was high enough to enable a sufficient martensite formation to ensure the high elongation. However, if the M.sub.d30-temperature was even lower, too little martensite would form and the elongation would be clearly lower.

(30) Based on the experiments, the chemical composition and the thermo-mechanical treatments shall be designed so that the resulting M.sub.d30-temperature of the steel ranges is between 0 and +50° C., preferably between 10° C. and 45° C., and more preferably 20-35° C.

(31) The tensile test data in Table 5 illustrates that the elongation at fracture is high for all steels according to the invention while the commercial lean duplex steel (LDX 2101) with a more stable austenite exhibits lower elongation values typical for standard duplex steels. FIG. 3a illustrates the influence of the measured M.sub.d30 temperatures of the austenite on the ductility. For the actual examples an optimum ductility is obtained for the M.sub.d30 temperatures between 10 and 30° C. In FIG. 3b the influence of the calculated M.sub.d30 temperatures on ductility is plotted.

(32) Both the diagrams, FIG. 3a and FIG. 3b, illustrate clearly that there is an almost parabolic correlation between the M.sub.d30 temperature values and the elongation regardless of how the M.sub.d30 temperature has been obtained. There is a clear discrepancy between the measured and calculated M.sub.d30 values in particular for alloy C. The diagrams show that the desired range of the M.sub.d30 temperature is much narrower than the calculations predict, which means that the process control needs to be much better optimized to obtain a desired TRIP effect. FIG. 4 shows that the austenite content for the optimum ductility ranges from about 50 to 70% for the used examples. In FIG. 5 the M.sub.d30 temperature of the alloy A is tested at 40° C. having in the microstructure 18% martensite (grey in image) and about 30% of austenite (black in image) the rest being ferrite (white in image).

(33) FIG. 6 shows the microstructures of the alloy B of the invention after annealed at 1050° C. The phases in FIG. 6 are ferrite (grey), austenite (white) and martensite (dark grey within the austenite (white) bands) In FIG. 6 the part a) relates to a reference material, the part b) relates to the M.sub.d30 temperature test performed at room temperature, the part c) relates to the M.sub.d30 temperature test performed at 40° C. and the part d) relates to the M.sub.d30 temperature test performed at 60° C.

(34) The control of the M.sub.d30 temperature is crucial to attain high deformation elongation. It is also important to take the material temperature during deformation into consideration as it largely influences the amount of martensite that can form. Data in Table 5 and in FIGS. 3a and 3b refers to room temperature tests but some increase in temperature cannot be avoided due to adiabatic heating. Consequently, steels with a low M.sub.d30 temperature may not show a TRIP effect if deformed at an elevated temperature while steels having an apparently too high M.sub.d30 temperature for optimum ductility at room temperature will show excellent elongation at elevated temperatures. The tensile tests with the alloys A and C at different temperatures (Table 7) showed the following relative changes in elongation:

(35) TABLE-US-00007 TABLE 7 The tensile tests with the Alloys A and C at different temperatures Temperature Alloy 20° C. 45° C. 65° C. A 100% 100%  85% C 100% 120% 115%

(36) The results show that the alloy A with lower M.sub.d30 temperature exhibits a reduction in elongation at elevated temperature, while the alloy C with the higher M.sub.d30 temperature demonstrates an increased elongation when the temperature is raised.

(37) Table 6 shows that the pitting corrosion resistance, expressed as pitting potential in 1M NaCl, is at least as good as that of the austenitic standard steel 304L.

(38) Prior art has not disclosed sufficient capability to design duplex steels with TRIP-effect properly as the predictions of the steel behaviour using established formulas are unsecure giving too wide ranges in the compositions and in other specifications. According to the present invention lean duplex steels can be more safely designed and manufactured with optimum ductility by selecting certain composition ranges and by using a special procedure involving measurement of the actual M.sub.d30 temperature and by employing special empirical knowledge to control the manufacturing processes. This new innovative approach is necessary to be able to utilize the real TRIP effect in the design of highly formable products. As illustrated in FIG. 7 a toolbox concept is used where empirical models for the phase balance and the austenite stability based on the measurements are used to select the alloy compositions that will be subjected to special thermal-mechanical treatments for designed formability (the austenite fraction and the M.sub.d30 temperature). By this model it is possible to design the austenite stability giving the optimum formability for a certain customer or solution application with a greater flexibility than for austenitic stainless steels exhibiting TRIP effect. For such austenitic stainless steels, the only way to adjust the TRIP effect is to choose another melt composition, while according to the present invention utilizing TRIP effect in a duplex alloy, the heat treatment such as the solution annealing temperature gives an opportunity to fine-tune the TRIP effect without necessarily introducing a new melt.