Precipitation-hardening austenitic iron-manganese-aluminum-carbon alloy welding filler wire

Abstract

To provide a weld filler wire for precipitation hardening type austenite iron-manganese-aluminum-carbon alloy fusion-welding.SOLUTION: A precipitation hardening type austenite iron-manganese-aluminum-carbon alloy fusion-welding contains: manganese of 23 to 34 wt.%; aluminum of 7.5 to 11.5 wt.%; carbon of 1.35 to 1.95 wt.%; and the remaining component comprising iron, Alternatively, a precipitation hardening type austenite iron-manganese-aluminum-carbon alloy fusion-welding contains: manganese of 23 to 34 wt.%; aluminum of 7.5 to 11.5 wt.%; carbon of 1.40 to 1.95 wt.%, and further, contains: any one of titanium of 0.1 to 2.5 wt.%; niobium of 0.1 to 3.0 wt.%; and vanadium of 0.1 to 2.5 wt.%; and the remaining component comprising iron.SELECTED DRAWING: Figure 2(a)

Description

[Technical Field]

[0001] The present invention relates to a new welding filler wire in the field of welding filler wire technology, particularly for joining parts manufactured from lightweight precipitation-hardening iron-manganese-aluminum-carbon alloys (Fe-Mn-Al-C alloys), and more specifically to joining using fusion welding utilizing gas tungsten arc welding (GTAW) and gas metal arc welding (GMAW). [Background technology]

[0002] Precipitation hardening (also known as age hardening) is a heat treatment technique used to increase the yield strength (YS) and hardness of ductile alloys. Examples of so-called ductile alloys include the highest-strength 7xxx series aluminum alloys (e.g., AA7075), precipitation-hardened stainless steels, and precipitation-hardened austenitic Fe-Mn-Al-C alloys. The precipitation hardening process involves three basic steps: solution heat treatment (SHT), quenching, and aging. The first step (solution heat treatment) involves heating the alloy to a solvus temperature and maintaining that temperature for a certain period of time until a uniform single-phase solid solution is formed. The second step (quenching) involves rapidly cooling the solid solution from the solution heat treatment temperature to room temperature to form a supersaturated solid solution. In the quenched state, the microstructure of the alloy is a single-phase supersaturated solid solution without precipitates. As a result, the alloy in the as-quenched state may have a good combination of ultimate tensile strength (UTS) and elongation (El), but the yield strength (YS) is always significantly lower. However, after heating the supersaturated solid solution to a temperature lower than the solid solution temperature and performing appropriate aging treatment at that temperature, high-density nano-sized precipitates are coherently (or semicoherently) formed and evenly distributed in the base.Next, these densely distributed, nano-order, coherent (or semi-coherent) second-phase precipitates act as effective obstacles to dislocation movement, significantly increasing the alloy's strength (especially its yield strength YS) and hardness, without significantly reducing its ductility. In other words, densely distributed nano-order precipitates coherently (or semi-coherently) deposited in the base are the most fundamental component for obtaining an ideal combination of mechanical strength, particularly a significantly improved yield strength YS, hardness, and elongation.

[0003] Recently, lightweight, precipitation-hardening, fully austenitic iron-manganese-aluminum-carbon alloys have been recognized as one of the most promising materials due to their exceptional combination of mechanical strength and ductility, making them widely applicable in various fields such as the automotive, armor, and mining industries. In particular, the addition of aluminum can produce another significant advantage: it can drastically reduce the density of conventional high-strength steel (approximately 1.3-1.5% reduction in density for every 1% by weight of aluminum added), which is generally expected to have a significant impact on both energy savings and reductions in carbon dioxide emissions.

[0004] According to various previous studies, for this type of alloy possessing an ideal combination of strength (especially yield strength YS) and ductility, the typical alloy microstructure basically consists of a complete austenitic iron base at room temperature and a large amount of nano-order iron-manganese-aluminum carbides ((Fe,Mn)3AlC carbides (κ-carbides)) precipitated on the austenitic iron base. These κ-carbides have an ordered face-centered cubic (FCC) L'12 crystal structure. Currently, there is ample evidence that manganese is a strong stabilizing element for austenitic iron, and that aluminum and carbon are the basic elements necessary for the formation of κ-carbides. The former (manganese) can give the alloy the face-centered cubic (FCC) crystal structure necessary for high ductility at room temperature, while the latter is the main strengthening precipitate in this type of alloy. It forms precipitates (κ-carbides). Therefore, in order to give the precipitation-hardening austenitic iron-manganese-aluminum-carbon alloy an excellent combination of strength and ductility, the manganese content in the precipitation-hardening austenitic iron-manganese-aluminum-carbon alloy is preferably greater than 18% by weight, so that the alloy can have a face-centered cubic perfect austenitic iron base at room temperature, and at the same time, the aluminum and carbon content is greater than approximately 7% by weight and 0.7% by weight, respectively. According to some recent studies, the chemical composition range is iron-(17.45~35.0)manganese-(7.1~12)aluminum-(0.7~1.2)carbon (Fe-(17.45~35.0) Mn-(7.1~12) Al-(0.7~1.2) C) (Note: Unless otherwise specified, the chemical composition of all alloys referred to below in this invention is in weight percent) is an austenitic iron-manganese-aluminum-carbon alloy obtained by hot-rolling or hot-forging, solution heat treatment at 1050-1200°C for 1-2 hours, rapid quenching in water or oil solution and cooling to room temperature, and the resulting alloy microstructure is a single austenite (γ-phase) without any precipitates.The yield strength, ultimate tensile strength, elongation, and microhardness of the alloy in this quenched state range from 350 to 540 megapascals (MPa), 700 to 1000 megapascals, 56 to 72%, and 200 to 230 Vickers hardness (Hv), respectively, and the actual values ​​are determined by the chemical composition of the alloy. While this quenched state alloy possesses excellent ductility and good tensile strength, unfortunately, its yield strength remains considerably low. Subsequent aging treatment is essential to improve the mechanical strength of the alloy, particularly its yield strength, and this aging treatment can promote the precipitation of nano-order κ-carbides within the γ-base. κ-carbides are rich in carbon and aluminum, and the precipitation process in a supersaturated austenitic iron base inevitably involves a diffusion process of various alloying elements associated with a large amount of carbon. Therefore, optimization and appropriate aging time and / or high aging temperature are usually required. Previous reports indicate that aging treatment at 550-600°C for 15-40 hours yields lightweight iron-manganese-aluminum-carbon alloys with the best combination of strength and ductility. The yield strength, ultimate tensile strength, and hardness of iron-manganese-aluminum-carbon alloys after optimized aging treatment reach 680-990 megapascals, 995-1180 megapascals, and 350-400 Vickers hardness, respectively, while the elongation remains within the range of 55-26.0%, with the actual value determined by the chemical composition of the alloy. According to conventional studies, compared to alloys in a solution-treated state followed by rapid quenching, the ultimate tensile strength of iron-manganese-aluminum-carbon alloys (carbon content ≤1.2%) after optimized aging treatment is improved by approximately 180-295 megapascals (18-42%), and the yield strength is increased by approximately 330-450 megapascals (83-94%). This clearly demonstrates that precipitation hardening plays a crucial role in improving the yield strength and microhardness of this type of alloy.It should be noted that further extending the aging time often leads to the formation of coarse κ-carbides, ferrite (α), regular DO3 phases, and β-manganese phases (β-Mn phases) at the austenite grain boundaries, and these products unfavorably affect the ductility of this type of alloy.

[0005] To explain in more detail the key factors related to the development of this type of lightweight, precipitation-hardening austenitic iron-manganese-aluminum-carbon alloy, we have carefully selected and described in detail several cases published in recent years, which show significant improvements in alloy mechanical strength, particularly yield strength, while maintaining good ductility (elongation exceeding 25%). According to a report by Gutierrez-Urrutia and Raabe, when an iron-30.5-manganese-8.0-aluminum-1.2-carbon (Fe-30.5Mn-8.0Al-1.2C) alloy is hot-rolled, solution-treated at 1100°C for 2 hours, then quenched in water, and subsequently aged at 600°C for 24 hours, the yield strength, ultimate tensile strength, and elongation of the alloy become 990 megapascals, 1180 megapascals, and 37%, respectively. The observed improvement in mechanical strength is attributed to the precipitation of nano-order κ-carbides with high volume fractions and uniform distribution on the austenitic iron base during aging treatment. Similarly, according to Wu et al., when an iron-26 manganese-10 aluminum-1.0 carbon (Fe-26Mn-10Al-1.0C) alloy is hot-rolled, solution-heat-treated at 1100°C for 1 hour, then cold-rolled, annealed at 1000°C for 15 minutes, and subsequently quenched, the resulting alloy microstructure is a single austenitic iron phase (γ-phase) without precipitation, and its yield strength, ultimate tensile strength, and elongation are 485 megapascals, 820 megapascals, and 72%, respectively. Furthermore, after aging treatment at 550°C for 40 hours, the resulting alloy microstructure is γ+κ-carbides, where κ-carbides are added to the γ phase. Within this microstructure, the volume fraction of κ-carbides precipitated on the austenitic iron base and uniformly distributed is approximately 43%. Notably, the mechanical properties of this alloy are further improved, with yield strength, ultimate tensile strength, and elongation of 955 megapascals, 1040 megapascals, and 38.2%, respectively.Compared to the mechanical properties of the alloy in its quenched state, it is clear that the precipitation of high-density κ-carbides, accelerated by aging at 550°C for 40 hours, can improve the ultimate tensile strength to 26.8%, and the yield strength is further significantly increased to 96.9%. In another recent example, Haase et al. showed that after hot-rolling an iron-29,8-manganese-7.65-aluminum-1.11-carbon (Fe-29,8Mn-7.65Al-1.11C) alloy, solution heat treatment at 1150°C for 5 hours, and quenching, the microstructure in the quenched state is a single austenite iron phase without precipitation. The yield strength, ultimate tensile strength, and elongation of this quenched alloy are 540 megapascals, 840 megapascals, and 56%, respectively. After aging this quenched alloy at 550°C for 15 hours, its microstructure consists mainly of densely distributed nano-order κ-carbides, with the γ phase added to the austenitic iron base. Its yield strength, ultimate tensile strength, and elongation are 880 megapascals, 995 megapascals, and 26%, respectively. With a good elongation of 26%, the improvement in yield strength and ultimate tensile strength reaches 63% and 18.5%, respectively. A common method to precipitate a sufficient amount of nano-order κ-carbides on the γ-phase base while maintaining the absence of precipitation and / or second phases at grain boundaries is to perform long-duration aging (reaching 15-40 hours) in the 550-600°C range. This process also produces the highly desirable combination of high strength (especially yield strength) and high ductility for this type of lightweight precipitation-hardening austenitic iron-manganese-aluminum-carbon alloy.

[0006] Compared to the above references, for each representative reference material in which the carbon content of the iron-manganese-aluminum-carbon alloy is less than 1.2% by weight, refer to two recent U.S. patents owned by the inventor of the present invention, Tzeng-Feng Liu: US9,528,177B2, granted in 2016, and US10,167,528B2, granted in 2019, which disclose a series of newly designed iron-manganese-aluminum-carbon alloys consisting of (23-34) manganese, (6-12) aluminum, and (1.4-2.2) carbon, with the remainder being iron, and having a relatively high carbon content. One of the outstanding properties of this series of alloys is that, after hot rolling, solution heat treatment at 980-1200°C for 1 hour, and then quenching in water or ice water to cool to room temperature, the microstructure of the alloy in this quenched state already contains high-density nano-order κ-carbides formed on the austenitic iron base by spinodal decomposition during quenching. This property is in stark contrast to the austenitic iron-manganese-aluminum-carbon alloys with a carbon content of less than 1.2%. In alloys with low carbon content, κ-carbides were not observed in either the solution heat treatment or quenched state. The high-density nano-order κ-carbides formed during quenching have at least two significant effects. The first effect is a significant improvement in the combined mechanical strength, particularly yield strength and ductility, of the alloy in the quenched state, with yield strength, ultimate tensile strength, and elongation being 865-925 megapascals, 1030-1155 megapascals, and 50-65%, respectively. The second effect is that, in the case of austenitic iron-based alloys, existing high-density nano-order κ-carbides can achieve the best combination of mechanical properties in this type of alloy using only considerably shorter aging times and lower aging temperatures compared to the iron-manganese-aluminum-carbon alloys with a carbon content of less than 1.2 wt%. Take the iron-28.6-manganese-9.84-aluminum-2.05-carbon (Fe-28.6Mn-9.84Al-2.05C) alloy disclosed by Liu as an example.In the quenched state, high-density nano-order κ-carbides were formed on the austenitic iron base by spinodal decomposition during quenching, and no precipitation was observed at the grain boundaries. The volume fraction and average size of the κ-carbides were approximately 45% and 12 nanometers (nm), respectively. The yield strength, ultimate tensile strength, and elongation of the quenched alloy were 912 megapascals, 1123 megapascals, and 52.5%, respectively. After aging the quenched alloy at 450°C for 6 hours, the volume fraction and average size of the κ-carbides increased to 53% and 25 nanometers, respectively. In this case, the yield strength, ultimate tensile strength, and elongation were 1179 megapascals, 1306 megapascals, and 39.8%, respectively.

[0007] From the representative and latest research examples in this field and the cited US patent, it is clear that high-density nano-order κ-carbides formed on a γ-phase base can improve the mechanical strength of alloys, particularly the yield strength, while simultaneously maintaining excellent malleability and ductility, thus exhibiting a significant effect. In particular, the inventor Mr. Liu's patent demonstrates that during quenching, the nano-order κ-(carbides) conventionally present on the austenitic iron base grow uniformly without causing long-distance diffusion between aluminum and carbon, and that this uniform growth of nano-order κ-carbides significantly improves the effectiveness of aging treatment. Clearly, this type of lightweight precipitation-hardening austenitic iron-manganese-aluminum-carbon alloy is a promising material in the materials industry and can be widely applied to high-performance structural materials required not only in the automotive industry but also in military vehicles, shipbuilding, and the aerospace industry. However, it should be noted that the most important thing to understand the potential of this type of alloy for promising applications is how to join the various parts made from this type of high-strength, highly ductile precipitation-hardening austenitic iron-manganese-aluminum-carbon alloy. Surprisingly, despite the very urgent needs of users, there is an extreme lack of information related to the weldability of precipitation-hardening austenitic iron-manganese-aluminum-carbon alloy. This is a common problem for all precipitation-hardening alloys, including high-strength 7-series aluminum alloys, and the reasons for this will be explained in detail later, but first, here is an overview. (1) Because fusion welding involves the remelting, resolidification, and redistribution of alloying elements, the microstructure of the weld metal (fusion zone, FZ) is completely different from that of the original base material (the material being welded). In particular, in addition to both dendritic microstructure and element segregation, precipitates obtained by solution heat treatment and aging treatment, which mainly have a strengthening effect, dissolve completely, and the mechanical strength of the as-welded weldments decreases significantly after welding. Therefore, the hardness of the fusion zone of the welded member is always significantly reduced after fusion welding. (2) Similarly, high heat input during welding dissolves and / or generates strengthening precipitates in the heat-affected zone (HAZ) of the welded member, causing grain growth and coarsening (HAZ), which significantly weakens the strength of the welded member. (3) Generally, precipitation-hardening alloys typically contain a large amount of alloying elements. Therefore, in addition to a significant decrease in mechanical strength, these types of metals are often prone to hot cracking during the welding process, such as solidification cracking along the weld path and liquefaction cracking near the boundary between the molten region and the weld heat-affected region. In fact, to our knowledge, there are only three relevant reports, which evaluate the microstructural changes that occur during the welding period and the ultimately obtained mechanical properties of alloys whose chemical composition falls within the range of the precipitation-hardening austenitic iron-manganese-aluminum-carbon alloys of the present invention.

[0008] The following literature provides a more detailed explanation and discussion of the above characteristics and advantages. [1] AJ Ardell, “Precipitation Hardening”, Metall. Trans., 16A (1985) 2131-2165.[2] “Precipitation hardening of aluminum alloys”, Totalmateria, 2010. [3] I. Gutierrez-Urrutia, D. Raabe, “Influence of Al content and precipitation state on the mechanical behavior of austenitic high-Mn low-density steels”, Scripta Mater., 68 (2013) 343-347. [4] Z. Q. Wu, H. Ding, X. H. An, D. Han, X. Z. Liao, “Influence of Al content on the strain-hardening behavior of aged low density Fe-Mn-Al-C steels with high Al content”, Mater. Sci. Eng. A, 639 (2015) 187-191. [5] W. Song, W. Zhang, J. von Appen, R. Dronskowski, W. Bleck, “κ-phase formation in Fe-Mn-Al-C austenitic steels”, Steel Res. Intern., 86 (2015) 1161-1169. [6] K. Lee, S. J. Park, J. Moon, J. Y. Kang, T. H. Lee, H. N. Han, “β-Mn formation and aging effect on the fracture behavior of high-Mn low-density steels”, Scripta Mater., 124 (2016) 193-197. [7] E. Welsch, D. Ponge, S. M. Hafez Haghighat, S. Sandlobes, P. Choi, M. Herbig, S. Zaefferer, D.Raabe, “Strain hardening by dynamic slip band refinement in a high-Mn lightweight steel”, Acta Mater., 116 (2016) 188-199. [8] H. Ding, D. Han, J. Zhang, Z. Cai, Z. Wu, M. Cai, “Tensile deformation behavior analysis of low density Fe-18Mn-10Al-xC steels”, Mater. Sci. Eng. A, 652 (2016) 69-76. [9] S. G. Peng, R. B. Song, Z. D. Tan, C. H. Cai, K. Guo, Z. H. Wang, “Abrasive wear behaviors of lightweight austenitic Fe-24Mn-7Al-1C steel and Mn13Cr2 steel”, J. Iron Steel Res. Int., 23 (2016) 857-866.

[10] J. Moon, S. J. Park, C. Lee, H. N. Han, T. H. Lee, C. H. Lee, “Microstructure evolution and age-hardening behavior of microalloyed austenitic Fe-30Mn-9Al-0.9C light-weight steels”, Metall. and Mater. Trans. A, 48 (2017) 4500-4510.

[11] Z. Y. Huang, A. L. Hou, Y. S. Jiang, P. Wang, Q. Shi, Q. Y. Hou, X. H. Liu, “Reitveld refinement, microstructure, mechanical properties and oxidation characteristics of Fe-28Mn-xAl-1C (x = 10 and 12 wt.%) low-density steels”, J. Iron and Steels Res. Intern., 24 (2017) 1190-1198.

[12] C. Haase, C. Zehnder, T. Ingendahl, A. Bikar, F. Tang, B. Hallstedt, W. Hu, W. Bleck, D. A. Molodov, “On the deformation behavior of κ-carbide-free and κ-carbide-containing high-Mn light-weight steel”, Acta Mater., 122 (2017) 332-343.

[13] J. Xing, Y. Wei, L. Hou, “An overview of the effects of alloying elements on the properties of lightweight Fe-(15-35)Mn-(5-12)Al-(0.3-1.2)C steel”, JOM, 70 (2018) 929.

[14] J. Lee, S. Park, H. Kim, S. J. Park, K. Lee, M. Y. Kim, P. P. Madakashira, H. N. Han, “Simulation of κ-carbide precipitation kinetics in aged low-density Fe-Mn-Al-C steels and its effects on strengthening”, Metals and Mater. Int., 24 (2018) 702-710.

[15] S. W. Park, J. Y. Park, K. M. Cho, J. H. Jang, S. J. Park, J. Moon, T. H. Lee, J. H. Shin, “Effect of Mn and C on age hardening of Fe-Mn-Al-C lightweight steels”, Metals & Mater. Int., 25 (2019) 683-696.

[16] J. Pang, Z. Zhou, Z. Zhao, D. Tang, J. Liang, Q.He, “Tensile behavior and deformation mechanism of Fe-Mn-Al-C low density steel with high strength and high plasticity”, Metals, 9 (2019) 897.

[17] Tzeng-Feng Liu, “Composition design and processing methods of high strength, high ductility and high corrosion resistance FeMnAlC alloys”, US 9,528,177 B2 / 2016.

[18] Tzeng-Feng Liu, “Composition design and processing methods of high strength, high ductility and high corrosion resistance FeMnAlC alloys”, US 10,167,528 B2 / 2019.

[19] L. Bartlett, D. Van Aken, “High manganese and aluminum steels for the military and transportation industry”, JOM, 66 (2014) 1770.

[20] W. Evans, A. J. Ramirez, K. Sebeck, “Investigation of hot cracking phenomena in lightweight armor steel based on the FeMnAlC alloy system”, 2018 NDIA GVSET symposium, Aug. 7-9, 2018, Novi, Michigan.

[0009] To our knowledge, CP Chou and CH Lee conducted perhaps the earliest relevant research on the welding of this type of alloy having the chemical composition of the precipitation-hardening austenitic iron-manganese-aluminum-carbon alloy of the present invention in 1989. Their research investigated the effect of changes in solidification microstructure within the molten region on two types of fully austenitic iron phase alloys (Fe-30Mn-8.9Al-1.29C and Fe-29Mn-9.0Al-1.0C) treated with carbon-content gas tungsten arc welding (GTAW). Autogenous weld is a form of welding in which the weld filler material (weld filler wire or welding rod) can be from the molten base material or other individual components. Both alloys were manufactured using the following methods: non-vacuum induction melting, casting, hot-forging at 1200°C, homogenization at 1050°C for 12 hours, cold-rolling, followed by annealing at 950°C for 1 hour. The experimental alloy thickness was 1 / 8 inch. Before welding, both base alloys had a fully austenitic structure. Their study found that the carbon content significantly influenced the microstructure and morphology of the molten region of the weld bead. In the iron-30 manganese-8.9 aluminum-1.29 carbon (Fe-30Mn-8.9Al-1.29C) alloy, the molten region of the weld has a complete austenite phase (ferrite less than 1 volume%), and further contains cellular γ-dendrite crystals and several κ-carbides present in the eutectic region. In this example, the microstructure of the molten region has a typical multi-branched dendritic crystal structure, with primary dendrites reaching several hundred micrometers in length (greater than 300 micrometers (μm)) and secondary dendrite arms approximately 20-40 micrometers in length.It should be noted that, with the exception of a small amount of eutectic κ-carbide, the primary and secondary dendritic crystal cells do not contain κ-carbide or other precipitates, which may be caused by rapid solidification during the welding process. The carbon content decreases to 1.0 wt%, and a large amount of ferrite (approximately 10.2 vol%) can be observed in the molten region, mainly located between the secondary austenitic iron dendritic crystal arms. Clearly, the fusion-welded welded member does not have a homogeneous structure and cannot fully maintain the properties of the base material. Unfortunately, this study does not mention the mechanical properties associated with the welded member.

[0010] In 1990, Chou and Lee reported on the microstructure and mechanical properties of austenitic iron alloys (Fe-29.3Mn-8.6Al-0.81C and Fe-29Mn-8Al-1.17C) welded using two types of butt-joint autogenous gas tungsten arc welding (GTAW) processes. Hereafter, these two alloys will be referred to as 0.81C and 1.17C, respectively. First, these base alloys are prepared in a vacuum induction furnace. Next, the cast alloys are hot forged to a thickness of 8 mm at 1200°C, then homogenized at 1050°C for 12 hours, and the homogenized pieces are cold-rolled into 4 mm thick plates. These are then annealed in argon at 950°C for 1 hour, followed by water quenching. The microstructure of the quenched base alloy is a complete austenite phase. The ultimate tensile strength and elongation of these two quenched base alloys, 0.81C and 1.17C, are 1040 megapascals, 50%, and 1080 megapascals, 54%, respectively. In comparison to the above, the combined values ​​of ultimate tensile strength and elongation for two iron-manganese-aluminum-carbon alloys with a carbon content of less than 1.2% in the solution heat treatment and quenched state are within the range. Unfortunately, the yield strength of these alloys is not mentioned. However, judging from the absence of κ-carbides observed in the austenitic iron base, a decrease in the yield strength of these two alloys can be reasonably expected, which is very similar to that of the iron-manganese-aluminum-carbon alloys with a carbon content of less than 1.2% after solution heat treatment and quenching, i.e., 350-540 megapascals as mentioned above. Following the spontaneous gas tungsten arc welding process, the microstructure of the molten region of the butt weld exhibits general characteristics very similar to the multi-branched dendritic structure described above. Specifically, the length of the primary dendritic crystals extends to several hundred micrometers (greater than 300 micrometers), and except for a small amount of eutectic κ-carbide, no κ-carbide or other precipitates appear in the primary and secondary austenite dendritic crystal cells. Simultaneously, the presence of approximately 5% by volume and 0.5% by volume of the ferrite phase is observed in the molten regions of the 0.81C and 1.17C alloys, respectively.After post-welding heat-treated (PWHT) at 1050°C for 5 to 240 minutes, no precipitates appeared in the primary and secondary austenite dendritic crystal cells. Tensile tests of the post-welding heat-treated butt joint samples showed that the 0.81C alloy had an ultimate tensile strength and elongation of 930 megapascals and 25%, respectively, while the 1.17C alloy had an ultimate tensile strength and elongation of 900 megapascals and 19%, respectively. Clearly, the ultimate tensile strength of the iron-manganese-aluminum-carbon welded members in the post-welded state retained nearly 90% of the ultimate tensile strength of the base alloy, but the elongation deteriorated significantly, decreasing from 54% (base metal) to approximately 20%. Furthermore, it should be noted that this study did not investigate the yield strength of samples of both the base alloy and the butt joint welds. However, as mentioned above, no nano-order precipitates were observed in the austenite dendritic crystal cells of the molten region, indicating that the yield strength of the resulting welded member was low. These results indicate that the above treatment cannot achieve the goal of improving the yield strength of the welded member in the post-weld state. Consequently, from 2017 to 2019, the development of precipitation-hardening austenitic iron-manganese-aluminum-carbon alloy welding filler wires for special fusion welding was considered the most urgent issue in this field.

[0011] Based on the two studies by Zhou and Li mentioned above, there are several points to note. (1) In their study, the base alloy used is a complete austenitic iron phase in the solution heat treatment and quenched state. In other words, the austenitic iron base does not contain the κ-carbides necessary to strengthen the precipitation-hardening austenitic iron-manganese-aluminum-carbon alloy. In particular, fusion welding clearly changes the microstructure of the molten region significantly, converting the homogeneous single-phase austenitic iron base into a dendritic crystalline structure, where the dendritic crystalline region is formed with a ferrite phase and a eutectic phase (e.g., γ + κ-carbide). Furthermore, within the molten region of the welded member, there is no κ-carbide precipitation in either the primary or secondary austenitic dendritic crystalline cells. This clearly indicates that the microstructure within the molten region and the microstructure of the base alloy are completely different, even if both have the same chemical composition. Therefore, the mechanical properties of the resulting welded member are still unpredictable to some extent if judged solely from the chemical composition of the base alloy and the weld filler wire. (2) As mentioned above, in this type of alloy, uniform dispersion of high-density nano-order κ-carbides based on austenitic iron is an essential requirement for obtaining a combination of excellent strength, particularly yield strength, and ductility. Therefore, during welding, the final mechanical properties of the welded member depend on how the fusion weld affects the changes and distribution of κ-carbides in both the molten and heat-affected zones. Intuitively, these possible situations are very similar to those encountered by other precipitation-hardening alloys (e.g., AA7075 aluminum alloy), and are expected to be severe softening phenomena, which generally occur in the molten and heat-affected zones. The main reason for severe softening is that the nano-order precipitates with strengthening properties dissolve, coarseen, and / or change into other incoherent phases with less effect. These types of problems remain a major challenge when joining various precipitation-hardening alloys using fusion welding, and a wide range of related research and development has been dedicated to solving these problems for over 50 years. However, as Howell and Gerth pointed out in 2017, to this day, direct welding of lightweight precipitation-hardening austenitic iron-manganese-aluminum-carbon alloys in the "aged state" still does not exist. The lack of a viable solution for joining aged precipitation-hardening austenitic iron-manganese-aluminum-carbon alloys using fusion welding is a major obstacle, preventing these types of alloys from being widely applied in various industrial fields.

[0012] As an initial test, in 2018, Sebeck et al. presented the results of welding a refined age-hardened iron-30 manganese-9 aluminum-1 silicon-0.9 carbon-0.5 molybdenum (Fe-30Mn-9Al-1Si-0.9C-0.5Mo) alloy using gas metal arc welding (GMAW), double vee groove butt, and commercial 316LSi austenitic stainless steel welding filler wire (composed of (11-14) nickel, (18-20) chromium, (1.0-2.5) manganese, (2.0-3.0), molybdenum, (0.65-1) silicon, with the remainder being iron). The base material was hot-rolled to approximately 12.7 mm (approximately 0.5 inches) at 1204°C, solution-treated at 1050°C for 2 hours, and then quenched. The report does not describe the yield strength, ultimate tensile strength, or elongation of the base material after quenching, but its Vickers hardness is approximately 210 Hv. After aging treatment at 538°C for 30 hours, the microstructure of the obtained base material was found to be a complete austenitic iron phase, and at the same time, κ-carbides were uniformly formed at the austenitic iron base and grain boundaries. The typical yield strength, ultimate tensile strength, elongation, and hardness of the base material after aging treatment were 800 megapascals, 827 megapascals, 36%, and approximately 360 Vickers hardness, respectively. In this study, because 316LSi metal solder was used, the microstructure of the molten region was not clearly described, resulting in a significant decrease in the yield strength and elongation of the resulting welded material to 350-400 megapascals and 23-5%, respectively. These are clearly insufficient and unsatisfactory. Therefore, this study particularly emphasizes one point: there is a need to develop a special metal solder for use with precipitation-hardening austenitic iron-manganese-aluminum-carbon alloys.

[0013] Based on the limited existing literature and research reports mentioned above, research in this field still lacks a complete and comprehensive plan, and a welding filler wire applicable to this long-awaited precipitation-hardening austenitic iron-manganese-aluminum-carbon alloy that can solve the relevant problems encountered during welding has yet to be found. However, based on the fragmentary data obtained from these previous studies, there is a fundamental problem that deserves particular attention. During fusion welding, alloy components have the properties of rapid heating, remelting, resolidification, and redistribution, so the design of the alloy component of the welding filler wire is necessary to control changes in the microstructure and avoid severe softening of the molten region, and all of these problems need to be solved. In particular for precipitation-hardening austenitic iron-manganese-aluminum-carbon alloys, the most important reinforcing component is high-density nano-order κ-carbide, and previous studies have shown that whether or not this carbide can form in the austenitic dendritic crystal cells of the molten region is severely suppressed due to the processes of remelting and rapid cooling. Therefore, finding a welding filler wire with an appropriate alloy component design to address this problem is important and urgent. The challenges we face in developing metal solders for fusion welding of precipitation-hardening austenitic iron-manganese-aluminum-carbon alloys are similar to the problems that such precipitation-hardening aluminum alloys face during fusion welding. Therefore, investigating and studying the significant problems that occur during fusion welding of precipitation-hardening aluminum alloys provides us with inspiration from past experience.

[0014] Among many aluminum alloys, precipitation-hardening aluminum alloy 7075 (AA7075) is the most preferred by the aerospace and military industries due to its excellent specific strength, fracture properties, and high formability. AA7075 typically contains a variety of alloying elements, with the chemical composition ranging from: Aluminum-(5-6.5)Zinc-(1.6-2.9)Magnesium-(1.2-2.0)Copper-(0-0.3)Manganese-(0-0.28)CChromium-(0-0.5)Iron-(0-0.15)Titanium-(0-0.4)Silicon (Al-(5-6.5)Zn-(1.6-2.9)Mg-(1.2-2.0)Cu-(0-0.3)Mn-(0-0.28)Cr-(0-0.5)Fe-(0-0.15)Ti-(0-0.4)Si). To obtain the ideal mechanical properties of AA7075 aluminum alloy, the most widely used heat treatment methods are T6 and / or T651. T6 is a heat treatment code commonly used for heat-treated alloys and includes solution heat treatment in the α-solid solution phase region having a face-centered-cubic (FCC) structure, followed by rapid quenching to room temperature, and then artificial aging. For AA7075 aluminum alloy, the most common artificial aging treatment for T6 is aging at 110-120°C for 16-24 hours. Another code, T651, involves further stress relief of the product after the T6 treatment. The precipitation hardening process of AA7075 aluminum alloy varies depending on the results of aging at a constant temperature for a certain period, and can be summarized as follows: Supersaturated α-solid solution → GP zones → metastable η'-phase → equilibrium η-phase. The GP zones and the face-centered cubic α-base are completely coherent, while the η' precipitates and the α-base are semi-coherent, but both are formed within aluminum crystal grains.In this alloy system, the uniformly distributed nano-order GP regions and η' precipitates, due to their fully coherent (or semi-coherent) nature with the base, are the primary factors causing age hardening and strengthening. While precise values ​​vary depending on the detailed composition, typical yield strength, ultimate tensile strength, elongation, and hardness of AA7075 aluminum alloy treated with standard T6 and / or T651 are 459–539 megapascals, 510–597 megapascals, 8.5–14.6%, and 157–180 Vickers hardness, respectively.

[0015] Extending the aging time or increasing the aging temperature (overaging, etc.) accelerates the equilibrium of η precipitates, which form non-uniform phases at the dispersoid / base interface and grain boundaries, at the cost of the GP region and η precipitates disappearing. Since η precipitates not only become rougher but also lose coherence with the base, overaging inevitably leads to a significant decrease in hardness, mechanical strength, and malleability. However, despite extensive research and development on fusion welding of precipitation-hardening AA7075 aluminum alloys over the past several decades, a method has not been found to thoroughly solve problems such as significant strength (hardness) loss in the molten region, solidification cracks, liquefaction cracks, and porosity occurring along the weld bead direction and the interface of the molten region / weld heat-affected zone, which are some of the common problems that occur when welding AA7075 aluminum alloys using conventional fusion welding. The following is a brief summary of the latest developments in how to address these issues, as mentioned in several recently published representative studies related to automated welding and / or dissimilar material-based welding filler wires.

[0016] As described above, after standard T6 and / or T651 treatment, the base of the AA7075 aluminum alloy matrix contains high-density nano-order GP regions and η precipitates. However, in the case of automatic welding and / or dissimilar material welding filler wires (e.g., ER5356 (Aluminum-(4.5-5.5 wt.%) Magnesium), ER4043 (Aluminum-(4.5-6.0 wt.%) Silicon), ER5356 (Al-(4.5-5.5 wt.%) Mg), ER4043 (Al-(4.5-6.0 wt.%) Si) after fusion welding, the microstructure of the molten region generally consists of a typical dendritic crystalline structure with eutectic regions and hardening nano-order GP regions. When a small amount of large eutectic precipitates (e.g., flake-like layers of ηMg(ZnCuAl)2 and θ(Al2Cu)) are added within the welding zones, the η' precipitates originally present in the base metal completely dissolve within the molten region. Within this, the hardened GP region and η' precipitates in the weld heat-affected zone also dissolve or are converted into softer η' precipitates. As a result, the molten and weld heat-affected zones of the welded member in the as-welded condition are significantly softened, with their hardness reduced to only 50-67% and 67-87% of the original base metal hardness, respectively. Clearly, the softening of the molten region is more pronounced than that of the weld heat-affected zone. Simultaneously, due to the redistribution of alloying elements during the fusion welding process, post-weld heat treatment clearly cannot effectively improve the micro-segregation of the alloy. Therefore, in both the as-welded and aged states, the weld molten region remains the most fragile region. Furthermore, thermal stress and solidification shrinkage due to high input heat also contribute to these issues. Due to shrinkage, solidification cracks are always detected within the weld area during the autogenous fusion welding period. In this regard, filler materials are often used to reduce the crack sensitivity of the alloy in order to reduce the occurrence of solidification cracks. However, while using dissimilar filler materials to reduce the crack sensitivity of the alloy, the concentration of reinforcing alloy material originally present in the AA7075 aluminum alloy also decreased.As a result, not only is the strength (or hardness) of the alloy significantly reduced after welding, but the alloy's reactivity to post-weld heat treatment also deteriorates, leading to the appearance of various defects such as liquefaction cracks and pores near the fusion line (e.g., the boundary between adjacent molten regions / weld heat-affected regions). Therefore, due to these three major obstacles associated with fusion welding, tensile testing of AA7075 aluminum alloys after welding is rarely performed. In fact, until recently (2019), precipitation-hardening AA7075 aluminum alloy was considered an alloy that could not be welded using conventional fusion welding methods.

[0017] To further address the cracking problem caused during fusion welding of AA7075 aluminum alloy, researchers recently (2019) submitted two investigative reports. Ipekoglu and Cam welded a 2cm thick AA7075-T6 alloy plate using ER5356 welding filler wire and cold metal transfer gas metal arc welding (CMT-GMAW). The main features of the cold metal transfer process are the controllability of material deposition and low heat input, which is presumably why the formation of hot cracks is suppressed. However, although the cracks disappeared, the researchers observed the appearance of large pores in the molten region. Furthermore, it was observed that nano-order GP regions and η precipitates in the molten region completely dissolved, and precipitates in the weld heat-affected zone also underwent an η'→η change. As a result, the yield strength, ultimate tensile strength, elongation, and hardness of the alloy decreased from 539 megapascals, 597 megapascals, 14.6%, and 175 Vickers hardness, respectively, to unmeasurable (e.g., no yield strength), 312 megapascals, 0.03%, and 65 Vickers hardness (only 37.1% of the base material). The numerous large holes that appear in welded members are generally considered to be the cause of brittle fracture. Therefore, although cold metal welding can repair cracks caused by fusion welding, this method has the problem of hole formation, a significant decrease in ductility (elongation), and serious softening problems, and a solution is still needed.

[0018] Meanwhile, in 2019, Sokoluk et al. used a newly developed welding filler wire for welding AA7075 alloy. This wire contains approximately 1.7 volume% titanium carbide (TiC) nanoparticles (size approximately 40-60 nanometers) mixed into the AA7075 alloy (aluminum-6.4Zn-3.2Mg-1.2Cu-0.15Cr). According to their report, the addition of titanium carbide nanoparticles to the welding filler wire results in the following very important properties: (1) During the solidification process after arc welding, the titanium carbide nanoparticles appear, significantly slowing down the solidification of the front surface and reducing the growth rate of dendritic grains. Due to this slow growth rate of dendritic grains, the solidification microstructure that develops in the molten region consists of fine spherical particles rather than the directional, long dendritic grains commonly seen after normal fusion welding. (2) During the solidification process, the titanium carbide nanoparticles remain in the second phase (eutectic) of magnesium (zinc, copper, aluminum) 2. As a result, the size, shape, and distribution of the second phase present in the eutectic region are effectively improved. Due to the absence of directional dendritic grain growth and the improvement of the second phase within the eutectic region, the hot cracking susceptibility of the AA7075 alloy can be completely eliminated. The welding filler wire treated with titanium carbide nanoparticles appears to have overcome two insurmountable challenges, namely hot cracking and porosity, problems encountered during more than 70 years of fusion welding of the AA7075 alloy. These remarkable results are said to be an unprecedented breakthrough in fusion welding of the AA7075 alloy, and at the same time, appear to make the AA7075 alloy a weldable alloy. However, in this study, tensile test results show that the ultimate tensile strength and elongation of the welded member in the post-welded state are only 392 megapascals and 1.5%, respectively. However, after solution heat treatment at 480°C and artificial aging at 120°C for 19 hours, the ultimate tensile strength and elongation of the post-weld heat-treated material were improved to 551 MPa and 5.21%, respectively.The elongation values of as-welded welded members and post-weld heat treated welded members (1.5% and 5.21% respectively) are apparently far lower than the typical elongation of the base material T6-AA7075 (~8.5-14.6%). Furthermore, it should be noted that this study does not mention the yield strength of as-welded welded members and post-weld heat treated welded members. According to the microstructure of the molten zone described in the present study, it is clear that the molten zones of both as-welded welded members and post-weld heat treated welded members are mostly composed of α-dendrites and a small amount of eutectic crystals. More importantly, there are no precipitates in the dendritic crystal cells of the molten zone. Based on the above-mentioned strengthening mechanism of precipitation hardening alloys, most areas of the molten zone are poor in precipitates, which inevitably leads to severe softening in this area, and at the same time greatly affects the yield strength of the welded member. Therefore, it is obvious that melting and welding precipitation-hardened AA7075 aluminum alloy is still the most challenging task in this field, since common problems such as softening in the molten zone, solidification cracking, liquefaction cracking and pore formation still exist.

[0019] However, the above specific case strongly suggests that a properly alloy-designed welding filler wire is essential for welding precipitation-hardened austenitic iron-manganese-aluminum-carbon alloys. The following embodiments disclosed in the present invention clearly show that all unsolved problems in fusion welding of precipitation-hardened austenitic iron-manganese-aluminum-carbon alloys can be simultaneously solved by appropriately designing the chemical composition of the welding filler wire. That is, by using the welding filler wire disclosed in the present invention, the as-welded welded member can obtain excellent hardness, strength (especially yield strength) and ductility in the molten zone, without formation of solidification cracks, liquefaction cracks and pores in said zone.

[0020]

[21] C. P. Chou and C. H. Lee, “The influence of carbon content on austenitic-ferrite morphology in Fe-Mn-Al weld metals”, Metall. Trans. A, 20 (1989) 2559-2561.

[22] C. P. Chou and C. H. Lee, “Effects of carbon on the weldability of Fe-Mn-Al alloys”, J. Mater. Sci., 25 (1990) 1491-1496.

[23] R. A. Howell, R. J. Gerth, “Fe-Mn-Al-C Alloy Steels-A New Armor Class”, SAE International, 2017; doi: 10.4271 / 2017-01-1703.

[24] J. Moon, S. J. Park, C. Lee, H. N. Han, T. H. Lee, C. H. Lee, “Microstructure evolution and age-hardening behavior of microalloyed austenitic Fe-30Mn-9Al-0.9C light-weight steels”, Metall. Mater. Trans. A, 48 (2017) 4500.

[25] K. Sebeck, I. Toppler, M. Rogers, R. Howell, K. Limmer, B. Cheeseman, W. Herman, “High Mn, High Al steels for thick plate armor applications”, 2018 NDIA GVSET symposium, Aug. 7-9, 2018, Novi, Michigan.

[26] G. Ozer and A. Karaaslan, “Properties of AA7075 aluminum alloy in aging and retrogression and re-aging process”, Trans. Nonferrous Met. Soc., 27 (2017) 2357-2362.

[27] J. Z. Liu, J. H. Chen, X. B. Yang, S. Ren, C. L. Wu, H. Y. Xu, and J.Zhou, “Revisiting the precipitation sequence in Al-Zn-Mg-based alloys by high-resolution transmission electron microscopy”, Scripta Mater., 63 (2010) 1061-1064.

[28] B. Cevik, “Gas tungsten arc welding of 7075 aluminum alloy: microstructure properties, impact strength, and weld defects”, Mater. Res. Express, 5 (2018) 066540.

[29] B. Hu and I. M. Richardson, “Microstructure and mechanical properties of AA7075(T6) hybrid laser / GMA welds”, Mater. Sci. Eng. A, 459 (2007) 94-100.

[30] G. Ipekoglu and G. Cam, “Formation of weld defects in cold metal transfer arc welded 7075-T6 plates and its effect on joint performance”, IOP Conf. Series: Mater. Sci. Eng., 629 (2019) 012007.

[31] M. Sokoluk, C. Cao, S. Pan, X. Li, “Nanoparticle-enabled phase control for arc welding of unweldable aluminum alloy 7075”, Nat. Comm., 10 (2019) 98; doi:10.1038 / s41467-018-07989-y. Summary of the Invention Problems to be Solved by the Invention

[0021] As described in the prior art, precipitation hardening (or age hardening) is one of the most effective means of improving the yield strength and hardness of ductile alloys. Generally speaking, precipitation hardening (or age hardening) alloys such as the strongest AA7xxx series aluminum alloys (AA7075), precipitation hardened stainless steels, and precipitation hardened austenitic iron-manganese-aluminum-carbon alloys, after solution heat treatment, quenching, and optimal aging, have their yield strength and hardness significantly improved, but their ductility is not severely impaired, because high-density nano-order precipitates are coherently (or semi-coherently) formed on the base. However, as described in the prior art in the background of the present invention, all nano-order precipitates present on the base due to previous aging treatment and having a strengthening effect are completely dissolved during fusion welding. After fusion welding, the typical microstructure of the molten region in the post-welded state consists of a majority of dendritic crystalline cells (dendrite cells) and a small number of eutectic regions. More importantly, dendritic crystalline cells do not contain precipitates. As a result, fusion welding inevitably leads to severe softening of the molten region. Furthermore, the molten region is the area most susceptible to high-temperature cracking reactions (e.g., solidification cracks and liquefaction cracks), and pores are formed simultaneously. To those skilled in the art, these three main problems are obstacles to overcome when fusion welding precipitation-hardening alloys; they are not insurmountable, but they are extremely difficult challenges.

[0022] Similarly, after solution heat treatment, quenching, and aging, nano-order κ-carbides with high-density strengthening properties are formed within the austenitic iron base of precipitation-hardening austenitic iron-manganese-aluminum-carbon alloys. Because high-density nano-order κ-carbides precipitate within the austenitic iron base, the alloy's strength (especially yield strength) and hardness are significantly improved without a substantial decrease in malleability. However, during the fusion welding process of precipitation-hardening austenitic iron-manganese-aluminum-carbon alloys, similar to other precipitation-hardening alloys, problems such as severe softening of the molten region, areas affected by high-temperature cracking effects (e.g., solidification cracking and liquefaction cracking), and pores are encountered, and solutions have been expected for decades. In 2017, Moon et al. pointed out that a high carbon content in precipitation-hardening austenitic iron-manganese-aluminum-carbon alloys makes welding difficult for automotive structures and poses a significant obstacle to the wider application of the alloy. A 2017 study by Howell and Gerth noted that, to date, no formally used welding solution exists for this type of (precipitation-hardening austenitic iron-manganese-aluminum-carbon alloy) plate. In 2018, Sebeck et al. further pointed out that "the high magnesium content of this type of alloy (precipitation-hardening austenitic iron-manganese-aluminum-carbon alloy) presents a significant challenge to the welding process...for example, the interchangeability of welding filler wires is limited..." Also in 2018, Evans et al. called for the development of special metal solders to precipitation-harden austenitic iron-manganese-aluminum-carbon alloys. From the above points, it can be considered that this remains a common understanding in this field.

[0023] To overcome all these unresolved problems encountered during fusion welding of precipitation-hardening austenitic iron-manganese-aluminum-carbon alloys, the inventors, based on decades of practical experience in the design and technological development of iron-manganese-aluminum-carbon alloys in the field of materials research, have conducted numerous experiments on the design of welding filler wires and have submitted this innovative invention. [Means for solving the problem]

[0024] The features of this invention are as follows: When using the precipitation-hardening austenitic iron-manganese-aluminum-carbon composite welding filler wire disclosed in this invention, the inventors disclose the following post-weld state characteristics of the microstructure of the welded region and the welding bead, which are not only unprecedented but also completely solve the problems of softening of the molten region and hot cracking. (1) In the post-welded state, high-density nano-order κ-carbides (approximately 3-5 nanometers) are present in the dendritic crystalline cells, and this unique structure is the most noteworthy feature disclosed by the present invention. As described in the prior art, after fusion welding, the fusion-welded region in the post-welded state has excellent hardness and strength, particularly yield strength, which is equivalent to or better than iron-manganese-aluminum-carbon alloys with a carbon content of less than 1.2 wt% after optimized aging treatment. This result is completely different from the results mentioned in the prior art relating to fusion welding of precipitation-hardening austenitic iron-manganese-aluminum-carbon alloys (e.g., AA7075 precipitation-hardening aluminum alloy, precipitation-hardening stainless steel, precipitation-hardening iron-manganese-aluminum-carbon alloy, etc.). In the prior art, after fusion welding of precipitation-hardening carbon alloys, the high-density strengthening (or hardening) of nano-order precipitates originally present in the base material always completely dissolves. In the post-weld state, the microstructure of the molten region consists mainly of a large number of dendritic crystal cells and a small number of eutectic regions. Furthermore, it is obvious that no precipitates are observed within the dendritic crystal cells of the molten region in the post-weld state, which causes severe softening of the molten region. (2) In the post-welded state, nano-order κ-carbides are present at the tips of dendritic crystal cells in austenite iron grains in the molten region, slowing the growth of dendritic crystal cells during solidification. As a result, the length and spacing of austenite dendrite cells formed in columnar austenite iron grains are only about 20-30 micrometers and 5-10 micrometers, respectively, which is not what is observed in fusion-welded AA7075 aluminum alloys or other iron-manganese-aluminum-carbon alloys. After fusion welding, the microstructure of the molten region of these alloys usually contains mostly long primary dendritic crystals (exceeding 200-300 micrometers). Furthermore, the amount of nano-sized (about 6-10 nanometers) κ-carbides present in the eutectic region is far greater than the amount of carbides observed in previous studies in iron-manganese-aluminum-carbon alloys with a carbon content of less than 1.29 wt% after fusion welding. (3) In the post-weld state in which appropriate amounts of titanium, niobium, and vanadium are added to the welding filler wire, in addition to the high-density nano-order κ-carbides present in the austenite dendritic crystal cells and eutectic regions, large amounts of titanium-rich titanium carbides (Ti-rich Ti-carbides), niobium-rich niobium carbides (Nb-rich Nb-carbides), and vanadium-rich vanadium carbides (V-rich V-carbides) are formed in the eutectic regions. These carbides have the same ductile face-centered cubic structure (FCC) as austenitic iron. Furthermore, the hardness of these carbides is very high, reaching a Vickers hardness of approximately 2000 to 3500. Therefore, the hardness of the molten region in the post-weld state increases significantly, but the ductility does not decrease significantly. (4) In the post-welded state, all phases present in the molten region consist of a base (austenite dendritic crystal cells with eutectic regions added) and all types of precipitates (κ-carbides, titanium-rich titanium carbides, niobium-rich niobium carbides, and vanadium-rich vanadium carbides), all of which have a ductile face-centered cubic (FCC) structure. Furthermore, the size of all types of precipitates is only about 3 to 10 nanometers. Therefore, the tensile test results of the welded material indicate excellent ductility in the molten region in the post-welded state. (5) By using the welding filler wire disclosed in the present invention, the problems of hot cracking and porosity (such as solidification and liquefaction cracking) that are often encountered in the fusion welding process of precipitation-hardening austenitic iron-manganese-aluminum-carbon alloys can be completely solved. These hot cracks and porosity often appear in the weld bead, the molten region, and / or in the region near the interface with the molten region / weld heat-affected region.

[0025] By using the welding filler wire disclosed in this invention, the microstructure of the molten region has the above-mentioned characteristics, and after fusion welding, the welded product in its molten region has high hardness, high strength (especially yield strength), and high malleability. For example, the average microhardness range of the malleable molten region obtained in this invention is 365 to 465 Vickers hardness, which is much larger than the average microhardness value (350 to 400 Vickers hardness) of iron-(17.45 to 35)manganese-(7.1 to 12)aluminum-(0.7 to 1.2)carbon (Fe-(17.45 to 35)Mn-(7.1 to 12)Al-(0.7 to 1.2)C) after hot rolling, solution heat treatment, quenching, and aging treatment optimized at 550 to 600°C as described in the prior art.

[0026] To achieve the goal of giving the molten region and weld bead the above-mentioned properties in the post-weld state, the present invention has conducted a comprehensive investigation and, through numerous experiments and analyses, has adjusted the alloy design of this innovative welding filler wire and modified various fusion parameters of the alloy. According to the present invention, the range of chemical composition components of each alloying element is as follows: Manganese (23-34 wt%, preferably 24-32 wt%), aluminum (7.5-11.5 wt%, preferably 8.0-11.0 wt%), carbon (1.35-1.95 wt%, preferably 1.40-1.95 wt%), titanium (0.0-2.5 wt%, preferably 0.1-2.5 wt%), niobium (0.0-3.0 wt%, preferably 0.1-3.0 wt%), vanadium (0.0-2.5 wt%, preferably 0.1-2.5 wt%), the sum of titanium, niobium and vanadium (0-3.0 wt%, preferably 0.1-3.0 wt%); and the remainder is iron. The novel properties and unprecedented features obtained in the molten region of the post-welded state disclosed in the present invention are mainly due to detailed studies of the individual effects of each alloying element, and further relevant details are as follows. (1) Manganese: Manganese is a strong austenitic iron stabilizing element and plays the most important role in maintaining the ductile face-centered cubic structure phase. Previous studies have demonstrated and confirmed that iron-manganese-aluminum-carbon alloys with a manganese content exceeding 17.5 wt% have a complete austenitic iron microstructure at room temperature after hot rolling, solution heat treatment, and quenching, as described in the prior art. However, in Example 9 of the present invention, a welding filler wire with a manganese content of 20.1 wt% in the post-welded state after fusion welding has a large amount of ferrite phase in the molten region of the weld bead, and its structure is body-centered cubic (BCC). In order to obtain a complete austenitic iron microstructure in the molten region in the post-welded state, the manganese content in the welding filler wire must exceed 22 wt%. Therefore, in the welding filler wire disclosed in the present invention, the manganese concentration is set in the range of about 23 to 34 wt%, preferably about 24 to 32 wt%. (2) Aluminum: Aluminum is one of the main elements that form (Fe, Mn)3AlC carbide (κ-carbide). As described in the prior art, κ-carbide is the most important and effective component in enhancing the properties of precipitation-hardening iron-manganese-aluminum-carbon alloys. Therefore, the aluminum content in the weld filler wire plays an important role in inducing the most unique properties of the present invention. That is, after fusion welding, in the post-welded state, high-density nano-order κ-carbide is present in the austenite dendritic crystal cells of the molten region. The present invention designed a series of weld filler wires with different aluminum concentrations and simultaneously conducted comprehensive observation and analysis of these weld filler wires. The results show that when the aluminum concentration is less than 7.0 wt%, κ-carbide is not formed in the austenite dendritic crystal cells of the molten region in the post-welded state. When the aluminum concentration in the welding filler wire increases to 7.5% by weight or more, high-density nano-order κ-carbides are readily observed in the austenite dendritic crystal cells of the molten region in the post-welded state. However, when the aluminum concentration in the welding filler wire reaches 12% by weight, in addition to the high-density nano-order κ-carbides formed in the austenite dendritic crystal cells and eutectic region, aluminum-rich particles appear at the austenite grain boundaries of the molten region. It is well known that the aluminum-rich secondary phase present at the austenite iron grain boundaries significantly reduces the ductility of the iron-manganese-aluminum-carbon alloy. In Example 8 of the present invention, it was further observed that aluminum-rich particles formed at the austenite iron grain boundaries of the molten region caused the formation of solidification cracks during the welding process. Clearly, the aluminum concentration of the welding filler wire of the present invention should be limited to the range of 7.5 to 11.5% by weight, preferably 8.0 to 11.0% by weight. (3) Carbon: Carbon is clearly one of the main elements that form iron-manganese-aluminum carbides ((Fe,Mn)3AlC,κ-carbides). Similar to aluminum as described above, the carbon content of the weld filler wire is also important in forming the most unique properties of the present invention. That is, after fusion welding, in the post-welded state, high-density nano-order κ-carbides are present in the austenite dendritic crystal cells of the molten region. The present invention designed a series of weld filler wires with different carbon concentrations and simultaneously conducted comprehensive observation and analysis of these weld filler wires. The results show that when the aluminum concentration is less than 1.30%, no κ-carbides are formed in the austenite dendritic crystal cells of the molten region in the post-welded state. As the carbon concentration in the weld filler wire increases, exceeding 1.35 wt%, high-density nano-order κ-carbides are readily observed in the austenite dendritic crystal cells of the molten region in the post-welded state. However, when the aluminum concentration in the welding filler wire exceeds 2.1 wt%, coarser κ-carbides appear at the austenite iron grain boundaries in the molten region, in addition to the high-density nano-order κ-carbides formed in the austenite dendritic crystal cells and eutectic regions. It is well known that the coarser κ-carbides present at the austenite iron grain boundaries and their associated precipitation-free zones have a serious adverse effect on the ductility of iron-manganese-aluminum-carbon alloys. In Example 7 of the present invention, in the case of a welding filler wire with a carbon content of 2.2 wt%, despite careful adjustment of parameters related to the welding process and the use of various different parameter combinations, the result was still that the alloy developed serious solidification high-temperature cracking in the weld bead. Therefore, the carbon content of the welding filler wire of the present invention should be limited to a range of about 1.35 to 1.95 wt%, preferably about 1.40 to 1.95 wt%. (4) Titanium, Niobium, and Vanadium: Titanium, niobium, and vanadium are very strong carbide-forming elements. The present invention also investigated the effect of adding these elements to the welding filler wire disclosed in the present invention on the microstructure and properties of the molten region of the welded member after welding. The results showed that by adding appropriate amounts of titanium, niobium, and vanadium to the welding filler wire, the microstructure of the molten region after welding had the following remarkable characteristics: (i) High-density nano-order κ-carbides are formed in the austenite dendritic crystal cells and their eutectic regions, forming titanium carbides rich in nano-order titanium, niobium carbides rich in niobium, and vanadium carbides rich in vanadium in the eutectic regions. (ii) The size of these eutectic carbides is only about 6 to 10 nanometers. (iii) All of these eutectic carbides have a ductile face-centered cubic structure, such as nano-order κ-carbides formed in austenite dendritic crystal cells and eutectic regions. (iv) The addition of these elements significantly reduces the size of the austenite dendritic crystal cells. (v) These titanium-rich titanium carbides, niobium-rich niobium carbides, and vanadium-rich vanadium carbides are all very hard (approximately 2000-3500 Vickers hardness), and the presence of these carbides brings about remarkable strengthening (hardening) of the molten region without significantly reducing ductility. It should be particularly emphasized that the most important feature of the present invention is the increase in yield strength and hardness of the molten region, and the condition for this feature to be met is the presence of high-density nano-order κ-carbides within the austenite dendritic crystal cells in the post-welded state. As described above, in order to obtain high-density nano-order κ-carbides formed in the austenite dendritic crystal cells of the molten region in the post-weld state, the carbon content of the welding filler wire must exceed 1.35% by weight.However, overall experimental results show that when approximately 1.0 wt% titanium is added to a welding filler wire with a carbon content of 1.35 wt%, κ-carbides are not formed within the austenite dendritic crystal cells after fusion welding in the post-welded state. In this case, to obtain high-density nano-order κ-carbides formed in the austenite dendritic crystal cells in the post-welded state, the carbon content in the welding filler wire must exceed 1.50 wt%. Further experiments have shown that to obtain high-density nano-order κ-carbides formed in the austenite dendritic crystal cells in the post-welded state, the carbon content in the welding filler wire must exceed 2.2 wt% when the total content of titanium, niobium, and vanadium exceeds 3.5 wt%. However, in that case, high-temperature cracking of the solidified bead is often observed. However, by using the novel alloy design of the present invention described above, the molten region of the resulting alloy has the expected unique microstructure and an excellent combination of microhardness, yield strength, ultimate tensile strength, and ductility. In other words, this demonstrates that this alloy design has unprecedented advantages, such as promoting mechanical strength, particularly yield strength, and improving the microhardness of the molten region in the post-welded state, while still retaining excellent malleability and ductility. (5) Chromium and Molybdenum: Chromium and molybdenum are also very strong carbide-forming elements. The present invention also investigated the effect of adding chromium and molybdenum to welding filler wire on the microstructure of the molten region in the post-weld state. The investigation found that when chromium and molybdenum are added to welding filler wire, coarse molybdenum-rich molybdenum carbide and chromium-rich chromium carbide, which are micrometer-sized, are formed at the austenite iron grain boundaries of the molten region after welding. At the same time, there are clearly visible non-deposit regions around the coarse carbides of the molten region, which adversely affect the ductility of the alloy after welding. Therefore, the present invention does not propose adding chromium and molybdenum to welding filler wire.

[0027] Finally, it is important to note that the base material used in the embodiments of the present invention is a hot-rolled, high-carbon-content (AS-Hot-Rolled) iron-manganese-aluminum-carbon alloy (with carbon exceeding 1.5% by weight). The microstructure of the base material before fusion welding consists of a complete austenitic iron base and nano-order κ-carbides with a high density of L'12 structures uniformly distributed within the austenitic base. Interestingly, after fusion welding, these high-density nano-order κ-carbides do not undergo significant changes in the weld heat-affected zone of the welded member or in the base material region. In other words, these high-density nano-order κ-carbides do not undergo significant dissolution or coarsening. In other precipitation-hardening alloys (e.g., AA7075), dissolution or coarsening of carbides is often observed after fusion welding. Therefore, as described in the embodiments disclosed below, no significant softening phenomenon occurs in the weld heat-affected zone of the alloy. [Effects of the Invention]

[0028] The precipitation-hardening austenitic iron-manganese-aluminum-carbon alloy welding filler wire of the present invention can solve the problems of softening of the molten zone and high-temperature cracking. [Brief explanation of the drawing]

[0029] [Figure 1(a)] This is a bright-field transmission electron microscope image of a hot-rolled iron-28.5 manganese-9.0 aluminum-1.55 carbon matrix (Fe-28.5Mn-9.0Al-1.55C). [Figure 1(b)] The electron diffraction pattern of a selected region in a transmission electron microscope image of a hot-rolled iron-28.5 manganese-9.0 aluminum-1.55 carbon matrix (Fe-28.5Mn-9.0Al-1.55C), with the region axis

[0001] , (hkl: austenitic iron; hkl: κ-carbide). [Figure 1(c)] This is a (100)κ dark-field image acquired from the same region as shown in Figure (a). [Figure 1(d)]This is a macroscopic image of a weld bead obtained by a carbon gas tungsten arc welding process using a welding filler wire with the composition iron-27,8 manganese-9,1 aluminum-1.86 (Fe-27,8Mn-9,1Al-1.86C). [Figure 1(e)] Scanning electron microscope images of welded components after welding (FZ: molten zone; HAZ: heat-affected zone; BM: base metal). The inset shows a magnified image of the marked area in the figure. [Figure 1(f)] The images show scanning electron microscope images of the molten region, the heat-affected zone of the weld, and the microhardness indentation applied to the base material. [Figure 1(g)] Figure (f) shows the corresponding minute hardness values ​​measured at each measurement point. [Figure 1(h)] This is a bright-field transmission electron microscope image obtained from the molten region of a welded component. The solid lines and white arrows indicate austenite dendritic crystal cells and κ-carbides in the eutectic region, respectively. [Figure 1(i)] This is an electron diffraction pattern of a selected region obtained from the molten region, with the region axis being

[0001] (hkl: austenitic iron; hkl: κ-carbide). [Figure 1(j)] Analysis results obtained from the molten region of welded components using a transmission electron microscope and X-ray energy dispersive analyzer (TEM-EDS). [Figure 1(k)] Macro images of the tensile test sample before and after the tensile test are shown. [Figure 2(a)] This image shows a macroscopic view of a weld bead obtained using gas metal arc welding and iron-28,2-manganese-9,2-aluminum-1.72-carbon (Fe-28,2Mn-9,2Al-1.72C) welding filler wire. [Figure 2(b)] The image shows scanning electron microscope images (FZ: molten zone, HAZ: heat-affected zone, BM: base metal) acquired from the welded member after welding. The embedded diagram is a magnified image of the area marked in the figure. [Figure 2(c)]These are scanning electron microscope images showing the molten region, the heat-affected zone of the weld, and the micro-hardness test indentations of the base material. [Figure 2(d)] These are the corresponding minute hardness values ​​measured at each measurement point in Figure (c). [Figure 2(e)] This is a bright-field transmission electron microscope image obtained from the molten region of a welded component. The solid lines and white arrows indicate κ-carbides dispersed in the austenite dendritic crystal cells and eutectic regions, respectively. [Figure 2(f)] This is an electron diffraction pattern of a selected region obtained from the molten region, with the region axis being

[0001] , (hkl: austenitic iron hkl: κ-carbide)). [Figure 2(g)] These are the results of analysis using a penetrant electron microscope and X-ray energy dispersion analyzer obtained from the molten region of the welded component. [Figure 2(h)] These are macroscopic images of tensile samples before and after tensile testing. [Figure 3(a)] This is a bright-field transmission electron microscope image of a hot-rolled iron-28.8 manganese-8.9 aluminum-1.62 carbon (Fe-28.8Mn-8.9Al-1.62C) matrix. [Figure 3(b)] This is the electron diffraction pattern of a selected region in a transmission electron microscope image of a hot-rolled iron-28,8-manganese-8,9-aluminum-1,62-carbon (Fe-28,8Mn-8,9Al-1,62C) matrix, with the region axis being

[0001] , (hkl: austenitic iron, hkl: κ-carbide). [Figure 3(c)] This is a (100)κ dark-field image acquired from the same region as shown in Figure 3(a). [Figure 3(d)] The images are scanning electron microscope images (FZ: molten zone, HAZ: heat-affected zone, BM: base metal) acquired from the welded member after welding. The welding was performed by gas tungsten arc welding, using iron-28,2-manganese-10,6-aluminum-1.62-carbon (Fe-28.2Mn-10.6Al-1.62C) welding filler wire. [Figure 3(e)] This is a scanning electron microscope image showing the molten region, the heat-affected zone of the weld, and the microhardness test indentations performed on the base material. [Figure 3(f)] These are the corresponding minute hardness values ​​measured at each measurement point in Figure 3(e). [Figure 3(g)] This is a bright-field transmission electron microscope image obtained from the molten region of a welded component. The solid lines and white arrows indicate κ-carbides dispersed in the austenite dendritic crystal cells and eutectic regions, respectively. [Figure 3(h)] This is an electron diffraction pattern of a selected region obtained from the molten region, with the region axis being

[0001] (hkl: austenitic iron, hkl: κ-carbide). [Figure 3(i)] These are the results of analysis using a penetrant electron microscope and X-ray energy dispersion analyzer, obtained from the molten region of the welded component. [Figure 4(a)] Scanning electron microscope images (FZ: molten zone, HAZ: heat-affected zone, BM: base metal) were acquired from the welded member in the post-weld state. The welding method used was gas tungsten arc welding, and iron-29.5 manganese-9.8 aluminum-1.48 carbon (Fe-29.5Mn-9.8Al-1.48C) welding filler wire. [Figure 4(b)] This is a scanning electron microscope image showing the molten region, the heat-affected zone of the weld, and the microhardness test indentations performed on the base material. [Figure 4(c)] These are the corresponding minute hardness values ​​measured at each measurement point in Figure 4(b). [Figure 4(d)] These are transmission electron microscope and bright-field images obtained from the molten region of a welded component. [Figure 4(e)] This is an electron diffraction pattern of a selected region obtained from the molten region, with the region axis being 001 (hkl: austenitic iron, hkl: κ-carbide). [Figure 4(f)] These are the analysis results obtained from the molten region of the welded component using a penetrant electron microscope and X-ray energy dispersion analyzer. [Figure 5(a)] Scanning electron microscope images (FZ: molten zone, HAZ: heat-affected zone, BM: base metal) were acquired from the welded member in the post-weld state. The welding method used was gas tungsten arc welding, and iron-33.8 manganese-10.6 aluminum-1.38 carbon (Fe-33.8Mn-10.6Al-1.38C) welding filler wire. [Figure 5(b)] This is a scanning electron microscope image showing the molten region, the heat-affected zone of the weld, and the microhardness test indentations performed on the base material. [Figure 5(c)] These are the corresponding minute hardness values ​​measured at each measurement point in Figure 4(b). [Figure 5(d)] This is a bright-field image obtained from the molten region of a welded component using a transmission electron microscope. [Figure 5(e)] This is an electron diffraction pattern of a selected region obtained from the molten region, with the region axis being

[0001] (hkl: austenitic iron, hkl: κ-carbide). [Figure 5(f)] These are the analysis results obtained from the molten region of the welded component using a penetrant electron microscope and X-ray energy dispersion analyzer. [Figure 6(a)] These are scanning electron microscope images (FZ: molten zone, HAZ: heat-affected zone, BM: base metal) acquired from the welded member in the post-weld state. The welding was performed using gas tungsten arc welding and iron-31.8Mn-9.2Al-1.25C (Fe-31.8Mn-9.2Al-1.25C) welding filler wire. The inset image shows a magnified view of the marked area within the image. [Figure 6(b)] This is a scanning electron microscope image showing the molten region, the heat-affected zone of the weld, and the microhardness test indentations performed on the base material. [Figure 6(c)] These are the corresponding minute hardness values ​​measured at each measurement point in Figure 6(b). [Figure 6(d)] These are macroscopic images of tensile samples before and after tensile testing. [Figure 7(a)] This image shows a macroscopic view of the weld bead after welding using gas tungsten arc welding and iron-30.5 manganese-8.9 aluminum-2.2 carbon (Fe-30.5Mn-8.9Al-2.2C) welding filler wire. Note the solidification cracks in the weld bead. [Figure 7(b)]The image shows scanning electron microscope images (FZ: molten zone, HAZ: heat-affected zone, BM: base metal) acquired from the welded member in the post-weld state. The inset image shows a magnified view of the marked area within the image. The arrows point to coarse κ-carbides formed at the austenite iron grain boundaries. [Figure 7(c)] This is a bright-field image obtained from the molten region of a welded component using a transmission electron microscope. [Figure 7(d)] The electron diffraction pattern type for the selected region was obtained from the coarse κ-carbide marked "K" in Figure 7(c), with the region axis being

[0001] , (hkl: austenite hkl: κ-carbide). [Figure 7(e)] These are the results of transmission electron microscopy and X-ray energy dispersion analysis of coarse κ-carbides. [Figure 8(a)] This image shows a macroscopic view of the weld bead after welding using gas tungsten arc welding and iron-30.5 manganese-12.5 aluminum-1.55 carbon (Fe-30.5Mn-12.5Al-1.55C) welding filler wire. Note the solidification cracks in the weld bead (indicated by the arrows). [Figure 8(b)] The image shows scanning electron microscope images (FZ: molten zone, HAZ: heat-affected zone, BM: base metal) acquired from the welded material after welding. The arrows point to aluminum-rich particles formed at the austenite iron grain boundaries (indicated by α). [Figure 8(c)] These are the results of scanning electron microscopy and X-ray energy dispersive spectroscopy (SEM-EDS) analysis of discrete aluminum-rich particles. [Figure 9(a)] These are scanning electron microscope images (FZ: molten region, HAZ: heat-affected zone, BM: base metal) acquired from the welded member after welding. The welding was performed using gas tungsten arc welding and iron-20,1 manganese-8,5 aluminum-1.52 carbon (Fe-20,1Mn-8.5 Al-1.52C) welding filler wire. The arrows indicate the discrete ferrite phase dispersed in the molten region (indicated by α). [Figure 9(b)]This is a bright-field image obtained from a transmission electron microscope in the molten region of a welded component. Discrete fatty iron phases are represented by α, and the circles indicate coarse κ-carbides formed at the ferrite / austenite boundary. [Figure 9(c)] The selected region electron diffraction pattern is shown in the circled area of ​​Figure 9(b) (hkl: austenitic iron, hkl: κ-carbide, (hkl): ferrite). [Figure 9(d)] These are the analysis results obtained from austenite regions using a transmission electron microscope and X-ray energy dispersion analyzer. [Figure 9(e)] These are the analysis results obtained from a transmission electron microscope and X-ray energy dispersive analyzer from the ferrite region. [Figure 10(a)] This is a bright-field transmission electron microscope image of a hot-rolled iron-29,3-manganese-9,3-aluminum-1.75-carbon (Fe-29,3Mn-9,3Al-1.75C) matrix. [Figure 10(b)] This is the electron diffraction pattern of a selected region in a transmission electron microscope image of a hot-rolled iron-29,3-manganese-9,3-aluminum-1.75-carbon (Fe-29,3Mn-9,3Al-1.75C) matrix, with the region axis being

[0001] , (hkl: austenitic iron, hkl: κ-carbide). [Figure 10(c)] This is a (100)κ dark-field image of the same region as shown in Figure 10(a). [Figure 10(d)] These are macroscopic images of weld beads obtained using gas tungsten arc welding and iron-28.5 manganese-9.3 aluminum-1.82 carbon-1.6 titanium (Fe-28.5Mn-9.3Al-1.82C-1.6Ti) welding filler wire. [Figure 10(e)] The images show scanning electron microscope images acquired from the welded material after welding (FZ: molten zone, HAZ: heat-affected zone, BM: base metal). [Figure 10(f)] This is a scanning electron microscope image showing the molten region, the heat-affected zone of the weld, and the microhardness test indentations performed on the base material. [Figure 10(g)] These are the corresponding minute hardness values ​​measured at each measurement point in Figure 10(f). [Figure 10(h)] This is a bright-field image obtained from the molten region of a welded component using a transmission electron microscope. The arrows point to κ-carbides and titanium-rich titanium carbides (Ti-rich Ti-carbides) dispersed in the eutectic region. [Figure 10(i)] The selected region electron diffraction pattern in the circled area of ​​Figure 10(h) is shown (hkl: austenite, hkl: κ-carbide, (hkl): titanium carbide). [Figure 10(j)] Figure 10(h) shows the analysis results obtained from the same region circled in the image, using both a transmission electron microscope and an X-ray energy dispersion analyzer. [Figure 10(k)] These are macroscopic images of tensile samples before and after tensile testing. [Figure 11(a)] These are macroscopic images of weld beads obtained by welding using gas metal arc welding and iron-29.5 manganese-8.2 aluminum-1.56 carbon-1.0 niobium (Fe-29.5Mn-8.2Al-1.56C-1.0Nb) welding filler wire. [Figure 11(b)] The images show scanning electron microscope images acquired from the welded material after welding (FZ: molten zone, HAZ: heat-affected zone, BM: base metal). [Figure 11(c)] This is a scanning electron microscope image showing the molten region, the heat-affected zone of the weld, and the microhardness test indentations performed on the base material. [Figure 11(d)] These are the corresponding minute hardness values ​​measured at each measurement point in Figure 11(c). [Figure 11(e)] These are transmission electron microscope (BEM) images taken from the molten region of a welded component. The arrows point to κ-carbides and niobium-rich niobium carbides (Nb-rich Nb-carbides) dispersed in the eutectic region. [Figure 11(f)] The selected region electron diffraction pattern (hkl: austenitic iron, hkl: κ-carbide, (hkl): niobium carbide) was obtained from the circled area in Figure 11(e). [Figure 11(g)] Figure 11(e) shows the analysis results obtained from the same region circled in the image, using both a transmission electron microscope and an X-ray energy dispersion analyzer. [Figure 11(h)] These are macroscopic images of tensile samples before and after tensile testing. [Figure 12(a)] These are macroscopic images of weld beads obtained by welding using gas tungsten arc welding and iron-31.2 manganese-9.8 aluminum-1.65 carbon-1.2 vanadium (Fe-31.2Mn-9.8Al-1.65C-1.2V) welding filler wire. [Figure 12(b)] This image shows a scanning electron microscope image (FZ: molten zone, HAZ: heat-affected zone, BM: base metal) obtained from a welded component after welding. [Figure 12(c)] This is a scanning electron microscope image showing the molten region, the heat-affected zone of the weld, and the microhardness test indentations performed on the base material. [Figure 12(d)] These are the corresponding minute hardness values ​​measured at each measurement point in Figure 12(c). [Figure 12(e)] This is a bright-field image obtained from the molten region of a welded component using a transmission electron microscope. The arrows point to κ-carbides and vanadium-rich vanadium carbides (V-rich V-carbides) dispersed in the eutectic region. [Figure 12(f)] Selective region electron diffraction pattern of the circled area in Figure 12(e) (hkl: austenitic iron, hkl: κ-carbide, (hkl): vanadium carbide). [Figure 12(g)] Figure 12(e) shows the analysis results obtained from the same region circled in the image, using both a transmission electron microscope and an X-ray energy dispersion analyzer. [Figure 12(h)] These are macroscopic images of tensile samples before and after tensile testing. [Figure 13(a)] These are macroscopic images of weld beads obtained by performing gas tungsten arc welding and welding using iron-30.6 manganese-9.2 aluminum-1.88 carbon-0.8 titanium-1.0 niobium (Fe-30.6Mn-9.2Al-1.88C-0.8Ti-1.0Nb) welding filler wire. [Figure 13(b)] The image shows a scanning electron microscope image obtained from a welded component after welding (FZ: molten zone, HAZ: heat-affected zone, BM: base metal). [Figure 13(c)] This is a scanning electron microscope image showing the molten region, the heat-affected zone of the weld, and the microhardness test indentations performed on the base material. [Figure 13(d)] These are the corresponding minute hardness values ​​measured at each measurement point in Figure 13(c). [Figure 13(e)] This is a bright-field image obtained from the molten region of a welded component using a transmission electron microscope. The arrows point to κ-carbides, titanium-rich titanium carbides, and niobium-rich niobium carbides dispersed in the eutectic region. [Figure 13(f)] Figure 13(e) shows the analysis results obtained from the same region as the selected area using a transmission electron microscope and X-ray energy dispersion analyzer. [Figure 13(g)] These are macroscopic images of tensile samples before and after tensile testing. [Figure 14(a)] The images show scanning electron microscope images (FZ: molten zone; HAZ: heat-affected zone; BM: base metal) obtained from welded members after welding. The welding was performed using gas tungsten arc welding and iron-28,2-manganese-10,6-aluminum-1.62-carbon (Fe-28.2Mn-10.6Al-1.62C) welding filler wire. Note that the base metal underwent hot rolling, solution heat treatment, and quenching, followed by aging at 550°C for 12 hours. [Figure 14(b)] This is a scanning electron microscope image showing the molten region, the heat-affected zone of the weld, and the microhardness test indentations performed on the base material. [Figure 14(c)] These are the corresponding minute hardness values ​​measured at each measurement point in Figure 14(b). [Figure 14(d)] These are macroscopic images of tensile samples before and after tensile testing. [Figure 15(a)]The images show scanning electron microscope images (FZ: molten zone; HAZ: heat-affected zone; BM: base metal) obtained from welded parts after welding. The welding was performed using gas tungsten arc welding and iron-28,2-manganese-9,1-aluminum-1.55-carbon-1.8-molybdenum (Fe-28.2Mn-9.1Al-1.55C-1.8Mo) welding filler wire. The arrows point to coarse molybdenum-rich molybdenum carbides formed at the austenite iron grain boundaries. [Figure 15(b)] These are the results of scanning electron microscopy and X-ray energy dispersive analysis of molybdenum carbides, which are rich in coarse molybdenum and formed at the austenite iron grain boundaries. [Figure 16(a)] The images show scanning electron microscope images (FZ: molten zone; HAZ: heat-affected zone; BM: base metal) obtained from welded members after welding. The welding was performed using gas tungsten arc welding and iron-29,3 manganese-8.8 aluminum-1.58 carbon-1.5 chromium (Fe-29,3Mn-8.8Al-1.58C-1.5Cr) welding filler wire. The arrows point to coarse chromium-rich carbides (Cr-rich chromium carbides) formed at the austenite iron grain boundaries. [Figure 16(b)] These are the results of scanning electron microscopy and X-ray energy dispersion analysis of coarse, chromium-rich chromium carbides formed at austenite iron grain boundaries. [Modes for carrying out the invention]

[0030] The structure and technical features of the present invention will be described in conjunction with the drawings, but each drawing is used solely to illustrate the structural relationships and related functions of the present invention, and the dimensions of the members in each drawing are not based on actual proportions and do not limit the present invention.

[0031] In the following, selected preferred embodiments will be described in detail to illustrate the innovative properties of the present invention. Unless otherwise specified, the iron-manganese-aluminum-carbon matrix materials with different compositional components investigated in the present invention are produced in an atmospheric induction furnace. After homogenization at 1150°C for 2 hours in a protective argon atmosphere, the metal ingot is hot-rolled from 80 mm thick to plates 8-12 mm thick, quenched in water, and cooled to room temperature. The hot-rolled plates are then machined to form single V-shaped joints, which are used as welding joints with a groove angle of 60 degrees. The welding filler wire is prepared in an air induction furnace based on the designed alloy compositional components. The molten metal is cast to form a steel mold measuring 80 mm × 80 mm × 1500 mm. Next, the metal ingot is heated to 1150°C for 2 hours, then hot-rolled into a coil with a diameter of approximately 5 mm, followed by pickling and cold drawing to form weld filler wires with diameters of approximately 1.2, 2.4, and 3.2 mm, respectively. Fusion welding is then performed using gas metal arc welding (GMAW) or gas tungsten arc welding (GTAW), followed by shielding gas welding. Pure argon or a mixture of argon and helium (He) is used as the shielding gas. Typical flow rates of the shielding gas for gas tungsten arc welding and gas metal arc welding are approximately 10 and 18 liters / min, respectively. In the gas metal arc welding process, the welding filler wire addition rate is approximately 200 mm / min, and the voltage and current used are 22-26 volts (V) and 140-170 amperes (A), respectively. In the gas tungsten arc welding process, the voltage and current used are 15-20 volts and 125-200 amperes, respectively, and are determined according to the diameter of the welding filler wire used. In all embodiments, the fusion welding is performed with a root opening of approximately 1 mm.The microstructure of welded components is analyzed using optical microscopes (OM), scanning electron microscopes (SEM), and transmission electron microscopes (TEM), respectively. For scanning electron microscopy, samples are mechanically polished before etching with 5% nitric acid. For transmission electron microscopy, film samples are prepared using a twin-jet electropolisher, with the electrolyte composition being 60% acetic acid, 30% ethanol, and 10% perchloric acid. If required, the composition of the alloy or sample is analyzed using an energy dispersive spectrometer (EDS) attached to the scanning electron microscope or transmission electron microscope. It should be noted that quantitative analysis of elements with atomic numbers less than 10 (such as carbon) using an energy dispersive spectrometer often yields unreliable results and is therefore frequently ignored. The sample for tensile testing was prepared according to the U.S. Service of Testing and Materials (ASTM) ASTME8 standard (Standard Test Method for Tensile Testing of Metallic Materials), with a gauge length of 25 mm, a gauge width of 6.25 mm, and a thickness of 4 mm. The tensile test was performed at room temperature using an Instron 8501 universal testing machine, at 6.7 × 10⁻⁶. -4 s -1 The test was conducted at the specified strain rate. Tensile strength was measured with an offset strain of 0.2%. Vickers microhardness measurements were used to summarize the characteristics of the hardness of each different molten region of the welded member, e.g., the molten region, the heat-affected zone of the weld, and the base metal. In the Vickers hardness test, a load of 100 grams (gf) was applied for a load duration of 15 seconds.

[0032] <Example 1> Figure 1(a) is a bright-field (BF) image of the iron-28.5Mn-9.0Al-1.55C (Fe-28.5Mn-9.0Al-1.55C) matrix after hot rolling, observed using a transmission electron microscope. It shows that high-density nano-order (approximately 3-5 nanometers) precipitates are uniformly distributed across the base. As indicated by the arrows in Figure 1(a), the precipitates at dislocations are slightly larger (approximately 5-8 nanometers), which means that these precipitates are formed during the hot rolling process. Figure 1(b) shows the matrix after hot rolling, and its selected-area diffraction pattern (SADP) clearly indicates that these high-density nano-order precipitates are L'12 structured iron-manganese-aluminum carbides ((Fe,Mn)3AlC, κ-carbides). Figure 1(c) shows the same region as Figure 1(a) (100) κ This is a dark-field (DF) image showing the austenitic iron base and the κ-carbides appearing in the transition. This figure clearly shows that the hot-rolled microstructure of the base material used in this embodiment is single-phase austenitic iron with high-density nano-order κ-carbides uniformly dispersed therein. Tensile tests revealed that the yield strength, ultimate tensile strength, and elongation of the hot-rolled base material were 952 megapascals, 1100 megapascals, and 56%, respectively.

[0033] Two hot-rolled base plate pieces, each measuring 80mm x 80mm x 8mm, are machined to form a single V-groove butt weld. In this embodiment, the nominal composition of the weld filler wire is iron-27.8 manganese-9.1 aluminum-1.86 carbon (Fe-27.8Mn-9.1Al-1.86C), with a diameter (Φ) of approximately 2.4mm. Fusion welding was performed using a gas tungsten arc welding process, with pure argon used as the shielding gas. Figure 1(d) is a macroscale image of the weld bead, showing that there are no visible solidification cracks or holes in the molten region of the welded sample.

[0034] Figure 1(e) is a typical scanning electron microscope image showing a region covering the molten area, the heat-affected zone of the weld, and the base metal within a welded sample. In this figure, several properties that are quite different from conventional fusion-welded precipitation-hardening iron-manganese-aluminum-carbon alloys (carbon content of 1.29 wt.% or less, C ≤ 1.29 wt.%) can be observed. (i) It is clear that the molten area contains typical columnar austenite iron grains consisting of austenite dendritic crystal cells and eutectic regions. More importantly, high-density nano-order κ carbides are present in the austenite dendritic crystal cells and eutectic regions. This is one of many remarkable features revealed for the first time by the present invention. Conventional techniques have shown that fusion-welded iron-manganese-aluminum-carbon alloys (carbon content of 1.29 wt.% or less) show no precipitates in either primary or secondary austenite dendritic crystals, with only small amounts of eutectic κ carbides present in the eutectic regions. Furthermore, in this example, the amount of nano-order κ-carbides present in the eutectic region is far greater than the amount of carbides observed in previous studies in alloys with a carbon content of 1.29 weight percent or less. (ii) From Figure 1(e), it is clear that the columnar austenite iron grains do not have any obvious long primary dendrites. The length and spacing of the austenite dendritic crystal cells formed in the columnar austenite iron grains are approximately 20-30 micrometers and approximately 5-10 micrometers, respectively. Clearly, the austenite dendritic crystal cells in this example are significantly refined. This is quite different from what can be observed anywhere in known or conventional fusion welding. In this context, the microstructure of the weld bead after welding is always composed of clearly visible primary dendritic crystals, with the length of the dendritic crystals reaching several hundred micrometers (e.g., 300 micrometers or more). Many of the factors that contribute to these differences are explained below. As shown in the magnified image of the marked area in the upper right corner of Figure 1(e), it is clear that precipitates are present not only at the tips of each austenite dendritic crystal cell, but also at the fine eutectic κ carbides formed between the austenite dendritic crystals.During the solidification period after fusion welding, the stability of the solid-liquid interface depends on the thermal state and constitutional supercooling near the interface, and the size of the dendritic crystals in the weld bead depends mainly on the cooling rate and alloy composition. Under these conditions, precipitates appearing at the tips of austenite dendritic crystal cells due to compositional supercooling appear to slow the solidification front and effectively hinder the growth of dendritic crystal cells, thus resulting in the appearance of much shorter dendritic crystal cells disclosed in Figure 1(e). Therefore, the composition of the weld filler wire is considered to be a more important factor in the short and thin dendritic structures observed in this example. (iii) Although the molten region in this example is a complete austenite iron phase, in the prior art, after autogenous fusion welding, even if the base material was complete austenite before welding, a certain amount of ferrite phase is often included in the solidified microstructure. Furthermore, Figure 1(e) shows that no significant microstructural changes were observed between the heat-affected zone (HZ) and the base metal region. Also, in this example, there were no discernible differences in the size and density of κ-carbides between the molten region, the HZ, and the base metal region. Therefore, there was no softening phenomenon in the HZ. It should also be noted that no traces of microcracks or porosity were observed in the molten region or near the boundary between the molten and HZ. This indicates that solidification and liquefaction cracks can be significantly eliminated by using the molten welding filler wire of the present invention, both of which are common problems in the melting process of precipitation-hardening alloys. Figure 1(f) shows the microhardness measurements performed in the molten region, the HZ, and the base metal region. The corresponding values ​​measured at each measurement point are shown in Figure 1(g). As shown in Figure 1(g), the average microhardness of the molten region (approximately 425 Vickers hardness) is actually slightly greater than the average microhardness of the weld heat-affected zone (approximately 417 Vickers hardness) and the base metal area (approximately 412 Vickers hardness). This is in stark contrast to the common belief that the molten region is always the weakest region after fusion welding.It is worth mentioning here that the microstructure of the base material region used in this example after hot rolling is a complete austenitic iron phase, uniformly dispersed in an austenitic iron base, with high density of nano-order κ-carbides, and the microhardness of the base material region is significantly greater than the hardness (350-400 Vickers hardness) of iron-manganese-aluminum-carbon alloys with optimized aging treatment and a carbon content of 1.2 wt% or less, as described in the prior art.

[0035] To fully understand the main factors leading to these unprecedented results, a more detailed analysis of the microstructure of the molten region is necessary. Transmission electron microscopy in this example shows that the base of the molten region is a complete austenite iron phase, and that both the austenite dendritic crystal cells and the nano-order precipitates in the eutectic region are κ-carbides. An example shown in Figure 1(h) is a bright-field electron micrograph (BF) acquired from the molten region. In this example, the κ-carbides of the austenite dendritic crystal cells (indicated by black arrows) and the associated κ-carbides in the eutectic region (indicated by white arrows) are clearly visible, with sizes of 3–5 nanometers and 6–10 nanometers, respectively. The length of the austenite dendritic crystal cells is approximately 100–300 nanometers. This indicates that the alloy design of the weld filler wire of the present invention can effectively suppress the growth rate of the front of the dendritic crystal cells during the solidification process after fusion welding. As shown in Figure 1(i), the electron diffraction pattern of the selected region confirms that only austenite iron phase and κ-carbide are present in the molten region of the welded member. Furthermore, as shown in Figure 1(j), analysis results obtained from the molten region using a transmission electron microscope and X-ray energy dispersive analyzer (TEM-EDS) showed that the compositional components of the molten region, iron, manganese, and aluminum, were 64.3%, 26.7%, and 9.0%, respectively.

[0036] Tensile tests demonstrate that welded members obtained using the weld filler wire disclosed in this example exhibit excellent ductile deformation behavior, with yield strength, ultimate tensile strength, and elongation of approximately 946 megapascals, 1086 megapascals, and 44%, respectively. Figure 1(k) shows macroscopic images of the sample before and after the tensile test. Clearly, the welded member as a whole exhibits typical characteristics of ductile plastic deformation. In particular, the molten region also exhibits typical characteristics of plastic deformation, with a corrugated surface and visible elongation. Because the tensile sample contains regions with various different mechanical properties (i.e., the molten region and the base metal region), it is usually difficult to directly infer the strengths of the molten and base metal regions from the measured yield strength. Nevertheless, it may be possible to distinguish the relative yield strength between the molten and base metal regions by judging from the change in width of each region of the sample fractured by the tensile test. For example, if the yield strength of the base metal region is high, and the tensile stress exceeds the yield strength of the molten region but is still lower than the yield strength of the base metal region, then the molten region will first undergo plastic deformation (e.g., becoming longer and thinner), and simultaneously experience a reduction in width. However, the base metal region remains within its elastic regime, so its width does not change, and vice versa. Therefore, if the yield strengths of the molten region and the base metal region differ significantly, there will be some change in width in each region where the sample breaks after a tensile test, and the edges of the sample will not be parallel. In contrast, if the widths of each region of the broken sample are approximately the same after a tensile test, and the edges of the sample remain basically parallel, it can be reasonably inferred that the molten region and the base metal region have approximately the same yield strength. In fact, the width of the entire region including the base metal region, the weld heat-affected zone, and the molten region (in Figure 1(k), d) BM d INT , and d FZThe (marked as) regions are essentially the same, and the edges of the sample in this region remain parallel, indicating that the yield strength of the molten region and the base material region of the tensile test sample are nearly identical. It should be particularly emphasized that the exceptional yield strength present in the molten region in this example is virtually entirely expected. It is well known that the presence of high-density nano-order κ-carbides significantly improves the yield strength of a material without significantly reducing its ductility, which is also a major reason why many precipitation-hardening alloys have excellent yield strength and elongation combinations. This is precisely the phenomenon seen in Figure 1(h). On the other hand, in Figure 1(k), it is clear that necking and final fracture (indicated by the white arrow) have occurred in the region opposite the base material region, and the ultimate tensile strength of the molten region is greater than that of the base material region, which may be consistent with the phenomenon of slightly lower hardness in the base material region, as explained in Figure 1(g).

[0037] In fact, the high-density nano-order κ-carbides present in the austenite dendritic crystal cells are considered to be the main factor contributing to the high hardness and strength (especially yield strength) obtained in the molten region of this example, while still maintaining high malleability. The microstructure of the molten region described in this example is undoubtedly the most unique and unprecedented property of the present invention, and this property stems from the novel compositional design of the weld filler wire disclosed by the present invention.

[0038] <Example 2> The purpose of this example is to clarify how changing the carbon content in the welding filler wire affects the characteristics of the molten region of a sample in the post-welded state. The base material used in this example is the same as that used in Example 1. However, in this example, a gas metal arc welding process is performed. The base plate size is 80 mm × 80 mm × 12 mm, and a single V-groove butt joint is formed after machining. The gas metal arc welding process is performed using a voltage of 26 volts and a current of 140-170 amperes. The flow rate of the shielding gas (25% argon + 75% helium) is approximately 18 liters / min, and the welding filler wire addition rate is approximately 200 mm / min. The standard composition of the welding filler wire used in this example is iron-28,2 manganese-9,2 aluminum-1.72 carbon (Fe-28,2Mn-9,2Al-1.72C), with a diameter of approximately 1.2 mm.

[0039] Figure 2(a) shows a macroscopic image of the weld bead obtained by gas metal arc welding. From Figure 2(a), it can be seen that the weld bead has a very smooth shape and there are no macrocracks or pores visible to the naked eye throughout the weld bead. Figure 2(b) shows a scanning electron microscope image of the region including the molten region, the heat-affected zone, and the base metal region of the weld sample. Similar to the phase shown in Example 1, the molten region is composed of typical columnar austenite iron grains, and at the same time, there are high-density nano-order κ-carbides in the austenite dendritic crystal cells and the eutectic region. As shown by the arrow in Figure 2(b) and the magnified image in the upper right corner of Figure 2(b), in addition to the fine eutectic κ-carbides formed in the eutectic region, there are κ-carbides present at the tips of each austenite dendritic crystal cell. This indicates that the microstructure of this molten region has properties very similar to those observed in Example 1. Furthermore, it should be noted that no traces of microcracks or pores were observed in the molten region or near the boundary between the molten region and the heat-affected zone. Similar to Example 1, Figure 2(b) also shows the fusion welding process, which does not appear to cause the dissolution and significant roughening of the κ carbide in the heat-affected zone of the weld in this example.

[0040] Figure 2(c) shows Vickers microhardness tests performed in the molten region, the heat-affected zone of welding, and the base metal region. The corresponding microhardness values ​​measured at each measurement point are shown in Figure 2(d). The average microhardness of the molten region, the heat-affected zone of welding, and the base metal region are approximately 410 Vickers, 412 Vickers, and 413 Vickers, respectively. Clearly, the microhardness of the molten region and the heat-affected zone of welding is almost the same as that of the base metal region, indicating that no significant softening occurs in the molten region and the heat-affected zone of welding. This is consistent with the unique microstructure shown in Figure 2(b), where the high-density nano-order κ-carbides, the most important reinforcing component, remain essentially intact in these three regions after fusion welding.

[0041] Examination of the molten region using a transmission electron microscope revealed that a large amount of κ-carbides were formed in the austenite dendritic crystal cells and the eutectic region. Figure 2(e) shows a typical example. In this transmission electron microscope bright-field image, it is clear that high-density nano-order κ-carbides (approximately 3-5 nm) are dispersed in the austenite dendritic crystal cells (indicated by black arrows) and the eutectic region (indicated by white arrows). Compared to the results described in Example 1 (Figure 1(h)), the density of κ-carbides formed in the austenite dendritic crystal cells and the eutectic region is slightly reduced. Figure 2(f) is the electron diffraction pattern of the selected region obtained from Figure 2(e), confirming that the base of the molten region is a single austenite iron phase, and that the precipitates present in both the austenite dendritic crystal cells and the eutectic region are κ-carbides with a regular L'12 structure. Figure 2(g) shows the results of transmission electron microscopy and X-ray energy dispersive analysis obtained from the molten region, clearly indicating that the iron, manganese, and aluminum components of the molten region were approximately 63.4%, 27.5%, and 9.1%, respectively.

[0042] In the tensile test, it can be seen that the welded member welded with the welding filler wire disclosed in this example also has excellent ductile deformation properties, and the yield strength, ultimate tensile strength and elongation are about 942 megapascals, 1075 megapascals and 36% respectively. Figure 2(h) shows macroscopic images of the sample before and after the tensile test. Obviously, the entire welded part has the typical characteristic of ductile plastic deformation. In particular, a zig-zag weaving fracture pattern and a corrugated fracture surface appear in the molten region, which are characteristics of ductile deformation. Furthermore, interestingly, as described in Example 1, the widths of the regions including the base material region, the molten region, and the interface therebetween (respectively d BM , d FZ , and d INT shown) are substantially the same, and it can be observed that the sample remains parallel at the edges of this region, which indicates that the molten region and the base material region of the tensile test sample have substantially the same yield strength. The main reason for the extremely high yield strength of the molten region is as described in Example 1. However, in this case, unlike the phenomenon observed in Example 1, slight necking and cracks occur in the molten region rather than the base material region, which means that the molten region has a slightly lower ultimate tensile strength compared with the base material region. In fact, this phenomenon and another phenomenon are consistent only in that the microhardness value of the molten region is only slightly lower (about 410 Vickers hardness) compared with that of the base material region (about 410 Vickers hardness). Based on the results shown in Example 1 and Example 2, samples welded using the welding filler wire having the composition design disclosed in the present invention still all have obviously excellent ductility, extraordinary yield strength and microhardness after undergoing remelting, redistribution of alloying elements and solidification during fusion welding.

[0043] <Example 3> In this example, we investigated how a slightly lower carbon content affects the properties of the molten region of a sample after welding. In this example, a base material with a standard composition of iron-28.8Mn-8.9Al-1.62C (Fe-28.8Mn-8.9Al-1.62C) was used after hot rolling. Figure 3(a) is a bright-field transmission electron microscope image of the base material after hot rolling. As is clear from Figure 3(a), the microstructure of the hot-rolled base material used in this example is very similar to that used in Examples 1 and 2. That is, high-density nano-order precipitates (approximately 3-5 nm) are uniformly dispersed in the base and dislocation sites (indicated by arrows). Figure 3(b) is the electron diffraction pattern of a selected region of the base material after hot rolling, clearly showing that the base of the base material is a complete austenitic iron phase and the high-density nano-order precipitates are κ-carbides with a regular L'12 structure. Figure 3(c) was obtained from the same region as Figure 3(a) (100) κ This is a dark-field image showing the presence of high-density nano-order κ-carbides. Transmission electron microscopy revealed that the microstructure of the hot-rolled matrix used in this example is single-phase austenitic iron containing uniformly distributed high-density nano-order κ-carbides. Tensile tests showed that the yield strength, ultimate tensile strength, and elongation of the matrix after hot rolling were 975 MPa, 1152 MPa, and 50%, respectively.

[0044] Two hot-rolled base metal plates, measuring 80mm x 80mm x 8mm, are machined to form a single V-groove butt weld. The nominal composition of the weld filler wire is iron-28.2 manganese-10.6 aluminum-1.62 carbon (Fe-28.2Mn-10.6Al-1.62C), with a diameter of approximately 2.4mm. The parameters for tungsten gas arc welding are the same as those used in Example 1. Figure 3(d) shows a typical scanning electron microscope image of the post-weld sample, including the molten region, the weld heat-affected region, and the base metal region. As seen in Examples 1 and 2, the molten region contains typical columnar austenite iron grains, which consist of austenite dendritic crystalline cells and eutectic regions. At the same time, the true properties of the high-density nano-order κ-carbides distributed in the austenite dendritic crystalline cells and eutectic regions can be observed. However, compared to Example 1 (Figure 1(e)) and Example 2 (Figure 2(b)), the amount of nano-order κ-carbides in this example appears to be slightly reduced, suggesting that the carbon content of the weld filler wire used in this example is relatively low. Furthermore, it should be noted that no traces of microcracks or pores were observed in the molten region or near the boundary between the molten region and the weld heat-affected zone. At the same time, the density of nano-order κ-carbides originally present in the weld heat-affected zone and the base metal region was clearly not affected by the welding process.

[0045] Figure 3(e) shows the Vickers microhardness test across the molten region, the heat-affected zone of welding, and the base metal region. The corresponding microhardness values ​​measured at each measurement point are shown in Figure 3(f). The average microhardness of the molten region, the heat-affected zone of welding, and the base metal region were found to be approximately 402 Vickers hardness, 417 Vickers hardness, and 420 Vickers hardness, respectively. Clearly, the softening in the molten region is very slight. In this example, the microhardness of the molten region is still considerably higher than that of a precipitation-hardening austenitic iron-manganese-aluminum-carbon alloy with a carbon content of 1.2 wt% or less after ideal aging treatment (350-400 Vickers hardness) (as described in the prior art).

[0046] Figure 3(g) is a bright-field image obtained from a transmission electron microscope from the molten region. Compared to those seen in Example 1 (Figure 1(h)) and Example 2 (Figure 2(e)), the density of nano-order κ-carbides in this figure is slightly lower. The electron diffraction pattern of the selected region shown in Figure 3(h) confirms that the microstructure of the molten region is composed of austenitic iron and κ-carbides. Figure 3(i) shows the analysis results from a transmission electron microscope and X-ray energy dispersive analyzer obtained from the molten region, showing that iron, manganese, and aluminum in the molten region are approximately 61.9%, 27.6%, and 10.5%, respectively. Tensile testing of the post-welded sample shows that the sample that fractured after the tensile test had a large amount of plastic deformation in the molten region, and that necking and fracture mainly occurred in the molten region. On the other hand, the amount of deformation in the base material region is relatively small. This may be due to the fact that the microhardness value of the molten region (approximately 402 Vickers hardness) is very low compared to the base material region (approximately 420 Vickers hardness). As a result, there is clear ductile deformation in the molten region, and the total elongation in this example is slightly reduced compared to the example above. The yield strength, ultimate tensile strength, and elongation of the welded member of this post-welded sample are approximately 916 megapascals, 1035 megapascals, and 33%, respectively.

[0047] From the results above, it is clear that the microstructure of the welded component in this example and the resulting mechanical properties are very similar to those observed in Examples 1 and 2.

[0048] <Example 4> In this example, the base material used is the same as that used in Example 3. The size of the base material used in this example is 80mm x 80mm x 8mm, and it is machined to form a single V-shaped recessed butt joint. The standard composition of the welding filler wire is iron-29.5 manganese-9.8 aluminum-1.48 carbon (Fe-29.5Mn-9.8Al-1.48C), and its diameter is approximately 2.4mm. Pure argon is used as the shielding gas, and fusion welding is performed by gas tungsten arc welding, with the welding parameters being the same as those used in Example 1.

[0049] Figure 4(a) shows a typical scanning electron microscope image of a post-weld sample, including the molten region, the heat-affected zone, and the base metal region. Clearly, the microstructure of the molten region is very similar to that seen in Examples 1-3. That is, typical columnar austenite iron grains consist of austenite dendritic crystal cells and eutectic regions. The dense distribution of nano-order κ-carbides in the austenite dendritic crystal cells and eutectic regions is still observable. However, the carbon content is further reduced by the welding filler wire used in this example, and the amount of nano-order κ-carbides dispersed in the austenite dendritic crystal cells and eutectic regions appears to be slightly reduced. Furthermore, it is noteworthy that no traces of microcracks or pores were observed in the molten region or near the boundary between the molten region and the heat-affected zone. At the same time, the density of nano-order κ-carbides originally present in the heat-affected zone and base metal region was clearly not affected during the welding process. Figure 4(b) shows the Vickers microhardness test across the molten region, the heat-affected zone of the weld, and the base metal region. The corresponding microhardness values ​​measured at each measurement point are shown in Figure 4(c). Figure 4(c) shows that the average microhardness of the molten region, the heat-affected zone of the weld, and the base metal region are approximately 385 Vickers, 414 Vickers, and 419 Vickers, respectively. The decrease in microhardness of the molten region is presumed to be due to a decrease in the carbon content of the weld filler wire used in this example. It should be noted that, as seen in the previous examples, there was no obvious softening phenomenon caused by fusion welding.

[0050] Figure 4(d) is a bright-field image obtained from a transmission electron microscope of the molten region. Nano-order κ-carbides present in the austenite dendritic crystal cells and eutectic region are clearly visible. However, when comparing the nano-order κ-carbides present in the austenite dendritic crystal cells and eutectic region with those observed in Examples 1 to 3, it is clear that the density is slightly lower. The electron diffraction pattern of the selected region obtained from the molten region shown in Figure 4(e) confirms that the microstructure of the molten region is indeed composed of austenite iron and κ-carbides. Figure 4(f) shows the analysis results from a transmission electron microscope and X-ray energy dispersive analyzer obtained from the molten region. It can be seen that the iron, manganese, and aluminum in the molten region are approximately 61.7%, 28.8%, and 9.5%, respectively. Tensile tests show that the yield strength, ultimate tensile strength, and elongation of the welded member in this example are approximately 875 megapascals, 1012 megapascals, and 29.8%, respectively. After detailed examination, the sample that fractured in the tensile test was found to exhibit typical characteristics of ductile-plastic deformation in the molten region. However, there was a significant difference in microhardness between the molten region (approximately 385 Vickers hardness) and the base metal region (approximately 419 Vickers hardness) (see Figure 4(c)). Almost all deformation occurred in the molten region, and the amount of deformation in the weld heat-affected zone and base metal region was relatively small, resulting in a lower overall elongation rate. Nevertheless, the severe necking observed in the molten region after the tensile test, without a wavy surface or meandering fracture pattern, indicates that the weld bead has good ductility.

[0051] It should be noted that the microhardness of the molten region in this example (approximately 385 Vickers hardness) is similar to that of the austenitic iron-manganese-aluminum-carbon alloy with a carbon content of 1.2% by weight or less of precipitation hardening after ideal aging treatment as described in the prior art (350-400 Vickers hardness). From the scanning electron microscope and transmission electron microscope results above, it is clear that the welding filler wire designed with the component composition used in this example retains most of the microstructural properties in the molten region. In other words, these properties refer to the austenitic dendritic crystal cells and high-density nano-order κ-carbides present in the eutectic region. Therefore, it can be reasonably expected that the molten region has an excellent combination of yield strength and ductility, which is due to the use of the new welding filler wire disclosed in this invention. In fact, the measured yield strength (approximately 875 megapascals) and elongation (approximately 29.8%) of the welded member in this example are very similar to the values ​​measured from precipitation-hardening austenitic iron-manganese-aluminum-carbon alloys with a carbon content of 1.2% by weight or less after ideal aging treatment (yield strength: approximately 680-990 megapascals, elongation: 55-26%).

[0052] <Example 5> In this example, the base material used is the same as that used in Example 1. The size of the base material used in this example is 80mm x 80mm x 8mm. Two base material plates are machined to form a single V-shaped grooved butt joint. The standard composition of the welding filler wire used in this example is iron-33.8 manganese-10.6 aluminum-1.38 carbon (Fe-33.8Mn-10.6Al-1.38C), with a diameter of approximately 2.4mm. 75% argon and 25% helium are used as shielding gases, and gas tungsten arc welding is used for fusion welding, with welding parameters the same as those used in Example 1.

[0053] Figure 5(a) is a scanning electron microscope image showing the microstructure of the post-weld sample, including the molten region, the weld heat-affected zone, and the base metal region. Similar to Examples 1-4, the molten region is composed of typical columnar austenite iron grains, and the characteristic uniform dispersion of numerous nano-order κ-carbides in austenite dendritic crystalline cells is still observed. However, the density of nano-order κ-carbides has decreased significantly. On the other hand, the density of nano-order κ-carbides originally present in the weld heat-affected zone and the base metal region was clearly unaffected during the welding process. Furthermore, it should be noted that no traces of microcracks or pores were observed within the molten region or near the boundary between the molten region and the weld heat-affected zone. Figure 5(b) shows the Vickers microhardness test performed across the molten region, the weld heat-affected zone, and the base metal region. The corresponding microhardness values ​​measured at each measurement point are shown in Figure 5(c). Figure 5(c) shows that the average microhardnesses in the molten region, the heat-affected zone of the weld, and the base metal region are approximately 369 Vickers hardness, 385 Vickers hardness, and 412 Vickers hardness, respectively. The decrease in microhardness in the molten region is presumed to be due to the low carbon content of the welding filler wire used in this example. This relatively low carbon content clearly further reduces the density of nano-order κ-carbides dispersed within the austenite dendritic crystal cells and eutectic region, resulting in an even lower hardness.

[0054] Figure 5(d) is a transmission electron microscope image obtained from the molten region. The bright-field image shows nano-order κ-carbides present in the austenite dendritic crystal cells and eutectic region. At the same time, it can be seen that the density of nano-order κ-carbides in the austenite dendritic crystal cells and eutectic region is significantly lower in this example compared to the nano-order κ-carbides observed in Examples 1 to 4 above. Figure 5(e) shows the electron diffraction pattern of a selected region obtained from the molten region, which further confirms that the microstructure of the molten region is indeed composed of austenite iron and κ-carbides. Figure 5(f) shows the analysis results from a transmission electron microscope and X-ray energy dispersive analyzer obtained from the molten region, showing that iron, manganese, and aluminum in the molten region are approximately 56.6%, 33.1%, and 10.3%, respectively.

[0055] Tensile tests showed that the yield strength, ultimate tensile strength, and elongation of the welded member in this example were approximately 725 megapascals, 1,005 megapascals, and 23.8%, respectively. After careful inspection, it was found that the sample fractured by tensile testing exhibited clear ductile-plastic deformation in the molten region. However, there was a considerable difference in microhardness between the molten region (approximately 369 Vickers hardness) and the base metal region (approximately 412 Vickers hardness) (see Figure 5(c)), and in the sample fractured after tensile testing, almost all deformation occurred in the molten region, with only slight deformation in the weld heat-affected zone and base metal region, resulting in a significant decrease in the overall elongation.

[0056] From the results above, it is clear that the weld filler wire with the designed composition used in this example still retains most of the microstructural properties of the molten region disclosed by the present invention. Furthermore, the microhardness obtained by the molten region (approximately 369 Vickers hardness) is still equivalent to the microhardness of a precipitation-hardening austenitic iron-manganese-aluminum-carbon alloy with a carbon content of 1.2 wt% or less after ideal aging (approximately 350 (400 million Pa)). As described in Example 4, because the microstructure and the resulting microhardness have a high degree of similarity, the molten region in this example can also be expected to have a satisfactory combination of yield strength and elongation. Of particular note here is the characteristic microstructure of the base material used in Examples 1 to 5 of the present invention. That is, in the state after hot rolling, the microstructure of the base material consists of complete austenitic iron and high-density nano-order κ-carbides uniformly distributed on an austenitic iron base. The density of nano-order κ-carbides is significantly higher than that of nano-order κ-carbides in precipitation-hardening austenitic iron-manganese-aluminum-carbon alloys with a carbon content of 1.2 wt% or less after ideal aging treatment. Consequently, the microhardness (approximately 412-420 Vickers hardness) of the base materials used in Examples 1-5 is all far higher than that of microhardness alloys of precipitation-hardening austenitic iron-manganese-aluminum-carbon alloys with a carbon content of 1.2 wt% or less after ideal aging treatment (approximately 350-400 Vickers hardness).

[0057] <Example 6> The purpose of this example is to clarify how further reduction in the carbon content of the welding filler wire affects the properties of the molten region of the post-welded sample. The base material used in this example is the same as that used in Example 1. The base plate size is 80 mm × 80 mm × 8 mm, and a single V-shaped grooved butt joint is formed after machining. Gas tungsten arc welding is performed, and the welding parameters are the same as those used in Example 1. The standard composition of the welding filler wire used in this example is iron-31.8 manganese-9.2 aluminum-1.25 carbon (Fe-31.8Mn-9.2Al-1.25C), and the diameter is approximately 2.4 mm.

[0058] Figure 6(a) shows typical scanning electron microscope images of a sample after welding, including the molten region, the heat-affected zone of the weld, and the base metal region. The microstructure of the molten region clearly exhibits several characteristics that are quite different from those of Examples 1 to 5 described above. First, the base is still complete austenitic iron, containing austenitic dendritic crystal cells and a eutectic region. However, κ-carbides are not dispersed in the austenitic dendritic crystal cells, and the amount of κ-carbides in the eutectic region is significantly reduced. To understand this characteristic more clearly, please refer to a magnified image of one of the marked regions (indicated by the arrow) inserted in the upper right corner of Figure 6(a). This characteristic is completely different from the microstructure described in Examples 1 to 5. In fact, this property is very similar to the phenomenon observed by Chou and Lee in the spontaneous fusion of iron-manganese-aluminum-carbon alloys with a carbon content of 1.29 weight percent or less, meaning that, unlike conventional techniques, the κ-carbides are not dispersed in the austenite dendritic crystal cells.

[0059] Figure 6(b) shows the Vickers microhardness test across the molten region, the weld heat-affected zone, and the base metal region. The corresponding microhardness values ​​measured at each measurement point are shown in Figure 6(c). Clearly, these two figures show that the microhardness in the weld heat-affected zone and the base metal region is approximately 412 Vickers hardness, but the microhardness in the molten region drops rapidly to 242 Vickers hardness. This Vickers hardness value is very similar to that of precipitation-hardening austenitic iron-manganese-aluminum-carbon alloys with a carbon content of 1.2% or less after solution heat treatment and quenching, and a Vickers hardness of 210-230, as in conventional techniques. Generally, austenitic dendritic crystal cells lack κ-carbides, and the significantly reduced κ-carbides in the eutectic region are the main factors contributing to the rapid decrease in microhardness in the molten region.

[0060] Figure 6(d) shows macroscopic images of the sample before and after the tensile test. As a result, the yield strength, ultimate tensile strength, and elongation of the welded member were approximately 462 megapascals, 938 megapascals, and 14%, respectively. Clearly, the yield strength and elongation of the welded member obtained using this example are not satisfactory. The main reason for the significant decrease in elongation is thought to be that the microhardness of the molten region (approximately 242 Vickers hardness) is significantly smaller than the microhardness of the heat-affected zone and the base metal region (approximately 412 Vickers hardness). Therefore, during the tensile test period, the overall deformation of the welded member is basically limited to the molten region, while the heat-affected zone and base metal region hardly deform and remain unchanged, resulting in a significant decrease in elongation. At the same time, in Figure 6(d), it can be observed that a fracture crack appears from the interface between the molten region and the heat-affected zone (indicated by the arrow), then penetrates the molten region laterally, and forms a curved, sawtooth-like fracture line. The elongation in the molten region is estimated to be approximately 27%. In fact, its microhardness is approximately 242 Vickers hardness, indicating that deformation within a certain range, the initiation of necking, and subsequent cracking all occur in the molten region, and that the strength of the molten region is much lower than that of the weld heat-affected zone and the base metal region. The main reason is thought to be that κ-carbides are not dispersed in the austenite dendritic crystal cells and the amount of κ-carbides in the eutectic region is significantly reduced. Weld filler wire with a carbon content of 1.25 wt.% does not adequately promote the formation of nano-order κ-carbides in the austenite dendritic crystal cells and eutectic region during the solidification process, and the formation of these nano-order κ-carbides is also the most unique property revealed by this invention.

[0061] <Example 7> The purpose of this example is to investigate the effect of a higher carbon content in the weld filler wire on the properties of the molten region of the welded sample. The base material used in this example is the same as that used in Example 3. The base plate size is 80 mm × 80 mm × 10 mm, and a single V-groove butt joint is formed after machining. Gas tungsten gas arc welding is used for fusion welding, and the welding parameters are the same as those used in Example 1. The standard composition of the weld filler wire used in this example is iron-30.5 manganese-8.9 aluminum-2.2 carbon (Fe-30.5Mn-8.9Al-2.2C), and the diameter is approximately 2.4 mm.

[0062] Figure 7(a) is a macroscopic image of the weld bead, clearly showing macrocracks penetrating the weld bead. Clearly, a high carbon content in the weld filler wire leads to a severe solidification high-temperature cracking reaction. Figure 7(b) is a scanning electron microscope image of a post-welded sample, including the molten region, the weld heat-affected region, and the base metal region. This figure clearly shows that the molten region is mainly composed of typical columnar austenite iron grains, while high-density nano-order κ-carbides are dispersed in the austenite dendritic crystalline cells and eutectic regions. This property is similar to that observed in Examples 1 to 5. However, coarse carbides of micrometer size (μm-sized) are observed at the grain boundaries (indicated by arrows in Figure 7(b)). Referring to the inserted upper right corner of Figure 7(b), the magnified image of the marked region (indicated by arrows) clearly shows whether there is a precipitate region near the coarse carbides formed at the grain boundaries. It is known that both grain boundary precipitates and associated non-precipitation regions negatively affect mechanical properties, particularly elongation. Furthermore, Figure 7(b) shows that the density of nano-order κ-carbides originally present in the weld heat-affected zone and the base material region is clearly unaffected during the fusion welding process.

[0063] Figure 7(c) is a transmission electron microscope image obtained from the molten region. The bright-field image clearly shows that, in addition to high-density nano-order κ-carbides, micrometer-sized coarse precipitates (labeled K) are also present at the austenite iron grain boundaries. Furthermore, there are clear non-precipitation regions near the coarse precipitates. Figure 7(d) is a selected-region electron diffraction pattern obtained from the coarse precipitates at the grain boundaries (labeled K), indicating that these coarse precipitates at the grain boundaries are L'12 structured κ-carbides. This structure is identical to the nano-order κ-carbide structure dispersed in the austenite dendritic crystal cells and eutectic regions. Figure 7(e) shows typical transmission electron microscope and X-ray energy dispersive analyzer results for coarse κ-carbides at the grain boundaries, showing that the iron, manganese, and aluminum content of the coarse κ-carbides is approximately 56.7%, 34.2%, and 9.1%, respectively.

[0064] Based on the observations described above, it is clear that when the carbon content of the welding filler wire is high (2.2 wt%), high-density nano-order κ-carbides are present in austenite dendritic crystal cells and eutectic regions, but coarse κ-carbides also appear at grain boundaries. Therefore, it is expected that the coarse κ-carbides at grain boundaries and the associated non-precipitation regions will not only have adverse effects on the welded member, such as ductility, but will also cause serious solidification and high-temperature cracking reactions in the weld bead, as shown in Figure 7(a).

[0065] <Example 8> This example illustrates how a high aluminum content in the welding filler wire affects the properties of the molten region of the welded sample. The base material used in this example is the same as that used in Example 1. The two base plate dimensions are 80mm x 80mm x 8mm, and after machining, a single V-shaped grooved butt joint is formed. The welding parameters used in gas tungsten arc welding are the same as those used in Example 1. The standard composition of the welding filler wire is iron-30.5 manganese-12.5 aluminum-1.55 carbon (Fe-30.5Mn-12.5Al-1.55C), and its diameter is approximately 2.4mm.

[0066] Preliminary inspection of the welded component revealed solidification hot cracking in a portion of the weld bead, as indicated by the arrow in Figure 8(a). The scanning electron microscope image shown in Figure 8(b) shows the microstructure of the molten region, which mainly consists of typical columnar austenite iron grains composed of austenite dendritic crystal cells and eutectic regions. Similar to what was observed in Examples 1 to 5, high-density nano-order κ-carbides are present in the austenite dendritic crystal cells and eutectic regions. However, in addition to the above characteristics, several discrete particles were observed in the austenite iron grains, as indicated by the arrows and symbol α in Figures 8(b) and 8(c). Transmission electron microscopy and X-ray energy dispersive analysis (TEM-EDS) of the discrete particles showed that the iron, manganese, and aluminum content of the compositional components was approximately 52.3%, 34.2%, and 13.5%, respectively. Clearly, the concentration of aluminum in the discrete particles reached as high as 13.5%. This means that during the welding process, particles rich in aluminum are formed within the austenitic iron grains. It has been confirmed that the presence of an aluminum-rich secondary phase in austenitic iron grains significantly deteriorates the ductility of iron-manganese-aluminum alloys. Therefore, if the welding filler wire contains excessive aluminum, not only will solidification cracks be formed during the welding process, but an aluminum-rich phase will also be formed that adversely affects the austenitic iron grains.

[0067] <Example 9> The main objective of this example is to determine how a high manganese content in the welding filler wire affects the properties of the molten region of the welded sample. The base material used is the same as that used in Example 1. The base plate size is 80 mm × 80 mm × 8 mm, and a single V-shaped grooved butt joint is formed after machining. The welding parameters used in gas tungsten arc welding are the same as those used in Example 1. The standard composition of the welding filler wire used in this example is iron-20.1 manganese-8.5 aluminum-1.52 carbon (Fe-20.1Mn-8.5Al-1.52C) with a diameter of approximately 3.2 mm.

[0068] Figure 9(a) is a scanning electron microscope image of the region including the molten area, the heat-affected zone, and the base metal region of the welded sample. The microstructure of the molten area is clearly different from that observed in Examples 1 to 5 above. High-density nano-order κ-carbides are still observed in austenite dendritic crystal cells and the eutectic region, but there is also a large amount of discrete ferrite phase (indicated by arrows and symbol a in Figure 9(a)) which is dispersed throughout the molten region. The volume fraction of the ferrite phase is estimated to be approximately 18%, and this value was calculated using the point-counting technique.

[0069] Figure 9(b) is a bright-field transmission electron microscope image obtained from the molten region, showing austenite dendritic crystal cells and nano-order κ-carbides dispersed in the eutectic region. Clearly, coarse, micrometer-sized κ-carbides are formed on the austenite iron / ferrite iron grains. Figure 9(c) is the selected-region electron diffraction pattern of the circled area in Figure 9(b), clearly showing the presence of three phases: austenite iron phase, ferrite iron phase, and κ-carbides. The orientation relationships between the austenite iron phase and the ferrite iron phase are as follows: TIFF0007923092000001.tif17123 and The filename is TIFF0007923092000002.tif21120. Figures 9(d) and 9(e) show typical transmission electron microscope and X-ray energy dispersive analyzer results obtained from the austenite dendritic crystal cell region and the ferrite region, respectively. Quantitative analysis results for the austenite dendritic crystal cell and ferrite revealed that their chemical compositions are iron-20% manganese-8.2% aluminum (Fe-20.6%Mn-8.2%Al) and iron-15.8% manganese-8.8% aluminum (Fe-15.8%Mn-8.8%Al), respectively.

[0070] Based on the results above, the interesting and noteworthy point of this example is that, in the case of an iron-manganese-aluminum-carbon alloy with the composition of iron-(17.5~35)manganese-(7.1~12)aluminum-(0.8~1.2)carbon (Fe-(17.5~35)Mn-(7.1~12)Al-(0.8~1.2)C), the microstructure of the alloy always exhibits single-phase austenite after hot rolling, solution heat treatment, and quenching, as described in the prior art. Therefore, this example further demonstrates that even materials with the same chemical composition can have significantly different microstructures after fusion welding compared to materials obtained through general heat-treatment processes. Clearly, in order to create a ferrite-free, full austenite microstructure in the molten region and disperse high-density nano-order κ-carbides in the austenite iron dendritic cells and eutectic regions, the manganese content of the welding filler wire used in this example must be basically higher than about 20% (by weight).

[0071] <Example 10> In this embodiment, the effect of adding titanium to the welding filler wire on the properties of the molten region of a welded sample is investigated. The base material used in this example is a hot-rolled base material with a standard composition of iron-29,3-manganese-9,3-aluminum-1.75-carbon (Fe-29,3Mn-9,3Al-1.75C). Figure 10(a) shows a bright-field image of the hot-rolled base material obtained by transmission electron microscopy, clearly showing that during the hot-rolling process, high-density nano-order precipitates (approximately 3-5 nanometers) are uniformly dispersed in the base and formed at dislocations (indicated by arrows). Figure 10(b) is an electron diffraction pattern of a selected region, confirming that the base of the base material is a complete austenitic iron phase and that the high-density nano-order precipitates are κ-carbides with a regular L'12 structure. Figure 10(c) shows (100) obtained from the same region as Figure 10(a). κ This is a dark-field image showing high-density nano-order κ-carbides appearing on an austenitic iron base and located at dislocation sites. Microstructural analysis using a transmission electron microscope confirmed that the microstructure of the hot-rolled matrix used in this example is single-phase austenitic iron with uniformly dispersed high-density nano-order κ-carbides. Tensile tests revealed that the yield strength, ultimate tensile strength, and elongation of the hot-rolled matrix were approximately 1020 megapascals, 1198 megapascals, and 45%, respectively.

[0072] Two hot-rolled base plate plates, measuring 80mm x 80mm x 12mm, are machined to form a single V-groove butt joint. The welding parameters used in gas tungsten arc welding are the same as those used in Example 2. The standard composition of the weld filler wire used in this example is iron-28.5manganese-9.3aluminum-1.82carbon-1.6titanium (Fe-28.5Mn-9.3Al-1.82C-1.6Ti) with a diameter of approximately 1.2mm. Figure 10(d) shows a macroscopic image of the welded sample, showing that the entire weld bead is very smooth and there are no large cracks or holes visible to the naked eye. Figure 10(e) is a scanning electron microscope image of the welded sample, including the molten region, the heat-affected region of the weld, and the base material region. From Figure 10(e), several important properties caused by the addition of titanium can be seen. Firstly, the size of the austenite iron particles in the molten region is clearly much smaller than that seen in Examples 1 to 5, and even smaller than in the weld heat-affected zone and the base metal region. Secondly, the morphology of the austenite iron particles becomes more spherical and the directionality decreases. However, it is clear that high-density nano-order κ-carbides uniformly dispersed in the austenite dendritic crystalline cells and eutectic region are still present. Also, although the austenite iron particles in the weld heat-affected zone have grown slightly, the high-density nano-order κ-carbides originally present in the base metal region remain unchanged. Finally, it should be noted that no traces of minute cracks and pores were observed in the molten region or near the boundary between the molten region and the weld heat-affected zone, which means that solidification cracks and liquefaction cracks commonly seen in precipitation-hardening alloys after fusion welding can be largely eliminated by using the weld filler wire disclosed in this invention. Figure 10(f) shows the Vickers microhardness test across the molten region, the weld heat-affected zone, and the base metal region. Figure 10(g) shows the corresponding microhardness values ​​measured at each measurement point, indicating that the average microhardness of the molten region, weld heat-affected region, and base metal region of the welded sample of the present invention is approximately 451 Vickers hardness, 435 Vickers hardness, and 436 Vickers hardness, respectively. It is clear that the most common softening phenomenon in most precipitation-hardening alloys welded in the molten region does not occur.This is mainly because, in the post-fusion state (Figure 10(e)), a high density of nano-order κ-carbides exists in the austenite dendritic crystal cells and eutectic region. Furthermore, in addition to the nano-order κ-carbides, numerous high-hardness (2200-3500 Vickers hardness) nano-order titanium-rich titanium carbides (Ti-rich Ti-carbides) are formed within the eutectic region (Figure 10(h)). As a result, not only is there no softening phenomenon in the molten region, but the micro-hardness is also higher than in the base material region.

[0073] Examination of a thin foil in the molten region using a transmission electron microscope revealed that the addition of titanium resulted in the formation of numerous nano-order titanium carbides in a regular face-centered cubic structure within the eutectic region. See Figure 10(h) for a typical example. In the bright-field image obtained from this transmission electron microscope, the length of the austenite dendritic crystal cells is only about 100-150 nanometers, and high-density nano-order κ-carbides (about 3-5 nanometers) are uniformly dispersed in the austenite dendritic crystal cells and the eutectic region. Furthermore, as indicated by the arrows in Figure 10(h), there are also slightly larger nano-order titanium carbides (about 6-10 nanometers) in the eutectic region. Figure 10(i) is the electron diffraction pattern of a selected region obtained from the circled area in Figure 10(h), showing the coexistence phenomenon of austenite iron base, κ-carbides, and titanium carbides. Analysis of the electron diffraction pattern of the selected region indicates that the crystal orientation relationship between the three phases is cubic-to-cubic. Figure 10(j) shows the results of transmission electron microscopy and X-ray energy dispersive analysis obtained from the same region, demonstrating that titanium-rich titanium carbides appear in the eutectic region. The standard compositions of the welding filler wires used in Example 1 and Example 10 have almost the same aluminum and carbon content, except that approximately 1.6 wt% titanium is added in Example 10. Therefore, it is necessary to compare the microstructure of the molten regions shown in Figure 1(h) and Figure 10(h). In Figure 1(h), the size of the austenite dendritic crystal cells is slightly larger than that of the austenite dendritic crystal cells in Figure 10(h), but the density of nano-order κ-carbides in the austenite dendritic crystal cells in Figure 1(h) appears to be slightly higher than in Figure 10(h) (titanium-added sample). Furthermore, both samples have numerous high-density nano-order precipitates in the eutectic region (Figure 1(h) shows κ-carbide, and Figure 10(h) shows titanium-rich titanium carbide). However, the microhardness of the molten region of Example 10 (approximately 451 Vickers hardness) is much greater than that of Example 1 (approximately 425 Vickers hardness).The main reason for this phenomenon may be that titanium carbides, which are rich in titanium, have a Vickers hardness of approximately 2200-3500, which is significantly higher than the hardness of (Fe,Mn)3AlC carbides (approximately 630-670 Vickers).

[0074] Tensile tests have shown that the yield strength, ultimate tensile strength, and elongation of the welded member are 1,006 megapascals, 1,145 megapascals, and 40%, respectively. Figure 10(k) shows macroscopic images of the sample before and after the tensile test. The corrugated deformation surface in the molten region clearly demonstrates the ductility of the weld bead. Furthermore, Figure 10(k) d FZ d BM , and d INT As shown, the edges of the tensile-fractured sample remain parallel, span the molten region and the base material region, and are found to be essentially the same width, where d INT This represents the width of the interface between the molten region and the base material region. As described in detail in Examples 1 and 2, such special deformation properties strongly suggest that the yield strengths of the molten region and the base material region are at approximately the same level. As mentioned above, the remarkable yield strength and microhardness obtained in the molten region may be due to the high density of nano-order κ-carbides and nano-order titanium-rich titanium carbides in the molten region. At the same time, it should be noted that a slight necking phenomenon occurs in the base material region, eventually leading to rupture, and furthermore, the tensile strength of the molten region is higher than that of the base material region, which is consistent with the phenomenon of high microhardness in the molten region, as shown in Figure 10(g).

[0075] Based on the above analysis, it is clear that adding titanium to the welding filler wire produces the following advantageous effects. First, the austenite iron grains in the molten region change from cylindrical to spherical (globular), and the grain size is significantly refined. Second, in the post-weld state, high-density nano-order κ-carbides exist in the eutectic region with significantly refined austenite dendritic crystal cells. Furthermore, numerous high-hardness, nano-order titanium-rich titanium carbides are formed in the eutectic region. Third, the austenite dendritic crystal cells, nano-order κ-carbides, and nano-order titanium-rich titanium carbides constituting the molten region all have a ductile face-centered cubic structure. Moreover, since the size of the κ-carbides and titanium-rich titanium carbides in the molten region is only 3 to 10 nanometers, mechanical properties, especially yield strength, can be effectively improved without causing a significant decrease in ductility. Therefore, after using the welding filler wire disclosed in the present invention, the molten region of the welded sample can simultaneously possess an excellent combination of hardness, yield strength, ultimate tensile strength, and ductility.

[0076] <Example 11> In this embodiment, the effect of niobium added to the weld filler wire on the properties of the molten region of the welded sample was investigated. The base material used is the same as that used in Example 10. The two base plate sizes are 80 mm × 80 mm × 8 mm, and after machining, a single V-shaped grooved butt joint is formed. The welding parameters used in gas tungsten arc welding are the same as those used in Example 1. The standard composition of the weld filler wire used is iron-29.5 manganese-8.2 aluminum-1.56 carbon-1.0 niobium (Fe-29.5Mn-8.2Al-1.56C-1.0Nb), ​​with a diameter of approximately 3.2 mm. Figure 11(a) shows a macro image of the welded sample, showing that the entire weld bead has a very smooth shape and there are no large cracks or holes visible to the naked eye. Figure 11(b) is a scanning electron microscope image of the welded sample, including the molten region, the heat-affected region, and the base material region. Unlike the sample before titanium addition (Example 10), the austenitic iron particles in the molten region of this example are a mixture of both globular and columnar-like forms. However, similar to the phenomenon observed in Example 10, high-density nano-κ-carbides can be observed uniformly dispersed in the austenitic dendritic crystalline cells and eutectic regions. Similarly, while the austenitic iron particles in the weld heat-affected zone grew slightly, the high-density nano-order κ-carbides originally present in the hot-rolled base material region remained essentially unchanged during the fusion welding period. Furthermore, it should be noted that no traces of microcracks or pores were observed in the molten region or near the boundary between the molten region and the weld heat-affected zone. This indicates that using the exposed weld filler wire in this invention can significantly eliminate solidification cracks and liquefaction cracks. Figure 11(c) shows the Vickers microhardness test across the molten region, weld heat-affected zone, and base material region. The corresponding microhardness values ​​measured at each measurement point are shown in Figure 11(d). This indicates that the microhardness of the molten region, weld heat-affected region, and base metal region of the welded state sample of the present invention are approximately 431 Vickers hardness, 435 Vickers hardness, and 437 Vickers hardness, respectively. It is clear that the most common softening phenomenon in most precipitation-hardening alloys that are fusion-welded is absent in the molten region.Similar to the phenomenon observed in Example 10, the main factors are the presence of high-density nano-order κ-carbides in the austenite dendritic crystal cells and eutectic region, and the presence of a large amount of high-hardness (2000-3200 Vickers hardness) nano-order niobium-rich niobium carbides (Nb-rich Nb-carbides) in the eutectic region (Figure 11(e)). Therefore, the microhardness of the molten region is at almost the same level as the microhardness of the weld heat-affected zone and the base material region.

[0077] Similar to the phenomenon observed in the titanium-added sample (Example 10), transmission electron microscopy analysis revealed that niobium carbides, rich in numerous nano-order niobium atoms with a regular face-centered cubic structure, were formed in the eutectic region. Figure 11(e) is a typical transmission electron microscope bright-field image obtained from the molten region of the welded member, clearly showing that the morphology of the austenite dendritic crystal cells and eutectic region in this example is entirely different from that observed in Example 10. The dendritic crystal cells and the surrounding eutectic region appear more roundish. However, this characteristic remains unchanged because high-density nano-order (approximately 3-5 nanometers) κ-carbides are dispersed in the austenite dendritic crystal cells and eutectic region. In addition to the eutectic κ-carbides, the eutectic region also contains numerous slightly larger nano-order precipitates (approximately 6-10 nanometers), as indicated by the arrows in Figure 11(e). Figure 11(f) was obtained from the circled region in Figure 11(e), and the electron diffraction pattern of the selected region shows the coexistence of austenitic iron base, κ-carbide, and niobium-rich niobium carbide. Analysis of the electron diffraction pattern of the selected region indicates that the crystal orientation relationship between the three phases is cubic allotropy. Figure 11(g) shows the analysis results from a transmission electron microscope and an X-ray energy dispersive analyzer obtained from the same region, confirming the presence of niobium-rich niobium carbide. Since the hardness of nano-order niobium-rich niobium carbide is approximately 2000-3200 Vickers hardness, it can be expected that the precipitation of nano-order niobium-rich niobium carbide will similarly affect the properties of the molten region, which simultaneously explains the microhardness results shown in Figure 11(d).

[0078] Tensile tests show that the yield strength, ultimate tensile strength, and elongation of the welded member are 980 megapascals, 1108 megapascals, and 46%, respectively. Figure 11(h) shows macroscopic images of the samples before and after the tensile test. In the tensile fracture sample, a wavy deformation surface and a bent, bypass fracture pattern are clearly observed in the molten region, indirectly indicating the ductility of the molten region. Furthermore, the region (d) including the base material region, the molten region, and the interface between them is also shown. BM d FZ , and d INT The widths (indicated by ) are essentially the same, and as shown in Figure 11(h), the edges of the sample throughout this region remain parallel. This peculiar property means that the yield strength of the molten region and the base material region are at approximately the same level, and that deformation under tensile stress continues uniformly until necking occurs in the molten region and eventually ruptures. In fact, the slight necking and eventual rupture in the molten region indicate that the ultimate tensile strength of the molten region is lower than the ultimate tensile strength of the base material region.

[0079] Based on the above results, it is clear that the addition of niobium to the weld filler wire has a similar effect to the addition of titanium (Example 10). That is, in addition to retaining high-density nano-order κ-carbides present in austenite dendritic crystal cells and eutectic regions, nano-order niobium-rich niobium carbides formed in the eutectic region are also advantageous in strengthening the mechanical strength and microhardness of the molten region of the welded sample, while simultaneously maintaining excellent ductility. As described in Example 10, all phases constituting the microstructure of the molten region have a ductile face-centered cubic structure, and the size of the strengthened nano-order κ-carbides and niobium-rich niobium carbides is approximately 3 to 10 nanometers, which is very effective in increasing the yield strength. Therefore, by using the exposed weld filler wire in the present invention, the molten region of the welded sample simultaneously possesses an excellent combination of hardness, yield strength, ultimate tensile strength, and ductility.

[0080] <Example 12> In this embodiment, the effect of adding vanadium to the weld filler wire on the characteristics of the molten region of a welded sample was investigated. The base material used in this example is the same as that used in Example 10. The two base plate dimensions are 80 mm × 80 mm × 10 mm, forming a single V-shaped grooved butt joint after machining. The welding parameters used in the gas tungsten arc welding are the same as those used in Example 1. The standard composition of the weld filler wire used is iron-31.2 manganese-9.8 aluminum-1.65 carbon-1.2 vanadium (Fe-31.2Mn-9.8Al-1.65C-1.2V), with a diameter of approximately 3.2 mm. Figure 12(a) shows a macroscopic image of the welded sample, showing that the entire weld bead has a very smooth shape and there are no large cracks or holes visible to the naked eye. Figure 12(b) is a scanning electron microscope image of one region of the post-welded sample, including the molten region, the weld heat-affected zone, and the base metal region. Similar to the pre-added sample (Example 11), the image shows austenite iron grains in the molten region with a mixture of spherical and columnar morphologies. Furthermore, as observed in Examples 10 and 11, high-density nano-κ-carbides are uniformly dispersed in the austenite dendritic crystalline cells and eutectic regions. Similarly, while the austenite iron grains in the weld heat-affected zone grew slightly, the high-density nano-order κ-carbides originally present in the hot-rolled base metal region remained essentially unchanged during the fusion welding process. Finally, it should be noted that no traces of microcracks and pores were observed within the molten region or near the boundary between the molten region and the weld heat-affected zone. Figure 12(c) shows Vickers microhardness tests performed in the molten region, the heat-affected zone, and the base metal region. The corresponding microhardness values ​​measured at each measurement point are shown in Figure 12(d), which indicates that the microhardness of the molten region, weld heat-affected region, and base metal region of the welded state sample of the present invention are approximately 438 Vickers hardness, 434 Vickers hardness, and 436 Vickers hardness, respectively. Clearly, no softening phenomenon appeared in the molten region and the weld heat-affected region, which is consistent with the unique microstructure shown in Figure 12(b), in which the high-density nano-order κ-carbides, the main reinforcing components in the three regions, remain essentially unchanged.Furthermore, the addition of vanadium leads to the formation of large amounts of high-hardness (2200-3000 Vickers hardness) vanadium-rich vanadium carbides (V-rich V-carbides) in the eutectic region (Figure 12(e)), and the microhardness of the molten region becomes almost the same as the microhardness of the weld heat-affected zone and the base material region, which is similar to the phenomenon observed in Example 11.

[0081] Similar to the phenomena observed in the titanium-doped sample (Example 10) and the niobium-doped sample (Example 11), transmission electron microscopy analysis indicates that vanadium addition has a similar effect. Specifically, there are numerous vanadium carbides rich in nano-order vanadium with a regular face-centered cubic structure formed in the eutectic region. Figure 12(e) is a bright-field transmission electron microscope image obtained from the molten region, clearly showing that the austenite dendritic crystal cells are only about 20-60 nanometers long, and their morphology differs from that observed in the titanium-doped sample (Figure 10(h)) and the niobium-doped sample (Figure 11(e)). However, it is still possible to observe the properties of high-density nano-order κ-carbides (about 3-5 nanometers) dispersed in the austenite dendritic crystal cells and the eutectic region. Furthermore, as indicated by the arrows in Figure 12(e), it is clear that there are numerous slightly larger precipitates (about 6-10 nanometers) in the eutectic region. Figure 12(f) was obtained from the region circled in Figure 12(e), and the electron diffraction pattern of the selected region shows the coexistence of austenitic iron base, κ-carbide, and vanadium-rich vanadium carbide. Analysis of the electron diffraction pattern of the selected region indicates that the crystal orientation relationship between the three phases is cubic allotropy. Figure 12(g) shows the analysis results from a transmission electron microscope and an X-ray energy dispersive analyzer obtained from the same region, confirming the presence of vanadium-rich vanadium carbide. Since the hardness of vanadium-rich vanadium carbide is similar to that of niobium-rich and titanium-rich vanadium carbide, it can be expected that vanadium-rich vanadium carbide will have a similar effect on the properties of the molten region, and the results for the microhardness shown in Figure 12(d) are also shown.

[0082] Tensile tests revealed that the yield strength, ultimate tensile strength, and elongation of the welded member were 998 megapascals, 1167 megapascals, and 44%, respectively. Figure 12(h) shows macroscopic images of the sample before and after the tensile test. Interestingly, the edges of the tensile fractured sample remained parallel, and the region crossing the molten and base metal regions was essentially the same width. In Figure 12(h), each region is d BM d FZ , and d INT This is shown by (where d INT (where represents the width of the interface between the base metal region and the molten region). As discussed in Examples 1 and 2, and in Examples 10 and 11, this special property means that the yield strengths of the molten region and the base metal region are at approximately the same level. Further examination of the tensile fracture sample revealed that the wavy deformation surface (an indicator of ductility) of the molten region of the weld bead was particularly pronounced. The elongation in the molten region is estimated to be as high as 28%. Furthermore, when the applied stress exceeded the yield strength, the plastic deformation of the entire sample, including the molten region, the heat-affected zone, and the base metal region, exhibited very uniform deformation up to the vicinity of the interface between the molten region and the heat-affected zone, resulting in slight necking and even rupture, as indicated by the arrows in Figure 12(h). As a result, even after rupture, the overall edges remained nearly parallel. In fact, the slight necking and eventual rupture near the interface between the molten region and the base metal region indicate that the ultimate tensile strengths of the molten region, the heat-affected zone, and the base metal region are very close to each other.

[0083] From the above results, it is clear that the addition of vanadium to the weld filler wire produces similar effects to the addition of titanium (Example 10) and niobium (Example 11). In any case, in addition to retaining high-density nano-order κ-carbides in the austenite dendritic crystal cells and eutectic regions, the nano-order vanadium-rich vanadium carbides formed in the eutectic regions are also advantageous in strengthening the mechanical strength and microhardness of the molten region of the welded sample.

[0084] Furthermore, nano-order κ-carbides (3-10 nanometers), titanium-rich titanium carbides, niobium-rich niobium carbides, and vanadium-rich vanadium carbides all possess the same ductile face-centered cubic structure as austenite dendritic crystal cells, thus achieving remarkable strengthening effects without a significant decrease in ductility. The detailed analysis described above clearly demonstrates that the molten regions obtained by using the welding filler wires disclosed in this invention all possess excellent combinations of minute hardness, yield strength, ultimate tensile strength, and ductility.

[0085] <Example 13> In this embodiment, the effect of adding titanium and niobium to the weld filler wire on the properties of the molten region of the welded sample was further investigated. The base material used in this example is the same as that used in Example 10. The two base plate sizes are 80 mm × 80 mm × 12 mm, and a single V-shaped grooved butt joint is formed after machining. The welding parameters used in gas tungsten arc welding are the same as those used in Example 1. The standard composition of the weld filler wire used is iron-30.6 manganese-9.2 aluminum-1.88 carbon-0.8 titanium-1.0 niobium (Fe-30.6Mn-9.2Al-1.88C-0.8Ti-1.0Nb), ​​with a diameter of approximately 3.2 mm. Figure 13(a) shows a macro image of the welded sample, showing that the entire weld bead has a very smooth shape and there are no large cracks or holes visible to the naked eye. Figure 13(b) is a scanning electron microscope image of the region of the weld sample, including the molten region, the heat-affected zone, and the base metal region. In this figure, all austenite iron grains in the molten region are spherical, and the size of these spherical austenite iron grains is significantly reduced, becoming much smaller than the austenite iron grains in the heat-affected zone and the base metal region. As described in Examples 10 to 12, high-density nano-order κ-carbides uniformly dispersed in the austenite dendritic crystalline cells and the eutectic region can be observed. Furthermore, the high-density nano-order κ-carbides originally present in the hot-rolled base metal region remained essentially unchanged during the fusion welding process. It should also be noted that no traces of microcracks or pores were observed in the molten region or near the boundary between the molten region and the heat-affected zone. Figure 13(c) shows the Vickers microhardness test across the molten region, the heat-affected zone, and the base metal region. Figure 13(d) shows the corresponding microhardness values ​​measured at each measurement point. This figure indicates that the average microhardness of the molten region, weld heat-affected zone, and base metal region of the welded sample is approximately 467 Vickers hardness, 444 Vickers hardness, and 435 Vickers hardness, respectively. Not only is no softening observed in the molten region, but it also has a much higher microhardness than the weld heat-affected zone and base metal region.Examples 10 to 12 are approximate, and this is mainly because, in the post-weld state (Figure 13(b)), a high density of nano-order κ-carbides is present in the austenite dendritic crystal cells and eutectic region. Furthermore, in addition to nano-order κ-carbides, the eutectic region contains many high-hardness nano-order titanium-rich titanium carbides and niobium-rich niobium carbides (Figure 13(e)). Consequently, the microhardness of the molten region is much greater than that of the heat-affected zone and the base metal region, and no signs of softening appear. Furthermore, when approximately 0.8% titanium and 1.0% niobium are added, the microhardness of the molten region appears to be higher than that of the previous Examples 10 to 12, which may be due to an increase in the total amount of both the high-hardness titanium-rich titanium carbides and niobium-rich niobium carbides in this example.

[0086] Figure 13(e) is a bright-field transmission electron microscope image obtained from the molten region, clearly showing that the morphology of the austenite dendritic crystal cells and eutectic region in this example is completely different from that observed in Examples 10-12. The dendritic crystal cells and the surrounding eutectic region appear to be more refined and more closely intertwined with each other. However, the high-density nano-order (approximately 3-5 nanometers) κ-carbides are dispersed in the austenite dendritic crystal cells and eutectic region, and this property remains unchanged. Furthermore, in the eutectic region, as indicated by the arrows in Figure 13(e), there is a significant increase in κ-carbides, titanium-rich titanium carbides of a slightly larger size (approximately 6-10 nm), and niobium-rich niobium carbides. Figure 13(f) was obtained from the circled region in Figure 13(e), and the electron diffraction pattern of this selected region confirms the presence of titanium-rich titanium carbide and niobium-rich niobium carbide.

[0087] Tensile tests show that the yield strength, ultimate tensile strength, and elongation of the welded member are 1015 megapascals, 1168 megapascals, and 38%, respectively. Figure 13(g) shows macroscopic images of the sample before and after the tensile test. The wavy deformation surface observed in the molten region indirectly indicates the ductility of the molten region. Furthermore, the edges of the tensile fractured sample remain parallel, and the distances from the base material region and molten region to both sides of the interface between the base material region and the molten region are shown in d in Figure 13(g). BM d FZ , and d INT As shown, they are essentially the same. This special property means that, as shown in Figure 13(g), the yield strength of the molten region and the base material region are at almost the same level, and deformation under tensile stress continues uniformly until necking occurs in the base material region and eventually ruptures. In fact, slight necking and final fracture occur in the base material region, indicating that the ultimate tensile strength of the molten region is greater than the ultimate tensile strength of the base material region. This phenomenon is consistent with the fact that the microhardness obtained in the molten region is much greater than the microhardness of the base material region, as shown in Figure 13(d).

[0088] Based on the above results, it is clear that adding titanium and niobium to the welding filler wire results in significant hardening due to morphological changes and refinement of the austenite iron particle size in the molten region. Furthermore, this addition also leads to a remarkable increase in the amount of high-density nano-order titanium-rich titanium carbides and niobium-rich niobium carbides in the eutectic region compared to the amounts of carbides observed in Examples 10 and 11. More importantly, the high-density nano-order κ-carbides present in the austenite dendritic crystal cells and eutectic region are still largely retained. As described above, since nano-order κ-carbides (approximately 3-10 nanometers), titanium-rich titanium carbides, niobium-rich niobium carbides, vanadium-rich vanadium carbides, and austenite dendritic crystal cells all possess the same ductile face-centered cubic structure, remarkable strengthening effects can be achieved without a significant decrease in ductility. Detailed analysis of the above examples clearly demonstrates that the molten regions obtained by using the welding filler wire disclosed in this invention all possess excellent combinations of minute hardness, yield strength, ultimate tensile strength, and ductility.

[0089] <Example 14> In the process of developing high-strength, particularly high-tensile strength and high-elongation, precipitation-hardening austenitic iron-manganese-aluminum-carbon alloys, significant progress has been made primarily through heat treatments including solution heat treatment, quenching, and ideal aging treatments, which have yielded high-density nano-order κ-carbides in the austenitic iron base. However, as mentioned in the prior art section, despite numerous related studies conducted by many over the past few decades, most of the problems related to the weldability of these alloys remain unresolved. Recently, in 2019, Jeong et al. attempted to disclose the microstructure of the weld heat-affected zone during welding and the influence of its mechanical properties in the weld heat-affected zone on a lightweight precipitation-hardening austenitic iron-31,4-manganese-11,4-aluminum-0.9-carbon (Fe-31,4Mn-11,4Al-0.9C) alloy using Gleeble simulation to confirm that high-density nano-order κ-carbides in the austenitic iron base play a crucial role in maintaining the strength, particularly tensile strength and ductility, of welded members. The alloy they used was hot-rolled to a thickness of 13 mm at 1200°C, water-quenched, solution-treated at 1050°C for 2 hours, then water-quenched again, and finally aged at 550°C for 100 minutes to obtain high-density nano-order κ-carbides within the austenite base. The aging treatment clearly increased the hardness of the alloy, raising the hardness from 298 Vickers in the solution-treated state to 349 Vickers after aging at 550°C for 100 minutes. The simulation tests were conducted using the Greeble simulation test machine. The results showed that when the simulation test samples were cooled to room temperature, softening was indeed observed in the samples after each aging treatment, and complete dissolution of the κ-carbides was observed in the sample with a peak temperature of 1150°C. From these observed phenomena, it can be inferred that there is indeed a relationship between the κ-carbide precipitates and the changes in the mechanical properties of the weld heat-affected zone.However, it should be noted that these Greeble simulation studies do not directly perform fusion welding, and therefore do not include the remelting and resolidification processes, and thus cannot directly obtain the actual microstructure and property values ​​of the molten region. Nevertheless, in their Greeble simulation test studies, it was observed that after aging-treated samples were heated to peak temperatures of 750°C to 1150°C and then cooled to room temperature, the strengthening κ-carbides that were originally present in the weld heat-affected zone dissolved, resulting in a severe softening phenomenon. In fact, this type of phenomenon is very similar to the situation that occurs during fusion welding of 7xxx-series precipitation-hardened aluminum alloys, as described in the prior art.

[0090] The above advantages and characteristics are discussed in detail in the following publication:

[32] S. Jeong, G. Park, B. Kim, J. Moon, SJ Park, C. Lee, “Precipitation behavior and its effect on mechanical properties in weld heat-affected zone in age hardened FeMnAlC lightweight steels”, Mater. Sci. Eng. A, 742 (2019) 61-68.

[0091] This example investigates the effect of fusion welding on the properties of welded components using an age-hardened iron-manganese-aluminum-carbon alloy as the base material. The standard composition of the base material is iron-31.5-manganese-8.5-aluminum-1.25-carbon (Fe-31.5Mn-8.5Al-1.25C). Before fusion welding, the hot-rolled base material is solution-heat-treated at 1050°C for 1 hour, water-quenched, and then age-hardened at 550°C for 12 hours. Two age-hardened base material plates, each measuring 80mm x 80mm x 8mm, are machined to form a single V-groove butt joint. The welding parameters used in gas tungsten arc welding are the same as those used in Example 1. The welding filler wire used in this example is the same as that used in Example 3, with a diameter of approximately 2.4mm.

[0092] Figure 14(a) is a scanning electron microscope image showing the microstructure of the molten region, the heat-affected zone, and the base metal region of the welded sample. In Figure 14(a), it is clear that the microstructure of the molten region is very similar to that seen in Figure 3(d), and a typical microstructure of columnar austenite iron grains is present, as well as austenite dendritic crystal cells and high-density nano-order κ-carbides dispersed in the eutectic region. However, in this example, it is clearly observed that the amount of nano-order κ-carbides in the heat-affected zone has decreased significantly, indicating that most of the nano-order κ-carbides originally present in the base metal region have dissolved through aging treatment during the fusion welding period. According to the investigation of this invention, this is thought to be due to insufficient carbon content in the base metal used in this example. These results are consistent with the 2019 research report by Jeong et al., who used Greeble simulation tests to investigate the effects of fusion welding on age-hardened iron-31.4-manganese-11.4-aluminum-0.9-carbon alloy (Fe-31.4Mn-11.4Al-0.9C). The effects of the weld heat-affected zone are observed in the dissolution of κ-carbides and significant softening phenomena within the weld heat-affected zone. In fact, this is very similar to the phenomena observed during fusion welding of precipitation-hardening AA7075 aluminum alloy, as described in prior art.

[0093] Figure 14(b) shows Vickers microhardness tests performed across the molten region, the heat-affected zone of the weld, and the base metal region. The corresponding microhardness values ​​measured at each measurement point are shown in Figure 14(c), which shows that the average microhardness of the molten region, the heat-affected zone of the weld, and the base metal region of the welded sample were approximately 406 Vickers hardness, 282 Vickers hardness, and 380 Vickers hardness, respectively. Clearly, the dissolution of nano-order κ-carbides, which have a reinforcing effect, within the heat-affected zone of the weld causes a severe softening phenomenon, as explained in many of the prior art referenced above. Figure 14(d) shows macroscopic images of the sample before and after tensile testing. From the tensile fracture sample, it is clear that, due to the severe softening phenomenon in the heat-affected zone of the weld, the fracture path propagates along the interface between the heat-affected zone and the base metal region, as indicated by the arrows.

[0094] From these results, it is clear that appropriate alloy design is fundamentally very important for both the weld filler wire and the base material in order to retain nano-order κ-carbides that have a reinforcing effect in the molten region and the weld heat-affected zone.

[0095] <Example 15> The purpose of this example is to clarify how the addition of molybdenum to the welding filler wire affects the properties of the molten region of the welded sample. The base material used in this example is the same as that used in Example 1. The two base plate dimensions are 80 mm × 80 mm × 10 mm, and a single V-shaped grooved butt joint is formed after machining. The welding parameters used in gas tungsten arc welding are the same as those used in Example 1. The standard composition of the welding filler wire used is iron-28.2 manganese-9.1 aluminum-1.55 carbon-1.8 molybdenum (Fe-28.2Mn-9.1Al-1.55C-1.8Mo), with a diameter of approximately 3.2 mm.

[0096] Figure 15(a) is a scanning electron microscope image showing the molten region, heat-affected zone, and base metal region of a welded sample. The microstructure of the molten region clearly contains mainly typical columnar austenite iron grains, and simultaneously, high-density nano-order κ-carbides are observed in the austenite dendritic crystal cells and eutectic regions, similar to the phenomenon observed in Examples 1 to 5. However, numerous micrometer-sized coarse particles and broad non-precipitation regions are also observed at the austenite iron grain boundaries (indicated by arrows). Figure 15(b) shows the results of scanning electron microscope and X-ray energy dispersive analysis of the coarse particles, indicating that these coarse particles are molybdenum-rich molybdenum carbides. Numerous studies have confirmed that the presence of coarse grain boundaries and associated non-precipitation regions severely impairs the ductility and strength of the alloy. Therefore, it is not recommended to add molybdenum alloy elements that strongly form carbides to the welding filler wire disclosed in the present invention.

[0097] <Example 16> The purpose of this example is to clarify how chromium added to the welding filler wire affects the properties of the molten region of the welded sample. The base material used in this example is the same as that used in Example 1. Two base plate pieces measuring 80 mm × 80 mm × 10 mm are machined to form a single V-shaped grooved butt joint. The welding parameters used in gas tungsten arc welding are the same as those used in Example 1. The standard composition of the welding filler wire used is iron-29,3 manganese-8,8 aluminum-1.5,8 carbon-1.5 Cr (Fe-29,3Mn-8,8Al-1.5,8C-1.5Cr), and the diameter is approximately 3.2 mm.

[0098] Figure 16(a) shows a typical scanning electron microscope image of a region including the molten region, the heat-affected zone of the weld, and the base metal region of a welded sample. The microstructure of the molten region mainly contains typical columnar austenite iron grains, and at the same time, high density of nano-order κ-carbides can be observed in austenite dendritic crystal cells and eutectic regions, which is similar to the phenomenon observed in Examples 1 to 5. However, at the same time, a large number of micrometer-sized coarse particles were observed at the austenite iron grain boundaries (indicated by arrows), and non-precipitation regions existed around these coarse particles. Figure 16(b) shows the results of scanning electron microscope and X-ray energy dispersive analyzer analysis of the coarse particles, indicating that these coarse particles are chromium-rich chromium carbides. As described in Example 15 above, the presence of coarse particles at grain boundaries, and the associated non-precipitation regions, causes significant fracture in the ductility and strength of the alloy. Therefore, it is not recommended to add chromium alloy elements, which may strongly form carbides, to the welding filler wire disclosed in the present invention.

[0099] In the present invention, it is interesting and noteworthy that although titanium, niobium, vanadium, molybdenum, and chromium are all alloying elements that strongly form carbides, in the above examples (Examples 10-13, 15, and 16), these metals have very different effects on the microstructure of the molten region of the weld sample. In the case of titanium, niobium, and vanadium, the present invention shows that during fusion welding, these metals form titanium carbides rich in nano-order (6-10 nm) titanium, niobium carbides rich in niobium, and vanadium carbides rich in vanadium within the eutectic region. These metal carbides formed in the eutectic region during fusion welding not only significantly improve microhardness and yield strength but also maintain the extremely good malleability of the weld alloy in the molten region. Conversely, our tests have shown that when molybdenum and chromium are added to the welding filler wire, coarse, micrometer-sized particles of molybdenum-rich molybdenum carbide and chromium-rich chromium carbide are formed at the austenitic iron grain boundaries. Simultaneously, distinct non-deposited regions appear around the coarse carbides in the molten area, and both of these adversely affect the ductility of the alloy in the post-welded state.

[0100] It should be particularly emphasized here that the above examples are intended to illustrate the unprecedented properties that the welding filler wires disclosed in this invention bring about by the new alloy composition design, and are merely preferred embodiments of the invention, and do not limit the scope of the invention. Equivalent variations and modifications made based on the claims of this invention should all be included within the scope of the claims of this invention.

[0101] The above description illustrates preferred embodiments illustrating the technical features of the present invention. Those skilled in the art can make modifications and modifications without departing from the spirit of the invention, and such modifications and modifications are included within the scope of the claims of the present invention.

Claims

1. A welding filler wire for fusion welding using a hot-rolled precipitation-hardening austenitic iron-manganese-aluminum-carbon alloy (with more than 1.5% by weight of carbon) as the base material, A precipitation-hardening austenitic iron-manganese-aluminum-carbon alloy welding filler wire for financial welding, containing 23-34% manganese, 7.5-11.5% aluminum, and 1.35-1.95% carbon, with the remainder being iron, in weight percent.

2. A welding filler wire for fusion welding using a hot-rolled precipitation-hardening austenitic iron-manganese-aluminum-carbon alloy (with more than 1.5% by weight of carbon) as the base material, A precipitation-hardening austenitic iron-manganese-aluminum-carbon alloy welding filler wire for financial welding, containing 24-32% manganese, 8.0-11.0% aluminum, and 1.40-1.95% carbon, with the remainder being iron, in weight percent.

3. A welding filler wire for fusion welding using a hot-rolled precipitation-hardening austenitic iron-manganese-aluminum-carbon alloy (with more than 1.5% by weight of carbon) as the base material, A precipitation-hardening austenitic iron-manganese-aluminum-carbon alloy welding filler wire for financial welding, containing 23-34% manganese, 7.5-11.5% aluminum, 1.40-1.95% carbon, and 0.1-2.5% titanium, with the remainder being iron, in weight percent.

4. A welding filler wire for fusion welding using a hot-rolled precipitation-hardening austenitic iron-manganese-aluminum-carbon alloy (with more than 1.5% by weight of carbon) as the base material, A precipitation-hardening austenitic iron-manganese-aluminum-carbon alloy welding filler wire for financial welding, containing 23-34% manganese, 7.5-11.5% aluminum, 1.40-1.95% carbon, and 0.1-3.0% niobium, with the remainder being iron, in weight percent.

5. A welding filler wire for fusion welding using a hot-rolled precipitation-hardening austenitic iron-manganese-aluminum-carbon alloy (with more than 1.5% by weight of carbon) as the base material, A precipitation-hardening austenitic iron-manganese-aluminum-carbon alloy welding filler wire for financial welding, containing 23-34% manganese, 7.5-11.5% aluminum, 1.40-1.95% carbon, and 0.1-2.5% vanadium, with the remainder being iron, in weight percent.

6. A welding filler wire for fusion welding using a hot-rolled precipitation-hardening austenitic iron-manganese-aluminum-carbon alloy (with more than 1.5% by weight of carbon) as the base material, A precipitation-hardening austenitic iron-manganese-aluminum-carbon composite welding filler wire for financial welding, comprising 23-34% manganese, 7.5-11.5% aluminum, 1.40-1.95% carbon, and at least two elements selected from the group consisting of titanium, niobium, and vanadium, in a total amount of 3% or less, with the remainder being iron, in weight percent.

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