Fe-Cr-Al alloy powder for additive manufacturing, Fe-Cr-Al alloy member, and method for manufacturing Fe-Cr-Al alloy member

Abstract

The present invention provides: an Fe-Cr-Al-based alloy powder for additive manufacturing that has improved high-temperature oxidation resistance and excellent malleability and is capable of suppressing cracking; an Fe-Cr-Al-based alloy member that comprises an additively manufactured article using the alloy powder; and a method for producing an Fe-Cr-Al-based alloy member.　Provided is an Fe-Cr-Al-based alloy member which is an additively manufactured article that contains, in terms of mass%, 20-30% Cr, 2.0-7.0% Al, more than 0.15% but 0.75% or less Ti, and 0.001-0.02% C, where 15≤(Ti / C)≤300 is met, with the remainder being Fe and unavoidable impurities, wherein carbides deposited at the grain boundaries of crystal grains are Ti carbides having an average diameter of 100 nm or less in equivalent circle diameter.

Description

[Technical Field]

[0001] The present invention relates to an Fe—Cr—Al alloy powder for additive manufacturing, an Fe—Cr—Al alloy member, and a method for producing an Fe—Cr—Al alloy member. [Background technology]

[0002] Fe-Cr-Al alloys have advantages such as high resistivity, good thermal stability, and long service life. They can also form aluminum oxide, chromium oxide, or composite oxide films that provide protection on their surfaces at high temperatures, providing resistance to oxidation, carburization, sulfur, and hydrogen gas corrosion at high temperatures. Therefore, they are used as components and parts in industrial diffusion furnaces, high-temperature furnaces, chemical plants, semiconductor manufacturing processes, and the like. Fe-Cr-Al alloys, which have excellent oxidation resistance at high temperatures (hereinafter referred to as high-temperature oxidation resistance), are particularly popular for components in highly corrosive high-temperature environments.

[0003] In addition, the designs of the above-mentioned components and parts have become more complex, and there is a growing demand for complex shapes that are difficult to produce using die forging or machining. Therefore, additive manufacturing has recently been applied to the production of components with complex shapes.

[0004] Patent Documents 1 and 2 disclose methods for producing Fe-Cr-Al alloy powder that are applied to additive manufacturing using vacuum gas atomization. Patent Document 2 also discloses additive manufacturing (AM) as a method for producing Fe-Cr-Al alloy members. Additive manufacturing involves repeatedly supplying a heat source to raw material powder to melt and solidify the raw material powder, thereby obtaining a three-dimensional additively shaped object. Additive manufacturing (hereinafter sometimes referred to as additive manufacturing) makes it possible to obtain three-dimensionally shaped members or products (hereinafter referred to as additively shaped objects) in net shape or near net shape, even if the shape is complex.

[0005] Additive manufacturing (AM) can produce complex-shaped components in net or near-net shape, and when applied to alloys with high-temperature oxidation resistance, it can achieve excellent properties in both strength and oxidation resistance. However, because Fe-Cr-Al alloy components are composed of a body-centered cubic (BCC) lattice structure, they are known to undergo a ductile-brittle transition (DBT) in the low-temperature range (100°C to 350°C). In particular, AM methods are prone to solid-state cracking (SSC) caused by DBT due to high thermal stress caused by rapid cooling rates. According to Non-Patent Document 1, AM objects obtained from Fe-Cr-Al alloy powders by AM have the problem of cracking and reduced ductility. [Prior art documents] [Patent documents]

[0006] [Patent Document 1] Publication number CN110125383 [Patent Document 2] WO2021-078885 publication

[0007] [Non-Patent Document 1] MATER.RES.LETT.2021,VOL.9,NO.8,350-357 Summary of the Invention [Problem to be solved by the invention]

[0008] Based on the above, the present invention aims to provide an Fe-Cr-Al alloy powder for additive manufacturing that has improved high-temperature oxidation resistance, excellent ductility, and can suppress cracking, as well as an Fe-Cr-Al alloy part made from an additive manufacturing body using this alloy powder, and a method for manufacturing an Fe-Cr-Al alloy part. [Means for solving the problem]

[0009] The Fe-Cr-Al alloy powder of the present invention contains, by mass%, Cr: 20% to 30%, Al: 2.0% to 7.0%, Ti: more than 0.15% to 0.75%, and C: 0.001% to 0.010%, and the ratio of Ti to C (Ti / C) satisfies 15≦(Ti / C)≦300, with the remainder consisting of Fe and inevitable impurities, and is characterized in that it is used for additive manufacturing. 76.6 ≦(Ti / C)≦ 200 It is preferable that the following is satisfied.

[0010] The Fe-Cr-Al alloy member of the present invention is an additively manufactured body containing, by mass%, Cr: 20% to 30%, Al: 2.0% to 7.0%, Ti: more than 0.15% to 0.75%, and C: 0.001% to 0.02%, and the ratio of Ti to C (Ti / C) satisfies 15≦(Ti / C)≦300, with the remainder being Fe and unavoidable impurities, characterized in that the carbides precipitated at the grain boundaries of crystal grains are Ti carbides with an average equivalent circle diameter of 100 nm or less. It is also preferable that the ratio of the amount of Ti to the amount of C (Ti / C) satisfies 76.6≦(Ti / C)≦200.

[0011] The method for producing an Fe-Cr-Al alloy member of the present invention comprises the steps of: applying a heat source energy density J of 20 to 120 J / mm according to the following formula (1) to Fe-Cr-Al alloy powder containing, by mass%, 20% to 30% of Cr, 2.0% to 7.0% of Al, more than 0.15% to 0.75% of Ti, and 0.001% to 0.02% of C, with the ratio of Ti to C (Ti / C) satisfying 15≦(Ti / C)≦300, the balance being Fe and unavoidable impurities; 3 The method is characterized in that a laser beam or an electron beam is irradiated onto the material so that the material melts and solidifies, thereby performing layered manufacturing. It is also preferable that the ratio of the amount of Ti to the amount of C (Ti / C) satisfies 76.6≦(Ti / C)≦200. J=P / (v×a×t) (1) J: Heat source energy density (J / mm 3 ), P: laser beam or electron beam power (W), v: scanning speed (mm / s), a: scanning pitch (mm), t: layer thickness (mm) [Effects of the Invention]

[0012] According to the present invention, it is possible to provide an Fe-Cr-Al alloy powder for additive manufacturing that has improved high-temperature oxidation resistance, excellent ductility, and can suppress cracking, an Fe-Cr-Al alloy part made from an additive manufacturing body using this alloy powder, and a method for manufacturing this Fe-Cr-Al alloy part. [Brief explanation of the drawings]

[0013] [Figure 1] 1 is an inverse pole figure map (IPF) image obtained by electron back-scattering diffraction (EBSD) for an additive manufacturing body of an example. [Figure 2] (a) is an image of the additive manufacturing body of the example observed with a scanning transmission electron microscope (STEM), (b) is an enlarged image of the black-framed area in (a), and (c) and (d) are composition maps of Cr and Ti observed with energy-dispersive X-ray spectroscopy (EDS) for (b). [Figure 3] 1 shows an IPF image obtained by EBSD for an additive manufacturing body of a comparative example. [Figure 4] (a) is an image of the additive manufacturing body of the comparative example observed by STEM, (b) is an enlarged image of the black framed area in (a), and (c) and (d) are composition maps of Cr and Ti observed by EDS for (b). [Figure 5] 1A and 1B show examples of cracks occurring in the additive manufacturing body of the comparative example, where (a) is an example of solid-phase cracking originating from the edge, and (b) is an example of cutting cracking originating from the cut surface. [Figure 6] FIG. 10 is a diagram showing the results of a high-temperature oxidation resistance test of layered manufactured bodies according to an example of the present invention and a comparative example. [Figure 7]FIG. 1 is a diagram showing an example of an additive manufacturing apparatus and an additive manufacturing method using a powder bed fusion (PBF) method. [Figure 8] 1A and 1B are diagrams showing examples of support members, including (a) a solid support, (b) a grid support, and (c) a cone support. DETAILED DESCRIPTION OF THE INVENTION

[0014] The Fe-Cr-Al alloy powder for additive manufacturing of the present invention, the Fe-Cr-Al alloy member made of an additive manufactured body, and the method for manufacturing the Fe-Cr-Al alloy member will be described below with reference to the drawings. First, the composition and particle size distribution of the alloy powder will be described, followed by an explanation of the additive manufacturing method, the additive manufactured body, and its structure. In this specification, % indicating the content of metal elements means mass %. Furthermore, a numerical range expressed using "to" means that the numerical values ​​written before and after "to" are included as the lower and upper limits. The upper and lower limits can be combined arbitrarily.

[0015] [Alloy powder composition] The Fe-Cr-Al alloy powder for additive manufacturing according to this embodiment has improved high-temperature oxidation resistance and ductility, resulting in excellent additive manufacturing properties. Specifically, the Fe-Cr-Al alloy powder of the present invention contains, by mass, 20% to 30% Cr, 2.0% to 7.0% Al, more than 0.15% to 0.75% Ti, and 0.001% to 0.02% C, satisfying the relationship 15≦(Ti / C)≦300, with the remainder consisting of Fe and unavoidable impurities. The above alloy composition is an Fe-based alloy with the highest Fe content, followed by Cr and Al, by mass. The main constituent elements are Cr, Al, and Fe, with Ti and C in appropriate and balanced amounts. Furthermore, unavoidable impurities such as Si, Mn, Ni, Zr, B, S, P, O, and N are permitted, but the Si and Mn contents are particularly limited. The chemical components and contents of the Fe-Cr-Al alloy powder are described in detail below. Note that the chemical components and contents of the layered manufactured body are the same as those of the alloy powder, so a detailed description is omitted.

[0016] (Cr: 20% to 30%) Cr is an effective element for improving high-temperature oxidation resistance by stabilizing the alumina coating formed on the surface of the additive manufacturing product. The lower limit of the Cr content is 20% to ensure sufficient high-temperature oxidation resistance. On the other hand, if the Cr content is too high, the formation of the alumina coating is inhibited, so the upper limit of the Cr content is 30%. Furthermore, Cr bonds with C (carbon) to form Cr carbides such as Cr7C3, which increase hardness. When these carbides segregate to crystal grains or grain boundaries, they can undesirably cause intergranular cracking. Therefore, the Cr content is set to 20% to 30%. From the viewpoint of improving high-temperature oxidation resistance, the lower limit of the Cr content is preferably 23%, more preferably 24%. Furthermore, from the viewpoint of stabilizing the alumina coating, reducing elongation, and preventing intergranular cracking, the upper limit of the Cr content is preferably 27%, more preferably 26%.

[0017] (Al: 2.0% to 7.0%) Al forms an alumina film on the surface of the additive manufacturing body, improving high-temperature oxidation resistance and reducing the occurrence of oxide scale. To ensure sufficient high-temperature oxidation resistance, the lower limit of the Al content is 2.0%. On the other hand, if the Al content is too high, the high thermal stress caused by the rapid cooling rate in the melting and solidification process can easily cause solid-state cracking due to the ductile-brittle transition (DBT), so the upper limit of the Al content is 7.0%. Therefore, the Al content is set to 2.0% to 7.0%. From the viewpoint of forming a sound alumina film, the lower limit of the Al content is preferably 3.0%, more preferably 4.0%. Furthermore, from the viewpoint of preventing solid-state cracking due to DBT, the upper limit of the Al content is preferably 6.0%, more preferably 5.5%.

[0018] (Ti: over 0.15% and up to 0.75%) Ti bonds with carbon (C) to form Ti carbides such as TiC, which can suppress the precipitation of Cr carbides. This reduces precipitation hardening caused by Cr carbides and improves the high-temperature oxidation resistance and ductility of the additive manufacturing body. To suppress the precipitation of Cr carbides, the lower limit of the Ti content is set to more than 0.15%. On the other hand, if the Ti content is too high, high-temperature oxidation resistance decreases, so the upper limit of the Ti content is set to 0.75%. Therefore, the Ti content is set to more than 0.15% but not more than 0.75%. The lower limit of the Ti content is preferably 0.25%, more preferably 0.3%, from the viewpoint of generating Ti carbides and suppressing the segregation of Cr carbides. Furthermore, the upper limit of the Ti content is preferably 0.5%, more preferably 0.4%, from the viewpoint of decreasing high-temperature oxidation resistance.

[0019] (C: 0.001% to 0.02%) C generally improves strength and hardness by precipitating Cr carbides. However, if added in excessive amounts, Cr carbides are formed near grain boundaries, resulting in a deterioration in corrosion resistance. Furthermore, the ductile-brittle transition temperature (DBTT) increases, increasing the risk of solid-phase cracking. However, C can also be used as a deoxidizer in the melting process, making it difficult to eliminate C from the alloy. For these reasons, in this embodiment, the C content is particularly limited, with a lower limit of 0.001%. High C content increases hardness, resulting in reduced crack resistance, and the formation of Cr carbides at grain boundaries increases the susceptibility to intergranular cracking. Therefore, the upper limit of the C content is set to 0.02%. Therefore, the C content is set to 0.001% to 0.02%. The lower limit of the C content is preferably 0.0015%, more preferably 0.002%, from the viewpoints of appropriate Ti carbide formation, suppression of Cr carbide formation, and ease of manufacturing. The upper limit of the C content is preferably 0.010%, more preferably 0.005%, from the viewpoint of ensuring crack resistance.

[0020] (15≦(Ti / C)≦300) Ti easily bonds with carbon (C) and forms Ti carbides, thereby preventing the precipitation of Cr carbides. To achieve this effect, the ratio of Ti to C is important in this embodiment. That is, to obtain the effect of preventing the precipitation of Cr carbides while precipitating the minimum amount of Ti carbides, the ratio of Ti to C (Ti / C) must be 15 or greater. Therefore, Ti / C is set to 15 or greater. However, if the Ti / C ratio is too large, too much Ti dissolves in the ferrite phase (matrix), increasing hardness and making the material more susceptible to cracking. Therefore, the upper limit is set to 300. Therefore, Ti / C is set to 15 to 300. The lower limit of Ti / C is preferably 40 or greater, more preferably 70 or greater, and even more preferably 100 or greater, from the viewpoints of forming Ti carbides by bonding with C dissolved in the ferrite phase within the crystal grains, thereby reducing the hardness of the ferrite phase, and preventing the precipitation of Cr carbides. The upper limit of Ti / C is preferably 250 or less, more preferably 200 or less, from the viewpoint of hardness and cracking.

[0021] In this embodiment, to eliminate solid-phase cracking in the Fe-Cr-Al alloy, the C content in the alloy composition is limited to an extremely low level, and the Ti / C blending balance is set to an appropriate value. This alloy composition allows Cr and Ti to be dispersed and precipitated in the ferrite phase (parent phase), while minimizing the formation of carbides. Even if carbides are formed, Ti carbides are preferentially precipitated, preventing the precipitation and segregation of Cr carbides. Even if carbides precipitate within grain boundaries, only a small amount of Ti carbides, not Cr carbides, appears. These effects are thought to suppress cracking (including solid-phase cracking, intergranular cracking, etc.) in the additively manufactured body, improving its high-temperature oxidation resistance and ductility.

[0022] (Fe: remainder) In the alloy of this embodiment, Fe constitutes the balance of the alloy composition defined above.

[0023] (unavoidable impurities) The inevitable impurities include Si, Mn, Ni, Zr, B, S, P, O, and N. The content of these inevitable impurities is preferably low, and is preferably 0%. In this embodiment, the content of Si, Mn, Ni, Zr, B, S, P, O, and N is preferably limited as follows. (Si:0.3% or less) If the Si content is high, the material becomes harder and more susceptible to cracking during additive manufacturing. For this reason, the Si content is limited to 0.3% or less. It is preferably 0.1% or less, more preferably 0.01% or less, and even more preferably 0.005% or less. (Mn:0.3% or less) Since an increased Mn content reduces high-temperature oxidation resistance, the Mn content is limited to 0.3% or less, preferably 0.1% or less, more preferably 0.01% or less, and even more preferably 0.005% or less.

[0024] (Ni: 0.3% or less) If the Ni content is too high, cracks are more likely to occur during solidification during additive manufacturing, so the Ni content is limited to 0.3% or less, preferably 0.1% or less, more preferably 0.01% or less, and even more preferably 0.005% or less. (Zr:0.1% or less) If the Zr content is too high, it will combine with carbon (C) to form Zr carbides, which will segregate at grain boundaries and make cracks more likely to occur from the grain boundaries. Therefore, the Zr content is limited to 0.1% or less. It is preferably 0.05% or less, more preferably 0.02% or less, and even more preferably 0.01% or less. (B: 0.01% or less) If the B content is too high, TiB2, Fe2B, and Cr2B will form near the grain boundaries, making cracks more likely to occur. For this reason, the B content is limited to 0.01%, preferably 0.005% or less, more preferably 0.003% or less, and even more preferably 0.002% or less.

[0025] (S, P: 0.01% or less) Since S and P segregate at grain boundaries and cause hot cracking, the S and P contents are limited to 0.01% or less, preferably 0.005% or less, and more preferably 0.002% or less. (O: 0.04% or less) O (oxygen) bonds with Al, Ti, etc. to form oxides inside or on the surface of the alloy powder. While these oxides have the effect of improving strength, they can act as fracture initiation points and reduce toughness. For this reason, the O content is limited to 0.04% or less, preferably 0.03% or less, and more preferably 0.02% or less. (N: 0.001% or less) N (nitrogen) is unavoidably mixed in during the general manufacturing process of powders for additive manufacturing, such as gas atomization, and a content of about 0.001% is acceptable.

[0026] The component composition of the Fe-Cr-Al alloy powder of this embodiment can be determined by the following measurement method. As described in the Examples below, the classified powder for additive manufacturing is dissolved in an appropriate aqueous solution, and the aqueous solution is subjected to inductively coupled plasma (ICP) analysis to measure the content of specific components. The content of C, N, and O can be determined by gas analysis using a combustion method.

[0027] [Method of manufacturing alloy powder] The Fe-Cr-Al alloy powder of this embodiment can be produced by atomization methods such as gas atomization, water atomization, and jet atomization. Gas atomization, which is particularly suitable for obtaining spherical powder, is preferred. In the atomization method, raw material powders formulated to a specified composition are melted in a crucible, and the molten metal is then discharged from the bottom of the crucible while high-pressure gas is sprayed onto the discharged molten metal. The kinetic energy of the high-pressure sprayed medium causes the molten metal to fly as droplets, which are then solidified to produce spherical Fe-Cr-Al alloy powder. The crucible used here is preferably made of Al2O3 to prevent the inclusion of Zr impurities.

[0028] (particle size distribution) The Fe-Cr-Al alloy powder of this embodiment is an alloy powder for additive manufacturing suitable for additive manufacturing. The particle size distribution is determined by laser diffraction using a laser diffraction particle size distribution analyzer. For example, in the alloy powder suitable for SLM described below, in the cumulative distribution curve showing the relationship between particle size and cumulative volume from the smallest particle size side, the particle diameter D10 corresponding to a cumulative value of 10 volume % is 7 μm to 35 μm, preferably 10 μm to 25 μm. The particle diameter D50 corresponding to a cumulative value of 50 volume % is 10 μm to 60 μm, preferably 20 μm to 40 μm. Furthermore, the particle diameter D90 corresponding to a cumulative value of 90 volume % is 20 μm to 100 μm, preferably 40 μm to 80 μm. Particles that are too small can cause the powder bed to wear out or become uneven, resulting in poor coatability. On the other hand, particles that are too large can cause concerns about unmelted particles due to insufficient laser power, which can affect defects and surface roughness. By having the above particle size distribution, or simply by satisfying the D50 value, a powder bed can be repeatedly laid down to a flat and uniform thickness.

[0029] Furthermore, when the Fe—Cr—Al alloy powder of this embodiment is used in the LMD described below, the particle size distribution suitable for the LMD is as follows: the particle diameter D10 corresponding to a cumulative 10% by volume is 15 μm to 100 μm, preferably 20 μm to 80 μm; the particle diameter D50 corresponding to a cumulative 50% by volume is 30 μm to 250 μm, preferably 60 μm to 200 μm; and the particle diameter D90 corresponding to a cumulative 90% by volume is 50 μm to 500 μm, preferably 100 μm to 300 μm. In the case of LMD, a small particle size can cause uneven powder flow during transportation, making stable powder supply to the molten pool difficult. Furthermore, uneven powder flow can also cause uneven powder clogging within the discharge head. On the other hand, a large particle size can lead to concerns about insufficient laser power, resulting in residual melt, which may result in defects and surface roughness.

[0030] [Layered manufacturing method] The Fe-Cr-Al alloy powder of this embodiment is suitable for additive manufacturing. Additive manufacturing methods for metal materials are generally divided into powder bed fusion (PBF) and directed energy deposition (DED). The PBF method involves spreading alloy powder on a substrate to form a powder bed, then irradiating the alloy powder spread in the target area with a beam to melt and solidify the alloy powder to create a shape. Meanwhile, the DED method, also known as laser metal deposition (LMD), involves supplying alloy powder onto a substrate or an already-formed object and irradiating it with a heat source beam, melting and solidifying the alloy powder supplied to the formation area to create a shape. The additively manufactured object of this embodiment can be produced by either method.

[0031] Both the PBF and DED methods are divided into methods that use a laser beam as a heat source and methods that use an electron beam. For example, methods that use a laser beam can be broadly divided into selective laser melting (SLM) and selective laser sintering (SLS). In SLM and SLS, which use a laser beam as the heat source, additive manufacturing is carried out in an inert atmosphere such as nitrogen gas. On the other hand, methods that use an electron beam are called selective electron beam melting (SEBM or EBM). In methods that use an electron beam as the heat source, additive manufacturing is carried out in a high vacuum.

[0032] Figure 7 shows an example of the schematic configuration of the laser additive manufacturing method, which is one of the PBF methods and uses a laser beam as a heat source for additive manufacturing. As shown in Figure 7, 10 is the alloy powder used as the raw material, 20 is the powder supply stage, 30 is the recoater, 40 is the laser oscillator, 50 is the laser light, 60 is the galvanometer scanner, 70 is the additively manufactured object, and 80 is the manufacturing stage. In the additive manufacturing process, the powder supply stage 20 is raised a predetermined distance, the modeling stage 80 is lowered a predetermined distance, and the recoater 30 moves in the X direction to supply alloy powder 10 onto the modeling stage 80, forming a powder bed. A laser beam 50 from a laser oscillator 40 is controlled by a galvanometer scanner 60 to irradiate a predetermined position on the powder bed, selectively melting and solidifying the powder to form a solidified layer. By repeating this process, a three-dimensional additively manufactured object 70 is manufactured.

[0033] Next, we will explain how the manufacturing method of the Fe-Cr-Al alloy component of this embodiment has led to a reduction in solid-phase cracking due to improved molding conditions, and how improvements in the provision of support during molding have led to a reduction in cutting cracks when cutting out the molded body.

[0034] (Modeling conditions) Generally, increasing the heat source energy density of a laser beam or electron beam increases the molten pool temperature, resulting in greater thermal stress and residual strain. Conversely, decreasing the heat source energy density makes it impossible to completely melt the alloy powder, resulting in a low density additively manufactured object and resulting in defects. Fe-Cr-Al alloys have a low-temperature ductile-brittle transition (DBT), and the high thermal stress caused by rapid cooling rates makes them more susceptible to solid-phase cracking (SSC) due to DBT. Furthermore, as crystal grains grow larger, the DBT temperature (DBTT) increases, increasing the risk of SSC. Here, the heat source energy density J is expressed by the following equation (1): J=P / (v×a×t) (1) J: Heat source energy density (J / mm 3 ), P: laser beam or electron beam power (W), v: scanning speed (mm / s), a: scanning pitch (mm), t: layer thickness (mm)

[0035] The heat source energy density J (hereinafter referred to as energy density) can be controlled by changing the laser or electron beam output P, ​​scanning speed v, scanning pitch a, and layer thickness t. The controlled modeling parameters are closely related to the laser or electron beam output P, ​​scanning speed v, and scanning pitch a, and it is preferable to control the energy density by adjusting these three parameters. In this embodiment, the relationship between the occurrence of cracks when the laser output P and scanning speed v are changed was investigated. Note that, hereinafter, the term "cracks" includes solid-phase cracks and cutting cracks.

[0036] As a result, in this embodiment, it is possible to select molding conditions that match the composition of the Fe-Cr-Al alloy powder, and it is therefore possible to reduce the DBT temperature. That is, with the Fe-Cr-Al alloy powder of this embodiment, it is possible to select a low energy density within an appropriate range, and the energy density is 20 to 120 J / mm 3 The preferred energy density is 40 to 100 J / mm 3 and more preferably 50 to 90 J / mm 3 is. The ability to manufacture at such a low energy density reduces the thermal load on the molten pool and enables the formation of a product with relatively small, uniform, columnar crystal grains. This reduces thermal stress and allows the DBT temperature to be lowered to near room temperature or even lower. It is believed that lowering the DBT temperature reduces the risk of solid-phase cracking.

[0037] (Support Grant) In additive manufacturing, a model is built on a base plate, and residual stress and distortion due to thermal stresses generated during the building process accumulate between the two. After the building process is complete, when the built object is cut from the base plate, the stress is suddenly relieved at the cut surface of the built object, which can lead to cutting cracks that differ from solid-state cracks (see Figure 5(b)). Therefore, to reduce the risk of cutting cracks, it is effective to insert support members between the base plate and the built object. The addition of support members can reduce the accumulation of residual stress and distortion during the building process and prevent cutting cracks. There are no particular limitations on the shape or type of support; solid supports, cone supports, grid supports, and other types can be used. The preferred range for the height of the support members is 0.5 mm to 5 mm. Support heights less than 0.5 mm are less effective at reducing residual stress and distortion. On the other hand, support heights greater than 5 mm can increase the building time and increase the cost of building. By using the support members described above, it is possible to reduce the thermal stress and residual strain caused by rapid cooling even if the input energy from the heat source is large, and it is possible to reduce the risk of cutting cracks when separating the additive manufacturing body. Note that providing supports is optional and is not a required step.

[0038] [Organization] Compared to additive manufacturing products made from commercially available alloy powder compositions, the additive manufacturing products according to this embodiment exhibit reduced solid-phase cracking and improved high-temperature oxidation resistance and ductility, as shown in the examples below. Additive manufacturing products with improved properties have relatively small and uniform crystal grains. Furthermore, it was found that Cr and Ti carbides do not precipitate in the ferrite phase of the matrix, and when carbides do precipitate at grain boundaries, only small amounts of tiny titanium carbides precipitate. Details are provided below.

[0039] The mechanism by which solid-phase cracking of the additive manufacturing body is suppressed will be explained. (Melted solidified structure) First, in this embodiment, the carbon (C) content is limited to an extremely low level, and the Ti / C blend balance is set to an appropriate value. Using this alloy powder enables additive manufacturing under building conditions with relatively low energy density. The building process involves extremely high melting and solidification rates, resulting in repeated stacking of rapidly solidified layers, resulting in the formation of equiaxially stacked columnar structures. Figure 1 shows an inverse pole orientation mapping (IPF) image obtained by electron backscatter diffraction (EBSD) of an additively manufactured object according to this embodiment. Numerous columnar crystal grains elongated in the stacking direction are formed in the structure of the manufactured object. Each columnar crystal is relatively small and approximately circular or scale-shaped, with an average equivalent circle diameter of 57 μm. The average diameter is preferably 100 μm or less, preferably 80 μm or less, and more preferably 60 μm or less. It is believed that this fine, approximately circular, uniform crystal structure reduces DBTT and suppresses the occurrence of solid-state cracking. The average grain size is the area-weighted average grain size per 100 crystals.

[0040] (crystal structure) Next, the microstructure was observed using a scanning transmission electron microscope (STEM). In Fe-Cr-Al alloys, C in the alloy composition easily combines with Cr and Ti to form the respective carbides. For example, in the comparative example shown in Figure 4(b) described below, numerous Cr and Ti carbides are dispersed and precipitated along the grain boundaries. On the other hand, in the example shown in Figure 2(b), only a small amount of tiny Ti carbides precipitated at the grain boundaries, and no Cr carbides were present. This is because the C content in the alloy composition was limited to an extremely small amount, preventing carbide precipitation in the ferrite phase (matrix). Furthermore, by setting the Ti / C blending balance to an appropriate value, Ti, which has a higher affinity than Cr, suppresses the formation of Cr carbides, thereby preventing carbide precipitation at the grain boundaries. Even if carbides precipitate at the grain boundaries, they remain as tiny, trace amounts of Ti carbides. The size of these Ti carbides is such that the average equivalent circle diameter is 100 nm or less, preferably 80 nm or less, and more preferably 50 nm or less. Furthermore, the number density of Ti carbides is 1 μm 2The number of particles per molecule is 50 or less, preferably 30 or less, and more preferably 10 or less. As described above, the matrix is ​​free of segregation of TiC and CrC carbides, with the composition uniformly dispersed, and carbides precipitating at grain boundaries are limited, which is thought to have suppressed solid-phase cracking and improved high-temperature oxidation resistance. Furthermore, because precipitation of Cr carbides at grain boundaries could be suppressed, the relative decrease in hardness of the AM body is thought to have led to improved ductility.

[0041] [High temperature oxidation resistance] The additive manufacturing product of this embodiment has excellent high-temperature oxidation resistance. In a high-temperature oxidation test performed in an atmospheric furnace at 900°C for over 0 hours up to 1000 hours, the amount of oxidation per unit area, expressed as [mass loss before and after oxidation test] / [surface area before oxidation test], was 0.25 mg / cm2 even after 1000 hours, as shown in Figure 6 described below. 2 The following good high-temperature oxidation resistance was demonstrated: The unit area oxidation amount was 0.2 mg / cm over 500 to 1000 hours. 2 Furthermore, the unit area oxidation amount can be maintained at 0.12 mg / cm for 250 hours. 2 The following is the result.

[0042] [Mechanical properties] Examples of indicators of mechanical properties include yield strength, tensile strength, elongation, reduction of area, and Vickers hardness (HV). Note that yield strength and tensile strength are indicators for determining strength, while elongation and reduction of area are indicators for determining ductility. Vickers hardness (HV) is, as the name suggests, an indicator for determining hardness. The layered manufactured product according to this embodiment can have a proof stress of 610 MPa or more, preferably 650 MPa or more, and more preferably 730 MPa or more. The tensile strength can be 710 MPa or more, preferably 800 MPa or more, and more preferably 810 MPa or more. The layered object according to this embodiment can exhibit an elongation of 30% or more, preferably 35% or more, and more preferably 38% or more, and an area reduction of 60% or more, preferably 70% or more, and more preferably 80% or more. Furthermore, the Vickers hardness (HV10) of the layered manufactured body according to this embodiment is 210 HV or more, but the hardness can be suppressed from the viewpoint of DBT. Furthermore, if the hardness is too high, cutting workability is poor and cracks due to processing are likely to occur, so the upper limit of the hardness is about 300 HV. The preferred hardness range is 230 to 280 HV, more preferably 240 to 270 HV.

[0043] [Application] The additive manufacturing object of this embodiment has excellent high-temperature oxidation resistance, particularly at temperatures above 800°C, and does not develop an oxide film on its surface even when used in high-temperature environments. Therefore, it can be used for components and parts installed inside oxidation furnaces for semiconductor manufacturing and baking furnaces for electronic component manufacturing. It can also provide additive manufacturing objects with complex shapes that have excellent high-temperature oxidation resistance in a wide range of fields, including parts for the aerospace and automotive industries, chemical plants, pharmaceutical manufacturing facilities, and energy fields such as oil and gas. It can also be used as heating wires and electrical resistors for home appliances. [Example]

[0044] The present invention will be specifically described below based on examples. (Alloy powder for additive manufacturing) Alloy powders a, b, and c for additive manufacturing were prepared, each having the chemical composition shown in Table 1. Alloy powders a and b are examples, and alloy powder c is a comparative example. These alloy powders (hereinafter referred to as "powder") were prepared by preparing raw materials for melting, melting them in a conventional high-frequency vacuum melting furnace to produce a master alloy, and then gas atomizing them in an argon atmosphere. Powders with particle sizes of 10 to 60 μm were classified from the atomized powder and used for additive manufacturing. The D10, D50, and D90 of the classified powders are also shown in Table 1. Comparing the component compositions of powders a and b (Examples) and powder c (Comparative Example) shown in Table 1, powders a and b have a significantly reduced C content compared to powder c, and the ratio of Ti to C (Ti / C) of powders a and b is set to 15 or more compared to powder c. At the same time, the contents of impurities Si and Mn are significantly reduced in powders a and b.

[0045] [Table 1]

[0046] (Production of alloy components) Next, a layered object was fabricated by additive manufacturing using the powder. Using an additive manufacturing device (EOS M290: manufactured by EOS Corporation), additive manufacturing was performed using the SLM method under the following conditions to produce additively manufactured objects (10 mm x 10 mm x 10 mm). (Hereinafter, the additively manufactured objects will be referred to as "models.") The models manufactured using powders a, b, and c will be referred to as models A, B, and C, respectively. The modeling conditions in this example were a scanning pitch of 0.11 mm, a layer thickness of 0.04 mm, a laser output of 200 W to 350 W, and a scanning speed of 700 mm / sec to 1500 mm / sec. As a result, the energy density J was 30 to 114 J / mm 3 It was possible to produce various shaped bodies within the range of

[0047] The molded bodies A, B, and C used in the crack observation, structure observation, high-temperature oxidation resistance, and mechanical property evaluation described below were laser-processed at a laser output of 325 W, a scanning speed of 800 mm / s, a scanning pitch of 0.11 mm, a layer thickness of 0.04 mm, and an energy density of 92 J / mm 3 It was made with.

[0048] (Support Grant) To reduce the residual stress generated in the base plate and the additive manufacturing object, a block material (10 mm x 10 mm x 10 mm) was fabricated by placing support members with a height of 1 mm between the base plate and the additive manufacturing object. The support members used were (a) solid, (b) grid, (c) cone, and square (not shown) supports, as shown in Figure 8. After manufacturing, the molded object was separated from the support members using a wire electric discharge machine.

[0049] The following crack observation, structure observation, high-temperature oxidation resistance, and mechanical properties were evaluated. (Crack evaluation) A cross section of each molded body (10 mm x 10 mm x 10 mm) was cut and embedded in resin using a hot resin embedding vacuum device (CitoPress-30: Struers). The molded bodies embedded in resin were polished with waterproof emery paper to a grit size of #1500, then polished with diamond paste in the order of grain sizes from 1 μm to 0.3 μm, resulting in a mirror finish, yielding test specimens for measuring cracks and defect rates. Cracks were visually confirmed and defect rates were measured using a microscope (VHX-6000: Keyence).

[0050] Table 2-1 shows the crack occurrence state when powders a, b, and c are used, the scanning pitch is fixed at 0.11 mm, the layer thickness is fixed at 0.04 mm, the laser output is 200 to 350 W, and the scanning speed is changed to 700 to 1500 mm / s. As shown in Table 2-2, the energy density under the above molding conditions is 30 to 114 J / mm 3 It varies within a range of.

[0051] [Table 2-1]

[0052] [Table 2-2]

[0053] As can be seen from the results in Table 2-1, Body A, which was fabricated using Powder a, did not experience any cracks under any of the conditions. Body B, which was fabricated using Powder b, did not experience any solid-state cracks, but under some conditions with higher energy densities, cutting cracks developed from the cut surface extending inward, as shown in Figure 5(b). This resulted in a 7% occurrence rate. Body C, which was fabricated using Powder c, experienced solid-state cracks that extended horizontally inward from the edge, as shown in Figure 5(a). While Body C did not experience any solid-state cracks under low energy densities with low power and high scanning speed, these bodies had internal cavities and were of low quality. Although not shown in Table 2-1, some cutting cracks also occurred. From the above results, it was found that cracking of the molded body can be suppressed by reducing the C content, optimizing the Ti / C ratio, and reducing the Si and Mn contents within the alloy powder composition. Furthermore, since molding using this alloy powder can be performed under relatively low energy density conditions, it can be said that the risk of cracking such as solid-state cracking is reduced.

[0054] For model B, the laser output was 350 W and the scanning speed was 700 mm / s (energy density was 99.4 J / mm 3 ) and 800mm / s (energy density is 113.6J / mm 3 ), cutting cracks occurred as shown in Figure 5(b) (indicated as "Partially present" in Table 2). When various supports with a height of 1 mm as described above were added between the additive manufacturing body and the base plate, the cutting cracks in Model B disappeared. Table 3 shows the relationship between the addition of supports and the occurrence of cutting cracks.

[0055] [Table 3]

[0056] It was shown that regardless of the type of support, adding appropriate support members prevented cracks from occurring in the AM object. This is thought to be because adding support members alleviates the thermal stress and strain that occurs during the melting process in AM, preventing cracks from occurring.

[0057] (structural observation) The crystal orientation of the microstructures of the molded bodies A and C was evaluated using electron backscatter diffraction (EBSD) (JSM-7900F: manufactured by JEOL). The results for molded body A are shown in Figure 1, and the results for molded body C are shown in Figure 3. The shaped body A (Example) in Figure 1 was composed of fine, roughly circular or scale-like columnar crystals, as shown in the partial outline. The columnar crystal structure was generated by the rapid solidification of a tiny molten pool, and the columnar crystals grew overall in the stacking direction due to the heat flux of the shaped body. The columnar crystal structure of shaped body A had a width of 2 to 150 μm and a stacking direction of 4.5 to 316 μm, with an average equivalent circle diameter of 57 μm. The average grain size is the area-weighted average grain size per 100 crystals.

[0058] On the other hand, the shaped body C (comparative example) in Figure 3 was composed of relatively coarse and elongated columnar crystals, as shown in the outline of a portion of the shaped body. The columnar crystal structure grew in the stacking direction, similar to that of shaped body A, but was clearly larger. The columnar crystal structure of shaped body C had a width of 2 to 182 μm, a length of 4.5 to 501 μm in the stacking direction, and many lengths exceeding 500 μm were observed. Although the formed bodies A and C were manufactured under the same manufacturing conditions, differences in the columnar crystal structure that occurred during melting and solidification were observed, which is thought to be due to differences in composition. It can be said that the composition of the example can suppress the growth of the structure due to melting and solidification, even at low energy densities, and can form a columnar crystal structure of appropriate size.

[0059] Next, samples were obtained from the same molded bodies A and C at the observation locations shown in Figures 1 and 3, and thin sections were cut using focused ion beam processing. The crystal structure of the slices was observed using a scanning transmission electron microscope (STEM) (JEM-ARM200F: JEOL). The results are shown in Figure 2(a) for molded body A and Figure 4(a) for molded body C. Furthermore, enlarged images of the grain boundaries in the black framed areas are shown in Figures 2(b) and 4(b), respectively. Next, the enlarged images of the grain boundaries in Figure 2(b) and Figure 4(b) were subjected to elemental analysis (EDS) using a STEM to obtain elemental mapping of Cr and Ti. The results are shown in Figures 2(c) and 2(d) for Body A and Figure 4(c) and 4(d) for Body C.

[0060] First, when observing the shaped body C in Figure 4, a large number of black and white particles can be seen dispersed and precipitated along the grain boundaries in Figure 4(b). Here, the black particles are Ti carbide (TiC) and the white particles are Cr carbide (CrC). The average circle equivalent diameter per 10 TiC particles was 35 nm, with some particles exceeding 50 nm. The number density of TiC particles was 1 μm 2 On the other hand, for CrC, the average size was about 50 nm and the number density was 1 μm 2 There were more than two winners. Furthermore, as shown in the elemental mapping in Figure 4(c)(d), it was confirmed that Cr and Ti precipitated in large quantities at the grain boundaries, and furthermore that TiC was also segregated in the ferrite phase (parent phase), and it is believed that CrC was also segregated to a small extent.

[0061] Next, when observing the shaped body A in Figure 2, only a few black particles, i.e., TiC, were observed to have precipitated at the grain boundaries in Figure 2(b), but no CrC was observed. Furthermore, the average diameter of the circle equivalent diameter of TiC was about 35 nm, and none exceeded 40 nm. Also, the number density was 1 μm 2 It can be said that there are less than three per person. Furthermore, the elemental mapping in Figures 2(c) and (d) confirmed that Cr and Ti were not found at the grain boundaries. Furthermore, it was confirmed that Cr and Ti were evenly dispersed in the ferrite phase, with no segregation. The carbide size and number density of bodies A, B, and C are shown in Table 4.

[0062] [Table 4]

[0063] Based on the above, in the case of the sintered body C, which has a high C content and an optimum Ti / C balance, both Ti and Cr carbides precipitate and tend to segregate. At the same time, as shown in Figure 4(d), Ti carbides also segregate in large quantities in the ferrite phase other than at the grain boundaries. In contrast, in the case of the sintered body A, which has a minimum C content and an optimum Ti / C balance, the contents of Ti and C, which readily bond with C, are limited and balanced, resulting in the preferential precipitation of a minimum amount of TiC and the suppression of the formation of Cr carbides. This is thought to result in the sintered body C being evenly dispersed and precipitated in the ferrite phase, with a small amount of TiC precipitating at the grain boundaries. The same can be said for the sintered body B. As described above, the matrix is ​​free of segregation of TiC and CrC carbides, with the composition uniformly dispersed, and carbides precipitating at grain boundaries are limited, which is thought to have suppressed solid-phase cracking and improved high-temperature oxidation resistance. Furthermore, because precipitation of Cr carbides at grain boundaries could be suppressed, the relative decrease in hardness of the additively manufactured body is reflected in elongation and reduction of area, leading to improved ductility.

[0064] (High temperature oxidation resistance evaluation) Plate samples were obtained for the shaped bodies A, B, and C, and their surfaces were polished to #1000 with waterproof emery paper. After that, they were degreased by applying ultrasonic vibrations to acetone and holding for 5 minutes to obtain oxidation-resistant test pieces. After measuring the dimensions and mass of each test piece, they were held at 900°C in a small atmospheric furnace, and the oxidation time was set to four stages: 250 hours, 500 hours, 750 hours, and 1000 hours. After that, they were cooled to room temperature, and the mass change of the test piece was measured at each stage. Then, the amount of oxidation per unit area (mg / cm) was calculated for each test piece using the following formula (2): 2 ) was calculated. The measurement results are shown in Figure 6. Amount of oxidation per unit area = [mass loss before and after oxidation test] / [surface area before oxidation test] (2)

[0065] According to the oxidation test shown in Figure 6, the unit area oxidation amount of the molded bodies A and B was 0.2 mg / cm even after 1000 hours. 2 The density of the comparative example, model C, was maintained at about 0.26 mg / cm2 The amount of oxidation per unit area is suppressed compared to the previous example. This result confirms that the high-temperature oxidation resistance is excellent.

[0066] (mechanical properties) Tensile test specimens (parallel diameter: 3 mm, gauge length: 7 mm) were prepared for the above-mentioned shaped bodies A, B, and C in accordance with the standard test (ASTM E8). A tensile test (INSTRON5982: manufactured by Instron) was conducted on these tensile test specimens at room temperature (22°C) to determine the tensile strength, 0.2% proof stress, elongation, and reduction of area. Furthermore, hardness was measured using a Vickers hardness tester (FM-110: manufactured by Future Tech) at the center of the test specimen and five locations 2–3 mm above, below, left, and right from the center, with a 10 kg load held for 30 seconds, and the average value was calculated. The reduction of area was expressed as a percentage by dividing the cross-sectional area of ​​the necked portion after fracture by the original cross-sectional area. The results are shown in Table 5.

[0067] [Table 5]

[0068] The hardness of the formed bodies A and B (Examples) was lower than that of the formed body C (Comparative Example). The lower hardness of the formed bodies A and B is thought to be due to the fact that only a small amount of TiC segregation was present at the grain boundaries, as shown in Figure 2, and the amount of Cr carbide precipitated was reduced. [Explanation of symbols]

[0069] 1: Ti carbide precipitation phase 2: Cr carbide precipitate phase 3: Edge crack 4: Cutting crack

Claims

1. In mass%, Cr: 20% or more and 30% or less, Al: 2.0% or more and 7.0% or less, Ti: more than 0.15% and not more than 0.75%; An Fe-Cr-Al alloy powder for additive manufacturing containing C: 0.001% or more and 0.010% or less, and having a ratio of Ti to C (Ti / C) satisfying 76.6≦(Ti / C)≦200, with the remainder being Fe and unavoidable impurities.

2. 2. The Fe-Cr-Al alloy powder according to claim 1, wherein the Si content of the unavoidable impurities is limited to 0.3 mass % or less, and the Mn content is limited to 0.3 mass % or less.

3. 3. The Fe-Cr-Al alloy powder according to claim 1, wherein the particle diameter D50, which is the particle diameter at 50 volume % of the integrated value measured by laser diffraction particle size distribution measurement, is 10 to 60 μm, and the Fe-Cr-Al alloy powder is used in a powder laser melting method.

4. The Fe-Cr-Al alloy powder according to claim 1 or 2, characterized in that the particle diameter D50, which is the particle diameter at 50 volume % integrated value as measured by laser diffraction particle size distribution measurement, is 30 to 250 μm, and the powder is used in directed energy deposition.

5. In mass%, Cr: 20% or more and 30% or less, Al: 2.0% or more and 7.0% or less, Ti: more than 0.15% and not more than 0.75%; An additive manufacturing body containing C: 0.001% or more and 0.02% or less, and having a ratio of the amount of Ti to the amount of C (Ti / C) satisfying 76.6≦(Ti / C)≦200, with the remainder being Fe and unavoidable impurities, An Fe-Cr-Al alloy member, characterized in that the carbides precipitated at the grain boundaries of crystal grains are Ti carbides having an average equivalent circle diameter of 100 nm or less.

6. The number density of the Ti carbides in cross-sectional structure observation is 1 μm ２ 6. The Fe-Cr-Al alloy member according to claim 5, wherein the number of defects is 50 or less per one.

7. The Fe-Cr-Al alloy member according to claim 5 or 6, characterized in that the molten solidified structure of the layered manufactured body is composed of columnar crystals, and the average equivalent circle diameter of the columnar crystals is 100 μm or less.

8. The Fe-Cr-Al alloy member according to claim 5 or 6, wherein the layered manufactured body has an elongation of 30% or more and a reduction in area of ​​60% or more.

9. The amount of oxidation per unit area, expressed as [mass loss before and after oxidation test] / [surface area before oxidation test], is 0.25 mg / cm ２ 7. The Fe—Cr—Al alloy member according to claim 5, wherein:

10. In mass%, Cr: 20% or more and 30% or less, Al: 2.0% or more and 7.0% or less, Ti: more than 0.15% and not more than 0.75%; An Fe-Cr-Al alloy powder containing C: 0.001% or more and 0.02% or less, a ratio of Ti to C (Ti / C) satisfying 76.6≦(Ti / C)≦200, and the balance being Fe and unavoidable impurities, The heat source energy density J according to the following formula (1) is 20 to 120 J / mm ３ and melting and solidifying the material to form an Fe—Cr—Al alloy member. J=P / (v×a×t)...(1) J: Heat source energy density (J / mm ３ ), P: laser beam or electron beam power (W), v: scanning speed (mm / s), a: scanning pitch (mm), t: layer thickness (mm)

11. The method for manufacturing an Fe-Cr-Al alloy member according to claim 10, characterized in that in the additive manufacturing, an additive manufactured body is formed on a base plate, and when the additive manufactured body is manufactured on the base plate, a support member is interposed between the base plate and the additive manufactured body.

Images