JP6112752B2 - Method for thermomechanical processing of tool steel and tools made from thermomechanically processed tool steel - Google Patents

Description

  The present invention relates to thermomechanical processing of tool steel, a method of forming a tool using thermomechanically processed tool steel, and a tool used in metal forming and metal cutting applications.

This application is a continuation-in-part of co-pending application 12/047532 filed on March 13, 2008, claiming the benefit of US Provisional Application No. 60/896729, each The disclosure is hereby incorporated by reference herein in its entirety. This application also claims the benefit of US Provisional Application No. 61/029236, filed Feb. 15, 2008, which is hereby incorporated by reference herein in its entirety.

  Among the various quality grades of commercially available carbon and alloy steels, the quality grade of tool steel is commonly used in applications where the tool is subjected to severe stress, impact and / or wear. Tool steel is generally characterized by a characteristic hardness, resistance to wear, the ability to hold the cutting edge, and resistance to deformation at high temperatures. As a result, tool steels are widely used in metal forming and metal cutting applications, inspection equipment and gauges, and wear / impact elements in machine tools.

  Various types of tools are used in metal forming and metal cutting applications such as machining, through drilling, casting, drawing, powder compression, metal engraving, pin stamping and the like. Specifically, punches and dies represent the type of metal forming tool used to pierce, drill and shape metal and non-metal workpieces. Cutting tools and inserts represent a type of metal cutting tool used in machining applications to form metallic and non-metallic workpieces. Plug gauges, thread gauges, pipe gauges, ring gauges and setting panels represent the types of tools used for inspection applications. Machine slides and gibs represent the types of wear and impact elements used in machine tools.

  Punches and dies are subjected to severe and repeated loads during their service life. Specifically, punches tend to fail during use due to catastrophic failure induced by significant stresses applied during their use. The requirements for metal forming tools are ultra high strength steel (UHSS), advanced high strength steel (AHSS), transformation induced plasticity (TRIP) steel, twinning induced plasticity (TWIP) steel, nano steel and martensitic (MART) With the introduction of workpieces made of steel having a higher strength to weight ratio, such as steel, it becomes more severe. For example, the automotive industry is shifting to more frequently use these types of high-strength, lightweight steels for body structures.

  Therefore, what is needed is a method of thermomechanically treating tool steel to improve its mechanical properties, and a tool formed by thermomechanical processing having improved mechanical properties.

  In one embodiment, a method for thermomechanically processing a preform made of tool steel is provided. The tool steel has a martensitic transformation start temperature and a stable austenite temperature. The preform has a region comprising austenite, the region comprising an outer surface and a plurality of outer dimensions with respect to the outer surface. The method comprises establishing at least a region of the preform at a processing temperature between the martensitic transformation start temperature and the stable austenite temperature. While the region of the preform is at the processing temperature, this region extends over a depth that extends at least one of the outer dimensions of the region and extends from the outer surface to a depth of 1 millimeter or more below the outer surface. It is deformed to modify the microstructure. After the region is deformed, the region is cooled to room temperature.

  In another embodiment, a tool for use in a machine for modifying a workpiece is provided. The tool has a member made of tool steel. The member has an outer surface that defines a first portion configured to be coupled to the machine and a second portion configured to contact the workpiece. The member has a first region extending from the outer surface to a depth greater than 1 millimeter and a second region separated from the outer surface by the first region. The first region comprises a plurality of particles having a misorientation angle distribution having an average misorientation angle greater than about 34 °, an average particle size that is at least 10% less than the second region, and in the second region It has a different particle orientation than the plurality of particles.

  The accompanying drawings, which are incorporated in and constitute a part of this specification, illustrate embodiments of the invention and, together with the brief description above and the detailed description of the embodiments below, of the embodiments of the invention. It plays a role in explaining the principle.

2 is a graphical representation of an exemplary time-temperature relationship for a thermomechanically processed M2 AISI tool steel according to an embodiment of the present invention. 2 is a graphical representation of an exemplary time-temperature relationship for a thermomechanical tool steel according to an embodiment of the present invention. 1 is a side view of a tool according to an exemplary embodiment of the present invention and a cross-sectional illustration of a corresponding die. FIG. FIG. 2B is an enlarged cross-sectional explanatory view of the tool and die of FIG. 2A. 1 is a perspective view of one embodiment of a preform having a shell and a core before deformation. FIG. 1 is a perspective view of one embodiment of a preform having a shell and a core after deformation. FIG. FIG. 3C is a perspective view of one embodiment of a tool made from the deformed preform of FIG. 3C. 2 is a graphical representation of the existing phase measurements of an exemplary embodiment of the invention made of M2 tool steel. 2 is a graphical representation of a measurement result of a particle misorientation angle distribution of an exemplary embodiment of the invention made of M2 tool steel. FIG. 3 is a pole figure of an exemplary embodiment of the present invention made of M2 tool steel. 2 is a graphical representation of the existing phase measurements of an exemplary embodiment of the invention made of M2 tool steel. 2 is a graphical representation of a measurement result of a particle misorientation angle distribution of an exemplary embodiment of the invention made of M2 tool steel. FIG. 3 is a pole figure of an exemplary embodiment of the present invention made of M2 tool steel. 2 is a graphical representation of the phase present in a heat treated M2 tool steel bar according to the prior art. 2 is a graphical representation of the distribution of misorientation angles of particles of a heat treated M2 tool steel bar according to the prior art. FIG. 2 is a pole figure of a heat treated M2 tool steel bar according to the prior art. 1 is a perspective view illustrating an exemplary preform configuration for thermomechanical processing tool steel according to one embodiment of the present invention. FIG. FIG. 8 is a plan view of the exemplary preform of FIG. 7 prior to processing according to an embodiment of the present invention. FIG. 8B is a partial cross-sectional view of the exemplary preform of FIG. 8A taken along section line 8B-8B of FIG. 7 following deformation. 5B is a schematic cross-sectional representation of an exemplary die and ram for thermomechanical processing the preform configuration shown in FIGS. 4 and 5A. FIG. FIG. 9 is a photomicrograph taken at 13 × magnification of a cross section taken through the elliptical cross section of FIG. 8B of a preform configured as shown in FIG. 8B. FIG. 10B is a schematic representation of the micrograph of FIG. 10A showing the preferential particle orientation drawn as a curve. It is a perspective view which shows one structure of the preform before a deformation | transformation. It is a perspective view which shows one structure of the preform after a deformation | transformation and machining. FIG. 11C is a set of perspective views of a tool made from the preforms shown in FIG. 11B in an operable position relative to each other that provides a shearing or trimming action to cut a sheet of steel material. FIG. 11D is a graphical representation showing a comparison of wear measurement results of an exemplary tool cutting edge profile of the present invention, and a tool cutting edge made of a reference material each having the configuration shown in FIG. 11C. FIG. 11B is a plan view of the tool of FIG. 11C showing measurement positions related to the wear profile given to the graph of FIG. 12A. FIG. 11B is a graphical representation showing a comparison of the wear measurement results of the contours of the cutting edges of an exemplary tool of the present invention, and a drawing showing cutting edges of a tool made of a reference material each having the configuration shown in FIG. 11C. FIG. 11B is a plan view of the tool of FIG. 11C showing measurement positions related to the wear profile given to the graph of FIG. 13A. FIG. 11B is a graphical representation showing a comparison of the wear measurement results of the contours of the cutting edges of an exemplary tool of the present invention, and a drawing showing cutting edges of a tool made of a reference material each having the configuration shown in FIG. 11C. FIG. 14 is a plan view of the tool of FIG. 11C showing measurement positions related to the wear profile given to the graph of FIG. 14A. FIG. 11B is a graphical representation showing a comparison of the wear measurement results of the contours of the cutting edges of an exemplary tool of the present invention, and a drawing showing cutting edges of a tool made of a reference material each having the configuration shown in FIG. 11C. FIG. 15B is a plan view of the tool of FIG. 11C showing measurement positions related to the wear profile given to the graph of FIG. 15A. FIG. 12 is a photomicrograph of a cross section taken at 17 × magnification of the region shown in FIG. 11B surrounding the cutting edge of one tool showing preferential particle orientation in the area surrounding the cutting edge. FIG. 16B is a schematic representation of the micrograph of FIG. 16A with lines drawn to indicate preferential particle orientation.

  According to one embodiment of the present invention, a method of making a tool includes manufacturing a preform from tool steel, wherein at least a region of the preform is thermomechanically processed. This region of the preform typically includes a substantial volume of tool steel or a majority of the preform. With respect to the cylindrical preform geometry, for example, a thermomechanically treated region that is processed in a radial or plane strain forging process encompasses the outer 60% of its volume, and the rest of the tool steel. The inner volume may be relatively unaffected by the process. Thus, for a simple preform geometry, the volume of this region may include at least some outer volume of one section of the preform. This region may extend at least partially or entirely across its cross-sectional area. Thus, in this embodiment, the outer volume or modified region extends from the outer surface of the region to a depth of at least greater than 0.039 inches (1 millimeter), but the volume dimension extends deeper within the preform. You may get something. However, the depth of the region need not be uniform; rather, the depth in one part of the region is less than 0.039 inch (1 millimeter), while the depth in the other part is 0.039 inch (1 millimeter). Extend more greatly.

  Although the modified region has been described above as an outer volume in the form of a layer around the inner volume, the modified region may be an irregularly shaped region. This may, for example, change at least one of the outer dimensions of the region so that the outer surface of the preform has a certain geometric shape before deformation, but then forms an object with a different shape. It is a case where it deform | transforms by. For example, the deformation may include a change in one or more of the cross-sectional area or other outer dimensions that may increase or decrease the length of the region. One skilled in the art will recognize that the volume of material being processed can depend on numerous other factors including, but not limited to, the size and shape of the preform and the capability and type of the deformation device. In general, as the load capacity of the forging device increases and the size of the preform decreases, the deformed region can include a larger, if not all, part of the preform. Thus, unlike surface treatment operations such as shot peening, embodiments of the present invention are not limited to forming a thin surface layer that is forced to follow a pre-established contour of a part. Furthermore, embodiments of the present invention deform a larger portion of the tool steel and, in some embodiments, determine the contour or outer surface dimensions of the preform. In this regard, the area of the preform can be measured over the majority of the thickness of the preform or tool. Also, the shape of the preform can be independent of the final shape of the tool.

  In addition to affecting the volume of tool steel being processed, the geometry or shape of the preform prior to thermomechanical processing can affect the final microstructure. For example, the shape of the preform can affect or determine the orientation of the particles as well as the properties of the microstructure in the thermomechanically treated region. Those skilled in the art will appreciate that the tool steel preform has multiple configurations with any number of cross-sectional shapes, such as rod stock with circular, rectangular or polygonal cross-sections, or stock materials with more complex shapes and cross-sections. You will understand that it can be one of Preform geometry determination can be advanced based on historical experience, tool requirements, and / or processing limitations. For example, the preform geometry may be selected based on the type of processing employed and the final geometry of the targeted tool.

  While the temperature of the region is maintained in the temperature range described below according to different embodiments of the present invention, the region undergoes deformation. In an embodiment of the invention, the amount of deformation is sufficient to improve the mechanical properties of the deformed area. The amount of deformation can be quantified by calculating a reduction ratio, defined as a relative decrease in cross-sectional area due to thermomechanical processing. The improvement in the characteristics of the region is considered to be proportional to the amount of deformation. By way of example and not limitation, a reduction factor of as little as 20% can result in a measurable improvement in the mechanical properties of the region. The amount of deformation that results in a measurable improvement in mechanical properties is believed to be limited only by dynamic recrystallization of the tool steel. In other words, the amount of deformation can be kept below an effective threshold to dynamically recrystallize the microstructure. When the deformed microstructure is recrystallized, a measurable decrease in mechanical properties can be observed compared to the non-recrystallized microstructure. The reduction in specific mechanical properties can be at least about 20%. However, despite a decrease may be observed, the mechanical properties are improved compared to tools prepared by heat treating the tool steel over a specified temperature range, as described in more detail below. Can be done. Those skilled in the art will appreciate that in addition to the amount of deformation, dynamic recrystallization is subject to the composition of the tool steel and the temperature at which the deformation occurs.

  As described above, thermomechanical processing includes plastically deforming the tool steel preform while the tool steel preform is held at a high temperature. Suitable processes that can plastically deform the preform include, but are not limited to, radial forging, ring rolling, rotary forging, swaging, thixoforming, ausforming and warm / hot up. Forging processes such as setting (warm / hot upsetting) are included, but other suitable deformation processes are available. For example, the technology may also include those where the primary deformation direction is not substantially perpendicular to the longitudinal axis of the preform. As mentioned above, other techniques such as shot peening at high temperatures can result in very shallow deformations, and therefore excluded if deeper plastic deformation is required to provide the required improvement in mechanical properties. Is done.

  One such process is primarily plane strain forging that results in plastic deformation in the radial and circumferential directions of the tool steel preform. Thus, plane strain forging can limit the elongation of particles in a direction that is perpendicular to the applied load. The preform can consequently exhibit a substantially uniform distribution of mechanical properties along and around its length. Accordingly, in one embodiment, plane strain forging includes a plastic deformation process that causes little, if any, elongation of particles in a particular direction. However, when thermomechanically processing a tool steel preform, any combination of the above-described processes that can plastically deform the preform can be used.

  In yet another embodiment, an existing tool serves as a preform. For example, in addition to unused tools, existing tools can include used tools, damaged tools, or broken tools. Existing tools are thermomechanically processed as described herein to remanufacture or reprocess the tools to restore their effectiveness.

  As previously defined, thermomechanical processing includes plastically deforming a region of the preform while maintaining the region at a high processing temperature. The preform temperature during deformation can be established by cooling the preform from a higher temperature. Such processing includes, by way of example only, casting a tool steel billet or preform from molten raw material, cooling the cast preform to a lower processing temperature, and transforming it at the processing temperature. May be included. Alternatively, the preform can be brought to a processing temperature at which deformation occurs by heating the preform at or near room temperature, as described in more detail below.

Specifically, also referring to FIG. 1, when the preform exceeds the martensitic transformation start temperature (Ms) of the tool steel (martensitic transformation start temperature) while the preform contains austenite. The tool steel is deformed at a processing temperature that is less than the stable austenite temperature (AC ₃ ). Ms is the temperature at which the transformation of austenite to martensite begins during cooling, and AC ₃ is the temperature at which the transformation of ferrite to austenite is completed during heating.

In addition, as is apparent from FIG. 1, the austenite transformation start temperature (AC ₁ ) represents the temperature at which austenite begins to form during heating. _One skilled in the art will appreciate that Ms, AC ₁ and AC ₃ each depend on the specific composition of the tool steel. Thus, any instance described herein in which Ms, AC ₁ or AC ₃ is referenced with a particular temperature is not intended to limit its definition to that particular temperature.

In view of the previously defined temperature, and according to one embodiment, when the tool steel preform is at a temperature between Ms and AC ₃ , the region is also austenite (e.g. metastable austenite). All or part of the tool steel preform is processed, i.e. the tool steel preform is plastically deformed or forged. As a result, the deformed region of the tool steel preform has some improved mechanical properties that are described below. For example, the improvement in impact strength or toughness of the deformed region can be at least about 20% greater, and in a further example the preform over AC ₃ if the microstructure is primarily stable austenite. Can be at least about 50% larger than deforming.

As introduced previously, in one embodiment, the method includes heating the tool steel preform to a temperature range such that at least a portion of the preform includes austenite. One skilled in the art will recognize that many different temperature profiles can be utilized to bring the tool steel preform within the aforementioned temperature range prior to deformation. By way of example only and with reference to FIG. 1, the tool steel preform may be heated from a temperature below Ms to a processing temperature above AC ₁ (denoted by reference numeral 10). In this example, the temperature is about 1530 ° F. (about 832 ° C.) and AC ₃ is about 2250 ° F. (about 1232 ° C.). Tool steel preform after which it is being held at the processing temperature between AC ₁ and AC _3, it may be deformed.

Heating the tool steel preform from other temperatures profile below Ms to a temperature between AC ₁ and AC ₃ and then processing temperature above Ms (at 11) before transforming it Cooling the tool steel preform to (as shown). In yet another embodiment shown in FIG. 1A, the temperature profile is between the AC ₁ and AC ₃ steps before heating the tool steel preform to a temperature above AC ₃ and then deforming it. Cooling to a processing temperature (indicated by reference numeral 12) or to a processing temperature between Ms and AC ₁ (indicated by reference numeral 13).

The processing temperature during the deformation may increase or decrease and remain substantially the same, but the temperature in the region remains between AC ₃ and Ms. As shown in FIGS. 1 and 1A, the temperature at which deformation occurs (eg, at 10, 11, 12, and 13) is shown as a horizontal line. Although the horizon may represent an isothermal condition, those skilled in the art will appreciate that some variation in actual process temperatures will occur. For example, the actual processing temperature of tool steel preforms can vary by ± 50 ° F. (± 28 ° C.) during deformation. Temperature control to maintain the region in a substantially isothermal state may require intentional addition or removal of heat through a closed loop temperature feedback control system.

However, an increase or decrease in temperature can occur during deformation. The increase or decrease in temperature can be intentional or the result of not controlling the temperature during deformation. For example, in some embodiments, the temperature of the preform can be increased by 150 ° F. (83 ° C.) due to the rate at which energy is applied to the preform by deformation. The additional energy is converted to heat, which increases the temperature of the region if not compensated for by reducing or removing heat. Therefore, the processing temperature is the temperature of the region, starting at temperatures above AC _1, may end at a temperature below AC _1, or starts at temperatures below AC _1, it may end at a temperature above the AC ₁ As such, it can rise or fall. In other embodiments, the region can be intentionally cooled to reduce the temperature of the region while deformation occurs. It should be noted, however, that dynamic recrystallization of the particles can reduce the impact strength and toughness of the region if the preform temperature changes substantially during the deformation process. Accordingly, isothermal processing, i.e., maintaining the actual processing temperature of the tool steel preform substantially constant during deformation, as described below, will increase the strength, toughness, and other mechanical properties of the region. Can be maximized.

With continued reference to FIGS. 1 and 1A, various heating and cooling processes may be employed, while the process temperature and process time are controlled to avoid carbide nose 14 or bainite nose 16. Those skilled in the art will understand that at temperatures below AC ₁ , tool steel can result in carbides or bainite if the region is held in these ranges for too long. As an example, the M2 AISI tool steel preform can be deformed over a period of at least 2 minutes without substantial carbide or bainite phase formation. However, the amount of time that the preform can be held at a temperature in this range depends on at least the composition and temperature of the tool steel, as well as other factors.

  Following thermomechanical processing, the preform is cooled to a lower temperature. Cooling or quenching can be achieved by forced air convection or by holding the region at an intermediate temperature before cooling the preform to room temperature. One skilled in the art will appreciate that quenching can have other cooling methods and media including, for example, water or oil quenching. As an additional example, the region may be subject to low temperature processing, where the region is about -150 ° in one or more stages to convert a larger percentage of retained austenite to martensite. Cool to a temperature between F (about -101 ° C) and about -300 ° F (about -184 ° C). Low temperature processing is achieved, for example, with liquid nitrogen, and other tool steels containing a substantial proportion of retained austenite can benefit from this type of processing, but can be employed primarily with A2 and D2 tool steels. The rate of quenching is greater than the critical cooling rate of the tool steel, i.e., the minimum rate of continuous cooling to prevent undesirable transformations such as carbide nose 14 and bainite nose 16. Thus, the cooling rate is sufficient to avoid substantial transformation of metastable austenite to undesirable decomposition products such as carbides or bainite. Faster cooling rates can also be used, but faster cooling rates are limited to those that do not cause thermal shock to the area or otherwise distort the tool steel preform.

Further, in one embodiment, one or more tempering processes are performed after cooling. For example, tempering may include heating the region to a temperature between about 850 ° F. (about 454 ° C.) and about 1000 ° F. (about 537 ° C.) for about 45 minutes to about 60 minutes. Tempering modifies the microstructure by converting martensite to retained austenite. As is known in the art, multiple tempering cycles can be used to convert residual austenite. Those skilled in the art will understand that tempering is shorter or longer depending on the composition of the tool steel, the geometry and size of the preform, the amount of retained austenite allowed, and the number of tempering processes utilized, It will be appreciated that the step of heating to a higher or lower temperature may be included. According to one embodiment, following quenching, before tempering, areas are not heat-treated at a temperature higher than in AC _3, or AC _3. Furthermore, the region may not be heated above any temperature experienced by the region during deformation. In other words, the preform can be reheated, but the temperature during any subsequent reheating is not the strain or the result of deforming the austenite in the region at a temperature between the stable austenite temperature and the martensite start temperature. Does not substantially suppress or change the dislocation increase.

  In another embodiment, the method further comprises the step of finishing the tool steel preform into a tool after the thermomechanical deformation process. The finish may have a material removal process to provide a final predetermined shape and / or surface finish. For example, conventional finishing processes may include machining, grinding, polishing / polishing, or a combination thereof to prepare the tool for use. However, the finish may require only a small amount of material removal to shape the preform into the tool. For example, the deformation can include a near net shape forging process such that if there is a small number of subsequent processing of the preform, it is required to make the tool.

  One or more secondary treatments may follow the cooling or finishing of the tool. Secondary processing includes forming a coating on the tool or further modifying the surface of the tool in some way. An exemplary secondary process includes spraying or cladding a deformed region of the tool or the entire tool with a wear resistant material. Other secondary treatments include applying a coating on the work surface of the tool by coating techniques including, but not limited to, physical vapor deposition (PVD), chemical vapor deposition (CVD) or salt bath coating. . Other surface modification techniques include ion implantation, laser or plasma hardening techniques, nitridation, or carburization that can be used to modify the surface layer on the work surface of the tool. It will be appreciated that a variety of different secondary processes can be used in any combination to further modify the tool.

  As described above, the preform is made of tool steel. Tool steel represents a class of steel from which tools used to cut, form, or otherwise shape other materials are made. Tool steel can exhibit quenching with heat treatment and can be tempered to achieve the desired mechanical properties. For example, the preforms can be made from a variety of different classes of tool steel, such as cold work, hot work, high speed tool steel grade material, or unique tool steel grade. Specifically, the tool steel is in the range of about 0.35 wt% to about 1.50 wt%, and in another example, along with other carbon content to be considered accordingly if any carbide phase is sought. An iron-carbon (Fe—C) alloy system with a carbon content in the range of about 0.85 wt% to about 1.30 wt%.

Tool steels often include the addition of carbide forming elements such as vanadium (V), tungsten (W), chromium (Cr), molybdenum (Mo), and combinations thereof. Depending on the added _{alloying, M 6 C, M 2 C} , one as _{M 23 C 6, M 7 C} 3 or M ₄ C or more carbide phases may precipitate, other types of carbides It can be formed as known in the art. With few exceptions, tool steel does not contain the intentional addition of nickel (Ni). Nickel is a known austenite phase stabilizer. However, the tool steel may contain trace amounts (0.3 wt% maximum) of this element.

Table 1 shows the nominal composition (balance of tool steel that is iron (Fe)) in weight percent of an exemplary tool steel that can be used to make tool steel according to embodiments of the present invention. As an example, the AC ₃ of the tool steel in Table 1 falls between about 2100 ° F (about 1149 ° C) and about 2400 ° F (about 1316 ° C), and the AC ₁ temperature is about 1300 ° F. (About 749 ° C) and about 1680 ° F (about 915.6 ° C), and Ms is between about 320 ° F (about 160 ° C) and about 480 ° F (about 249 ° C) Enter the range.

  In addition, the preform can also have a powdered metal material or, in particular, a powdered metal tool steel. Powder metal tool steel preforms physically break the tool steel bulk piece, or otherwise into many small individual particles, inject powder metal into the mold or weakly agglomerate It is generally made by passing a powder metal through a die to create a compact and sintering the compact as is known in the art. Tools formed from powdered metal tool steel are often characterized as having isotropic properties as a result of their manufacturing process. However, when processed according to the embodiments disclosed herein, the properties of the tool are improved compared to powder metal tools processed by conventional sintering and / or hot isostatic pressing methods. The

By processing the tool steel as disclosed herein, the microstructure of the tool steel is modified. As described above, the tool steel is deformed while it contains austenite. As is known in the art, austenite has a face centered cubic (fcc) crystal structure and martensite has a body centered tetragonal (bct) crystal structure. Due to its greater sliding surface, austenite is considered to have higher ductility than martensite and by those having ordinary skill in the art. Any austenite that forms above AC ₃ is generally recognized by those skilled in the art as stable. That is, at temperatures above AC _3, austenite typically does not decompose into other phases. At temperatures below AC _3, austenite, if a long time is maintained at a temperature between A _C 3 and Ms, to decompose into other phases, are known to be unstable, often ₃ metastable and be called. Austenite present in the temperature range described here is metastable. Although not wishing to be limited by theory, metastable austenite is believed to possess a strain history despite having the same crystal structure as austenite.

The plastic deformation of a preform containing metastable austenite has a different microstructure than when only quenching is performed between these temperatures, or when the preform is forged at a temperature exceeding AC ₃ and then quenched. Bring. The resulting microstructure and material properties of the deformation zone were determined by the type of tool steel, the type of thermomechanical treatment, the magnitude of strain induced in the austenite, the rate at which the strain was induced, and the thermomechanical treatment performed. Can depend on temperature. For example, thermomechanical processing of metastable austenite at temperatures between Ms and AC ₁ can result in a different microstructure than thermomechanical processing of metastable austenite at temperatures between AC ₁ and AC ₃ . In any case, however, the deformation region exhibits improved mechanical properties.

  As a result of deforming austenite in these temperature ranges, in one embodiment, the microstructure is fine grained. For example, the average size of the particles or crystals in the deformation region can be at least 10% less than that observed with tools made by conventional processing, and in another example can be at least about 25% smaller. In some embodiments, the fine microstructure of the particles facilitates uniform precipitation of the carbide phase along more particle boundaries during quenching or other processing.

In addition, other microstructure features may have an increase in dislocation density. As is known in the art, dislocations are linear defects in crystalline solids, such as in austenite. While other types of dislocations are known, and many types of dislocations are known to form simultaneously in a single crystal, one typical dislocation is an extra half-plane of atoms within the crystal (extra half-plane). -plane). Furthermore, the grain boundaries can be represented by one or more dislocations. In polycrystalline materials, like the tool steel material of a preform, the grain boundary that exists between adjacent crystals is the region of mismatch between the crystal lattice of one particle and the crystal lattice of the adjacent particle. As the degree of misalignment or misorientation angle between adjacent grains increases from 0 degrees at which the crystal structure of the adjacent grains aligns, the dislocation density at the grain boundaries increases. Therefore, the measurement of the misorientation angle between grains is a measurement of the dislocation density, particularly the dislocation density at the grain boundary. By deforming the region of the tool steel preform, the misorientation angle between the particles increases to an angle that exceeds the case where the region of the same composition is deformed by hot forging or heat treatment exceeding AC ₃ by a conventional method. The martensite particles after deformation, quenching and tempering are misoriented, for example, with an average angle of greater than about 34 °, and in another example, the martensite particles average on at least about 40 ° It can be misoriented. In addition, in one embodiment, the dislocation density in the region is at least 25% greater than the hot forged or heat treated portion of the conventional process. Dislocation density and particle size can be measured, for example, by using electron backscatter diffraction (EBSD) or X-ray diffraction (XRD) techniques. In addition to improving the impact strength of the deformation region, high dislocation density locations can provide nucleation points for the precipitation of the carbide phase during deformation or in subsequent heating or cooling operations.

  The deformation region may also indicate a suitable orientation of the particle structure. Specifically, in a cross-sectional view of the deformation region, the particles are collectively such that when the particles are placed or oriented relative to each other, the particles provide a preferential flow or direction to the microstructure. It can be stretched or have other shapes. The direction of preferential orientation can be in one direction relative to one of the tool faces, relative to the tool axis, or to other regions that also have preferential orientation. In essence, the preferential orientation can be in any direction. In one embodiment, the preferential orientation of the particles in the deformation region follows the surface contour of the work surface of the tool. For example, the preferential orientation may follow a surface contour formed by two intersecting surfaces that define an edge. The particle structure may be substantially parallel to each surface, transitioning from a first direction parallel to one surface to a second direction parallel to the second surface in an area adjacent to the edge. . The initial shape of the preform, any carbide or alloy banding present in the preform prior to processing, and processing techniques are key factors in determining the preferential orientation of the particles in the deformation region. obtain.

  Thus, in one embodiment, the deformation region is characterized by a combination of two or more of the microstructure characteristics. For example, the deformation region may have a particle size distribution with a small average particle size and the particles may be preferentially oriented with respect to the work surface or tool axis of the tool. Furthermore, the region can be characterized as having a relatively high dislocation density. In one embodiment, the region may be further characterized by having one or more finer, more uniformly distributed carbide phases located at grain boundaries and at high dislocation density locations. In addition, the features cannot vary greatly from one location to another in the deformation region, although large variations can exist between two or more separately formed regions. For example, a portion of the preform can have a relatively high dislocation density region separated by a relatively low dislocation density region. The variation in dislocation density between regions can be due to the different processes used (eg, radial forging compared to plane strain forging), different forging speeds or different forging strengths, different temperatures, and the like.

Without being bound by theory, the inventors may observe that external energy from thermomechanical processing may form a fine grain structure, provide orientation to the grain structure, increase dislocation density, Alternatively, it may be utilized to provide that combination within the metastable austenite phase. Following quenching, the deformed metastable austenite beneficially affects the microstructure that ultimately forms. In addition, external energy from thermomechanical processing can facilitate the precipitation of carbide phases in the microstructure. For example, thermomechanical treatment at a temperature below AC ₁ is believed to reduce the solubility of carbon in metastable austenite and consequently promote carbide precipitation. In related embodiments, the carbide phase may precipitate at grain boundaries and / or dislocation sites during deformation or during cooling, or both during deformation and cooling. Thus, tool steel preforms processed at less than AC ₁ exhibit greater strength among other improved properties compared to tool steel preforms processed at temperatures above AC ₁ . Furthermore, the increase in dislocation density in this temperature range is believed to be substantially higher compared to preforms that are thermomechanically processed at temperatures higher than AC ₁ .

As previously mentioned, the deformation region of the preform is characterized by improved properties compared to conventional processing (eg, heat treatment and / or hot forging above AC ₃ ). Thus, a tool made from a tool steel preform may, for example, exhibit a longer service life. The improved properties may have an improvement in one or more of impact strength, toughness, hardness or wear resistance (according to Charpy test), or a combination thereof. As a comparison, the impact strength of the deformed region of the preform M2 AISI tool steel processed by an embodiment of the present invention, from the tool of the heat-treated the same composition without modified or forged beyond AC ₃ It can be at least 50% larger. In any embodiment, longer tool life may be due to improved resistance to impact, improved resistance to other stresses, or improved resistance to polishing conditions experienced during use. it can.

  Referring to FIGS. 2A and 2B, and according to another embodiment of the invention, a tool 18 is connected to or coupled to a machine (not shown), a first portion 24, and the tool 18 is metal formed. And a member 20 with an outer surface 22 generally having a second portion in the typical shape of a work surface 26 that contacts the workpiece 28 when used in metal cutting applications. Furthermore, the outer surface 22 surrounds and defines the outer boundary of the bulk volume or mass of tool steel. As best shown in FIG. 2B, at least one region 30 is formed within the enclosed bulk volume, as described herein. If the region 30 is not constituted by the entire bulk volume of the tool 18, the member 20 differs in one or more of the features of the microstructure and thus other regions that differ in the properties described above compared to the region 30. May have 32.

  In one embodiment, referring again to FIG. 2A, the member 20 is elongated and the outer surface 22 is a barrel or shank 34, a head 36 disposed at one end of the shank 34, and , Defining a nose or body 38 with a tip 40 disposed at the end of the shank 34 opposite the head 36. The work surface 26 supported by the tip 40 is connected to the side wall 42 of the tip 40 along the cutting edge 44. The cutting edge 44 defines a corner along which the side wall 42 and the work surface 26 meet. Cutting edge 44 and work surface 26 collectively define the portion of tool 18 that contacts the surface of workpiece 28. The workpiece 28 may have material to be processed by the tool 18 in metal forming or metal cutting applications.

  When viewed along the longitudinal axis or centerline 50 of the tool 18, the shank 34 and body 38 of the elongate member 20 may be of any suitable cross-sectional shape, such as, for example, a round, rectangular, square, or elliptical cross-sectional shape. Have The shank 34 and the body 38 may have the same area cross-sectional shape, or the body 38 may have a smaller cross-sectional area to provide a buffer area 52 between the shank 34 and the body 38. In some embodiments, the shank 34 and the body 38 are symmetrically disposed about the centerline 50 and may specifically have a circular or round cross-sectional shape centered on the centerline 50.

  The head 36 of the tool 18 has a structure suitable for being held by a tool holding device used with a metal working machine such as a machine tool or a press (not shown). In the exemplary embodiment, head 36 is a flange having a diameter that exceeds the diameter of shank 34. However, instead of the head 36, the tool 18 may be a ball-lock, wedge lock, turret, or other for connecting the shank 34 of the tool 18 to the tool holding device. This type of holding structure may alternatively be provided.

  A tool 18 having a punch structure in the exemplary embodiment typically forms the components of a die set 54. The die set 54 further includes a die 56 that includes an opening 58 that receives a portion of the tip 40 of the tool 18. The die 56 and tool 18 cooperate to form a hole formed in the workpiece 28 or to deform the workpiece 28 in some desired manner when pressed together. The tool 18 and die 56 can be removed from the metalworking machine while the tool 18 is temporarily attached using a tool holding mechanism to the end of a ram (not shown).

  Tool 18 generally moves in direction 61 toward workpiece 28 and with the load perpendicular to the contact point between work surface 26 and workpiece 28. The metalworking machine can be driven mechanically, hydraulically, pneumatically, or electrically to apply a load that pushes the tool 18 into the workpiece 28. The tip 40 of the tool 18 is pushed through the thickness of the workpiece 28 or into the thickness and into the die opening 58 under a high load applied by the metalworking machine. The workpiece 28 is deformed and / or cut at and around the contact zone between the work surface 26 of the tool 18 and the workpiece 28.

  The tool 18 may have other punch structures that are different from the structure of the exemplary embodiment. By way of example, the tool 18 may be configured as a blade, heel punch, pedestal punch, round punch, or the like. In an exemplary embodiment, the tool 18 is shown as having a structure that matches the punch, but those skilled in the art will recognize that the tool 18 may have other structures such as a die 56 (FIGS. 2A and 2B) or a die such as a stripper. Will be understood. Specifically, a tool 18 in the form of a punch, die or stripper can be applied in metal stamping and forming processes such as through-hole drilling and drilling, precision punching, forming and extruding or casting.

  The tool 18 may also have a cutting tool structure such as a rotary broach, a non-rotary broach, a tap, a reamer, a drill, a milling machine, a trimming tool or the like. Tool 18 may be used in conventional die casting, high pressure die casting, and casting and molding applications such as injection molding. Tool 18 can also be used in powder compression applications used in the formulation, nutritional, battery manufacturing, cosmetic, confectionery and food and beverage industries, as well as household products and nuclear fuel, tablets, explosives, ammunition, pottery, and others It can be used in the manufacture of products. Tool 18 can also be used in automation and part fixing applications, such as placing details or part-touching the details.

  Referring to FIG.2B, a preform (not shown) in which the region 30 of the tool 18, the region 62 of the die 56, or both the region 30 of the tool 18 and the region 62 of the die 56 have been thermomechanically processed as described above. ) Or machined. For example, the region 30 is often disposed in close proximity or includes the work surface 26 such that the region 30 is in close proximity to or in direct contact with the workpiece 28 during operation of the tool 18. Similarly, region 62 of die 56 is in close proximity to or in direct contact with workpiece 28 when tool 18 and die 56 are used. Region 30 extends from outer surface 22, eg, work surface 26, to a depth d1 greater than 0.039 inches (1 millimeter). Similarly, in the die 56, the region 62 can be irregularly shaped but also extends from the outer surface 63 to a depth d2 greater than 0.039 inches (1 millimeter).

  However, beneficial performance can be observed when regions 30 or 62 are formed at other locations within the tool steel preform. These positions can be determined by factors associated with the work in which the tool 18 is used, or by cost considerations used to balance the use of the tool 18 against its manufacturing costs. In any respect, the thermomechanically treated region 30 is characterized by a high dislocation density, fine grain structure, suitable orientation of the grains, or a combination thereof, as defined above. In one embodiment, high dislocation density, fine grain structure, suitable orientation of the grains, or combinations thereof can be related to the primary deformation direction during thermomechanical processing.

  Tool 18 may have multiple regions of high dislocation density, fine grain structure, suitable orientation of grains, or combinations thereof. In embodiments with two or more regions, each region may be adjacent to an adjacent region in the tool steel preform. It will be appreciated that the orientation of the particles in one region may or may not be substantially aligned with any of the other regions or axes of the tool 18. In yet another embodiment, the region of high dislocation density, preferred orientation of the fine grain structure or grain, or a combination thereof is substantially limited to the tool 18 rather than limited to one or more portions thereof. Extend. In other words, the tool 18 may be machined or formed from a tool steel preform that has been previously thermo-machined according to embodiments herein.

  Referring to FIGS. 3A and 3B, embodiments of the present invention are described and illustrated herein with reference to a preform that is substantially entirely made of tool steel, while in other embodiments, a preform 64 is illustrated. May take the form of a shell 66 made of tool steel with a core 68 made of dissimilar steel. As shown in FIG.3A, the core 68 fills the entire gap in the shell 66 or only a portion thereof, among other variables, depending on the application for the tool (not shown) made therefrom. Can be. The volume of tool steel in shell 66 can be small when compared to the volume of dissimilar steel, but shell 66 is 0.039 inches (1 mm) so that the deformation area is at least 0.039 inches (1 millimeter) thick. Mm). The shell 66 is designed to form the work surface 26 of the tool (see FIG. 1A). The core 68 can form the rest of the tool and can be designed to impart complementary mechanical properties to the tool. By way of example only, the shell 66 may be a tool steel tube, as shown in FIG. 3A. The core 68 can be a cylinder of other steel, such as low carbon or cold worked steel like D2, which is more economical. After insertion of the cylindrical core 68 into the tubular shell 66, the preform 64 is heated and at least the shell 66 is deformed by swaging or radial forging in the temperature range described above. For example, a deformed preform 69 and core 68 after radial forging of the shell 66 is shown in FIG. 3B. A tool 18a formed from a deformed or forged preform 69 or machined, for example, to improve tool life, such as a gear or gear rotation or screw (as shown in FIG. 3C) Although it may include a rotating die and can be used in applications where lateral strength is required, the material cost of the tool is significantly reduced.

  Further details of the invention will be described with reference to the following examples.

[Example 1]
1.5 inch (3.81 centimeters) diameter and 48 inches (121.9 centimeters) long, made up of as-rolled bars known in the art by the symbolic AISI M2, D2 and M4 Eight tool steel preforms were prepared according to one embodiment of the method disclosed herein.

To that end, the bar was heated to 2100 ° F. (1149 ° C.) above AC ₁ in a gas powered furnace. The temperature measurement results were recorded using an infrared pyrometer calibrated within the working range. At this temperature, each microstructure of the bar is considered to be composed of austenite. After the bars reached the target temperature, they were transferred individually (to avoid temperature loss during parts transfer) to the 200 ton-4 hammer radial forging machine inlet roll. Bars 1.500 inches (3.81 centimeters) in diameter by 48 inches (121.9 centimeters) long were each radially forged with four reductions to bars having a diameter of 0.875 inches (2.222 centimeters). Each reduction took between about 15 and about 20 seconds (up to a total of 80 seconds per bar). The calculated effective reduction ratio was 66%. The treated bar was air cooled to room temperature by forced convection.

  It is known that during thermomechanical processing, hot metal will lose heat due to losses due to convection and heat dissipation. Therefore, in order to maintain the temperature of each bar within a narrow temperature range near the target temperature of 2100 ° F (1149 ° C), the external and internal heat from the deformation process should compensate for any heat loss. Used. As a result, forging was performed in a substantially isothermal state. In addition, the temperature was monitored to ensure that any temperature change could be ignored.

Small sections were cut from each bar for analysis during intermediate pressure. None of the samples were observed to show recrystallization. In addition, the phase present in each sample is determined, misorientation between particles is measured, and pole figures are developed for the [001] face of martensite in the transverse (TD) and radial (RD) directions. It was. Measurements were taken after deformation and subsequent tempering at locations that were half the radius of the bar cross-section, or approximately 0.22 inches from the center of the M2 tool steel bar. Phase identification was performed on a Philips X 'Pert X-ray diffractometer. A phase analysis of one M2 bar of Example 1 is shown in FIG. 4A. In FIG. 4A, the fraction of each phase was 0.771473 iron martensite, 0.00419837 chromium-vanadium carbide (658741), 0.219877 iron-tungsten carbide (892579), and 0.00445168V4C3. An EBSD scan was performed on a FEI / Phillips XL30 ESEM-FEG equipped with a field emission controlled environmental scanning electron microscope (ESEM) -EBSD detector. Data was collected using orientation image microscope (OIM) data collection software and mapped with XRD data. A misorientation graph was generated by OIM analysis software . A typical distribution of misorientation angles measured for martensite particles for one of the M2 tool steel bars of Example 1 is shown in FIG. 4B. The pole figure developed for this M2 bar is shown in FIG. 4C.

[Example 2]
Some of the 0.875 inch (2.222 centimeter) diameter bars from Example 1 were reheated above AC ₁ to a temperature of 2100 ° F. (1149 ° C.). After the bars were heated above AC _1, the microstructure was believed to be composed of austenite. When the bars reached the target temperature, they were individually transferred to the inlet roll of a 200 ton-4 hammer radial forging machine. Each bar was radially forged while at 2100 ° F. (1149 ° C.). The bar diameter was reduced from 0.875 inches (2.222 centimeters) to 0.640 inches (1.626 centimeters) under four presses. This reduction in cross-sectional area reached an effective reduction ratio of 47% in addition to the 66% reduction from the first four reductions of Example 1. The treated bar was air cooled to room temperature by forced convection. Several samples were cut from one bar under intermediate pressure to record the effect of strain. Similar to the sample from Example 1, no recrystallization was observed in any of the samples.

  As before, the heat lost to the surroundings and the heat generated from the deformation was balanced in an attempt to maintain the bar at a constant temperature during thermomechanical processing. The temperature was monitored during the process and between the rolls to ensure that the temperature change was negligible. Thus, it is believed that all of the external energy has been transferred to the preform in order to increase the dislocation density and reduce the austenite grain size.

  The bar was then stress relieved at 1400 ° F. (760 ° C.) for 4 hours in a gas powered furnace and subsequently successfully processed through a bar straighter to minimize distortion.

[Example 3]
A 1.5 inch (3.81 centimeter) diameter and 48 inch (121.9 centimeter) length, with AISI M2, D2 and M4 symbolic representations of the as-rolled bar construction known in the art A tool steel preform was prepared.

The bar was heated to a temperature of 2050 ° F. (1121 ° C.) in a gas powered furnace. The microstructure of the bar is considered to be composed of metastable austenite. As before, the temperature measurement results were recorded using an infrared pyrometer calibrated within the working range. After the bars reached the target temperature, each of the bars was pulled out of the furnace and placed on the 200 ton-4 hammer radial forging machine inlet roll. The bar was then air cooled to a processing temperature between about 1100 ° F. (about 593 ° C.) and about 1200 ° F. (about 649 ° C.) (AC ₁ or less). A temperature drop occurred in about 1 minute. The bar was radially forged to a diameter of 1.000 inches (2.54 centimeters) under 7 rolls. The calculated reduction ratio was 56%. The 1.000 inch (2.54 centimeter) diameter bar was air cooled to room temperature by forced convection.

  Similar to the temperature control described in Examples 1 and 2, the bar was kept as constant as possible. The temperature of each bar was monitored during processing and during rolling to ensure that temperature changes were negligible.

Small sections were cut from each bar for analysis during intermediate pressure. None of the samples showed the microstructural characteristics of dynamic recrystallization. Phases are determined and misorientation measurements between particles are taken, and martensite [at a location that is half the radius of the cross section of the bar or about 0.25 inches from the center of the bar. The pole figure was developed for the [001] plane. A phase analysis of one M2 bar of Example 3 is shown in FIG. 5A. The phase fractions in FIG. 5A were 0.737644 iron martensite, 0.0111572 chromium-vanadium carbide (658741), 0.240541 iron-tungsten carbide (892579), and 0.0106579V ₄ C ₃ . A typical distribution of misorientation angles between martensite grains for one of the M2 tool steel bars of Example 3 is shown in FIG. 5B. The pole figure developed for this M2 bar is shown in FIG. 5C.

Comparative Example 1
The as-rolled AISI M2 bar stock is heated to about 1000 ° F, held for about 45 minutes to 1 hour, and cooled to achieve the same hardness as Examples 1 and 3, i.e. HRC61-63. One standard tempering cycle was followed by heat treatment in a 2 bar vacuum furnace using a standard heat treatment cycle by heating the bar to a temperature in excess of 2250 ° F. (about 1232 ° C.). The heat treated bar was then ground to the same outer dimensions as the bar of Example 3.

The pole figure measurement results for the phase, misorientation angle, and comparative bar are shown in FIGS. 6A, 6B and 6C. The fractions of the phases shown in FIG. 6A were 0.660257 iron martensite, 0.00451285 chromium-vanadium carbide (658741), 0.330886 iron-tungsten carbide (892579), and 0.00434446V ₄ C ₃ . The phase present in each bar was substantially the same as provided by the comparative analysis of FIGS. 4A, 5A and 6A.

However, the dislocation density of each bar of Examples 1 and 3 is substantially higher than that of the bar of Comparative Example 1. Specifically, by comparing FIGS. 4B and 5B with FIG. 6B, the misorientation angle of each of the M2 bars of Examples 1 and 3 is significantly higher than the comparative M2 bar shown in FIG. 6B. The average misorientation angle distribution for the bar of Example 1 (Figure 4B) is about 36 degrees, the average misorientation angle distribution for the bar of Example 3 (Figure 5B) is about 42 degrees, The average misorientation angle distribution for the bar of Comparative Example 1 (FIG. 6B) was about 34 degrees. The high average misorientation angle in the M2 tool steel bar of Examples 1 and 3 relative to the comparative heat treated M2 bar indicates higher dislocation density and strain. It is believed that deformation at temperatures below AC ₁ can tolerate an increase in the misorientation angle of the particles relative to deformation at higher temperatures as the particles have less thermal energy and recover from deformation at a slower rate.

  The improved dislocation density for the M2 bar of Examples 1 and 3 is also shown by the pole figures shown in FIGS. 4C and 5C, respectively, when compared to the pole figure of the M2 bar of Comparative Example 1 shown in FIG. 6C. Proved. The pole figure shows that the dislocation density or number of dislocations for the bars of Examples 1 and 3 is significantly higher than the dislocation density for the bar of Comparative Example 1 that was only heat treated. The relative dislocation density is indicated by the density of dots in each graph. Thus, Example 1 (FIG. 4C) has the highest number of dislocations, followed by Example 3 (FIG. 5C), and Comparative Example 1 (FIG. 6C) has the minimum number of dislocations.

[Example 4]
Some of the 1.000 inch (2.54 centimeter) diameter bars from the treatment of Example 3 were reheated to 2050 ° F. (1121 ° C.) (greater than AC ₁ but less than AC ₃ ). The bar was removed from the furnace and air cooled to a processing temperature between about 1100 ° F. (about 593 ° C.) and about 1200 ° F. (about 649 ° C.). After reaching the processing temperature, the bars were each radially forged into bars having a diameter of 0.700 inches (1.778 centimeters) under 7 presses. The calculated reduction ratio was 51%.

  The treated bar was air cooled to room temperature. Several samples were cut from each bar under intermediate pressure. As with the sample of Example 3, none of the bars showed the dynamic recrystallization microstructure characteristics.

  As before, temperature was monitored during processing and during rolling to ensure that temperature changes were negligible.

  The bar was then tempered three times for about 3 hours between about 950 ° F. (about 510 ° C.) and about 1000 ° F. (about 538 ° C.) in a vacuum furnace. It was confirmed that the tempering treatment transformed any retained austenite into martensite. In Examples 1-4, it was noted that the treated bar contained particles that were elongated and preferentially oriented along the longitudinal axis of the bar.

  Although Examples 1 to 4 employ radial forging, other forging techniques known in the art can be used to thermomechanically process the preform as described above. Thus, in the following examples, a near-plane-strain forging was reproduced on a horizontal hot upsetting machine. The preform 65 was developed to be a cylindrical bar when forged using this machine (see FIGS. 7 and 8A and 8B). The cylindrical bar could then be used as a preform for machining or forming metal cutting and metal forming tools.

  With reference to FIGS. 7 and 8A and 8B, in the near plane strain forging process, the geometry of the preform 65 completely composed of tool steel has an oval portion 70 and a cylindrical portion 72. The cylindrical portion 72 does not suffer any deformation and is primarily used to position and hold the preform 65 in the machine during forging. The oval portion 70 or region is heated and undergoes deformation during processing so that a tool can be formed therefrom. After deformation, the deformed preform 75 has a deformed oval portion 73 or region, as best shown in FIG. 8B.

  Referring now to FIG. 9, in the near plane strain forging process, the tool cavity 74 and ram 76 were each designed to form a semi-circular cavity. The resulting circular shape formed by the closure of the tool cavity 74 and ram 76 collectively allows the tool steel to flow in both the radial and circumferential directions while the tool in the elliptical portion 70 in one direction. Designed to block the movement of steel.

[Example 5]
The geometry AISI M2 tool steel preform shown in FIGS. 7 and 8A was machined from an as-rolled mill bar stock. The rolling direction or main carbide direction in a conventional mill bar stock was always concentric with the axis of the cylindrical portion, as indicated by the arrows in FIG. The direction of carbide banding before treatment can determine the orientation of the carbide after thermomechanical treatment. Subsequently, the preform is 1400 ° in a vacuum furnace for 45 minutes-60 minutes to relieve any remaining stress and to achieve a near-equiaxed grain structure. It was first annealed at F (760 ° C).

After annealing, the ellipsoidal portion of each preform was heated above AC ₁ to a temperature of about 1850 ° F. (about 1010 ° C.) using an induction coil. At this processing temperature, the microstructure was thought to be composed of austenite. The temperature was monitored using an infrared pyrometer incorporated into a 50 ton horizontal upsetting machine used to simulate the near plane strain forging process. After the ellipsoidal portion of the preform reached 1850 ° F. (1010 ° C.), each preform was individually forged into a near semicircular cross section (see, eg, FIG. 8B).

  After forging, each bar was quenched to room temperature by convection air cooling. The microstructure after forging was composed of fine-grained austenite. After quenching, austenite transformed to martensite and carbides precipitated. This microstructure is considered unstable and stresses in vacuum furnaces at temperatures between about 950 ° F (about 510 ° C) and about 1000 ° F (about 538 ° C) at a pressure of about 2 bar. Relaxed. After stress relaxation, the preform is placed between about 1200 ° F (about 649 ° C) and 1400 ° F (760 ° C) for 45-60 minutes per cycle for 3 to convert residual austenite to martensite. Processed through multiple tempering cycles. Thereafter, furnace cooling was performed to convert the retained austenite in the microstructure to martensite.

  The impact strength gain from near plane strain forging was due to increased dislocation density and reduced austenite grain size. However, unlike radial forging, in near plane strain forging, deformation occurs almost instantaneously along the entire length of the elliptical portion, so heat loss to the surroundings can be ignored.

[Example 6]
An AISI M2 tool steel preform with the geometry shown in FIG. 8A was machined from the as-rolled mill bar stock and then processed. Similar to the previous preform, the carbide rolling direction before treatment was oriented in the conventional direction (see FIG. 7). Prior to heating and deformation, the preform is 45 minutes at 1400 ° F (760 ° C) in a vacuum furnace to relieve any residual stress in the preform and to obtain a near equiaxed grain structure. Annealed for -60 minutes.

Each preform was heated to a temperature of 2050 ° F. (1121 ° C.) using an induction coil. This temperature was greater than AC ₁ but less than AC ₃ . The temperature was monitored using an infrared pyrometer. Both the coil and pyrometer were incorporated into the ACMA 50 ton horizontal upsetting machine. The microstructure at the temperature between AC ₁ and AC ₃ was composed of austenite. After heating to 2050 ° F. (1121 ° C.), the oval portion was air cooled to a temperature between about 1100 ° F. (about 593 ° C.) and about 1200 ° F. (about 649 ° C.). A temperature drop occurred in about 1 minute. The microstructure was composed of metastable austenite. Thereafter, the elliptical portion was forged to have a circular cross-sectional structure while being maintained at a processing temperature between 1100 ° F. (593 ° C.) and 1200 ° F. (649 ° C.).

  Thereafter, the forged preform was cooled to room temperature. Upon cooling, martensitic transformation and carbide precipitation occurred, resulting in a homogeneous fine-grained microstructure in the elliptical part of the preform. However, the microstructure was considered unstable for most applications due to the presence of residual austenite. Subsequently, the preform was tempered 3 times for 45 minutes-60 minutes at a temperature between 950 ° F. (510 ° C.) and 1000 ° F. (538 ° C.).

Impact strength gain was observed in each of the deformed oval portions. The impact strength gain was due to an increase in dislocation density, a reduction in austenite grain size, and onset of carbide precipitation. Similar to the results observed during radial forging trials, the mechanical properties of the preforms forged at lower than AC ₁ temperature, was improved over those forged beyond AC _1. It is believed that the dislocation density in preforms forged at lower temperatures is significantly higher than the dislocation density produced by forging at higher temperatures.

  Referring to FIGS. 10A and 10B, while the thermomechanical process in the previous exemplary embodiment improves impact strength, it is relatively high in each elliptical section due to the inherent nature of the near plane strain forging process. There are areas of strength and relatively low intensity. The areas of maximum deformation and minimum deformation are oriented substantially perpendicular to each other. For clarity, the preferential forged particle orientation is shown by the curve in FIG. 10B. The areas of relatively low impact strength are typically those that will be in contact with or close to the tool cavity and ram. The region of relatively high impact strength is related to the region of maximum deformation. The cross-sectional dimensions shown in FIG. 10A are about 13.11 millimeters high and about 11.03 millimeters wide, where the width is from the end of the preform (left) and the surface of the deformed oval portion 73 is cylindrical. Measured up to the place where it moves to the shape part 72 (right).

  In preforms where maximal improvement in near uniform material strength is required, a multi-step plane strain forging process can be used to continuously complete the strength of regions of relatively low impact strength. For example, to obtain a cylindrical bar for thermo-mechanically processed metal forming and metal cutting tools, a preform configured with a bar with a rectangular or square cross-sectional geometry is a bar with an oval cross-section. It could be thermomechanically processed using near plane strain forging into the material. Subsequent thermomechanical processing of the oval cross-section to form a bar with a circular cross-section can result in a more uniform distribution of deformation.

  Specifically, referring to FIG. 10B, as a result of the first thermomechanical treatment using plane strain forging, the relatively low strength region is aligned along or close to the region of minimal deformation, In addition, the relatively high strength region is aligned with the relatively high deformation region. Accordingly, a rectangular or square bar that is forged into an oval cross section can be used as a preform for subsequent near plane strain forging processes. In subsequent processing, regions of relatively low intensity can be aligned along the direction of highest deformation. This orientation can be, for example, perpendicular to the initial deformation direction. Thus, a relatively low intensity region will be strengthened as a result of deformation in that region. Conversely, a relatively high strength region from the first forging operation will observe a minimum deformation strength and thus a minimum improvement.

[Example 7]
Two tools were prepared from powder metal preforms of T15 tool steel. The preform was machined from T15 powder metal pressed to annealed thermal equilibrium. It was noted that the microstructure of the plate was almost isotope as a result of its preparation method. The preform had the structure shown in FIG. 11A. As shown, one end of the preform 76 had a pyramid shape. The overall length of the preform was measured to be 5.75 inches (14.6 centimeters), and the pyramidal portion had a total length of 1.75 inches (4.445 centimeters).

Preform 76 is heated using an induction heater in about 4 minutes to a processing temperature between 2000 ° F (1093 ° C) and 2050 ° F (1121 ° C) (between AC ₁ and AC ₃ ). It was. A hot preform is forged in one cycle to a near net shape using a 1000 ton horizontal mechanical AJAX upsetter with a 500 ton die clamping force. It was. A forged preform 78 is shown in FIG. 11B. Specifically, a 1.75 inch (4.445 centimeter) pyramidal end was forged into a 1 inch (2.54 centimeter) rectangular end 80 as shown.

  After forging, the forging 78 is stress relieved at 1400 ° F. (760 ° C.) for 45-60 minutes in an oven. The forged preform 78 was cooled to room temperature in an oven.

  The stress relaxed preform was triple tempered to convert residual austenite to martensite. The final hardness was measured between 63HRC and 66HRC. The three-stage tempered part was machined to remove scale, remove carbon, and provide the final tool shape. A set of two tools 18b, 18c was made from the preform structure shown in FIG. 11B by cutting the preform shown in half.

The two tools 18b, 18c, namely the upper tool and the lower tool, acted on each other (as indicated by the arrows in FIG. 11C) to cut the sheet steel workpiece (not shown). The gap between tools was 0.006 inches (0.01524 centimeters). The workpiece was 22MnB5 steel with AlSi coating sold under the name USIBOR 1500P . The workpiece steel was press hardened to UTS 1500MPa (50HRC). The sheet was measured to be 1.85 millimeters (0.07283 inches) thick. The test was conducted at about 68 ° F. (about 20 ° C.). Wear on the cutting edge was monitored at four locations. Measurements were taken at 5000 impacts or edge profile per cycle.

The measurement results of the cutting edge profile for each upper and lower T15 tool are shown in Figures 12A, 13A, 14A and 15A, which also provide the cutting edge profile for the reference material and CPM M4 powder metal tool doing. (Tools made of CPM M4 powder metal are fully described below in Example 8.) Wear measurements are made at four locations on both the upper and lower tools, while the upper and lower tools Only the two highest wear locations in are provided in the figure. Contour measurements were taken at the locations shown in FIGS. 12B, 13B, 14B and 15B, respectively.

  More specifically, FIGS. 12A and 13A are graphs of the cutting edge contour of the cutting edge of the upper tool at the locations specified in FIG. 12B (position 1) and FIG. 13B (position 4), respectively. 14A and 15A are graphs of the edge profile of the lower tool at the locations specified in FIG. 14B (position 1) and FIG. 15B (position 4), respectively. The cutting edge profiles at locations 1 and 4 as shown in the figure illustrate the wear measurement results at the remaining two unreported locations.

  Referring to FIGS. 12A, 13A, 14A and 15A, the line labeled “Starting Cutting Edge Shape” represents the cutting edge shape before any use. The line labeled “reference” represents the results of measurements made on a tool made of reference material processed according to industry standards.

  The cutting edge profiles at locations 1 and 4 for the T15 tool at 10,000 and 20000 strikes are named “T15 ... 1000 impact” and “T15 ... 2000 impact”, respectively. As illustrated by the graph, the cutting edge of the T15 tool made by the above procedure wears at 10,000 impacts, both the upper and lower tools wear at each location rather than the reference material wears at 10,000 impacts. There were few. At 20000 impact, the T15 tool was worn by an amount comparable to the reference material tool at 10,000 impact. Thus, the T15 tool according to one embodiment of the present invention is almost twice as resistant to wear and impact as the reference material.

[Example 8]
Two tools were prepared from powder metal preforms of CPM M4 tool steel. ( CPM is a US trademark of Crucible Materials, Inc., New York.) The preform was machined from annealed CPM M4 powder metal bulk material. It was noted that the microstructure of the CPM M4 plate had a major carbide banding as a result of the rolling direction used to prepare the bulk CPM M4 material. The preform has the structure shown in FIG. 11A. As shown, one end of the preform had a pyramid shape. The total length of the preform was measured at 5.75 inches (14.6 centimeters), and the pyramidal portion was 1.75 inches (4.445 centimeters) out of 5.75 inches (14.6 centimeters).

The preform was heated using an induction heater in about 4 minutes to between 2000 ° F. (1093 ° C.) and 2050 ° F. (1121 ° C.) (between AC ₁ and AC ₃ ). A hot preform is forged in one cycle to a near net shape using a 1000 ton horizontal mechanical AJAX upsetter with a 500 ton die clamping force. It was. The forged preform is shown in FIG. 11B. Specifically, a 1.75 inch (4.445 centimeter) pyramid shaped end (shown in FIG. 11A) was forged into a 1 inch (2.54 centimeter) rectangular end as shown.

  After forging, the preform is stress relieved at 1400 ° F for 45-60 minutes in an oven. The preform was cooled to room temperature in an oven.

  The stress-relieved preform was tempered in three stages so that any residual austenite was converted to martensite. The final hardness was measured between 62HRC and 64HRC.

The preferential grain orientation in the region of the cutting edge of the forged preform of FIG. 11B was similar to that shown in FIG. 16A. The sample dimensions shown in FIG. 16A were 17.98 millimeters from top to bottom and 13.82 millimeters from side to side. From the measurement results of the cutting edges shown in FIGS. 12A, 13A, 14A, and 15A, the CPM M4 forged tool showed less wear than the reference material at 10,000 impacts. Again, a substantial improvement in tool life was observed.

  Although the invention has been illustrated by descriptions of various embodiments, and these embodiments have been described in considerable detail, it limits the scope of the appended claims to such details, Or it is not the applicant's intention to limit it somewhat. Additional advantages and modifications will readily occur to those skilled in the art. Accordingly, the invention in broader embodiments is therefore not limited to the specific details, exemplary apparatus and methods, and illustrative examples shown and described. Accordingly, departures may be made from such details without departing from the scope of applicants' general inventive concept.

18 tools
20 parts
22 Exterior surface
28 workpieces
30 First area
32 Second area
65 preform
70 areas

Claims (19)

A method for manufacturing a tool by thermomechanically processing a preform composed of tool steel having a martensitic transformation start temperature and a stable austenite temperature, wherein the preform includes a first region, the first region, A second region adjacent to the region, each region having a region containing austenite, the first region having an outer surface and a plurality of outer dimensions with respect to the outer surface,
Establishing at least the first region of the preform at a processing temperature between the martensitic transformation start temperature and the stable austenite temperature;
While the first region of the preform is at the processing temperature, to change at least one of the outer dimensions of the first region and from the outer surface to less than 1 millimeter below the outer surface. Plane strain forging the first region without deforming the second region so as to modify the microstructure of the first region over a depth extending to the depth;
Cooling the first region to room temperature after the first region is deformed;
Including
The microstructure of the first region includes martensite grains having a misorientation angle distribution, the misorientation angle being characterized by an average misorientation angle that is greater than an average misorientation angle from a heat treatment that does not deform. The microstructure of the first region is different from the second region in at least one microstructure characteristic;
Said first region, Ru Tei extends from the outer surface to a depth greater than 1 mm method.   The method of claim 1, wherein after the first region is deformed, the outer dimension of the first region is approximately equal to a near net shape of a tool used in metal forming or metal cutting applications.   The method of claim 1, wherein the first region has a cross-sectional area and the depth extends across the cross-sectional area.   The method of claim 1, wherein the first region has a cross-sectional area, and a change in the at least one of the outer dimensions reduces the cross-sectional area.   The method of claim 1, wherein the first region has a length and a change in the at least one of the outer dimensions increases or decreases the length of the first region. .   The method of claim 1, wherein the microstructure modified by the deformation comprises martensite grains having a misorientation angle distribution characterized by an average misorientation angle greater than 34 °.   The method of claim 1, wherein the processing temperature is kept isothermal while the first region is deformed.   The method according to claim 1, wherein the treatment temperature exceeds an austenite transformation start temperature of the tool steel.   9. The method of claim 8, wherein establishing a preform at the processing temperature comprises heating the first region to a temperature that does not exceed the stable austenite temperature.   Before plane strain forging the first region, the first region is heated to a temperature exceeding the austenite transformation start temperature of the tool steel, and the first region is heated from the temperature exceeding the austenite transformation start temperature to the treatment temperature. The method of claim 1, further comprising the step of cooling one region. The treatment temperature exceeds the austenite transformation start temperature of the tool steel,
The method of claim 1, further comprising maintaining the processing temperature above the austenite transformation start temperature while the first region is deformed. The treatment temperature is between the martensitic transformation start temperature and the austenite transformation start temperature of the tool steel;
The method of claim 1, further comprising maintaining the treatment temperature between the martensitic transformation start temperature and the austenite transformation start temperature while the first region is deformed.   The method of claim 1, wherein the microstructure of the first region does not recrystallize.   The method of claim 1, further comprising tempering the first region, the tempering step comprising heating the first region to a temperature that does not exceed the processing temperature.   Prior to plane strain forging the first region, further comprising assembling a shell made of tool steel with a core made of dissimilar steel, and establishing the first region at the processing temperature; The method of claim 1, comprising establishing at least the shell at the processing temperature and deforming at least a portion of the shell while the shell is at the processing temperature. A tool for use in a machine for deforming and / or cutting a workpiece,
Including members made of tool steel,
The member has an outer surface defining a first portion configured to be coupled to the machine and a second portion adapted to contact the workpiece;
The second portion includes a first region extending from the outer surface to a depth greater than 1 millimeter, and a second region separated from the outer surface by the first region;
The first region has a misorientation angle distribution characterized by an average misorientation angle greater than 34 °, has an average grain size that is at least 10% smaller than the second region, and the second region A tool including a plurality of crystal grains having a crystal grain orientation different from the plurality of crystal grains in the region.   The tool according to claim 16, wherein the average misorientation angle is at least 40 °.   The tool of claim 16, wherein the microstructure of the first region is not recrystallized.   The member includes a shell made of tool steel and a core made of dissimilar steel, the shell including the outer surface defining the first portion, and at least the second region. The tool according to claim 16, comprising a core that forms a part.