Method for producing an ausferritic steel, austempered during continuous cooling followed by annealing

Abstract

A method for producing an austempered steel is provided. The method includes subjecting a steel alloy having a silicon content of 1.5 to 4.4 weight percent and a carbon content of 0.3 to 0.8 weight percent to continuous cooling followed by annealing. The cooling rate is initially sufficiently fast to prevent predominant formation of proeutectoid ferrite or pearlite, while subsequently at intermediate temperatures, the cooling rate is sufficiently slow to allow a transformation of the austenite to mainly ausferrite during cooling. The annealing is able to complete the transformation of carbon enriched austenite to ausferrite and to temper any martensite previously formed. The method results in the cost-efficient production of one or more continuously cooled and annealed austempered steel components or semi-finished products having mainly an ausferritic microstructure.

Claims

1. A method for producing an austempered steel, comprising: subjecting a steel alloy that has the following composition in weight percent: TABLE-US-00003 C 0.3-0.8.sup.  Si 1.5-4.4.sup.  Mn 0-2.0 Cr 0-2.0 Cu 0-0.4 Ni 0-3.5 Al 0-1.0 Mo 0-0.5 V 0-0.5 Nb 0-0.2 balance Fe and normally occurring impurities, to casting one or more steel components, or hot forging or hot rolling one or more semi-finished steel products; continuous cooling without quenching; and annealing, wherein the continuous cooling begins at a fully austenitic temperature, wherein the cooling rate during the continuous cooling is initially sufficient to result in less than 50% of proeutectoid ferrite and pearlite, wherein the continuous cooling subsequently includes an intermediate temperature range below the proeutectoid ferrite/pearlite region in a continuous cooling transformation (CCT) diagram but above the initial M.sub.s temperature, wherein in the intermediate temperature range, the cooling rate is sufficient to allow a transformation of austenite to at least 50% of ausferrite during cooling, wherein the transformation of austenite to at least 50% of ausferrite occurs before any carbon enriched austenite has reached a temperature below its M.sub.s temperature, thereby limiting the amount of martensite being formed, and wherein the annealing, which is carried out after the continuous cooling, completes the transformation of carbon enriched austenite to ausferrite and tempers any martensite previously formed, the method resulting in production of one or more continuously cooled and annealed austempered components or semi-finished products, whereby the austempered steel contains at least 50% of ausferrite.

2. The method according to claim 1, wherein the continuous cooling comprises cooling naturally in air and/or accelerated cooling and/or decelerated cooling in different temperature ranges.

3. The method according to claim 1, wherein the austempered steel has a microstructure that contains less than 10 volume percent of proeutectoid ferrite.

4. The method according to claim 1, wherein the austempered steel has a microstructure that contains less than 40 volume percent of tempered martensite.

5. The method according to claim 1, wherein the austempered steel has a microstructure that contains less than 1 volume percent of carbides.

6. The method according to claim 1, wherein the austempered steel has a silicon content of 3.1 to 4.4 weight percent and a carbon content of 0.4 to 0.6 weight percent.

Description

BRIEF DESCRIPTION OF THE DRAWINGS

(1) The present invention will hereinafter be further explained by means of non-limiting examples with reference to the appended figure where;

(2) FIG. 1 schematically shows the steps of a method for producing an austempered steel during continuous cooling followed by annealing according to an embodiment of the invention. The dashed M.sub.s line schematically illustrates that during formation of ausferrite, the nucleation and growth of acicular ferrite enriches the surrounding austenite with carbon, thus reducing its M.sub.s temperature during both continuous cooling and during annealing,

(3) FIG. 2 shows the microstructure by light optical microscopy after Nital etching of Example 2 as rolled (a) and after rolling followed by annealing in air for 6 hours (b); the scale bar corresponds to 50 μm in both micrographs,

(4) FIG. 3 shows the fracture surfaces by light optical microscopy of the complete tensile bar cross-sections initially 010.0 mm of Example 2 as rolled (a) and same steel after rolling followed by annealing in air for 6 hours (b),

(5) FIG. 4-5 show the fracture surfaces by scanning electron microscopy of Example 2 as rolled (a) and same steel after rolling followed by annealing in air for 6 hours (b); scale bars correspond to 50 μm in FIGS. 4a and 5a, and 10 μm in FIGS. 4b and 5b, and

(6) FIG. 6 shows stress-strain curves and mechanical properties for as-rolled (curve #1 and the first two rows in the legend) versus rolled and annealed steels using four different combinations of annealing temperature and annealing time (curves #2-5 and corresponding rows 3-10 in the legend). Curves #1 and #3 with their mechanical properties in this figure correspond to microstructures shown in FIGS. 2-5.a and 2-5.b, respectively.

DETAILED DESCRIPTION OF EMBODIMENTS

(7) FIG. 1 shows the steps of a method for producing an ausferritic steel according to an embodiment of the invention.

(8) The method comprises the steps of: (a) continuous cooling from an austenitic state passing the pearlite nose; (b) entering into the austempering intermediate temperature range during cooling; (c) nucleation and growth of acicular ferrite and carbon enriching of austenite with reducing M.sub.s; (d) incomplete transformation into ausferrite stops before cooling to ambient temperature; (e) heating to an annealing temperature; (f) completing the transformation to ausferrite with stabilized austenite having further reduced M.sub.s; (g) cooling to ambient temperature.

(9) The method comprises the step of subjecting a steel alloy having a preferred silicon content of 3.1 to 4.4 weight percent and a preferred carbon content of 0.4 to 0.6 weight percent to either casting of a steel component, or hot forging or hot rolling of a semi-finished steel product.

(10) After either casting of one or more steel components, or hot-working, i.e. hot forging or hot rolling of one or more semi-finished steel products, during which the one or more steel components or semi-finished steel products reach the fully austenitic temperature, the one or more steel components or semi-finished steel products is/are then continuously cooled from the fully austenitic temperature followed by annealing at one or more temperatures to produce one or more continuously cooled and annealed ausferritic steel components or semi-finished steel products. A hot-worked semi-finished product may be continuously cooled on a cooling bed, such as on the cooling bed of a hot-rolling mill for example, and subsequently annealed, in a belt oven or a batch oven for example.

(11) The cooling rate can, especially further down in the austempering temperature range, be decreased (but not prevented) by insulation, such as in the case of a cast component by keeping the cast component in the mould until it has reached a lower temperature before shake-out or even by insulating the mould by covering with a thermally insulating material, such as a blanket comprising refractory ceramic fibre (RCF) or high-temperature insulating wool (HTIW), and in the case of a hot-worked semi-finished product, a plurality of semi-finished hot-worked products may be stacked or placed adjacently to one another during the continuous cooling step and/or even insulated by covering them with a thermally insulating material, such as a blanket comprising refractory ceramic fibre (RCF) or high-temperature insulating wool (HTIW).

(12) The cast steel component, hot forged or hot rolled semi-finished product may be continuously cooled by natural cooling, forced cooling (but not quenching) or delayed cooling in an ambient atmosphere such as air. The continuous cooling may either reach asymptotically one or more temperatures for isothermal treatments, for example by cooling slower in an oven, or continue down to ambient temperature, or be cooled further to lower temperature to deliberately form some amount of martensite.

(13) If cooled to ambient temperature or lower, the steel is thereafter heated and annealed at one or several low austempering temperatures where austenite areas not yet transformed to ausferrite, but having carbon contents intermediate between the initial medium carbon austenite and the films of carbon stabilized austenite in ausferritic areas, will transform to new ausferritic areas having a microstructure similar to ausferrite formed isothermally at same temperature after quench. Concurrently any amount of martensite formed at earlier stages will be tempered and contribute to the strength of the austempered steel.

(14) The method according to the present invention results in the production of austempered steel that has a predominantly ausferritic microstructure. An ausferritic structure is well known and can be determined by conventional microstructural characterization techniques such as, for example, at least one of the following: Optical microscopy, transmission electron microscopy (TEM), scanning electron microscopy (SEM), Atom Probe Field Ion Microscopy (AP-FIM), and X-ray diffraction.

(15) According to an embodiment of the invention the microstructure of the ausferritic steel is substantially carbide-free, or contains less than 1 vol-% of carbides.

EXAMPLE 1

(16) Austempered steel having the following composition in weight percent was produced using a method according to an embodiment of the present invention:

(17) TABLE-US-00002 C 0.5 Si 3.3 Mn 0.5 Cr 0.3 Cu 0.2 Ni 1.6 Mo 0.2 V 0.3
balance Fe and normally occurring impurities, such as 0.012 weight percent P and 0.006 weight percent S.

(18) A 1400 kg rolling ingot having the chemical composition described above was cast vertically in a permanent cast iron mould having an internal height of 1690 mm, top and bottom sections having the dimensions 255×230 mm and 440×350 mm respectively and a conicity of 6.3°×4.1°.

(19) The ingot was subsequently forged into a rolling billet 165×165×4560 mm. Thereafter the billet was hot rolled into round bar having a diameter of Ø53 mm. The cast and forged billet was namely preheated in a furnace at a temperature of 1200° C. for two hours, rough rolled three times and then rolled continuously to a final bar diameter of Ø53 mm. After hot rolling finished at 1040° C., the Ø53 mm round bar was transferred to a walking beam cooling bed next to Ø53 mm round bars previously hot rolled and left to cool continuously during 18 minutes to 460° C., whereafter the bar was cut into 6 m lengths. A few minutes later the resulting nine bars from this billet were bundled together, followed by further air cooling to ambient temperature.

(20) The average cooling rate at 700° C. in Ø53 mm round bars is about 0.7° C./s in still air, but due to the surrounding hot rolled bars at the cooling bed (and no cooling fans) the actual mean cooling rate was 0.5° C./s. This cooling rate resulted in about 2-3% of proeutectoid ferrite formed near the bar surface and about 8-10% of proeutectoid ferrite in the center, while only occasional small areas of pearlite nucleated on the proeutectoid ferrite, since the high silicon content delays cementite formation. These microstructures indicate that the alloy had in this case a slightly too low hardenability for this cooling rate to result in ausferrite only, but if the bar dimension had been smaller and/or the cooling rate around 700° C. had been increased by cooling fans or water spray, the austenite would have been completely preserved for transformations to ausferrite at lower temperatures.

(21) The continuously cooled hot-worked semi-finished austempered Ø53 mm round steel bar had a Vickers hardness of 412±4.7 HV30, where the variations in hardness are mainly reflecting the difference in minor amounts of proeutectoid ferrite as earlier described. This hardness level can be compared with 369±5.2 HV30 in the previously cast and forged mainly pearlitic rolling billet.

(22) When the continuously cooled austempered bar was studied by microscopy, it was found that the mainly ausferritic microstructure (with small amounts of proeutectoid ferrite) also contained some austenitic areas being much thicker than the mainly submicron austenite films within ausferrite. From earlier experiences during development of SiSSADI™ it was concluded that although these austenite areas had been enriched with carbon sufficiently to avoid their transformation to martensite during cooling to ambient temperature (by decreasing M.sub.s temperature below ambient), these areas had not been able to transform completely into ausferrite during the short time within the austempering temperature range during continuous cooling, probably due to compositional variations from segregation since enrichment of carbon and some of the substitutional alloying elements are known to delay the otherwise surprisingly rapid transformation into ausferrite in high silicon medium carbon steels.

(23) Initial mechanical tensile testing verified the conclusions from microstructural observation. The results were as follows: R.sub.p0.2=820.5±7.8 MPa; R.sub.m=1269±19 MPa; A.sub.5=2.71±0.02%. In stark contrast to typical behaviors for fully ausferritic steels, fracture occurred far before necking, indicating the presence of areas of austenite being too low in carbon and too thick to resist their premature strain-induced transformation into martensite, before efficient strain hardening within the ausferritic microstructure has been able to increase plastic elongation and contraction before fracture.

(24) To investigate if the unfinished transformation into ausferrite could be completed, tensile testing bars were subject to an annealing heat treatment at 250° C. for 6 h. This long duration at elevated temperature was permitted since the high silicon content in the steel (3.3% Si) efficiently stabilizes the already formed ausferrite by delaying/preventing any destructive transformation of its high carbon austenite films within the ausferrite into brittle bainite. The hardness of the steel increased by annealing from 412±4.7 HV30 to 431±3.5 HV30. Microstructural observation confirmed that the previous thicker austenitic areas having intermediate carbon content were during the annealing replaced with ausferrite, being much finer than most of the ausferrite earlier formed during continuous cooling (that was mainly nucleated and grown in the beginning of the cooling at higher temperatures when carbon diffusion is more rapid), thereby increasing the hardness.

(25) Tensile testing verified also in this case the conclusions from microstructural observation. The results were as follows: R.sub.p0.2=1118±3.5 MPa; R.sub.m=1447±5 MPa; A.sub.5=23.1±0.9%. Compared to the previous results, the yield strength was much higher, followed by efficient strain hardening within the ausferritic microstructure that resulted in a considerable isotropic plastic elongation up to 18% where an increased ultimate tensile strength was reached, followed by necking and considerable contraction (Z=26.5±0.6%) before fracture.

EXAMPLE 2

(26) An alloy consisting of 0.45 wt % C, 3.33 wt % Si, 1.57 wt % Ni, 0.60 wt % Mn, 0.30 wt % V, 0.29 wt % Cr, 0.21 wt % Cu and 0.20 wt % Mo, was cast into a conical 1.4 tonne ingot with dimensions 1690×(440−255)×(350−230) mm. The ingot was then forged to a cross-section having an area of 165×165 mm, followed by hot-rolling to round bar having a diameter of Ø53 mm.

(27) The surface temperature of the bar was 1010° C. when entering the cooling bed after rolling, and 18 minutes later the surface of the bar had cooled to 461° C., when it was cut/sheared into nine bars having a length of six meters and immediately bundled together for further handling. The cooling time after bundling within the temperature range 460-320° C. was estimated to be about 10 minutes, by using data from the Atlas of Continuous Cooling Transformation Diagrams of Engineering Steels, by M. Atkins, ASM and British Steel Corporation 1980.

(28) Initial hardness testing surprisingly revealed a hardness level much higher than anticipated for a ferritic-pearlitic microstructure, in spite of the considerable substitutional solution hardening of its ferrite by Si and Ni. The first tensile tests of the as-rolled steel showed, however, only a few percent elongation at fracture for tensile strengths varying between 1040-1350 MPa.

(29) Later metallographic work and low temperature annealing of the rolled bar in air revealed some astonishing effects on the mechanical properties which initiated a deeper investigation of the causes. Microstructures before and after annealing were investigated using light optical microscopy and fracture surfaces by SEM (JEOL IT300).

(30) Finally the annealing treatment of the as-rolled bar was investigated for different combinations of temperature and time, followed by tensile testing (DARTEC M1000/RK) at room temperature according to EN ISO 6892-1:2016. The tensile testing bars were 120 mm long with 010 mm cylindrical parts between 022 mm heads. A 50 mm extensometer measured the A.sub.5 elongation during a cross-head speed of 2 mm/min.

(31) Conventional optical metallography after Nital etching revealed a predominantly ausferritic structure, although not fully developed, see FIG. 2.a. Remaining larger bright austenitic “islands” (“blocky” shape in contrast to “film” shape in ausferrite) are carbon enriched (since thermally stable at room temperature), but they have not reached their final carbon content or fine size.

(32) In the core of the rolled Ø53 mm bar about 5% of proeutectoid ferrite (but no pearlite due to the high silicon content) could be seen, indicating that the hardenability of the steel alloy was slightly lower than required for the air cooling experienced for the Ø53 mm diameter steel bar on the cooling bed. After low temperature annealing in air for six hours at a temperature level 30 K below the initial M.sub.s temperature of the austenite, the bright austenite islands had been transformed into fine ausferrite, see FIG. 2.b.

(33) FIGS. 3-5 show fracture surfaces from tensile testing bars. The stereo microscope photos show small “mirrors” in the as-rolled fracture (see FIG. 3.a) and shear lips after necking before fracture in the annealed sample (see FIG. 3.b).

(34) At higher magnification in SEM the as-rolled fracture is dominated by ductile dimple areas but does also contain probably weakening areas (corresponding to the bright austenitic islands in FIG. 2.a) of quasi-brittle fracture mixed with cleavage fracture, see SEM micrographs in FIG. 4. At the higher magnification in FIG. 4.b, the different types of fracture are indicated by arrows: The arrow in the middle for cleavage, the arrow on the right hand-side for quasi-brittle and the arrow on the left hand side for ductile dimples.

(35) After annealing for 6 hours at a temperature level 30 K below the initial M.sub.s temperature of the austenite before entering the cooling bed, the fracture becomes completely ductile, see FIG. 5.

(36) The stress-strain curves and resulting mechanical properties are presented in FIG. 6, together with three additional different combinations of annealing temperature and annealing time after rolling of the same steel.

(37) The as-rolled steel in curve #1 (see mechanical properties presented in row 2 in the legend) yielded early, presumably due to plastic deformation in the softer austenite islands, followed by fracture occurring far before necking. This indicates the presence of austenite being too low in carbon and too thick to resist premature strain-induced transformation into martensite, before the efficient strain hardening within the ausferritic microstructure has been able to increase plastic elongation and contraction before fracture. The scatter in properties was also high, especially for the ultimate tensile strength.

(38) After annealing for 6 h at T={M.sub.s initial−30 K}, the mechanical response became totally different, (see curve #3 and mechanical properties presented in row 6 in the legend). Both the yield strength .sub.Rp0.2 and the ultimate tensile strength R.sub.m increased by about 275 MPa and with very low scatter (standard deviation±4-5 MPa). Furthermore the elongation was isotropic beyond 18% (to R.sub.m) and finally ruptured at 23.7±2%.

(39) Finally, the annealing treatment of the as-rolled bar was investigated for different combinations of temperature and time. Due to difficulties in cutting the as-rolled bar with its mechanically unstable austenite using a HSS band saw (thus requiring expensive sectioning with EDM), only single bars were evaluated (without determination of standard deviation for properties). Both a lower annealing temperature for longer time (see curve #2 and the resulting mechanical properties presented in row 4 in the legend) and two higher annealing temperatures for shorter times (see curves #4 and #5 with mechanical properties presented in row 8 and row 10 in the legend) give similar results, namely substantial improvements of both yield strength and ultimate tensile strength, concurrently with very high ductility.

(40) The hardness of the ferritic-pearlitic forged ingot before rolling was 369±5 HV30. In the as-rolled Ø53 mm bar the hardness in the predominantly ausferritic microstructure formed during continuous cooling (with some less stable austenite islands remaining) increased 35 to 415±5 HV30.

(41) In the Ø53 mm bar rolled+annealed for 6 h at T={M.sub.s initial−30 K} the hardness increased further to 431±4 HV30.

(42) The small hardness increase during annealing corresponds well with microstructural observations that the as-rolled microstructure was already predominantly ausferritic. The very fine ausferrite subsequently formed during annealing can therefore only raise the hardness slightly, in spite of its high hardness being probably far above 500 HV, since the very fine ausferrite represents only a few volume percent.

(43) This is also the reason why ausferrite formed at various temperatures (see FIG. 6) results in similar mechanical properties. If less ausferrite has time to form during the previous continuous cooling, the influence from annealing temperature would be larger.

(44) To find out in which temperature range ausferrite was mainly formed during the continuous cooling of the hot rolled bar, a comparison was made with the same steel alloy after conventional austempering by complete austenitization followed by quenching and isothermal transformation in a salt bath held at T={M.sub.s initial+20 K}. The resulting hardness of the isothermally formed ausferritic steel was 490±5 HV30.

(45) Based on the inventor's experience of hardness dependence on isothermal transformation temperature, this implies that the as-rolled ausferrite structure established but not completed during continuous cooling would correspond to a far higher salt bath temperature of T={M.sub.s initial+95 K}. Furthermore, the strength levels of ausferritic steels isothermally transformed at such high salt bath temperatures are similar to levels in these rolled+annealed steels, in which mechanically unstable austenite areas have been eliminated.

(46) The advantages offered by the method for producing ausferritic steels according to the present invention can be summarized as follows:

(47) Quenching followed by isothermal transformation in salt baths are not necessary, on condition that the cooling rate of the steel around the eutectoid temperature is sufficiently rapid relative to the hardenability of the alloy to preserve most of the austenite for consecutive transformation to predominantly ausferrite during continuous cooling within the austempering temperature range.

(48) Continuous cooling in air (instead of quenching in liquids) followed by annealing at low temperatures reduces both residual stresses and production costs, while enabling very strong, ductile and tough ausferritic steels to be delivered in lengths exceeding 20 meters directly from rolling mills combined with low temperature belt ovens.

(49) The annealing is able to complete the transformation of austenite to predominantly ausferrite, on condition that carbon diffusion during the previous continuous cooling has sufficiently stabilized the remaining larger areas of austenite against transformation to more than minor amounts of martensite if cooled to ambient temperature, or cooled further to deliberately form martensite before annealing, where the transformation into ausferrite is completed concurrently with low-temperature tempering of any martensite, avoiding temper embrittlement in this temperature range due to the high silicon content.

(50) The annealing thus reduces the need to decrease cooling rates within the austempering temperature range in order to complete the transformation into ausferrite within current production processes such as casting, forging and rolling, while the subsequent annealing at low temperature in air in batch ovens or belt ovens may result in extremely good mechanical properties with small scatter.

(51) If martensite is formed during the continuous cooling to ambient temperature or deliberately lower temperatures it becomes tempered during the annealing, thus contributing to even higher strength of the predominantly ausferritic steel.

(52) Further modifications of the invention within the scope of the claims would be apparent to a skilled person. For example, it should be noted that any feature or method step, or combination of features or method steps, described with reference to a particular embodiment of the present invention may be incorporated into any other embodiment of the present invention.