Method of Heat Treatment of Medium- and High-Carbon Alloy Steel and Use Thereof

The subject matter of the exemplary arrangements is a method of heat treatment of medium- and high-carbon alloy steel, the aim of which is to produce a fragmented, multiphase microstructure consisting of tempered bainitic ferrite and martensite, separated by layers of stable residual austenite and finely dispersed alloy carbide precipitates responsible for secondary hardness. The exemplary microstructure may also contain primary and/or secondary carbides. The use of the exemplary method of heat treatment to tool steels makes it possible to achieve enhanced resistance to brittle fracture, high hardness, ductility, high strength parameters and resistance to abrasive wear.

To date, the commonly used heat treatment for tool steels to impart the desired mechanical and functional parameters is martensitic quenching combined with single or multiple tempering. This process involves gradually heating the steel to austenitizing temperature, holding it at this temperature and then cooling, usually in oil, fluidized bed or air. Thus quenched steel is subjected to tempering, i.e. annealing at an elevated temperature, usually from the range of 400-600° C., depending on the grade of steel. The effect of tempering is the reduction of martensite embrittlement, microstress relaxation, as well as carbide precipitation, which contribute significantly to the strengthening of the steel. In tool steels with higher carbon content, the precipitation processes result in destabilization of residual austenite, which undergoes martensitic transformation during cooling. In this case, it is necessary to carry out several tempering runs to eliminate the steel embrittlement associated with the presence of the newly formed martensite and to further reduce the fraction of residual austenite. As a result of this treatment, a microstructure is produced consisting of primary carbides that were not dissolved during austenitization, tempered martensite, secondary and tertiary carbides and a small content of residual austenite, usually up to 10% by volume. Such a treatment is presented, in patent PL101241 and, in combination with mechanical treatment, in patents PL116046, PL78487, and PL123120.

Tool steels after toughening are characterized by high hardness, good abrasion resistance and retention of mechanical properties at elevated operating temperatures. However, the main disadvantage of this solution is the occurrence of deformations, or quenching cracks resulting from the introduction of internal stresses into the structure as a result of martensitic transformation. In addition, the relatively low ductility and resistance to fracture of such treated pieces significantly limit their service life.

A method of steel strengthening by employing the phenomenon of secondary and tertiary carbide precipitation was used in WO99/02747 which corresponds to U.S. Pat. No. 5,900,075 which is incorporated herein by reference in its entirety. This method is used to shape the properties of high-strength steel with the following chemical composition (% by weight): 0.03-0.12% C, 0.1-0.5% Si, 0-2.0% Cu, 0.5-2.0% Ni, 0.03-0.12% Nb, 0.03-0.15% V, 0.2-0.8% Mo, 0.3-1.0% Cr, 0.03-0.05% Ti, 0.01-0.05% Al. In this steel, as a result of continuous cooling to temperature above 400° C., followed by high tempering, a microstructure is produced that includes tempered martensite, granular bainite and finely dispersed carbides and carbonitrides. A similar microstructure can be obtained by using the heat treatment described in application CN109207693A, which consists of continuous cooling from austenitization to room temperature followed by freezing or low/medium tempering in the temperature range of 150-550° C. to eliminate the residual austenite and high tempering, the parameters of which depend on the grade of steel. This method of treatment is intended for low-alloy steels with a carbon content of less than 0.5% by weight and a fraction of alloying additives in the form of (% by weight): Cr≤4%, Mo≤1.5%, V≤0.5%, Mn≤1.5%, and Ni, Si and traces of Nb, Ti, B, the total content of which does not exceed 5% by weight.

EP2126150, which corresponds to US Patent Publication 20110085930 which is incorporated herein by reference in its entirety, on the other hand, disclosed cold work tool steel with a fully or at least partially martensitic and/or bainitic structure strengthened by primary carbide, secondary carbide and intermetallic phase precipitations. It was also produced by heat treatment consisting of quenching under continuous cooling and one or more tempering steps. The steels described are characterized by improved weldability, as well as high hardness, resistance to abrasive wear and good impact strength. In these inventions, the aim is to eliminate as much residual austenite as possible, since it is treated as an undesirable component of the microstructure and partial bainitic transformation is not used as a tool to fragment the original austenite grain and control the content of phase components in the final microstructure.

From GB999866A there is known a heat treatment method for high-speed steel containing (in % by weight): 0.55-1.06% C, 3.5-5.5% Cr, 5.0-24.0% W, max. 14.0% Co, 0.75-6.00% V, max. 8.00% Mo. The described method involves four consecutive steps: (I) gradual austenitization, (II) partial bainitization by isothermal quenching at the temperature at which the bainitic transformation has the shortest incubation period ±5° C. for the time necessary to produce about 10% by volume bainite, (III) martensitic quenching to −80° C. or lower, (IV) followed by one or more high tempering at 550° C. The result of this treatment is a multiphase microstructure consisting of tempered bainite, martensite and secondary carbides. It is characterized by high hardness, impact strength and resistance to abrasive wear. The heat treatment method proposed in GB999866A is dedicated to high-speed steel that does not contain silicon or aluminum, which cause carbides to precipitate from the bainitic ferrite during the bainitic transformation and results in lower and/or upper bainite, which is characterized by higher embrittlement compared to carbide-free bainite. The high-speed steels do not contain Si or Al, i.a. to avoid having austenite in the final microstructure. In the cited reference, cooling to −80° C. is used to transform to martensite min. 95% of the remaining austenite. Additional reduction of the austenite content is achieved by using multiple tempering resulting in the carbide precipitation and destabilization of the remaining austenite.

CN111876569A discloses a heat treatment method in which, as above, the strengthening of the steel is realized through the secondary carbide precipitation, bainitic transformation and significant fragmentation of the multiphase microstructure. In the process presented, however, the bainitization step follows the martensitic quenching and tempering operation, making it possible to eliminate the disadvantages of lower bainite tempering, thereby increasing the impact strength of H13, 42CrMo, Cr12MoV steels while maintaining their very high hardness. In the cited invention, martensitic quenching is used after austenitizing, so that the size of the martensite packets formed, especially in the initial phase of the transformation are limited only by the size of the original austenite grain. As a result, the possibilities for martensite fragmentation and further improving the strength and resistance to fracture of the steel are limited.

A description of chemical compositions and a method for obtaining a multiphase microstructure composed of tempered bainite, martensite and carbides can also be found in the patent applications CN111593258 and CN101724736A related to low-carbon steels.

Widely described in the cited literature, the bainitic microstructure consists of two phases: plates or laths of carbon-supersaturated ferrite and cementite precipitates. Depending on the temperature and duration of the bainitic transformation, which affect the kinetics of the diffusion processes, cementite can precipitate inside the bainitic ferrite grains, at an angle of 60° to the lath/plate axis (bottom bainite) or at the borders of the laths (top bainite). The bainite structure is obtained by continuous cooling or isothermal annealing in bainite formation range. Compared to martensite, bainite is characterized by lower strength. At the same time, lower bainite shows much better ductility than martensite due to reduced microstresses.

Thus, the result of the introduction of a certain fraction of bainite into the martensitic microstructure due to multi-step heat treatments, as described in the state of the art presented above, is an increase in the impact strength of the steel, while maintaining preferable strength parameters, high hardness and abrasion resistance. However, the presence of cementite precipitates in bainite, capable of coagulation under annealing at elevated temperatures, can be a factor determining the low resistance to brittle fracture of the steel.

In order to make it possible to use bainitic transformation as a steel strengthening mechanism and eliminate embrittlement, steels with a carbide-free bainite structure were developed. Studies have shown that with the appropriate content of silicon and/or aluminum, it is possible to significantly delay the processes of cementite precipitation. As a result, a microstructure consisting of plates or laths of bainitic ferrite separated by thin layers of stable residual austenite is produced. Certain chemical compositions of the steels and the well-defined parameters of isothermal quenching, which leads to a carbide-free bainite structure with nanometer-sized components, are disclosed in U.S. Pat. No. 8,956,470 (which is incorporated herein by reference in its entirety), EP2410070A1 (which corresponds to US Patent Publication 20110126946 which is incorporated herein by reference it is entirety), WO2010013054A2, US20110126946A1, PL228168, and PL219414. However, these inventions are mainly limited to low- and medium-alloy steels with a carbon content in the range of 0.4-1.1% by weight, and an isothermal annealing process at a temperature slightly above the Ms, in the range of 190-300° C. is carried out to complete the bainitic transformation and can take up to several weeks.

From the patent PL234049 appears that it is possible to obtain a microstructure consisting of primary alloyed carbides in a matrix of carbide-free bainite and residual austenite in tool steels with a chemical composition (in % by weight): 0.6-2.3% C, 0.4-2.5% Mn, 0.5-3.0% Si, at least one of the following carbide-forming elements: Cr up to 17.0%, Mo up to 10.0%, V up to 4.0% and at least one alloying additive from the group: W up to 18.0%, Ti up to 0.2%, Nb up to 0.1%, Al up to 2.0%, Ni up to 4.5%, Co up to 10.0%, Cu up to 1.2%, while satisfying: (% Cr+% Mo+% V) 2.25% and (% Si+% Al) 1.5%. As a result of isothermal quenching at a temperature above Ms but lower than 320° C., a bainitic structure with nanometric and/or submicron dimensions of at least 50% by volume is produced, while residual austenite achieves thermal stability down to −40° C. In the cited invention, a binary microstructure of bainitic ferrite with residual austenite as a matrix for carbides is sought. Neither martensitic transformation nor secondary hardness is used to increase the strength and hardness of the steel.

A heat treatment combining isothermal annealing with martensitic quenching and partitioning (B-Q&P) is known from PL234490. After the step of full austenitization, the steel is subjected to isothermal quenching for the time necessary to obtain 10-80% by volume of bainitic ferrite in the structure, then it is cooled to a temperature between the temperatures of the start (Ms) and finish (Mf) of the martensitic transformation of the residual austenite, to finally undergo annealing to stabilize the residual austenite and reduce the martensite carbon supersaturation. The result of this heat treatment is a fragmented, multiphase microstructure consisting of carbide-free bainite, martensite and residual austenite, which results in high hardness, strength, good ductility and resistance to fracture. This method effectively reduces treatment time, compared to simple isothermal quenching, but is limited to low- or medium-alloy steels and steels with silicon and/or aluminum concentrations adequate to inhibit the precipitation processes, but not greater than (Si+Al) 3.2%. In the cited reference, the aim is to inhibit any precipitation processes during the heat treatment steps used. The partitioning temperature is selected so that only diffusive redistribution of carbon from martensite to the remaining austenite occurs. The method described does not use precipitation strengthening to increase the hardness and strength of the steel. An analogous treatment scheme is described in patent application CN102828011A, with the partitioning step replaced by tempering at temperatures above 300° C. to increase ductility, resistance to fracture and resistance to high-cycle fatigue, while maintaining high strength in steels used for production of drilling tools. However, this method of treatment is only for low-carbon steels with higher manganese and silicon content.

Steels for tool applications should be characterized by high hardness, resistance to abrasive wear, good strength parameters and resistance to elevated temperatures during operation. However, it is also desirable that they show increased resistance to brittle fracture and relatively good ductility in order to extend the service life of cyclically and/or dynamically loaded pieces. The combination of these properties is difficult to achieve by conventional heat treatment application. Controlling the fraction and morphology of the phase components through appropriate design of the heat treatment, allows the steel to achieve the desired parameters. The use of an additional mechanism of strengthening by bainitic transformation, in addition to strengthening by martensite and alloy carbide precipitations, in conventional tool steels results in a significant increase in ductility. However, the presence of cementite precipitates in the bainitic structure can reduce the material's resistance to brittle fracture. The addition of sufficient amount of Si and/or Al makes it possible to inhibit the processes of precipitation during isothermal quenching, resulting in an increase in the volume fraction of stable residual austenite in the final structure. From research work [H.K.D.H. Bhadeshia, Materials Science and Engineering, A 481-482 (2008), 36-39, the disclosure of which is incorporated herein by reference in its entirety] it appears that an appropriate austenite morphology that ensures its percolation can significantly increase the plasticity, impact strength and resistance to fracture of the steel. However, to achieve the most favorable results this plastic phase should be in the form of narrow layers or finely dispersed grains, evenly distributed in the structure, and its volume fraction should exceed 10%.

An object of exemplary arrangements is to produce steels with a nanometric or submicron multiphase structure, providing a sufficient level of hardness and resistance to abrasive wear, while increasing ductility and resistance to fracture. In addition, the exemplary method of heat treatment of alloyed medium- and high-carbon steel leads to a reduction in quenching deformations compared to conventional martensitic quenching and tempering treatment, while maintaining high temperature stability.

The exemplary method comprises heat treatment of articles comprised of medium- and high-carbon alloy steel, designated by the abbreviation BQT, in which a bainitic transformation takes place. The exemplary method, which includes steel heating, cooling and tempering, includes consecutive steps:

In step (A), the steel article is heated to an austenitizing temperature T>A(the critical temperature for austenitizing) and held at this temperature for a time t.

In the next step (B), the steel article in which a bainitic transformation takes place is cooled to an isothermal annealing temperature T, where 350° C.≤T≤Ms(γ)+10° C. and held at this temperature for a time t, until a 10-70% volume fraction of bainitic ferrite is formed.

In a subsequent step, step (Q), the steel article is cooled to a temperature Tthat is between the temperature of the start M(γ) and the temperature of the finish M(γ) of the martensitic transformation of the residual austenite that remains after partial bainitization (step B) and held for a time t, until the temperature equalizes throughout the piece.

In a subsequent step (T), the steel article is tempered one or more times for a time tand at a temperature T, at which a secondary hardness effect occurs in the treated steel until achieving the precipitation strengthening effect by alloy carbides.

As a result of the above exemplary treatment method, a structure consisting of a mixture of bainitic ferrite, tempered martensite, residual austenite and carbides can be obtained.

In some exemplary methods, the tempering step (T) is performed during surface engineering processes such as nitriding or physical vapor deposition (PVD). The PVD process results in a coating of a nature compatible with the intended use of the steel. These can be protective, metalized or even decorative coatings.

In some exemplary methods a carbon partitioning step (P) instead of the tempering step (T) is carried out.

In some exemplary methods in the first tempering, the tempering temperature Tis in the range of 400° C.≤T≤600° C., and the tempering time tis in the range of 1 h≤t≤3 h.

When a martensitic transformation occurs during cooling after the last tempering, the exemplary methods may include carrying out an additional tempering step at T≤400° C. to reduce the embrittlement of the newly formed martensite.

The exemplary method may be characterized by the fact that in the step (A) at the temperature Tand austenitizing time tthere is obtained Ms(γ) 340° C.

The exemplary method may be characterized by the fact that the step (B) of partial bainitization is carried out by means of one or more isothermal stops, or by continuous cooling in the temperature range of 350° C.≥T≥Ms(γ)+10° C., where the temperature Tis lower than the temperature of the bainitic transformation start Bs of the steel, and at least 10° C. higher than the temperature Ms(γ), wherein T≥200° C.

After BQT treatment in accordance with the exemplary method, there is at least 10% by volume of residual austenite in the final microstructure of the steel.

The exemplary method may be characterized by the fact that the treated steel contains alloying additives (in % by weight): 0.4-2.3% C, 0.4-2.5% Mn, 0.5-3.0% Si, at least one of the following carbide-forming elements: Cr up to 17.0%, Mo up to 10.0%, V up to 4.0%, while maintaining the relation: (% Si+% Al)≥0.8% and optional additives from the group: up to 18.0% W, up to 0.2% Ti, up to 0.1% Nb, up to 2.0% Al, up to 4.5% Ni, up to 10.0% Co, up to 1.2% Cu, the remaining components being iron and trace amounts of unavoidable admixtures.

The exemplary methods may be used for producing tools and parts of machinery and equipment that work under harsh operating conditions, such as industrial knives, coin punches, and punching dies for plastics.

The exemplary heat treatment methods may be used to shape the alloying properties of the starting material, which are medium- and high-carbon steels in which a bainitic transformation occurs, wherein the chemical composition should ensure the temperature of the martensitic transformation start in the steel Ms≤340° C. For purposes of this disclosure medium to high carbon alloy steels include steels having at least 0.4% C by weight. The chemical composition of the treated steel should be within the following ranges (in % by weight): 0.4-2.3% C, 0.4-2.5% Mn, 0.5-3.0% Si, at least one of the following carbide-forming elements: Cr up to 17.0%, Mo up to 10.0%, V up to 4.0% while maintaining the relation: (% Si+% Al)≥0.8% and optional additives from the group: W up to 18.0%, Ti up to 0.2%, Nb up to 0.1%, Al up to 2.0%, Ni up to 4.5%, Co up to 10.0%, Cu up to 1.2%. The remaining components are iron and trace amounts of unavoidable admixtures.

The addition of silicon and/or aluminum is used in the exemplary arrangements to inhibit precipitation processes during isothermal holding (B), allowing formation of carbide-free ferrite during isothermal quenching and for better control over carbide morphology during tempering. Thanks to the inhibition of precipitation processes during the bainitic transformation, carbon diffuses from the bainitic ferrite into the surrounding austenite instead of forming carbides. Increasing the carbon content of austenite increases its thermal stability and allows it to remain in the microstructure in the form of thin layers between plates of bainitic ferrite. Replacing carbides with residual austenite in the form of thin layers significantly improves the resistance to fracture of the steel. Elements such as chromium, molybdenum, niobium and vanadium—due to their tendency to form alloyed carbides—in turn have the greatest influence on the secondary hardness achieved during tempering. Additives that increase hardenability, such as nickel and manganese, are also desirable due to the ability to machine pieces with larger dimensions. The exemplary steel alloy may also contain alloying additives that form high-melting carbides or carbonitrides. Their presence reduces the grain size of the primary austenite, increasing the number of bainitic ferrite nucleation sites and causing a homogenization of the structure, which often has a beneficial effect on the properties of the alloy steel.

In the exemplary methods the content of impurities such as phosphorus, sulfur, oxygen, tin and lead in the steel should be reduced to reduce their segregation within the boundaries, and to avoid intergranular embrittlement and temper embrittlement during heat treatment.

In an exemplary method a first step—austenitizing (A)—is transforming the initial structure of the steel article into austenite or austenite with primary carbides. By selecting the appropriate austenitizing temperature (T) and time (t) of austenitization, the chemical composition of the austenite can be controlled, controlling the dissolution processes of the alloying carbides. Increasing the alloying of austenite results in a shift in the area of phase transformation occurrence toward longer times and lower temperatures. On the other hand, the presence of evenly distributed, finely dispersed primary carbides can positively affect the strengthening of the steel, and also affect the morphology of the phase components and the kinetics of transformations, providing additional sites for phase nucleation. Therefore, the parameters of subsequent BQT treatment method steps are largely determined by the austenitization conditions. Due to the risk of decarburization and/or oxidation resulting from the high temperatures of this operation, austenitization in a vacuum furnace or with a protective atmosphere is recommended in exemplary methods.

In an exemplary method a second step—isothermal quenching (B)—is producing a certain volume fraction of carbide-free bainite in the steel article in the form of bundles composed of nanometric or submicron plates of bainitic ferrite separated by thin layers of residual austenite. Nano- or submicron carbide-free bainite is characterized by a favorable combination of high strength, resistance to fracture and fatigue. An objective of bainite bundle formation is also dividing the primary austenite grain into smaller areas of so-called blocks, in which partial martensitic transformation will take place in a next treatment step (Q). The division of the primary austenite grain into smaller areas results in significant fragmentation of the martensite formed in them. The fragmentation of martensite increases the strength and plasticity of the microstructure. This step involves cooling the steel from the austenitization temperature (T) to the isothermal annealing temperature (T) at a rate higher than the critical rate, in order to avoid areas of diffuse ferritic and/or perlitic transformation. The temperature (T) and the isothermal annealing time (t) of the stop in the bainitic range may be determined experimentally for the particular steel depending on the austenite chemical composition obtained. However, in exemplary methods this step is performed in the lower temperature range of the occurrence of bainitic transformation T≤320° C., due to the possibility of obtaining a higher degree of microstructure fragmentation. Due to the lowered temperature of isothermal quenching, it is possible to produce a fine-grained structure with submicron thickness of bainite plates, and in some cases even on the order of tens of nanometers. Adequate fragmentation of the bainite structure results in the strengthening of the material, resulting in an increase in hardness and strength limit, without a decrease in ductility. It should also be remembered that the Tisothermal annealing temperature should be at least 10° C. higher than the Ms(γ) martensitic transformation start temperature, as the Ms temperature is the average temperature for a given alloy and may locally take higher values. The Ttemperature should therefore be selected to avoid the premature occurrence of martensitic transformation. In addition, too low isothermal stop temperature results in slowing down the kinetics of the transformation, which significantly increases the treatment time, so it is preferable that T≥200° C. The use of an isothermal stop above the Ms temperature further reduces the quenching deformations of the workpiece. As the bainitic transformation progresses, austenite is enriched in carbon, and thus its thermal stability increases, manifested, among other things, by a decrease in Ms(γ) martensitic transformation start temperature of the austenite remaining. Hence, it is important to control the advancement of the bainitization step so that it effectively affects the microstructure and properties of the steel, and so that the temperature Ms(γ) is not lowered excessively, which would prevent the next step, which is martensitic quenching.

In an exemplary method a third step—quenching (Q)—is producing martensite, in the form of very fine plates, laths or needles, in the steel structure. This step involves cooling the steel article to a temperature between the temperatures of the martensitic transformation start M(γ) and the martensitic transformation finish M(γ) of the martensitic transformation of austenite remaining in the material in the form of blocks. Martensite, as a hard phase, increases the hardness and resistance to abrasive wear of the machined steel. Adjustment of the Tquench temperature makes it possible to control the content of residual austenite and martensite, since with overcooling in the microstructure the fraction of martensite increases, at the expense of the austenite fraction. The phase composition of the steel can be determined depending on the application of the workpiece, although it may be desirable that at least 10% by volume of residual austenite remain in the steel. The quench holding time tdepends on the cooling medium and the size of the charge, and should be selected to allow the quench temperature Tto be reached throughout the workpiece.

In an exemplary method a fourth step—tempering (T)—is often important to the obtained properties of the steel as a result of the exemplary heat treatment method. The step includes annealing the charge one or more times at a temperature at which the precipitation processes and diffusion of carbon from martensite to austenite take place simultaneously, so that embrittlement is reduced and austenite increases its thermal stability. Thanks to the precipitation processes during tempering, the steel is strengthened due to the uniform precipitation of finely dispersed carbides from the matrix. The effect of this step is an increase in the hardness of the workpiece, known as secondary hardness. The finely dispersed alloy carbides, cohesive with the matrix, effectively contribute to increasing the mechanical and functional parameters of the steel article. In addition, in the case of high-carbon steels, in which the content of residual austenite remaining after the bainitization and quenching steps is too high, the precipitation of carbides can lead to the destabilization of austenite and cause its partial martensitic transformation during cooling to room temperature, resulting in a reduction in the volume fraction of this soft phase. Further tempering operations are then desirable to eliminate the embrittlement of the fresh martensite. The time, temperature and number of tempering steps should be selected individually for the workpiece, as the tempering step offers a wide range of possibilities for controlling the final properties of the particular steel and the workpiece. Typically, the tempering temperature Tis in the range of about 400-750° C., and sometimes in the range of 400-600° C., while the tempering time Tof a single segment should not exceed about 3 hours. For purposes of this disclosure when a value is referred to as about a particular value, the value shall be deemed to be the value stated±5%.

The content of residual austenite in the final microstructure of steel is crucial to the properties obtained—it is a plastic phase and, in the right amount and morphology, can significantly increase the impact strength, ductility and resistance to brittle fracture of the steel. At the same time, carbon-supersaturated austenite in the form of thin films and fine grains can undergo the TRIP (Transformation-Induced Plasticity) effect, increasing strength and resistance to abrasive wear. However, stabilization of an excessive amount of austenite during tempering can lead to unsatisfactory values of hardness and resistance to abrasive wear, while destabilization of too much austenite, especially when martensitic transformation occurs after the final step of tempering, may lead to embrittlement of the workpiece. Tempering parameters also directly affect the precipitation and growth of carbides, which also affects the final properties of the steel. Too high a temperature or too long a time for this treatment step can lead to excessive carbon depletion of martensite and bainitic ferrite, which translates into a decrease in the hardness of the material. At the same time, coagulation of precipitates will occur, resulting in reduced strengthening of the steel and a further decrease in hardness. Improperly selected tempering parameters can also lead to the growth of carbides at grain boundaries and the formation of the so-called carbide mesh significantly deteriorating the resistance to brittle fracture of the steel. It is also important to keep in mind the known embrittlement effects of tempering of the first and second kind, which can occur during tempering processes. Properly carried out tempering process steps provide the possibility to control the content and morphology of secondary carbides and the fraction of residual austenite, allowing achieving high values of hardness and abrasion resistance along with increased impact strength and resistance to fracture.

The exemplary heat treatment methods for medium- and high-carbon alloy steel with higher silicon and/or aluminum content and carbide-forming alloying elements makes it possible to produce a very fine, multiphase microstructure consisting of carbide-free bainite, tempered martensite, residual austenite and carbides. This type of microstructure allows the combination of high hardness, strength and abrasion resistance with ductility and resistance to fracture. This is made possible in exemplary arrangements by combining the mechanisms of precipitation strengthening by carbides, strengthening by hard martensite and bainitic ferrite phase, and strengthening by grain boundaries. Strengthening by grain boundaries is the result of a high fragmentation of the phase components to submicron and even nanometer sizes. On the other hand, the addition of sufficient amounts of silicon and/or aluminum inhibits the precipitation of cementite especially in bainite, making it more resistant to fracture. In turn, the presence of residual austenite in the form of thin layers and finely dispersed grains that exhibit percolation, thanks to their content of more than 10% by volume, is responsible for the increase in plasticity.

In addition, the presence of carbides that did not dissolve during austenitization can result in the formation of bainite with a star-shaped morphology (acicular type) during the partial bainitization step. This type of structure is characterized by a large disorientation of bainitic ferrite plates that nucleate on carbide precipitations, thus further increasing the resistance to fracture of the steel.

The steel produced according to the exemplary methods has a different microstructure and chemical composition than the materials disclosed in the state of the art. The exemplary heat treatment method simultaneously utilizes all the presented strengthening mechanisms in steels for tools, resulting in unique mechanical and functional properties. The method according to exemplary arrangements makes it possible to control the phase fraction and morphology of the microstructure, allowing achieving the desired properties of the steel in terms of its application and requirements for tools and structural components. The parameters of each heat treatment step can be determined by calculations, computer simulations and experimental tests using a dilatometer.

The method according to exemplary arrangements which may be designated by the abbreviation BQT, can be used to shape the properties of pieces that are required to have higher impact strength and resistance to fracture compared to conventional martensitic quenching and tempering treatment, while maintaining comparable hardness and strength. These pieces can include industrial knives, coin punches and punching dies for plastics, among other things. This treatment allows for longer life of the pieces and increases the failure-free life of equipment.

When the service conditions of the workpieces require steels with even higher resistance to fracture and ductility compared to BQT-treated steels, at the expense of lower hardness, a modified BQT treatment process may be used by what may be referred to herein as a BQP process. In an exemplary BQP treatment, a carbon partitioning (P) step is used instead of a tempering (T) step. During the partitioning step (P), the article is heated to a partitioning temperature Tand held at the partitioning temperature for a partitioning temperature time tduring which carbon diffuses from martensite and bainitic ferrite to austenite increasing thermal and mechanical stability. During the exemplary partitioning step (P), unlike the tempering step (T), there is no carbide precipitation and no secondary hardness effect.

The following exemplary methods caried out with exemplary steels are illustrative only and in no way limit the claimed invention.

Specimens made of K360 steel, with the chemical composition shown in Table 1, were subjected to BQT type method treatment. This steel is characterized by a high concentration of carbide-forming elements, i.e. Cr, Mo, V, Nb, W, as well as an increased content of silicon and aluminum (% Si+% Al)=1.74% by weight, at which there is an inhibition of precipitation processes during the isothermal quenching step. The chemical composition of the steel meets the requirements of the material for this type of treatment.

An austenitizing temperature about T=1050° C. and an austenitizing time about t=15 min were used as the parameters for the austenitizing step, which allow for partial dissolution of the carbides. As a result of carrying out the exemplary method step, the structure is left with carbides of the type MC, MCand MC, with a volume fraction not exceeding 10% by volume, evenly distributed in the matrix, while the alloying of austenite is sufficient to apply the developed technology. Based on dilatometric studies, the temperature of martensitic transformation start Ms(γ)=240° C. was determined.

An isothermal quenching step was carried out at an isothermal annealing temperature about T=260° C., which is 20° C. higher than Ms(γ). Cooling to the isothermal stop was carried out in hot oil, which allows a cooling rate higher than the critical one. Dilatometric studies showed that the completion of the isothermal transformation at this temperature occurs after 16 h 30 min. Based on the dilatometric curve showing the kinetics of the bainitic transformation, the durations of the stops (isothermal annealing stop times) that lead to 20% and 60% advancement of the bainitic transformation were determined. They are an isothermal annealing time for 20% bainitic transformation t=2 h 3 min and an isothermal annealing time for 60% bainitic transformation t=3 h 33 min, respectively.

As a result of the formation of bainitic ferrite during isothermal quenching, the temperatures of the martensitic transformation start of the remaining austenite fell to Ms(γ20%)=210° C. and Ms(γ40%)=190° C., respectively, while the martensitic transformation finish temperatures M(γ) were recorded below room temperature. The quenching temperature was determined at T=30° C. for both series of specimens. According to exemplary methods the quench temperature, Tis within the temperature range of M(γ) and M(γ). The tquench holding time was assumed to be 10 min.

On the basis of experimental studies and material safety data sheets, the tempering temperature and time were determined: T=500° C., t=2 h. These parameters allow observing the effect of secondary hardness in the steel. Only one tempering step segment was used, as no martensitic transformation was observed during cooling, so tempering did not destabilize the residual austenite. The fraction of residual austenite was estimated to be about 15% by volume in both series of specimens.