Patent Application
20240141465
This disclosure relates to a martensitic steel, and in particular to a martensitic steel alloyed with high aluminum, manganese and silicon contents.
Steel producers strive to produce steels with high mechanical resistance and high toughness. Such steels are for instance suitable as tool steels. For many applications of tool steels high hardness is required to fight the abrasive wear mechanism, while a reasonable toughness is desired to avoid the development of cracks during the mechanical work.
Known steels, however, always feature a tradeoff between the mechanical resistance and the toughness. That is, higher toughness is usually obtained by lowering the mechanical resistance due to the intrinsic behavior of the plasticity of the materials. For applications where it is necessary to achieve higher toughness at a targeted mechanical resistance, usually the thickness of the steel part is increased to reach the targeted mechanical resistance. However, the increase of thickness adds weight and cost to the steel part and may result in some loss on the hardenability due to higher thickness.
Further, there are tool steel applications such as knifes and saws where the design of the parts is fixed. In these cases it is necessary to improve the properties of the steel instead of performing changes in the design of the part.
Conventional tool steels typically use a martensitic matrix with high Cr content (>1.0 wt. %) and high additions of Mo (>0.5 wt. %), Ni (>0.5 wt. %) and carbide former elements (V, Nb, Ti). The mechanism of these steels relies on a solid solution strengthened martensite with precipitation of primarycarbides, i.e. formed in liquid phase, and a high volumetric fraction of secondary hardening carbides, which precipitate during subsequent tempering heat treatment.
One of the common steels employed in knife applications is DIN 1.2360 steel, which presents higher hardness at a reasonable toughness. Other examples of steels like AISI D2, AISI S7 and TENAX300®, a modified AISI H11 steel (DIN 1.2365), are usually applied for molds and dies. In order to develop high hardness these steels include considerable amounts of expensive alloying elements (e.g. Cr, Mo, W, V, Nb, Ti . . . ) to allow secondary precipitation of carbides as well as to increase the hardness of the martensite due to solid solution strengthening. High amounts of alloying elements, however, lead to expensive tool steel products.
According to an aspect of the disclosure a martensitic steel consists of, in % in weight: C: 0.30 to 0.80%, Si: 2.50 to 4.50%, Mn: 1.00 to 2.50%, Al: 0.40 to 1.50%, Cr: 0.10 to 2.00%, V: 0.01 to 0.40%, Ti: 0.005 to 0.35%, and optionally one or more of Nb: less than 0.35%, Zr: less than 0.35%, Ta: less than 0.35%, P: less than 0.25%, S: less than 0.25%, Co: less than 0.50%, Mo: less than 0.90%, W: less than 0.90%, Ni: less than 0.50%, Cu: less than 0.50%, N: less than 0.050%, Ca: less than 0.10%, Mg: less than 0.10%, Ce: less than 0.10%, La: less than 0.10%, B: less than 0.10%, the balance Fe and impurities, and comprising one or more intermetallic phases based on an Al—Fe—Mn—Si system.
According to another aspect of the disclosure a method of manufacturing a martensitic steel comprises providing a hardened and quenched steel having a composition as set out above; and tempering the hardened and quenched steel at a temperature preferably in a range between 300° C. and 600° C.
The steel disclosed herein uses high amounts of silicon, manganese and aluminum at the same time. According to the literature, such combination is believed to cause embrittlement of the steel and consequently a reduction of its toughness. The present disclosure teaches that balanced amounts of these elements with addition of some grain boundary stabilizers allow a high hardness coupled with a very high toughness.
In other words, the steel as disclosed herein goes against the expected behavior of the traditional steels. Traditional steels always have a tradeoff between toughness and hardness. Considering the hardened, quenched and tempered state, increasing the hardness of traditional steels always reduces the toughness of the steel. However, the steel as disclosed herein, after hardening, quenching and tempering heat treatments, increases the toughness with the increase of the hardness.
The importance and properties of the constituent chemical elements as well as their compositional ranges in the claimed steel are described in the following. Throughout this description and the claims, all percentages of the chemical composition are given in percentage in weight (wt. %). The upper, intermediate and lower limits of the individual elements can be freely combined within the compositional ranges set out in the claims.
Carbon (C: 0.30-0.80%) is responsible for improving the strength and the hardenability of the steel. During hardening, the matrix is mainly composed of austenite phase, which, after quenching, will transform into the martensite phase leading to a high hardness matrix but with lower toughness. This martensite, after tempering heat treatment, will be conditioned and the steel will present a higher toughness. The increase of carbon contents has the effect to increase the martensite start temperature (Ms). However, too high carbon contents deteriorate the weldability and the elongation of the steel. Carbon is important for the precipitation of carbides (e.g. VC, TiC), which enhance the wear resistance but can cause a reduction of the toughness of the steel of the present disclosure. For the steel of the present disclosure, carbon is desired between 0.30% and 0.80% being preferable between 0.40% and 0.60%. The upper limit for carbon may be set to 0.80% or 0.70% or 0.60%. The lower limit may be set to 0.30% or 0.35% or 0.40%. It is to be noted that the martensitic steel according to the disclosure does not comprise carbon in the form of graphite.
Silicon (Si: 2.50-4.50%) is usually present in the steels due to the deoxidation processes. For the steel of the present disclosure, the silicon is added with the aim of improving the oxidation resistance as well as retarding the eutectoid decomposition of the austenite. Further, silicon inhibits the precipitation of M3C type carbides and thus keeping the carbon in solid solution. The synergic effect of silicon with aluminum enhances the oxidation and corrosion resistance of the steel. With the increase of silicon content, the nitridability is reduced due to the effect of the silicon over the atomic mobility of interstitial elements, in special to the nitrogen. However this effect is counteracted by the high aluminum addition which enhances the chemical potential of nitrogen and reduces the deleterious effect of the high silicon addition. Silicon may form intermetallic phases such as Al2Mn2Si3. The silicon content of the steel of the present disclosure is desired to be between 2.50% and 4.50%, being preferable between 3.00% and 4.00%. The upper limit may be set to 4.50% or 4.40% or 4.30% or 4.20% or 4.10% or 4.00%. The lower limit may be set to 2.50% or 2.60% or 2.70% or 2.80% or 2.90% or 3.00%.
Manganese (Mn: 1.00-2.50%) is an important element of the steel of the present disclosure due to its ability to improve the hardenability coupled with an enhancement of the hot workability and toughness. Further, the addition of manganese improves the mechanical strength through a solid solution mechanism and stabilizes the residual sulfur in the matrix as MnS. The addition of manganese in excess is desirable to increase the mechanical resistance of the matrix and allows precipitation of intermetallic phases such as Al2Mn2Si3. Manganese shall therefore be present in a minimum content of 1.00%, preferably at least 1.1% or 1.2% or 1.3%. For the steel of the present disclosure, manganese is desired between 1.00% and 2.50%, preferable between 1.30% and 1.80%. The upper limit for the manganese content may be 2.50% or 2.30% or 2.20% or 2.10% or 2.00% or 1.90% or 1.80%.
Aluminum (Al: 0.40-1.50%) is an indispensable element of the steel of the present disclosure. The addition of aluminum promotes the formation of a passive oxide layer on the surface to enhance the oxidation resistance of the steel coupled with a better nitridability due to the precipitation of the AIN nitrides during the nitriding processes. Due to the addition of aluminum coupled with silicon and manganese, the martensite of the steel attains high hardness with a high toughness due to precipitation of intermetallic phases such as, e.g., Al2Mn2Si3 phase. For the steel of the present disclosure, aluminum is desired between 0.40% and 1.50%, being preferably between 0.70% and 1.20%. The upper limit may be set to 1.50% or 1.45% or 1.40% or 1.35% or 1.30% or 1.25% or 1.20%. The lower limit may be set to 0.40% or 0.45% or 0.50% or 0.55% or 0.60% or 0.65% or 0.70%.
Chromium (Cr: 0.10-2.00%) is an important element for the steel of the present disclosure to perform the fine-tuning of the martensite start temperature. An excess in the chromium addition will promote precipitation of chromium carbides of M3C, M23C6 and M7C3 types and this is not desirable for the steel of the present disclosure since this reduces the toughness of the martensite matrix. Chromium contents lower than 0.30% will not allow the fine-tuning of the martensite start temperature. For the steel of the present disclosure, chromium is desired between 0.10% and 2.00%, being preferably between 0.30% and 0.80%. The upper limit may be set to 2.00% or 1.50% or 1.00% or 0.95% or 0.90% or 0.85% or 0.80%.
Vanadium (V: 0.01-0.40%), Titanium (Ti: 0.005-0.35%), Niobium (Nb: <0.35%), Zirconium (Zr: <0.35%), Tantalum (Ta: <0.35%) are strong carbide formers that improve the hot mechanical resistance and the wear resistance of the steel. Higher contents of these elements are not desirable due to the precipitation of large MC carbides and the reduction of the steel toughness. In lower contents these elements are desirable due to its effect of grain boundary pinning and therefore to reduce grain coarsening. For the steel of the present disclosure, vanadium is desired between 0.01% and 0.40%, being preferably between 0.01% and 0.35%, most preferably between 0.01% and 0.15%. The upper limit for vanadium may be set to 0.40% or 0.30% or 0.25% or 0.20% or 0.15%. Titanium is desirable between 0.005 and 0.35%, preferably between 0.05 and 0.35%, more preferably between 0.08 and 0.25%, most preferably between 0.10% and 0.15%. Niobium is desired to be lower than 0.35%, preferably lower than 0.15%. Niobium is optional and may not be deliberately added. For zirconium and tantalum, the same applies as for niobium.
Sulfur (S: <0.25%) improves the machinability of the steel of the present disclosure and is a residual of the steelmaking process. The addition of sulfur for the alloy of the present disclosure is incidental and the sulfur content must be lower than 0.25%, preferably lower than 0.10%. For the steel of the present disclosure, the sulfur content is desired to be below 0.10% and preferably lower than 0.05%.
Phosphorous (P: <0.25%) is effective in strengthening of the steel by solid solution. However it reduces the toughness of the steel and must be controlled to be below 0.25%. For the steel of the present disclosure, the phosphorous content is desired to be below 0.050%, preferably lower than 0.035%.
Nitrogen (N<0.05%) as carbon is intended to improve the solid solution strengthening and the mechanical strength of the steel of the present disclosure. Nitrogen, when added in amounts higher than 0.05%, will provide the precipitation of Cr2N, AIN and TiN that are not desirable in the matrix of the steel according to this disclosure. Addition of nitrogen between 0.001% and 0.050% will enhance the matrix without promoting a high volumetric fraction of nitrides but with sufficient amount in order to allow the grain growth control by the mechanism of pinning of the grain boundaries, improving consequently the fatigue resistance of the steel. Lower additions than 0.0010% are impracticable due to the higher cost of melting, refining and processing of the steel. For the steel of the present disclosure, nitrogen is preferably desired to be lower than 0.050% and most preferable lower than 0.020%.
Cobalt (Co: <0.50%) presents very similar properties compared to nickel, i.e. causes the same effects and the same intermetallic compounds that can be formed, i.e. of Co3Al and Co3Ti types. Additionally, cobalt also is an impurity commonly present in nickel ores, being frequently found as a residual of the main sources of nickel for the alloys production. For the steel of the present disclosure, cobalt is desired to be lower than 0.50%, more preferably lower than 0.20%, most preferably lower than 0.05%. Co is optional and may not be deliberately added.
Molybdenum (Mo: <0.90%) and tungsten (W: <0.90%) are optional. They are responsible for the improvement of the hot mechanical properties and promote the precipitation of M2C carbides during the tempering heat treatments. Higher molybdenum and tungsten contents over 0.90% are not desirable due to the reduction of the hot workability of the steel, precipitation of M2C carbides and higher cost of the steel. The upper limits may be set to 0.90% or 0.50% or 0.30% or 0.20% or 0.10%. However, molybdenum and tungsten contents lower than 0.01% may be costly due to the use of scrap of steels in the elaboration process. For the steel of the present disclosure, molybdenum and tungsten are desired to be lower than 0.20%, being preferably lower than 0.10%. Most preferably, no Mo or W additions are made.
Nickel (Ni: <0.50%) is intended to be added lower than 0.50% in order to avoid the precipitation of intermetallic phases like Ni3Al and Ni3Ti combined with aluminum and residual titanium present in the steel. Nickel contents lower than 0.01% are not desirable due to the characteristic of the scrap and iron-alloys used to compose the composition of the steel. For the steel of the present disclosure, nickel is desired to be lower than 0.50%, preferably lower than 0.20%, most preferably lower than 0.05%. Ni is optional and may not be deliberately added.
Copper (Cu: <0.50%) is responsible for the enhancement of the corrosion resistance of the steels and for the improvement of the machinability. Higher copper contents than 0.50% are not desirable because of the precipitation of spherical copper precipitates, which reduce the hot mechanical strength. However, copper contents lower than 0.01% may be costly due to the use of scrap of steels in the elaboration process. Copper is desired to be lower than 0.50%, more preferably lower than 0.20%, most preferably lower than 0.10%. Copper is optional and may not be deliberately added.
Calcium (Ca: <0.10%), magnesium (Mg: <0.10%), cerium (Ce: <0.10%) and lanthanum (La: <0.10%) are unavoidable impurities in the steelmaking process of production of the steel as disclosed herein. Their concentration is desirably below 0.10%, preferably below 0.05%, even more preferably below 0.01% in order to avoid intermetallic phase precipitation or the interference with the desired properties of the steel of the present disclosure. These elements are commonly used as deoxidizer and desulfurizer of the steel during the melting refinement. For the steel as disclosed herein, calcium, cerium, lanthanum and magnesium are desired below 0.10%, preferably below 0.05%.
Boron (B: <0.10%) may be used in order to increase the hardness and the hardenability of the steel of the present disclosure. The amount of boron is limited to 0.10%, preferably 0.01%, more preferably to 0.0050%.
The steel as disclosed herein is based on a new concept that allows a cost reduction through the reduction of the amount of expensive alloying elements contents in comparison with traditional steels. The steel of the present disclosure contains comparatively high amounts of silicon, manganese and aluminum, which were discovered to increase the hardenability of the matrix in specific balanced amounts. This combination of aluminum, silicon and manganese also allows the precipitation of intermetallic phases which considerably improve the mechanical resistance without compromising the toughness. The literature, though describing to use high amounts of each one of these elements, teaches that the use of these elements shall be performed individually to avoid the embrittlement of the steel. Within the new concept taught herein, it was discovered that these elements have a very interesting synergic effect that allows achieving higher hardness but keeping a high toughness at the same time.
This allows to produce tool steels used, e.g., in knifes, e.g. knifes for sugar cane harvesting, saws, molds, dies, valves for hot or cold work applications with a martensitic matrix containing significantly lower contents of one or more of the elements of the group consisting of chromium, molybdenum and carbide former elements (vanadium, niobium, titanium, etc.) than in traditional work steels.
Differently stated, the concept of the steel of the present disclosure is to use a martensitic steel substantially without secondary hardening carbide precipitation to increase the toughness of the steel (since carbide precipitations are the main mechanism for embrittlement of steels). In other words, while some carbides (e.g. Ti-carbides and/or Nb-carbides) may be present in the steel as primary carbides, i.e. formed in the liquid phase during cooling/solidification of the molten steel, the martensitic steel according to the disclosure is (substantially) free from secondary hardening carbides, i.e. carbides formed during tempering heat treatment, such as, e.g., V- and/or Ti- and/or Nb-secondary hardening carbides. The reduction of cost of the steel as disclosed herein is mainly achieved by using high amounts of cheaper elements for the steel matrix, namely manganese, aluminium and silicon, in addition to carbon. Low chromium addition may be used to fine-tune the hardenability of the steel. Low additions of carbide former elements (vanadium, niobium, titanium) may only be used for the purpose of grain boundary pinning. Other elements such as molybdenum, tungsten, nickel, copper, etc. may be kept as low as possible.
The steel described herein was conceived essentially in the hardening of the martensite assuring a high hardness of the matrix, i.e. the steel material surrounding the carbides, with a minimum amount of carbides to guarantee the toughness of the steel and to control the grain boundary pinning during hot working of the steel. To add some extra hardness, intermetallic precipitation of phases rich in manganese and silicon coupled with aluminum were observed to improve the mechanical resistance of the steel. These precipitated phases were found to be of rich in Mn, Al and Si. Especially, ternary Al—Mn—Si intermetallic phases such as, e.g., Al2Mn2Si3 phases were identified by X-ray diffraction as will be described in more detail further below. Generally, the steel disclosed herein includes one or more precipitated intermetallic phases based on the Al—Fe—Mn—Si system, e.g. Al4Mn1Si2, Al4Mn1Si1, Al9Mn3Si1, Al2Mn2Si3, Al17Fe3.2Mn0.8Si2, α-Al8.36Mn2Si1.14, α-Al4.01Mn1.0Si0.74 or Al17Fe3.2Mn0.8Si2.
The steel as disclosed herein features an inversion of the tradeoff between hardness and toughness. While known steels show a reduction of the toughness with the increase of hardness, the steels disclosed herein increase the toughness when increasing the hardness.
Further, common steels usually rely on secondary hardening by which high amounts of alloying elements M have precipitation M23C6 and M2(C,N) and M6C carbides. The steel as disclosed herein may substantially not feature alloying elements carbide precipitation during the tempering heat treatment. Only martensite conditioning and at most precipitation of iron based carbides (E-type) may be found.
The martensitic steel with high hardness and high toughness according to the present disclosure can be produced through conventional (electric arc furnace or induction air furnace) or especial (vacuum induction melting) melting process being conventionally or continuously casted. The ingots or billets are heated up to the temperature of forging and/or hot rolling process depending upon the size of the ingot or billet and hot worked to the desired shape and diameter for the final product. The resulting bars are heat treated, finished and inspected. Components for hot or cold work such as, e.g., knifes, saws, molds, dies, internal combustion valves etc. . . . can be produced with the steel of the present disclosure. Differently put, the steel of the present disclosure may, e.g., be a hot-work steel and/or a cold-work steel, in particular a hot-work tool steel and/or a cold-work tool steel.
The steel of the present disclosure may be heat treated for different hardness depending upon the heat treatment cycles used and the desired application.
Hardening can be performed at temperatures between 850° C. and 1100° C. for a time commensurate with the thickness of the part followed by quenching in air, oil or water. This heat treatment is intended to produce a fully austenitic structure during the high temperature exposure. During quenching, this austenite will transform mainly in martensite and some retained austenite. Depending upon the cooling rate employed during quenching, some bainite may form, however for applications which requires high hardness a fully martensitic structure is preferable.
Accordingly, the claimed “martensitic steel” may either have a fully martensitic structure or a predominantly martensitic structure with some retained austenite and/or generated bainite contributions. Hardening temperatures are preferably between 850° C. and 1050° C., and most preferably between 900° C. and 1050° C. Hardening times may range from 2 minutes to several hours, about 1 hour hardening time is preferred.
After hardening and quenching the steel can optionally be tempered in temperatures between 100° C. and 600° C. for a time commensurate with the thickness of the part. For instance, a tempering time of about 2±1 h (hours) may be used, wherein the time count starts when the part is uniformly heated at the tempering temperature. Lower tempering temperatures are preferable to increase the toughness of the steel and keep the high hardness, the upper limit being preferably 500° C., more preferably 450° C., even more preferably 400° C., 350° C. or 300° C. Most preferably, the tempering temperature is between 150 to 300° C. to achieve higher hardness coupled with toughness. Due to the balanced amounts of alloying elements, the steel of the present disclosure may not present a significant secondary hardening, therefore exhibiting usually only a reduction of the hardness with the increase of the tempering temperature (and also a decrease of its toughness).
Tempering may be carried out at the customer's location, i.e. after the hardened and quenched but not yet tempered steel has been shipped to the customer. Optionally, annealing (also referred to as “soft annealing”) can be performed before hardening and tempering. Annealing softens the steel and makes it easier to machine processing (machining, turning, milling, drilling, . . . ). An annealing heat treatment can be performed at temperatures between 650° C. and 900° C. for a time commensurate with thickness of the part (e.g. 2 hours after uniform heating) followed by air-cooling or even a slower cooling rate.
After hardening and quenching, the steel of the present disclosure achieves a hardness between e.g. 29 HRC (Hardness Rockwell scale C—also referred to as Rockwell C hardness) and 65 HRC depending upon the hardening temperature and the quenching media. It is preferable that the hardness in the as quenched state is e.g. between 45 HRC and 65 HRC, being most preferable to be e.g. between 55 HRC and 65 HRC.
After tempering heat treatment the steel of the present disclosure may present a hardness between e.g. 30 HRC and 58 HRC being preferable to be between e.g. 40 HRC and 58 HRC and most preferable to be e.g. between 45 HRC and 58 HRC.
In the hardened, quenched and tempered state, the steel of the present disclosure presents impact energy, measured in accord with VDG M82 standard for an unnotched impact specimen, higher than 120 J/cm2, preferable higher than 140 J/cm2, preferably higher than 160 J/cm2, preferably higher than 200 J/cm2 and most preferably higher than 250 J/cm2. For instance, it was observed that when the steel of the present disclosure is hardened between 950° C. and 1050° C. followed by oil quenching and tempering between 300° C. and 450° C., the steel of the present disclosure is able to deliver impact energy higher than 250 J/cm2 with hardness higher than 55 HRC.
For the hardened, quenched and tempered condition, the steel of the present disclosure may present yield strength (YS) higher than 900 MPa, preferably higher than e.g. 1200 MPa and most preferably higher than e.g. 1500 MPa for room temperature tensile tests in accord with ASTM A370 standard. The ultimate tensile strength (UTS) may be higher than 1000 MPa, preferably higher than e.g. 1300 MPa and most preferable higher than e.g. 1700 MPa. The elongation in 4D (A4D) may be higher than 4%, preferably higher than e.g. 6% and more preferably higher than e.g. 10%. The reduction in area (RA) may be higher than 10%, preferably higher than e.g. 15% and more preferably higher than e.g. 20%.
The bending proof strength, evaluated in accord with the ASTM E855 standard, using specimens with 5 mm by 7 mm cross section, for the steel of the present disclosure in the hardened, quenched and tempered condition may be higher than 3000 MPa, preferably higher than e.g. 3500 MPa, and more preferably higher than e.g. 4000 MPa.
The steel of the present disclosure can also be coated through conventional process(es) such as CVD (Chemical Vapor Deposition), PVD (Physical Vapor Deposition), or by forming a diffusion layer via gas nitriding, plasma nitriding, carbonitriding, case hardening, oxidation followed by nitriding or nitriding followed by oxidation and similar deposition processes improving its surface properties. Due to its high aluminum content, it is expected that the steel of the present disclosure develops an outstanding behavior and achieve very high surface hardness after nitriding process due to precipitation of aluminum nitrides.
In the following examples, exemplary steels (which can be used as tool steels) according to the present disclosure are compared to known reference steels. The chemical compositions of the exemplary steels (Examples 1-7) and reference steels (DIN 1.2360 and TENAX300®) are presented in Table 1. All of the compositions were vacuum induction melted and conventionally casted into 25 kg ingots under vacuum. The ingots were heated up to 1180° C. and hot rolled into 40 mm square bars. The bars were cut in order to obtain specimens for heat treatments, metallographic characterization, Rockwell C hardness tests, tensile tests, impact tests and four point bending tests.
For each steel of Table 1, the hardening curve was determined with different quenching media. Specimens with cross section of 20 mm×20 mm and thickness of 10 mm were hardened in temperatures between 800° C. and 1100° C. for 1 h in temperature followed by water or oil quenching. The Rockwell C hardness of these specimens was measured in accord with ASTM A370 standard and is presented in Table 2.
As can be observed from Table 2, all of the steels of Examples 1-7 of the present disclosure achieved hardness over 58 HRC in at least one hardening temperature between 850° C. and 1100° C. From the highest hardness of each steel, its hardening temperature was defined. The defined hardening temperature (e.g. between 950° C. and 1000° C. for Examples 1-7) and quenching media were used as previous state for the following tempering heat treatments.
That is, after hardening and quenching in accord with the defined hardening cycles of Table 2, the steels were tempered at different temperatures between 300° C. and 600° C. for 2 h in temperature followed by air cooling. The Rockwell C hardness was measured in accord with ASTM A370 standard and the results are also presented in Table 2. The results indicate that the steels of the present disclosure achieved hardness over 55 HRC for tempering temperatures between 300 and 400° C. for all of the compositions with silicon content over 2.70%. With exception of steels of Examples 3 and 6 that have a silicon content less than 2.70% and did not achieve hardness over 55 HRC for tempering at 400° C., however, achieved hardness over 55 HRC for tempering at 300° C. It can be also observed that the reference steels (1.2360 and TENAX300®, a modified AISI H11 steel similar to DIN 1.2365) exhibit secondary precipitation indicated by the increase of the hardness for tempering temperatures over 400° C. This effect is a consequence of the higher alloying element content of these steels that promotes the secondary precipitation of carbides.
Some of the intermetallic phases presented in Table 3 exhibit variations of compositions due to some solubility of other elements like, e.g., iron. However, low amount of ICSD data is available for the Al—Fe—Mn—Si system for compositions rich in iron, since in literature the Al—Fe—Mn—Si system was mainly studied for aluminum alloys.
From the X-ray diffraction patterns of
Table 4 presents the results of tensile tests performed in accord with ASTM A370 standard and Charpy impact tests without notch at room temperature. Both tests were made in accord with VDG M82 standard for the steels of Examples 2, 3 and 4. From these results, it can be observed the synergic effect of the aluminum, silicon and manganese additions on the mechanical properties of the steels of the present disclosure.
The steel of Example 4 presents lower aluminum content (0.415%) in comparison with the other examples of the present disclosure. Aluminum improves the mechanical resistance due to precipitation of intermetallic compounds and hence indicates the reason because for the steel of the present disclosure, the aluminum content is desirable to be higher than 0.50% in weight percent. Further, it can be seen that Charpy impact tests without notch for steels of Examples 2, 3 and 4 yielded an absorbed energy over 160 J/cm2.
Further, a high yield strength coupled with a reduction of area higher than 10% is desired for the steels as disclosed herein. As set out in Table 4, only the steel of Example 4, when tempered at 400° C., had a yield strength less than 900 MPa which is desired as a lower limit in many applications.
As a way to validate the results of Charpy impact tests without notch, the steels were also evaluated by four point bending tests in accord with ASTM E855 using specimens with cross section of 5 mm by 7 mm. The results of the bending proof strength (in MPa) of Table 5 indicate that the material exhibits a high toughness and high hardness in comparison with the traditional steels presenting higher bending proof strength values for the higher hardness specimens. The bending proof strength is directly proportional to the toughness evaluated by unnotched Charpy impact tests. However it is more accurate than the unnotched Charpy impact tests for materials with high hardness and limitations on the elastoplastic behavior, and hence may be a better reference for the toughness of the steels as disclosed herein.
Although specific embodiments have been illustrated and described herein, it will be appreciated by those of ordinary skill in the art that a variety of alternate and/or equivalent implementations may be substituted for the specific embodiments shown and described without departing from the scope of the present invention. This application is intended to cover any adaptations or variations of the specific embodiments discussed herein. Therefore, it is intended that this invention be limited only by the claims and the equivalents thereof.
| Number | Date | Country | Kind |
|---|---|---|---|
| 21159990. 7 | Mar 2021 | EP | regional |
| Filing Document | Filing Date | Country | Kind |
|---|---|---|---|
| PCT/BR2022/050066 | 2/26/2022 | WO |
