Production of high-carbon tool steel components by powder-based additive manufacturing and thermal treatment

Abstract

Methods are disclosed for heat treating a high-carbon steel component, including a high-carbon tool- steel component, printed by powder-based full-melt additive manufacturing, wherein the printed component contains at least 0.2 wt% C and comprises at least 30 vol% of an austenitic phase and wherein the composition of the component may be selected according to a modified Schaeffler diagram. The methods comprise applying a thermal treatment to decompose at least a fraction of the austenitic phase into at least a ferritic phase, austenitization, rapid cooling or quenching to a temperature below which austenite transforms to martensite, and tempering.

Description

[0001] PRODUCTION OF HIGH-CARBON TOOL STEEL COMPONENTS BY POWDER-BASED ADDITIVE MANUFACTURING AND THERMAL TREATMENT

[0002] FIELD

[0003] The invention relates to powder-based additive-manufacturing of high-carbon steels, more particularly to the production of high-carbon steel and tool-steel components by full-melt powder-based additivemanufacturing and thermal treatment.

[0004] BACKGROUND

[0005] Applying conventional manufacturing techniques to produce high-carbon steel components, including high-carbon tool-steel components, with desired combinations of properties — hardness, toughness, abrasion resistance, corrosion resistance, etc. — and shape can require many processing steps and involve significant wastage of material.

[0006] Additive manufacturing ("AM”) offers the possibility of reducing both the number of processing steps required and material wastage, while enabling production of high-carbon steel components, including high-carbon tool-steel components, with complicated shapes, such as internal cooling channels, that would be difficult or even impossible to form with conventional techniques.

[0007] In full-melt powder-based AM, selected regions of a powder are melted using a directed heat source and solidified to successively build up, or print, a three-dimensional component having the same nominal composition as the powder. In Powder Bed Fusion ("PBF"), the selected regions of powder are located within a succession of powder layers formed on a build-plate located in a powder bed. In Directed Energy Deposition (DED), the selected regions of powder are located where the powder is delivered through a nozzle to the growth surface. In full-melt powder-based AM, the directed heat source can be, for example, a laser or an electron beam, as in L-PBF or E-PBF, respectively.

[0008] However, applying full-melt powder-based AM to produce tool steel components has proved difficult, because of the tendency of such components to warp and / or crack during and / or after printing. (WO 2020 / 221689 A1 , Deutsche Edelstahlwerke Speciality Steel GmbH & Co. KG, at page 5, lines 10 to 16). As such, it has been recognized that, for a given tool steel, good weldability is prerequisite for good printability in full-melt powder-based AM. (WO 2020 / 221689 A1 at page 5, lines 3 to 9). In particular, steels with a carbon content above about 0.2 wt%, which have proved difficult to weld without preheating, have been widely believed to be incompatible with full-melt powder-based AM (See, e.g.,Turk et al., Berg Huettenmaenn Monatsh (2019) Vol. 164 (3) pages 112-116, at pages 112, L.H. col., lines 4 to 7, and 113, L.H. col., lines 5 to 7). Similar difficulties with the application of fullmelt powder-based AM have been documented specifically for tool steels, which typically contain more than 0.6 wt% C, such high carbon content being advantageous for producing harder, more wear-resistant tools.

[0009] Attempts to overcome the difficulties of producing high-carbon tool steel components by full-melt powder-based AM have involved application of one or more thermal treatments known in the production of tool-steel components by conventional manufacturing techniques. In austenitization, austenite is formed at elevated temperatures, typically at around 1 ,000 °C. After austenitization, rapid cooling, such as air cooling, or quenching, such as rapid immersion into water, can transform austenite to mainly martensite with, possibly, some retained austenite. Austenitization followed by rapid cooling or quenching to form martensite is typically referred to as hardening. In tempering of martensite, typically above 300 °C, and cryo-treatment, retained austenite can be transformed into martensite.

[0010] US 2021 / 0040591 A1 (Hitachi Metals, Ltd. and Toyota Jidosha Kabushiki Kaisha), at paras. [0019-22, 61-67], reports the production of components with hardness of 40 to 50 HRC from tool-steel powders comprising 0.3 to 0.5 wt% C, 3 to 6 wt% Cr, <1.0 wt% Ni, < 2 wt% Si and <1.5 wt% Mn by L-PBF printing followed by optional austenitizing at 900 to 1100 °C and tempering at 500 to 700 °C.

[0011] WO 2021 / 262707 A1 (MacLean-Fogg Co.), at paras. [0044-73], reports the production of components with hardness of up to around 50 HRC from tool-steel powders comprising 0.13 to 0.21 wt% C, 9 to 15 wt% Cr, 2 to 3 wt% Ni, 0.2 to 0.8 wt% Mo, 0.9 to 1.7 wt% Mn, 0.3 to 0.7 wt% B and at least one boride former by L-PBF printing with build-plate preheating to around 300 °C, followed by austenitizing at 900 to 1 ,200 °C, cryo-treatment at -85 °C and tempering between 100 and 600 °C, yielding a component consisting of boride particles in a martensitic matrix.

[0012] Roetger et al., Additive Manufacturing 161 (2023) 103292 pp 1 to 13, reports the production of components with hardness of around 55 HRC from tool-steel powders comprising around 0.4 wt% C and falling within the austenite + martensite region of a modified Schaeffler diagram (Fig. 6, page 6, L.H. col.) by L-PBF printing with build-plate preheating to 300 °C, followed by austenitization at around 1 ,000 °C and tempering between around 200 and 700 °C. Saewe et al., Additive Manufacturing 46 (2021) 102125, reports cold cracking of components produced by L-PBF printing with build-plate preheating to 350 °C from high-speed steel powder and carbon powder mixtures comprising > 0.6 wt% C and falling within the austenite region of the Schaeffler diagram.

[0013] WO 2020 / 221689 A1 , at page 8, line 9 to page 10, line 24, and claims 1 and 9, discloses the production of components with a hardness of at least 45 HRC by full-melt AM from tool-steel powders comprising from 0.5 to around 5 wt% C and amounts of at least one carbide former selected from among the group “Cr, Mo, V, Ti, Nb, W” by printing and cooling to room temperature, to form a component consisting of a steel matrix with at least 25 vol% austenite, followed by tempering at 400 to 700 °C and / or deep cooling to below -50 °C, to form a component consisting of a steel matrix with at least 80 vol% martensite and 5 to 50 vol% hard-phase particles, such as carbides.

[0014] Kosiba et al., J. Mater. Sci. & Tech. 156 (2023) 1-9, discloses the production of components with high wear resistance and hardness of 59 HRC from tool steel powders with 1 wt% C by L-PBF printing followed by cooling to room temperature, to form a components with 32 wt% martensite.

[0015] Aubert & Duval (https: / / www.aubertduval.com / wp- media / uploads / sites / 2 / 2022 / 11 / X15TN_Stellar_AM_V0_GB.pdf) discloses the production of components with hardness of up to around 56 HRC from a tool steel powder comprising around 0.37 to 0.45 wt% C, 15.0 to 16.5 wt% Cr, 1.5 to 1.9 wt% Mo, 0.2 to 0.4 wt% V and 0.13 to 0.25 wt% N by L-PBF printing with build-plate preheating to 160 °C, followed by stress relieving at 400- 500 °C, austenitization at 1 ,075 °C, cryo-treatment at -80 °C and tempering at 525 °C.

[0016] Asberg et al., Mater. Sci. 742 (2019) 584-89, discloses the production of components with a hardness of 515 to 560 HV20 from H13 tool steel powder comprising 0.39 wt% C, 0.42 wt% Mn, 1.12 wt% Si, 5.36 wt% Cr, 0.09 wt% Ni and 1.05 wt% V by L-PBF printing with build-plate preheating to 200°C, followed by stress relieving at 650 °C, optional Hot Isostatic Pressing (“HIPping”), and standard hardening and tempering.

[0017] Deirmina et al., Mater. Sci. & Eng. A 753 (2019) 109-121 , discloses production of components with a hardness of up to 650 HV1 from H13 tool steel powder by L-PBF, with as-built components having up to 19 vol% austenite, followed by austenization at 1 ,020 °C and tempering at 500 to 650 °C.

[0018] SUMMARY

[0019] Accordingly, there is a need for a method for manufacturing components by powder-based full-melt AM methods having desired properties, such as hardness, toughness, abrasion resistance, corrosion resistance, as good or superior to those achievable with conventionally manufactured high-carbon steel or high-carbon tool-steel components, and having desired shapes and sizes, using high-carbon steel powders and high-carbon tool-steel powders with a range of compositions.

[0020] A method is disclosed for heat treating a high-carbon steel component, including a high-carbon toolsteel component, printed by powder-based full-melt additive manufacturing, wherein the printed component contains at least 0.2 wt% C and comprises at least 30 vol% of an austenitic phase. The method comprises applying a thermal treatment to decompose at least a fraction of the at least 30 vol% austenitic phase into at least a ferritic phase, austenitization, rapid cooling or quenching to a temperature below which austenite transforms to martensite, and tempering.

[0021] It is an advantage of the present disclosure to reduce cracking and / or warping during printing of high- carbon steel components, including high-carbon tool-steel components, formed by full-melt powderbased AM, by reducing the fraction of martensite formed during printing, potentially to zero.

[0022] It is another advantage of the present disclosure to reduce cracking during removal from the build plate of high-carbon steel components, including high-carbon tool-steel components, formed by fullmelt powder-based AM, by reducing the presence of internal stresses introduced during full-melt powder-based AM.

[0023] It is a further advantage of the present disclosure to form high-carbon steel components, including high-carbon tool-steel components, by full-melt powder-based AM without the need for Hot Isostatic pressing (“HIP”, “HIP” processing or “HIPping”). It is yet another advantage of the present disclosure to facilitate the machining of high-carbon steel components, including high-carbon tool-steel components, formed by full-melt powder-based AM.

[0024] It is another advantage of the present disclosure to form high-carbon steel components, including high-carbon tool-steel components, by full-melt powder-based AM with properties, such as hardness, toughness and wear resistance, equivalent or superior to components formed from the same high- carbon steel or high-carbon tool-steel by conventional manufacturing techniques.

[0025] It is a further advantage of the present disclosure to form high-carbon steel components, including high-carbon tool-steel components, by full-melt powder-based AM with a hardness value above 700 HV5 or 60 HRC, superior to components formed from the same high-carbon steel or high-carbon toolsteel by conventional manufacturing techniques.

[0026] It is a further advantage of the present disclosure to allow formation of high-carbon steel components, including high-carbon tool-steel components, by full-melt powder-based AM with a finer microstructure, and therefore a better trade off between toughness and hardness, than components formed from the same high-carbon steel or high-carbon tool-steel by conventional manufacturing techniques.

[0027] It is a further advantage of the present disclosure to allow formation of high-carbon steel components, including high-carbon tool-steel components, by full-melt powder-based AM with homogeneously dispersed ultra-fine carbides, and therefore better wear resistance and lower tendency for cracking, than components formed from the same high-carbon steel or high-carbon tool-steel by conventional manufacturing techniques.

[0028] It is yet another advantage of the present disclosure to allow formation of high-carbon steel components, including high-carbon tool-steel components, by full-melt powder-based AM having more complex shapes than can be produced readily or at all by conventional manufacturing techniques.

[0029] It is yet another advantage of the present disclosure to make use of waste powder that is a byproduct of conventional manufacturing of high-carbon steel or high-carbon tool steel components.

[0030] In an exemplary embodiment, a component is formed by full-melt L-PBF printing of 440C martensitic stainless steel powder — manufacturer specified content of C 0.95 to 1 .20 wt%, Cr 16 to 18 wt%, Mn < 1 wt%, Mo < 0.75 wt%, P < 0.04 wt%, Si < 1 wt%, and S < 0.03 wt% — followed by austenite decomposition performed at 750 °C or 800 °C for 8 hours, and austenitization, performed at 1 ,050 °C for 1 hour, quenching to room temperature, resulted in a component exhibiting a hardness of 747 or 728 HV5, respectively, and a microstructure with an average martensitic grain size and carbide diameter of less than 1 micrometer.

[0031] BRIEF DESCRIPTION OF THE DRAWINGS The above and other features and advantages of the present disclosure will become readily apparent to those skilled in the art by the following detailed description of examples thereof with reference to the attached drawings, in which:

[0032] Fig. 1 shows photographs of as-printed samples of high-carbon tool-steel formed by L-PBF: as- printed and after the austenite decomposition step of the presently disclosed method.

[0033] Fig. 2 shows a Schaeffler diagram with examples of candidate high-carbon tool steels.

[0034] Fig. 3 plots hardness measurements for high-carbon tool-steel-components L-PBF printed with the same composition and under the same process conditions but with different thermal treatments applied, as shown in Table 1 , including the austenite decomposition, austenitization and quenching / rapid cooling steps of the presently disclosed method.

[0035] Fig. 4A plots the hardness of L-PBF-printed high-carbon tool-steel samples as a function of the duration of the austenite decomposition step of the presently disclosed method. Figs. 4B, 4C and 4D are scanning electron microscope (“SEM”) images obtained from the samples used to generate the plot in Fig. 4A.

[0036] Fig. 5 shows SEM images of specimens of L-PBF-printed high-carbon tool-steel samples: Fig. 5A, as- printed; Fig. 5B, after the austenite decomposition step of the presently disclosed method.

[0037] Fig. 6 shows orientation maps obtained by electron back-scatter diffraction (“EBSD”) in a SEM from an L-PBF-printed high-carbon tool-steel component: Fig. 6A, as-printed; and, Fig. 6B, after the austenite decomposition, austenization and quenching / rapid cooling steps of the presently disclosed method.

[0038] Fig. 7 is a schematic representation of a method for producing, using full-melt powder-based AM and thermal treatment, a high-carbon tool-steel component.

[0039] DETAILED DESCRIPTION

[0040] The present disclosure relates to the thermal treatment, also referred to herein as heat treatment, of high-carbon steel and tool steel components formed by full-melt powder-based AM. Tool steels are used in the general sense to refer to any of various carbon steels and alloy steels that are particularly well-suited to be made into tools and tooling, including cutting tools, dies, hand tools, knives, and others, wherein their suitability comes from their distinctive hardness, resistance to abrasion and deformation, their ability to hold a cutting edge at elevated temperatures and, optionally, their resistance to corrosion.

[0041] For present purposes, a component formed by full-melt powder-based AM includes an intermediate or final step in the production of a finished component ready for its intented use. For example, a component formed by full-melt powder-based AM may need to be removed from a build plate or to undergo surface finishing before it can be put to its intended use or, if printed directly onto another component, which may itself have been formed by AM, including by full-melt powder-based AM, the component formed by full-melt powder-based AM may form only a part of a finished component that is ready for its intended use. For present purposes, a component formed by full-melt powder-based AM can include a shaped component with a range of geometries, including a component formed by fullmelt powder-based AM that is a surface layer applied to another component.

[0042] A high-carbon steel, for present purposes, means a steel with a carbon content, including in the form of carbides and carbonitrides, etc., of 0.2 wt% or greater. High-carbon steels include steels having greater than 0.3, 0.4, or 0.5 wt% C. For present purposes, the term high-carbon steel can include tool steels and high-carbon tool steels. A high-carbon tool steel for present purposes means a tool-steel with a carbon content, including in the form of carbides or carbonitrides, etc., of 0.6 wt% or greater. High-carbon tool-steels include those having greater than 0.7, 0.8, 0.9, 1.0, 1.1 , 1.2, 1.3, 1.4, 1.5, 1.6, 1.7, 1.8, 1 .9, 2.0, 2.1 , 2.2, 2.3, 2.4 or 2.5 wt% C.

[0043] For present purposes, decomposing at least a fraction of the austenitic phase, or decomposing austenite or the austenite decomposition step, can mean decomposing at least 10, 20, 30, 40, 50, 60, 70, 80 or 90 vol.% or 100 vol.% of the austenite phase and, into at least a ferritic phase, can mean decomposing solely into a ferritic phase, into a ferritic phase and a carbide, or into a ferritic phases and a carbide and / or other phases.

[0044] Examples of high-carbon tool-steels to which the presently disclosed heat treatment may be applied include, but are not limited to, AISI classifications D2, D6, the 420 and 440 high-C variations (420D, 440B and 440C), M2 and A2, and Uddeholm® 8% Cr Sleipner® and Caldie® and high-alloyed Vanadis® variations.

[0045] Fig. 1 is a photograph that illustrates the difficulties of printing high-carbon tool-steels by full-melt powder-based AM using conventional methods and the advantages of the presently disclosed thermal treatment methods. Fig. 1 shows samples 101 and 103 after their detachment from the build plate and a cm scale marked in mm intervals. Samples 101 and 103 were prepared from the same 440C martensitic stainless steel powder — manufacturer specified content of C 0.95 to 1.20 wt%, Cr 16 to 18 wt%, Mn < 1 wt%, Mo < 0.75 wt%, P < 0.04 wt%, Si < 1 wt%, and S < 0.03 wt% — by L-PBF printing in the same EOS GmbH M290 system using the same processing parameters — a volumetric energy density (“VED”) of 101.8 J / mm3with with no build-plate preheating. Before removal from the build plate, sample 101 was not subjected to any heat treatment and sample 103 was subjected to an austenite decomposition heat treatment — step S106 shown in Fig. 7 and as discussed in detail herein — in this instance, heat treatment at 750 °C for a holding time of 4 hours. A crack 105 is visible in sample 101 , extending from middle of surface 107, the surface of sample 101 that had been closest to the build plate, and towards but not reaching most recently grown surface 109, which appeared during removal of the sample from the build plate, may be attributed to the stresses introduced in the sample during L-PBF printing. No such cracking is evident in sample 103 that was subjected to the austenite decomposition thermal treatment before removal from the build plate.

[0046] Though the presence of predominantly austenite in the as-printed component, 101 in Fig. 1 , as shown in the SEM image of Fig. 5A and EBSD orientation map of Fig. 6A, discussed herein, might be expected to mitigate the tendency to form cracks, any such mitigation is evidently insufficient to prevent cracking. The absence of cracking in sample 103, which is shown in the same orientation as sample 101 , with the surface originally closest to the build plate at the bottom of the photograph, next to the scale, and the most recently grown surface at the top of the photograph, may be attributed to stress relief induced by the austenite to ferrite decomposition step of the presently disclosed method, resulting in predominantly ferrite in the sample after austenite decomposition, as shown in the SEM images in Figs. 4B, 4C, 4D and 5B and further discussed herein.

[0047] The Schaeffler diagram, originally developed for assessing the weldability of stainless steels, provides an empirically derived representation of the phases present as a function of composition and processing conditions after AM. A.L. Schaeffler, Constitution diagram for stainless steel weld metal, Metal Progress 56(11): 680-680B (1949).

[0048] The Schaeffler diagram shown in Fig. 2, following that in the original 1949 publication, plots chromium equivalent wt%, also referred to herein as “Or eq.”, on the horizontal axis against a nickel equivalent wt%, also referred to herein as “Ni eq.”, on the vertical axis, where:

[0049] Cr equivalent wt% = Cr wt% + Mo wt% + 1.5 x Si wt%+ 0.5 x Nb wt%

[0050] And where:

[0051] Ni equivalent wt% = Ni wt% + 30 x C wt%+ 0.5 x Mn wt%

[0052] In Fig. 2, the regions of interest for purposes of the presently disclosed methods, shown as shaded, consist of purely austenite, labeled “Austenite”, but also referred to herein as the “A region”, and those regions with austenite and ferrite, labeled “A+F”, with austenite, martensite and ferrite, labeled “A+M+F”, and with austenite and martensite, labeled “A+M”. Of particular interest are those portions of the A+F, A+M and A+M+F regions that represent more than 50% austenite by volume, those portions that lie closer to the A region.

[0053] The shaded regions of interest are bounded by the lines:

[0054] (I) Ni eq. = 19.3 - 0.8*Cr eq.

[0055] (II) Cr eq.= 14.1 + 0.33*Ni eq.

[0056] Line (I) delineates the boundary between the Martensite region and the shaded A+M region of interest and between the M+F region and the shaded A+M+F region of interest, while line (II) dilneates the boundary between the Ferrite region and the shaded A+M+F and A+F regions of interest.

[0057] Compositions of interest for purposes of the presently disclosed method may be further limited by carbon content, such that:

[0058] C > 0.6 wt% Examples of tool-steel grades with compositions that fall within the shaded regions of interest for purposes of the presently disclosed method are identified with a star symbol and tool-steel grades that fall outside the region of interest for purposes of the presently disclosed method are marked with a triangle.

[0059] The presence of a given tool-steel composition in the regions of interest of the Schaeffler diagram shown in Fig. 2 does not necessarily mean that a single powder of exactly that composition would be readily available, let alone in commercially viable quantities, for example, as a byproduct of conventional high-carbon tool-steel manufacturing. However, the desired tool-steel composition may be available as a mixture of powders of different compositions that, on average, correspond to the given tool-steel composition.

[0060] The samples shown in Fig. 1 were formed from high-carbon tool-steel 440C, as already noted, L-PBF printed under the same conditions using the same process parameters, corresponding to the A subregion of the Schaeffler diagram, as indicated by the presence of predominantly austenite phases in the as-printed state, as shown in the SEM image of Fig. 5A and the EBSD orientation map of Fig. 6A as further discussed herein.

[0061] In general, the hardness of a given component will be a function of factors that include the relative proportions of the Austenite, Martensite and Ferrite phases, the presence of dispersed hard phases, particular carbides but also carbo-nitrides and nitrides if sufficient nitrogen is available, and the grain size or paricle size of such features, factors that can change after heat treatment.

[0062] Fig. 3 plots, on the vertical axis, hardness measurements obtained from samples — printed using L- PBF from high-carbon tool-steel 440C under the same conditions as those detailed in connection with samples shown in Fig. 1, as-printed sample 101 that cracked upon removed from the build plate and sample 103 subjected to austenite decomposition that did not crack upon removal — and subjected to thermal treatments 1 through 14, as indicated on the horizontal axis, and as detailed in Table 1.

[0063] Hardness was measured using Vickers hardness indentation.

[0064] Thermal treatments below 950°C were carried out in a Netzsch Thermal Analyzer with a controlled heating and cooling rate (HT1-4, HT9-12). The remaining high temperature thermal treatments were carried out in a horizontal Kanthal tube furnace followed by subsequent gas quenching. All thermal treatments were carried out in an inert gas atmosphere.

[0065] TABLE 1

[0066]

[0067] As shown in Fig. 3 and Table 1 , thermal treatments at 450, 550 or 650 °C for 1 hour resulted in hardness values, HV5 of between 400 and 500 for thermally treated samples 1, 2 and 3, while thermal treatments at 750 °C for 1 , 2, 4 and 8 hours resulted in a strong decrease in hardness values, HV5 of between 300 and 330 for thermally treated samples 4, 5, 6 and 7. Thermal treatments at 850 °C for 1 and 8 hours, resulted in recovered / in creased hardness values, HV5 slightly above 500 for thermally treated samples 8 and 9, while thermal treatments at 950, 1050 or 1150 °C for 1 hour resulted in yet further increases in hardness values, HV5 between 700 and 800 in thermally treated samples 10, 11 and 12.

[0068] Samples of high-carbon tool-steel 440C printed using L-PBF, including those from which hardness data was obtained, as plotted in Fig. 3 and Fig. 4A and shown in Table 1 , were further investigated by X-ray diffraction (“XRD”), performed on 2.5 x 2.5 x 4.5 cm samples, electron back-scatter diffraction (“EBSD”) performed in a scanning electron microscope (“SEM”) using electropolished cross sections, light optical microscopy (“LOM”) performed on appropriately etched cross sections, and energy dispersive x-ray diffraction (“EDS”) elemental analysis performed in an SEM using electropolished cross sections.

[0069] As discussed herein, such investigatons of the microstructure and composition of these components allow for an improved understanding of the observed relationship between specific thermal treatments applied and hardness values obtained from components L-PBF printed from a particular high-carbon tool steel, 440C, as plotted in Figs. 3 and 4A and shown in Table 1, and of the underlying mechanisms that may give rise to such relationships. The results of such investigations of the L-PBF printed high-carbon tool steel 440C components thermally treated under the conditions detailed herein can be generalized to components formed by full-melt powder-based AM from a range of high- carbon steels and high-carbon tool steels subjected to similar thermal treatments. Based on these investigations of high-carbon tool steel 440C printed by L-PBF, certain relationships between thermal treatment and observed hardness for full-melt powder-based AM components may be understood, as detailed in the following.

[0070] Initial hardening, upon thermal treatment at relatively low temperature for a fixed time period — for the exemplary 440C samples, as plotted in Fig. 3 and shown in Table 1, as observed in samples 1, 2 and 3, after thermal treatment for 1 hour at 450, 550 and 650 °C, respectively — which may be understood in general terms as resulting from the known effects of secondary hardening, due at least in part to the precipitation of carbide particles.

[0071] Considerable softening, upon thermal treatment at moderate temperature for increasing time periods — for the 440C samples, as plotted in Figs. 3 and 4A and shown in Table 1, as observed in samples 4, 5, 6 and 7, after thermal treatments at 750 °C for 1 , 2, 4 and 8 hours, respectively and, more generally, corresponding to step S106 shown in Fig. 7 and described herein — which can be understood as resulting from austenite decomposition to ferrite and possibly carbides.

[0072] Hardness recovery, to that of the as-printed component, and then further hardening, upon moderate to high temperature thermal treatment for different time periods — for the 440C samples, as plotted in Fig. 3 and shown in Table 1, as in for samples 8, 9 and 10, after thermal treatment at 850 °C for 1 hour, 850 °C for 8 hours and 950 °C for 1 hour, respectively — which can be understood as resulting from the transformation of reverted austenite to ferrite and possibly carbides.

[0073] Further hardening, upon high temperature thermal treatment followed by rapid cooling or quenching — for the 440C samples, as plotted in Fig. 3 and shown in Table 1, as observed in samples 11 and 12 after thermal treatment at 1 ,050 °C and 1 ,150 C, respectively, for 1 hour followed by quenching to room temperature and, more generally, corresponding to step S108 shown in Fig. 7 and described herein — which can be understood as resulting from full re-austenitization during the high temperature thermal treatment and martensite formation during the rapid cooling or quenching.

[0074] Surprising additional hardening, upon application of a combined thermal treatment — for the 440C samples, as plotted in Fig. 3 and shown in Table 1, as observed in samples 13 and 14 after thermal treatment at 750 and 850 °C, respectively, for 8 hours, followed by thermal treatment for 1 hour at 1050 °C with quenching to room temperature and, more generally, corresponding to step S106 followed by step S108 shown in Fig. 7 and described herein — which can be attributed to the surprising effect of introducing an austenite decomposition step, as described further herein, before full re-austenitization and martensite formation.

[0075] As discussed herein, the particular combined thermal treatment that resulted in surprising additional hardening in the 440C samples — those applied to samples 13 and 14, corresponding to the austenite decomposition treatment applied to samples 7 and 9, respectively, followed by the high temperature thermal treatment and quenching to room temperature applied to sample 11 — may be applied to fullmelt powder-based AM components formed from high-carbon steels and high-carbon tool steels with a range of different compositions. Regarding the effects of the austenite decomposition step on measured hardness, Fig. 4A is a plot, based on the same batch of L-PBF printed 440C samples that was used to generate the data plotted in Fig. 3 and shown in Table 1 , of hardness measured by Vickers indentation as a function of duration of the austenite decomposition step, corresponding to step S106 shown in Fig. 7 and described herein. As shown in Fig. 4A, upon thermal treatments at 750 °C lasting for up to around 1 h, hardness drops sharply, from 500 to around 350 HV0.5, and, for thermal treatments at 750 °C lasting from 1 h up to around 100 h, hardness declines more gradually, from 350 to around 300 HVo.5.

[0076] Regarding the effects of the austenite decomposition step on microstructure, Figs. 4B, 4C and 4D are SEM images obtained from electropolished specimens of samples thermally treated at 750 °C for 0.5, 1.0 and 10 hours, respectively, as shown in Fig. 4A, and overlapping with the thermal treatment conditions of samples 4 to 7 plotted in Fig. 3 and shown in Table 1. The SEM images shown in Figs. 4B, 4C and 4D demonstrate the successive formation of an increasingly fine microstructure, indicative of austenite decomposition to ferrite. The SEM images in Figs. 4B, 4C and 4D were obtained at an operating voltage of 20 kV, using a back scatter electron detector, a working distance of 6.6 mm, and an aperture size of 60 pm. The scale bars are 2 pm, corresponding to an image width of 45.73 pm.

[0077] Regarding the effects of the austenite decomposition step on microstructure, Fig. 5 shows higher resolution SEM images obtained from electropolished specimens prepared from the same batch of L- PBF printed 440C samples used to generate the data plotted in Figs. 3 and 4A and shown in Table 1 and the SEM images shown in Figs. 4B to 4D. The SEM images in Figs. 5A and 5B were obtained at an operating voltage of 20 kV and magnification of 5,000 X, using a back scatter electron detector, a working distance of 7.5 mm and an aperture size of 120 pm. The scale bars are 2 pm, corresponding to an image width of 22.87 pm. The SEM image in Fig. 5A was obtained from an as-printed sample, while that in Fig. 5B was obtained from a sample that had been subjected to austenite decomposition to at least ferrite, by thermal treatment at 750 °C for 1 hour, corresponding to sample 4 plotted in Fig. 3 and shown in Table 1 and, more generally, to step S106 shown in Fig. 7 and described herein.

[0078] In the SEM image of Fig. 5A, the as-printed sample shows austenite dominating as a cellular dendritic structure, with average grain size > 10 micrometer, with lighter regions at the dendrite boundaries, the last regions to solidify, corresponding to mainly carbides, specifically M7C3, as shown by Transmission Electron Microscopy (“TEM”).

[0079] In the SEM image of Fig. 5B, the sample after austenite decomposition, performed by heating to 750 °C for 1 hour, lighter regions, again carbides, but very finely dispersed, nanometer sized, embedded in darker regions of finely grained ferrite and, possibly, martensite, are evident.

[0080] Regarding the effects of the combined austenite decomposition and hardening steps on microstructure, Figs. 6A and 6B show orientation maps obtained in an SEM using EBSD from electropolished specimens prepared from the same batch of L-PBF printed 440C samples used to generate the data plotted in Figs. 3 and 4A and shown in Table 1 , and the SEM images shown in Figs. 4B to 4D and 5. In the EBSD orientation map shown in Fig. 6A, obtained from an as-printed sample, coarse grains of characteristic cellular dendritic austenite are evident, with average grain size > 10 micrometer, and certain grains as large as 50 micrometer, consistent with the SEM image of an as-printed sample shown in Fig. 5A.

[0081] In the EBSD orientation map shown in Fig. 6B, obtained from a sample subjected to austenite decomposition (750 °C for 8 hours), austenitization (1 ,050 °C for 1 hour) and quenching to room temperature, corresponding to sample 13 plotted in Fig. 3 and shown in Table 1 that exhibited an extremely high measured hardness value, 747 HV5, a distribution of ultrafine, < 1 micrometer diameter, carbide particles are evident, embedded in, at least 50% by volume, martensite grains with an average diameter of less than < 1 micrometer. The average martensite grain size was determined from the area-weighted mean in the EBSD orientation map. The average carbide size of < 1 micrometer was determined from high-resolution SEM images, such as those shown in Fig. 5.

[0082] The lower hardness that results from the austenite decomposition heat treatment step, as illustrated in the hardness values plotted in Figs. 3 and 4A, has benefits that include facilitating sample removal from the build plate without the sample cracking, as illustrated by the photographs of cracked, without the austenite decomposition treatment step, and uncracked, with the austenite decomposition treatment step, samples 101 and 103, respectively, shown in Fig. 1 and described herein, and of increasing the ease with which such a thermally treated component may be machined. At the same time, the austenite decomposition into finer grained ferrite and possibly carbides acts as a stress relieving treatment, that minimizes the risk of cracking after printing.

[0083] Fig. 7 is a schematic representation of a method 700 for producing and heat-treating a printed high- carbon tool-steel component that results in the surprising effects of increased hardness and other advantages as detailed herein.

[0084] Step S102 corresponds to selection of a suitable powder composition, such as a powder composition that contains greater than 0.2 wt % C or a powder composition that lies within a region of interest of the Schaeffler diagram that contains austenite, which can be the region shown in Fig. 2 and labeled Austenite, or the at least portions of the austenite with ferrite region, A + F, the austenite with martensite and ferrite region, A + M + F. or the austenite with martensite region, A + M in Fig. 2 , as bounded the lines:

[0085] (I) Ni eq. = 19.3 - 0.8*Cr eq.

[0086] (II) Cr eq.= 14.1 + 0.33* Ni eq.

[0087] The powder composition falling within the above-identified regions can be further selected to satisfy the relation:

[0088] C > 0.6 wt%. Step S104 corresponds to full-melt powder-based AM using the selected austenite-containing powder composition, such as using L-PBF, under process conditions that can yield a partially or mainly austenitic microstructure. Those portions of the A + M, A + F and A + M + F regions that correspond to a mainly austenitic structure, one with more than 50 vol% austenite, lie closer to the boundary with the neighboring austenite region. Though selecting a powder composition within one of such defined regions of the Schaeffler diagram may allow for the possibility of printing a sample with partially, mainly or even fully austenitic structure, for example, more than 25, 30, 40, 50, 60, 70, 80 or 90 vol%, for a particular high-carbon steel or tool-steel powder composition, optimization of printing conditions — e,g., powder layer thickness, scan speed, laser power, hatch distance — may be needed to realize such a volume fraction of austenite. Such optimization would be expected to be within the capacity of the skilled person.

[0089] Step S106 corresponds to austenite decomposition to at least ferrite, for example, in higher carbon content austenite, to ferrite and carbides. As discussed herein, this can be performed at around 700 °C or 800 °C for a period as short as 1 minute or up 0.5 hour, or up to around 10 hours, or up to around 100 hours. More generally, the austenite decomposition may be performed at a temperature between around 650 °C and 850 °C, or between 650 °C and 950 °C, depending on the composition and structure of the particular steel and component, so long as the temperature is kept sufficiently high to decompose at least a fraction of the austenite to at least ferrite and kept sufficiently low so as not to generate a microstructure where austenite is the majority structure. For a given high-carbon steel or high-carbon tool-steel, at least a fraction of the austenite in the as-printed sample can be decomposed to at least a ferritic phase by thermal treatment, according to step S106 described herein, at temperatures above those that result in secondary hardening and softening but lower than those that at which austenite is stable. Such austenite decomposition can lead to further softening, as shown for the measured hardness of the high-carbon 440C components plotted in FIG. 4A and discussed herein, and is thus distinct from tempering of tool steel that contains a significant fraction of retained austenite, a process that can lead to minor softening if conducted at sufficiently low temperatures, or hardening if conducted in the secondary hardening regime. The austenite decomposition step may be performed by intercritical annealing.

[0090] For a given tool steel, such austenite decomposition differs from austenitization conducted at higher temperatures, higher than those required for decomposition of austenite to ferrite and possibly other phases, such as those described herein in connection with the hardness data plotted in FIG. 3. The skilled person will appreciate that the thermal treatment conditions required to decompose at least a fraction of the austenite to at least ferrite may vary, depending on the type of high-carbon steel or tool-steel, and that austenite decomposition of carbon containing steels and tool-steels can result in carbide formation.

[0091] As already noted, as illustrated by the high-carbon 440C samples shown in Fig. 1, after the austenite decomposition step, S106, the component may be more easily removed from the build plate without cracking. After the austenite decomposition step, S106, because of the increased softness of the component, for example, softer than 400 HV5, the component may also be more amenable to machining, including before or after removal from the build plate.

[0092] Step S108 corresponds to a hardening process that comprises austenitization followed by rapid cooling or quenching to a temperature typically lower than around 200 °C. During austenitization, austenite is formed at high temperature, typically between around 1 ,000 °C and around 1 ,250 °C, e.g., at or around 1 ,000 °C, 1 ,050 °C, 1 ,100 °C, 1 ,200 °C or 1 ,250 °C, depending on the composition of the tool-steel grade. The component may be maintained at such temperature for between around 5 minutes to 2 hours, depending on the tool-steel grade.

[0093] During rapid cooling or quenching, such as air cooling, or quenching, such as rapid immersion into water, austenite may be transformed to martensite, with, possibly, some retained austenite, by cooling at a sufficiently high rate from the austenitization temperature to a sufficiently low temperature, typically lower than around 200 °C. The skilled person will appreciate that such an austenite to martensite transformation causes hardening. Rapid cooling or quenching may involve cooling rates of more than 1 °C / s, such as 10 °C / s or 100 °C / s. The skilled person will appreciate that sufficiently high cooling rates may be required to drive martensite formation in place of decomposition reactions. Quenching to a temperature below around 200 °C may involve cooling to a temperature of at or around 100 °C, ambient temperature, 0 °C, -80 °C, or -196 °C, depending on the quenching media used, water, air, oil, dry ice, etc. Such quenching may involve cooling directly to sub-zero Celsius temperature, including to a cryogenic temperature below 120 K, or may involve a cooling to an intermediate temperature followed by cooling to a sub-zero Celsius temperature, such as in the deep cooling or cryogenic treatment step discussed herein. The skilled person will appreciate that rapid cooling or quenching from the austenitization temperature to below around 200 °C may be performed under a variety of conditions, using different cooling rates provided by different quenching media, depending on the martensite start temperature, the value of which strongly depends on the composition of the particular steel type.

[0094] Step S110 corresponds to a classical tempering process. A skilled person will appreciate that such tempering can be performed under a variety of conditions, examples being between around 100 and 700 °C for a period of between around 30 and 240 minutes. The skilled person will appreciate that the tempering step may be performed under a variety of conditions. Tempering step 110 may be repeated, as may the hardening by austenitization and quenching step S108.

[0095] An optional sub-zero Celsius treatment step or steps may be performed after hardening, austenitization and quenching, step S108, as already noted, and / or before / after tempering, step 110. Such a sub-zero Celsius treatment step or steps may involve cooling at a rate of between around 0.01 °C / s and around 1 ,000 °C / s, e.g., by cooling at a lower rate in a mechanical fridge or at a higher rate by immersion in cryogenic liquid, e.g., liquid nitrogen, or a solid, e.g., dry ice, with lower cooling rates being generally favored for larger components and higher cooling rates being generally more feasible for smaller parts. The cooling rate selected may also depend on the tool-steel grade and the shape of the component. The sub-zero Celsius treatment temperature may be in the deep cooling region below - 50 °C, e.g., around - 80 °C, for dry ice, or in the cryogenic region below 120 K, e.g., around - 196 °C, for liquid nitrogen, or - 269 °C, for liquid helium, and may be maintained for a period up to around 96 hours, including up to around 48 hours. The sub-zero Celsius step may be repeated before or after tempering and in the event the tempering step is repeated.

[0096] One or more of the steps shown in Fig. 7, and related steps described herein, may be performed in a single piece of equipment, for example, a component may be printed by L-PBF and thermally treated, to induce austenite decomposition to ferrite and possibly carbides, inside the build volume of the L- PBF printer, or in different pieces of equipment, including equipment at different locations. For example, a component may be printed by L-PBF at one location and transported with sufficient care to avoid detrimental mechanical and / or thermal stresses to another location for austenite decomposition to ferrite and possibly carbides step S106, including transportation with the component remaining attached to the L-PBF build plate or, if the printed component is built onto another component, still attached to the other component. After austenite decomposition to ferrite and possibly carbides in step S106, the component may be removed from the build plate, further machined, and hardened by austenitizing and quenching step S108 at the same or another location or locations. After hardening by austenitizing and quenching, the component may be tempered S110 and / or cryogenically treated at the same or another location or locations.

[0097] The processes described herein may be applied to a range of different high-carbon steels, including tool steels, with carbon content above 0.2 wt%, including those with a carbon content above 0.3, 0.4 or 0.5 wt% C, and, in particular, to high-carbon tool steels with carbon content between 0.6 and 2.5 wt%, such as tool steels with around 2.2 wt% C and around 25 wt % Cr. The process described herein maybe used with injection molding tool steels, cold-work tool steels, hot-work tool steels and high-speed steels (HSS), that remain fully or partially austenitic when consolidated via full-melt Additive Manufacturing techniques such as Powder Bed Fusion or Directed Energy Deposition, including Laser Cladding.

[0098] The disclosed methods for heat treating high-carbon steel components, including high-carbon toolsteel components, formed by powder-based full-melt AM disclosed herein including components having a composition defined according to the regions or portions of regions of the Schaaffler diagram shown in Fig. 2 and as discussed herein, can be used to form components with characteristic microstructures and properties, such as hardness, toughness, abrasion resistance, corrosion resistance, equal to and, in some instances, superior to the porperties of com HV5 HVsponents formed using conventional manufacturing techniques from the same high-carbon steel or tool steel.

[0099] The disclosed methods can be used to form high-carbon steel, including high-carbon tool-steel, components with a hardness that exceeds 700, 710, 720, 730, 740, 750, 760, 770, 780, 790 or even 800 HV5.

[0100] The disclosed methods can be used to form tool-steel components containing at least 0.6 wt% C with a microstructure comprising at least 50 vol% martensite, carbides and nitrides, in which all the microstructural features are smaller than 5, 3, 4 or 1 micrometers. The disclosed methods can be used to form tool-steel components with a microstructure comprising at least 50 vol% martensite and containing at least 0.6 wt% carbon for which the carbide and nitride phases do not together exceed 15 vol% with a hardness that exceeds 62 HRC.

[0101] The disclosed methods can be used to form a tool-steel component with a microstructure comprising at least 50 vol% martensite and containing at least 0.6 wt% carbon, for which the average size of the carbide and nitride phases does not exceed 2 micrometers and the average martensite grain size does not exceed 10 micrometers, with a hardness that exceeds 650 HV5.

[0102] The disclosed methods up to the austenite decomposition step, S106 in Fig. 7, can be used to form a tool-steel intermediate component with a eutectoid microstructure comprising ferrite and carbides, having a carbon content of at least 0.6 wt%, and with a maximum hardness of 400 HV0.5, such a product being removable being machinable without risk of cracking, as a result of stability against additional phase transformations during machining.

Claims

CLAIMS1. A method for heat treating a high-carbon steel component printed by powder-based full-melt additive manufacturing, wherein the printed component contains at least 0.2 wt% C and comprises at least 30 vol% of an austenitic phase, the method comprising: decomposing at least a fraction of the austenitic phase into at least a ferritic phase; austenitization; rapid cooling or quenching to a temperature below which austenite transforms to martensite; and tempering.

2. The method according to claim 1 , wherein the high-carbon steel is a high-carbon tool-steel having a composition as plotted on a Schaeffler diagram:Ni eq.= Ni+30*C+0.5*Mn andCr eq.= Cr+ Mo + 1 ,5*Si + 0.5*Nb, and lies within a full austenite or an austenite-containing region bounded by the lines:Ni eq. = 19.3 - 0.8*Cr eq.Cr eq.= 14.1 + 0.33*Ni eq. and contains at least 0.6 wt% C.

3. The method according to claims 1 or 2, wherein the austenitic phase comprises more than 50 vol% of the printed component.

4. The method according to any of the preceding claims, wherein decomposing at least a fraction of the austenitic phase is performed at a temperature between 650 and 950 °C and / or for a time interval of between 1 minute and 100 hours.

5. The method according to any of the preceding claims, wherein the austenization is performed at a temperature between 950 and 1 ,250 °C and / or for a time interval of between 5 minutes and 2 hours.

6. The method of any of the preceding claims, wherein the cooling rate in the rapid cooling or quenching step is between 1 °C / min and 3,000° C / min.

7. The method according to any of the preceding claims, wherein the tempering is performed at a temperature between 100 and 700 °C and / or for a time interval of between 30 minutes and 240 minutes.

8. The method of any of the preceding claims further comprising a sub-zero Celsius treatment of the component wherein the component is cooled to at least -78 °C after the austenitizing step and before the tempering step.

9. The method according to claim 8 or 9, wherein the sub-zero Celsius treatment step lasts between 5 minutes and 96 hours and / or the cooling rate is no less than 0.01 °C / s and / or the heating rate terminating the sub-zero Celsius treatment step is no more than 1 ,000 °C / s .

10. The method of any of the preceding claims, further comprising the step of printing of the high- carbon tool-steel component by powder-based full-melt additive manufacturing.11 . The method of claim 10, wherein the printing and at least one of the austenite decomposition, austenization and tempering steps are performed in the same equipment.

12. The method of any of the preceding claims, wherein the printed component remains attached to a build plate or another component during at least the austenite decomposition step thermal treatment.

13. A tool steel component having a composition as plotted on a Schaeffler diagram:Ni eq.= Ni+30*C+0.5*Mn andCr eq.= Cr+ Mo + 1 ,5*Si + 0.5*Nb and lies within a purely austenite or austenite-containing region bounded by the lines:Ni eq. = 19.3 - 0.8*Cr eq.Cr eq .= 14.1 + 0.33*Ni eq. and contains at least 0.6 wt% C, wherein the tool steel component comprises martensite, carbides and nitrides, wherein all the microstructural features are smaller than 1 micrometer.

14. The component of claim 13 wherein the martensite phase comprises at least 50 vol.%.

15. The component of claim 13 wherein the carbide and nitride phase do not exceed 15 vol.% and the hardness of the component exceeds 62 HRC

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