Electrode for fusion welding of austenitic Fe-Mn-Al-C alloy for precipitation hardening

Abstract

The present invention discloses a series of welding electrodes for fusion welding of precipitation hardening austenitic iron-manganese-aluminum-carbon alloys. The first type of welding electrode, by weight percentage, contains: 23-34 wt.% manganese, 7.5-11.5 wt.% aluminum, 1.35-1.95 wt.% carbon, and the balance is iron. After welding, it will produce high-density nanoscale (Fe,Mn)3AlC iron-manganese-aluminum carbide, which is uniformly dispersed in the austenite dendritic cells in the fusion zone, and the amount of nanoscale (~6-10 nm) κ carbide appearing in the eutectic zone will increase significantly, and the size of the austenite dendritic cells will decrease significantly. The second type of welding electrode, by weight percentage, contains: 23-34 wt.% manganese, 7.5-11.5 wt.% aluminum, 1.40-1.95 wt.% carbon, and 0.1-2.5 wt.% titanium and / or 0.1-3.0 wt.% niobium and / or 0.1-2.5 wt.% vanadium, and the balance is iron. There will be a large number of nanoscale (about 6-10 nm) titanium-rich titanium carbide and / or niobium-rich niobium carbide and / or vanadium-rich vanadium carbide with a face-centered cubic structure (FCC) and good ductility in the eutectic zone of the microstructure in the fusion zone. The hardness of these carbides is very high (2000~3500 Hv), thereby increasing the hardness of the fusion zone.

Description

Technical Field

[0001] The present invention relates to the technical field of welding electrodes, and particularly to a new welding electrode for joining parts made of lightweight precipitation-hardening Fe-Mn-Al-C alloys, especially for joining by fusion welding using gas tungsten arc welding (GTAW) and gas metal arc welding (GMAW). Background Art

[0002] Precipitation hardening (or age hardening) is a heat treatment technique used to increase the yield strength (YS) and hardness of ductile alloys. Examples of ductile alloys include the 7xxx-series aluminum alloys with the highest strength (such as AA7075), precipitation-hardening stainless steels, and precipitation-hardened austenitic Fe-Mn-Al-C alloys, etc. The precipitation hardening process involves three basic steps: solution heat treatment (SHT), quenching, and aging. The first step (solution heat treatment) is to heat the alloy above the solvus temperature and maintain it for a period of time until a uniform single-phase solid solution is formed. The second step (quenching) is to rapidly cool the solid solution from the solution heat treatment temperature to room temperature to form a supersaturated solid solution. In the as-quenched state, the microstructure of the alloy is a single-phase supersaturated solid solution without any precipitates. Therefore, although the alloy in the as-quenched state may have a good combination of ultimate tensile strength (UTS) and elongation (El), the yield strength (YS) is always significantly low. However, when the supersaturated solid solution is heated to a temperature below the solvus temperature and aged at this temperature for an appropriate period of time, high-density nano-sized precipitates will be coherently (or semi-coherently) generated and uniformly distributed in the matrix. Then, these high-density nano-sized and uniformly distributed coherent (or semi-coherent) second-phase precipitates will become effective obstacles to dislocation movement, thus greatly increasing the strength (especially the yield strength YS) and hardness of the alloy, while not significantly reducing the ductility. In other words, in order to obtain ideal mechanical strength, especially a significantly improved combination of yield strength YS, hardness, and elongation, these high-density nano-sized precipitates precipitated coherently (or semi-coherently) in the matrix are the most basic components.

[0003] Recently, lightweight precipitation-hardening fully austenitic Fe-Mn-Al-C alloys have been regarded as one of the most promising materials because of their excellent combination of mechanical strength and ductility, and can be widely used in various fields, such as the automotive industry, the armor industry, the mining industry, and so on. In particular, the addition of aluminum can further produce another significant advantage, that is, it can significantly reduce the density of traditional high-strength steel (about 1.3-1.5% reduction in density for every 1 wt% addition of aluminum). Generally, it is expected that this will have a significant impact on both energy conservation and reduction of carbon dioxide emissions.

[0004] According to previous various studies, for such alloys with an ideal combination of strength (especially yield strength YS) and ductility, the typical alloy microstructure basically consists of a fully austenitic matrix at room temperature and a large number of nanoscale Fe-Mn-Al carbides ((Fe,Mn) 3 AlC carbides (κ-carbide) precipitated in this austenitic matrix. The κ-carbide has an ordered face-centered cubic (FCC) L'1 2Crystal Structure. There is sufficient evidence showing that manganese is a strong austenite stabilizing element, while aluminum and carbon are essential elements for the formation of κ-carbides. The former (manganese) enables the alloy to have the face-centered cubic (FCC) crystal structure necessary for achieving high ductility at room temperature, while the latter forms the strengthening precipitates (κ-carbides) that play a major strengthening role in this type of alloy. Therefore, in order for the precipitation-hardened austenitic Fe-Mn-Al-C alloy to have an excellent combination value of strength and ductility, the manganese content in the precipitation-hardened austenitic Fe-Mn-Al-C alloy should preferably be greater than 18% by weight percentage. In this way, the alloy can have a face-centered cubic fully austenite substrate at room temperature, while at the same time, the aluminum and carbon contents are respectively greater than about 7% and 0.7% by weight percentage. Some recent studies have shown that for austenitic Fe-Mn-Al-C alloys with a chemical composition range of iron-(17.45 to 35.0) manganese-(7.1 to 12) aluminum-(0.7 to 1.2) carbon (Fe-(17.45 to 35.0)Mn-(7.1 to 12)Al-(0.7 to 1.2)C) (note: the chemical composition of all alloys cited in the present invention hereinafter is in weight percentage unless otherwise specified), after hot-rolled or hot-forged and solution heat-treated at 1050 - 1200 °C for 1 - 2 hours, and then rapidly quenched in water or oil solution to room temperature, the resulting alloy microstructure is a single austenite phase (γ-phase, γ-phase) without any precipitates. The yield strength, ultimate tensile strength, elongation, and microhardness ranges of the alloy in this quenched state are 350 - 540 megapascals (MPa), 700 - 1000 megapascals, 56 - 72%, and 200 - 230 Vickers hardness (Hv) respectively, and the actual values depend on the chemical composition of the alloy. Although the alloy in this quenched state has excellent ductility and good tensile strength, unfortunately, its yield strength is still quite low. Subsequent aging treatment is indispensable for improving the mechanical strength (especially the yield strength) of the alloy, and this aging treatment can promote the precipitation of nano-scale κ-carbides in the γ-substrate. Since the κ-carbides are rich in carbon and aluminum, the precipitation process of the κ-carbides in the supersaturated austenite substrate inevitably involves a large amount of diffusion processes of carbon and various related alloying elements. Therefore, it is usually necessary to optimize and moderate the aging treatment time and / or a higher aging treatment temperature. Previous reports have indicated that by performing aging treatment at 550 - 600 °C for 15 - 40 hours, a lightweight Fe-Mn-Al-C alloy with the best combination value of strength and ductility can be obtained.The optimized age-hardened Fe-Mn-Al-C alloy has a yield strength, ultimate tensile strength, and hardness that can reach 680 - 990 MPa, 995 - 1180 MPa, and 350 - 400 Vickers hardness, respectively, while the elongation remains in the range of 55 - 26.0%, and the actual value depends on the chemical composition of the alloy. According to existing research, compared with the alloy in the solution heat-treated state and after rapid quenching, the optimized age-hardened Fe-Mn-Al-C alloy (carbon content ≤ 1.2%) has an increase in ultimate tensile strength of approximately 180 - 295 MPa (18 - 42%) and an increase in yield strength of approximately 330 - 450 MPa (83 - 94%). This clearly shows that precipitation hardening plays an important 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 DO. 3 phases and β-Mn phases at the austenite grain boundaries, and these products will have an adverse effect on the ductility of this type of alloy.

[0005] To more specifically illustrate the important factors related to the development of such lightweight precipitation-hardened austenitic Fe-Mn-Al-C alloys, the present invention lists some cases published in recent years and elaborates on them in detail. These examples show that the alloys have significant improvements in mechanical strength (especially yield strength), while still maintaining good ductility (elongation greater than 25%). As reported by Gutierrez-Urrutia and Raabe, when alloys such as Fe-30.5Mn-8.0Al-1.2C are hot-rolled, solution heat-treated at 1100 °C for 2 hours, then quenched in water, and then aged at 600 °C for 24 hours, the yield strength, ultimate tensile strength, and elongation of the alloys are 990 MPa, 1180 MPa, and 37%, respectively. Such remarkable improvement in mechanical strength is attributed to the precipitation of nanoscale κ-carbides with a high volume fraction and uniform distribution in the austenite matrix during aging treatment. Similarly, as reported by Wu et al., after the 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 then quenched, the obtained alloy microstructure is a single austenite phase (γ-phase) without any precipitation, and its yield strength, ultimate tensile strength, and elongation are 485 MPa, 820 MPa, and 72%, respectively. After aging treatment at 550 °C for 40 hours, the microstructure of the obtained alloy becomes γ-phase plus κ-carbides (γ+κ-carbides). The volume fraction of κ-carbides precipitated and uniformly distributed in the austenite matrix can reach about 43%. More remarkably, the mechanical properties of the alloy are further improved, and its yield strength, ultimate tensile strength, and elongation are 955 MPa, 1040 MPa, and 38.2%, respectively. Compared with the mechanical properties of the alloy in the quenched state, it is obvious that the high-density κ-carbide precipitation promoted by aging treatment at 550 °C for 40 hours can improve the ultimate tensile strength by 26.8%, and the yield strength is even greatly increased, up to 96.9%. In another recent case, Haase et al. pointed out that after the Fe-29.8Mn-7.65Al-1.11C alloy is hot-rolled, solution heat-treated at 1150 °C for 5 hours, and quenched, the microstructure in the quenched state is a single austenite phase without any precipitation. The yield strength, ultimate tensile strength, and elongation of this quenched-state alloy are 540 MPa, 840 MPa, and 56%, respectively.When this quenched alloy is solution treated at 550 degrees Celsius for 15 hours, its microstructure mainly consists of γ-phase, along with nano-sized κ-carbides densely distributed in the austenite matrix. Its yield strength, ultimate tensile strength, and elongation are 880 MPa, 995 MPa, and 26% respectively. When maintaining a good elongation of 26%, the improvement rates of the yield strength and ultimate tensile strength reach 63% and 18.5% respectively. To allow a sufficient amount of nano-sized κ-carbides to precipitate in the γ-phase matrix while keeping no precipitation and / or second phases at the grain boundaries, a common practice is to conduct a long solution treatment (up to 15 to 40 hours) in the range of 550 to 600 degrees Celsius. This treatment process also promotes the combination of high strength (especially yield strength) and high ductility that is highly desired for this type of lightweight precipitation-hardened austenitic Fe-Mn-Al-C alloy.

[0006] Compared with the reference data cited above, for each representative reference data in which the carbon content in the Fe-Mn-Al-C alloy is less than 1.2 weight percent, please refer to two recent U.S. patents owned by the inventor of the present invention, Tzeng-Feng Liu, US 9,528,177 B2 approved in 2016, and US10,167,528B2 approved in 2019, which disclose a series of newly designed Fe-Mn-Al-C alloys, which are composed of (23-34) manganese, (6-12) aluminum, and (1.4-2.2) carbon, and the rest is iron, and its carbon content is relatively high. One of the excellent characteristics of this series of alloys is that after hot rolling, solution heat treatment at 980 to 1200 degrees Celsius for 1 hour, and then quenching treatment in water or ice water to rapidly cool to room temperature, the alloy microstructure in this quenched state already contains a high density of nanoscale κ-carbides formed in the austenite matrix by spinodal decomposition during quenching. This characteristic is in sharp contrast to the austenitic Fe-Mn-Al-C alloy with a carbon content below 1.2%. In the alloy with a lower carbon content, no κ-carbides were observed in the solution heat treatment state and the quenched state. The formation of a high density of nanoscale κ-carbides during quenching brings at least two major effects: The first effect is that the combination value of the mechanical strength (especially the yield strength) and ductility of the alloy in the quenched state is significantly improved, and its yield strength, ultimate tensile strength, and elongation are 865-925 MPa, 1030-1155 MPa, and 50-65% respectively; The second effect is that in the austenite matrix, the pre-existing high density of nanoscale κ-carbides can enable this type of alloy to obtain the best combination of mechanical properties with only a relatively short aging treatment time and a lower aging treatment temperature compared to the Fe-Mn-Al-C alloy with a carbon content below 1.2 weight percent described above. Taking the Fe-28.6Mn-9.84Al-2.05C alloy disclosed by Liu Zengfeng as an example, in the quenched state, high density nanoscale κ-carbides are formed in the austenite matrix by spinodal decomposition during quenching, and no precipitation is observed at the grain boundaries. The volume fraction and average size of the κ-carbides are approximately 45% and 12 nanometers (nm) respectively. The yield strength, ultimate tensile strength, and elongation of the quenched state alloy are 912 MPa, 1123 MPa, and 52.5% respectively. When the quenched state alloy is aged at 450 degrees Celsius for 6 hours, the volume fraction and average size of the κ-carbides increase to 53% and 25 nanometers respectively. In such a case, its yield strength, ultimate tensile strength, and elongation are 1179 MPa, 1306 MPa, and 39.8% respectively.

[0007] As can be seen from the above-described representative latest research examples in the art and the cited U.S. patents, the high-density nanoscale κ-carbides formed in the γ-phase substrate play an important role in improving the mechanical strength of the alloy (especially the yield strength), while still maintaining excellent ductility. In particular, as disclosed in the Liu Zengfeng patent, during quenching, the nanoscale κ-carbides that already existed in the austenite substrate have been proven to grow uniformly without triggering long-distance diffusion of aluminum and carbon, and the uniform growth of these nanoscale κ-carbides has also significantly improved the effectiveness of the aging treatment. Obviously, such lightweight precipitation-hardened austenitic Fe-Mn-Al-C alloys are the highly anticipated stars of tomorrow in the materials field and can be widely used in the automotive industry and high-performance structural materials required for military vehicles, shipbuilding, and the aviation industry. However, it is worth noting that in order to understand the promising application potential of such alloys, the most important thing is how to join the various different parts made of such high-strength and high-ductility precipitation-hardened austenitic Fe-Mn-Al-C alloys. Unexpectedly, despite the very urgent needs of users, there is extremely scarce information related to the weldability problems of precipitation-hardened austenitic Fe-Mn-Al-C alloys. This seems to be a common problem for all precipitation-hardened alloys (such as high-strength 7-series aluminum alloys), the reasons for which will be discussed in more detail later and are briefly summarized as follows:

[0008] (1) Since fusion welding involves the re-melting, re-solidification, and redistribution of alloying elements, the microstructure of the weld metal (fusion zone, FZ) will be very different from the microstructure of the original base material (the material being welded). In particular, in addition to dendritic microstructure and element segregation, the main strengthening precipitates obtained through solution heat treatment and aging treatment will completely dissolve, and the mechanical strength of the as-welded weldments will be significantly weakened. Therefore, after fusion welding, the hardness in the fusion zone of the as-welded weldments will always drop severely.

[0009] (2) Similarly, the high heat input during welding will cause the strengthening precipitates to dissolve and / or coarsening accompanied by grain growth in the heat-affected-zone (HAZ) of the weldment, all of which will significantly weaken the strength of the weldment.

[0010] (3)Generally speaking, precipitation hardening alloys usually contain a large amount of alloying elements. Therefore, in addition to a significant reduction in mechanical strength, such metals often tend to be very prone to hot cracking during welding, such as solidification cracking that occurs along the welding path, and liquation cracking that forms near the junction of the molten zone / weld heat-affected zone. In fact, as far as we know, there are only three relevant reports that evaluated the microstructural changes generated during fusion welding and the final obtained mechanical properties of alloys with chemical compositions falling within the scope of the precipitation hardening austenitic Fe-Mn-Al-C alloy of the present invention.

[0011] The following documents provide more detailed descriptions and discussions on the above-mentioned properties and characteristics.

[0012] [1]A.J. Ardell, “Precipitation Hardening”, Metall. Trans., 16A (1985) 2131 - 2165.

[0013] [2] “Precipitation hardening of aluminum alloys”, Totalmateria, 2010.

[0014] [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.

[0015] [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.

[0016] [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.

[0017] [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.

[0018] [7]E. Welsch, D. Ponge, S. M. Hafez Haghighat, S. 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.

[0019] [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.

[0020] [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.

[0021]

[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.

[0022]

[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 12wt.%) low-density steels”, J.Iron and Steels Res. Intern., 24(2017)1190-1198.

[0023]

[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.

[0024]

[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.

[0025]

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

[0026]

[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 lightweightsteels”,Metals&Mater.Int.,25(2019)683-696.

[0027]

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

[0028]

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

[0029]

[18] Tzeng-Feng Liu,“Composition design and processing methods of highstrength,high ductility and high corrosion resistance FeMnAlC alloys”,US10,167,528B2 / 2019.

[0030]

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

[0031]

[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.

[0032] To our knowledge, the earliest relevant research on the welding of alloys with the chemical composition of the austenitic Fe-Mn-Al-C alloys with precipitation hardening of the present invention was probably carried out by C.P. Chou and C.H. Lee in 1989. In their research, they examined the effect of carbon content on the changes in the solidification microstructure in the fusion zone of two fully austenitic alloys (Fe-30Mn-8.9Al-1.29C and Fe-29Mn-9.0Al-1.0C) processed by gas tungsten arc welding (GTAW). Autogenous weld is a form of welding in which the welding filler material (welding wire or electrode) can be from the melted base metal or other individual components. Both alloys were manufactured by the following methods: air induction melting, casting, hot-forged at 1200 °C, homogenized at 1050 °C for 12 hours, cold-rolled, and then annealed at 950 °C for 1 hour. The thickness of the experimental alloy was 1 / 8 inch. Before welding, both base alloys had a fully austenitic structure. In their research, it was found that the carbon content had a significant impact on the microstructure and morphology in the weld fusion zone. For the Fe-30Mn-8.9Al-1.29C alloy, the heat-affected zone of the weld had a fully austenitic phase (ferrite less than 1 vol.%), and there were cellular γ-dendrites and some κ-carbides in the eutectic zone. In this case, the microstructure in the fusion zone had a typical multi-branched dendritic structure, where the length of the primary dendrite extended up to several hundred micrometers (greater than 300 μm), and the length of the secondary dendrite arms was about 20 - 40 μm. It is worth noting that except for a small amount of eutectic κ-carbides, there were no κ-carbides or other precipitates in the primary and secondary dendritic cells, which might be due to rapid solidification during the welding process. As for the carbon content, it decreased to 1.0 wt.%, and a large amount of ferrite (about 10.2 vol.%) was observed in the fusion zone, mainly located between the secondary austenitic dendrite arms. Obviously, the weldment processed by fusion welding no longer had a homogeneous structure and could not fully maintain the characteristics of the base metal.Unfortunately, in this study, the mechanical properties related to the welded parts were not mentioned.

[0033] In 1990, Mr. Chou and Mr. Lee subsequently reported on the microstructure and mechanical properties of two austenitic alloys (Fe-29.3Mn-8.6Al-0.81C and Fe-29Mn-8Al-1.17C) welded by two butt-joint autogenous GTAW processes. Hereinafter, these two alloys will be referred to as 0.81C and 1.17C, respectively. These base alloys were first prepared in a vacuum induction furnace, and then the cast alloys were hot forged to a thickness of 8 mm at 1200 °C, followed by homogenization treatment at 1050 °C for 12 hours. The homogenized small pieces were cold rolled into 4-mm-thick plates, and then annealed in argon at 950 °C for 1 hour, followed by water quenching. The microstructure of the base alloys in the quenched state was completely austenite phase. The ultimate tensile strength and elongation of these two base alloys (0.81C and 1.17C) in the quenched state were 1040 MPa and 50%, and 1080 MPa and 54%, respectively. Compared with the above-mentioned Fe-Mn-Al-C alloys with a carbon content of less than or equal to 1.2% in the solution heat-treated and quenched state, the combined values of the ultimate tensile strength and elongation of these two types of alloys just fall within the range. Unfortunately, the yield strength of these alloys was not mentioned in the article. However, judging from the fact that no κ-carbide was observed in the austenite substrate, we can reasonably expect that the yield strength of these two types of alloys will be very similar to those of the solution heat-treated and quenched Fe-Mn-Al-C alloys with a carbon content of less than or equal to 1.2%, that is, 350-540 MPa cited above. After the autogenous GTAW process, the microstructure of the fusion zone of the butt weld has very similar general characteristics to the multi-branched dendritic structure described above. That is to say, the primary dendritic length extends up to several hundred micrometers (greater than 300 μm), and in addition to a small amount of eutectic κ-carbide, no κ-carbide or other precipitates appear in the primary and secondary austenite dendritic cells. At the same time, ferrite phases with about 5 and 0.5 volume percentages were observed in the fusion zones of 0.81C and 1.17C alloys, respectively. After post-welding heat treatment (PWHT) at 1050 °C for 5 to 240 minutes, no precipitates appeared in the primary and secondary austenite dendritic cells. Tensile tests were carried out on the butt-welded samples after post-welding heat treatment. The results showed that the ultimate tensile strength and elongation of the 0.81C alloy could reach 930 MPa and 25%, respectively, while those of the 1.17C alloy were 900 MPa and 19%, respectively.Obviously, the ultimate tensile strength value of the Fe-Mn-Al-C welded joint in the as-welded state can retain nearly 90% of the ultimate tensile strength value of the base alloy, while the elongation rate deteriorates severely, dropping from 54% (base metal) to about 20%. Additionally, it is worth noting that the yield strength of samples of both the base alloy and the butt-welded joint was not investigated in this study. However, as mentioned above, the absence of nano-scale precipitates in the austenite dendritic cells in the fusion zone indicates that the resulting welded joint has a low yield strength. These results show that the above treatment cannot achieve the goal of improving the yield strength of the welded joint in the as-welded state. As a result, until 2017 - 2019, developing special precipitation-hardening austenitic Fe-Mn-Al-C alloy electrodes for fusion welding was still regarded as the most urgent task in this field.

[0034] Based on the above-mentioned two studies by Zhou and Mr. Li, several key points are worth noting:

[0035] (1) In their studies, the initially used base alloy was in a fully austenite phase in the solution heat treatment and quenched state. In other words, there were no κ-carbides required for precipitation hardening of the austenitic Fe-Mn-Al-C alloy in the austenite substrate. In particular, fusion welding obviously significantly changes the microstructure of the fusion zone, transforming the homogeneous single-phase austenite substrate into a dendritic structure, with the formation of ferrite phase and eutectic phase (such as γ plus κ-carbide, γ+κ-carbid) in the inter-dendritic regions. Additionally, there are no κ-carbide precipitates in either the primary or secondary austenite dendritic cells in the fusion zone of the welded joint. This clearly indicates that the microstructure in the fusion zone is very different from that of the base alloy, even though both have the same chemical composition. Therefore, the mechanical properties of the resulting welded joint are somewhat unpredictable if judged only from the chemical compositions of the base alloy and the electrode.

[0036] (2) As described above, for this type of alloy, to obtain an excellent combination value of strength (especially yield strength) and ductility, a high-density nanoscale κ-carbide uniformly dispersed in the austenite matrix is an indispensable prerequisite. Therefore, during welding, the final mechanical properties of the welded component depend on how the fusion welding affects the change and distribution of κ-carbide in both the fusion zone and the heat-affected zone. Intuitively, one can expect that these possible situations would be very similar to those encountered in other precipitation-hardened alloys (such as AA7075 aluminum alloy), like the widespread severe softening phenomenon that commonly occurs in the fusion zone and the heat-affected zone. The main reason for the occurrence of severe softening is that the strengthening nanoscale precipitates dissolve, coarsen, and / or transform into other less effective incoherent phases. Such problems still pose a major challenge when using fusion welding to join various precipitation-hardened alloys, despite extensive research and development efforts to address these issues having been ongoing for over 50 years. However, as pointed out by Howell and Gerth in 2017, until now, there is still no formal method for direct welding of this type of "aged" lightweight precipitation-hardened austenitic Fe-Mn-Al-C alloy. The lack of a viable solution for using fusion welding to join the age-hardened precipitation-hardened austenitic Fe-Mn-Al-C alloy has become a major obstacle, preventing the widespread application of this type of alloy in various industrial fields.

[0037] As a preliminary experiment, in 2018, Sebeck et al. proposed the results of welding a refined age-hardening Fe-30Mn-9Al-1Si-0.9C-0.5Mo alloy by gas metal arc welding (GMAW), double vee groove butt, using a commercial 316LSi austenitic stainless steel electrode (composition: (11-14) nickel, (18-20) chromium, (1.0-2.5) manganese, (2.0-3.0) molybdenum, (0.65-1) silicon, and the rest iron). The base metal was hot-rolled to approximately 12.7 mm (about 0.5 inches) at 1204 °C and solution heat-treated at 1050 °C for 2 hours, followed by quenching. The yield strength, ultimate tensile strength, and elongation of the base metal after quenching were not mentioned in the report. However, its hardness was approximately 210 Hv Vickers hardness. After aging treatment at 538 °C for 30 hours, the microstructure of the resulting base metal was a fully austenitic phase, and κ-carbides were uniformly formed at the austenite substrate and grain boundaries. The typical yield strength, ultimate tensile strength, elongation, and hardness of the aged base metal were 800 MPa, 827 MPa, 36%, and approximately 360 Vickers hardness, respectively. In this study, since the metal solder used was 316LSi, the microstructure of the fusion zone was not clearly described, and the yield strength and elongation of the resulting welded joints decreased significantly to 350-400 MPa and 23-5%, respectively, which was clearly far from satisfactory. Therefore, this study particularly emphasized the necessity of developing a special metal solder for precipitation-hardening austenitic Fe-Mn-Al-C alloys.

[0038] From the limited existing literature and research reports described above, it can be determined that to date, the research conducted in this field still lacks a complete and comprehensive plan, and has not yet found a welding rod applicable to such highly anticipated precipitation-hardened austenitic Fe-Mn-Al-C alloys that can solve the related problems encountered during welding. However, based on the fragmentary information obtained from these previous studies, there is a particularly notable fundamental problem: since the alloy composition has the properties of rapid heating, remelting, solidification, and redistribution during fusion welding, through the design of the alloy composition of the welding rod, controlling the change of the microstructure to avoid severe softening of the fusion zone, these problem points must be solved. Especially for precipitation-hardened austenitic Fe-Mn-Al-C alloys, the most important strengthening component is the high-density nanoscale κ-carbide. Previous studies have found that whether this carbide can form in the austenite dendritic cells of the fusion zone will be severely inhibited due to the remelting and rapid cooling processes. Therefore, it is crucial and urgently needed to find a welding rod with an appropriate alloy composition design to address this issue. The challenges faced in developing a metal solder for fusion welding of precipitation-hardened austenitic Fe-Mn-Al-C alloys are similar to those encountered during fusion welding of precipitation-hardened aluminum alloys. Therefore, investigating the major problems that occur during fusion welding of precipitation-hardened aluminum alloys will inspire us based on past experience.

[0039] Among various aluminum alloys, the precipitation-hardening aluminum alloy 7075 (AA7075) is most favored by the aviation industry and the military because of its excellent specific strength, fracture characteristics, and high formability. AA7075 usually contains various alloying elements, and its chemical composition ranges as follows: aluminum - (5 - 6.5) zinc - (1.6 - 2.9) magnesium - (1.2 - 2.0) copper - (0 - 0.3) manganese - (0 - 0.28) chromium - (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 desired mechanical properties of AA7075 aluminum alloy, the most widely used heat treatment methods are T6 and / or T651. T6 is a common heat treatment code for heat-treating alloys and includes the following procedures: solution heat treatment within the α-solid solution phase field with a face-centered-cubic (FCC) structure, followed by rapid quenching to room temperature, and subsequent artificial aging treatment. For AA7075 aluminum alloy, the most common artificial aging treatment for T6 is aging treatment at 110 - 120 °C for 16 - 24 hours. Another code, T651, refers to further stress relief of the product after T6 treatment. The precipitation-hardening process of AA7075 aluminum alloy changes with the result of aging treatment at a certain temperature for a period of time, and its summary is as follows: supersaturated α-solid solution to GP zones to metastable η′-phase to equilibrium η-phase. The GP zones are completely coherent with the α-substrate with a face-centered-cubic structure, while the η′ precipitate and the α-substrate are only semi-coherent, but both are formed in aluminum grains. In this alloy system, due to the fully coherent (or semi-coherent) characteristics with the substrate, the uniformly distributed nano-scale GP zones and η′ precipitates are the main factors causing age hardening and strengthening. Although the exact values depend on the detailed composition, for AA7075 aluminum alloy treated with standard T6 and / or T651, its typical yield strength, ultimate tensile strength, elongation, and hardness are 459 - 539 MPa, 510 - 597 MPa, 8.5 - 14.6%, and 157 - 180 Vickers hardness, respectively.

[0040] Prolonging the aging treatment time or increasing the aging treatment temperature (e.g., overaging) promotes the equilibrium of η precipitates to form heterogeneously at the interfaces between dispersoid particles / substrate and grain boundaries, at the expense of the disappearance of GP zones and η precipitates. Since the η precipitates not only become coarser but also become incoherent with the substrate, overaging inevitably causes a significant decrease in hardness, mechanical strength, and ductility. However, despite extensive research and development on fusion welding of precipitation-hardened AA7075 aluminum alloy in the past few decades, no method has been found to completely solve the following problems, such as the significant reduction in strength (hardness) in the fusion zone, solidification cracks and liquation cracks occurring along the weld bead direction and at the interface between the fusion zone / weld heat-affected zone, porosity, etc., which are some common problems that occur when welding AA7075 aluminum alloy using traditional fusion welding methods. In the following text, we will briefly summarize the latest progress in dealing with these problems mentioned in some recently published and representative studies related to the use of automatic welding and / or electrodes made of heterogeneous materials.

[0041] As mentioned above, after being processed by the standard T6 and / or T651 procedures, the substrate of the AA7075 aluminum alloy base material contains a high density of nano-sized GP zones and η precipitates. However, regardless of whether automatic welding and / or electrodes made of heterogeneous materials (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)) are used, after fusion welding, the microstructure of the fusion zone generally consists of a typical dendritic crystal structure, plus a small amount of large eutectic precipitates (such as lamellar ηMg(ZnCuAl) in the eutectic zone and hardening nano-sized GP zones) 2 and θ(Al 2Cu)), and the η′ precipitates originally present in the base metal will completely dissolve in the molten zone. Among them, the hardened GP zones and η′ precipitates in the welding heat-affected zone will also dissolve or transform into less hard η′ precipitates. As a result, in the as-welded condition of the welded part, the molten zone and the welding heat-affected zone will be severely softened, and their hardnesses will only be 50 - 67% and 67 - 87% of the original base metal hardness respectively. Obviously, the softening degree of the molten zone seems more obvious compared with that of the welding heat-affected zone. At the same time, due to the redistribution of alloying elements during the fusion welding process, post-weld heat treatment obviously cannot effectively improve the micro-segregation of the alloy. Therefore, in the as-welded state and the post-weld aging state, the welding molten zone is still the most vulnerable area. In addition, due to the thermal stress and solidification shrinkage caused by the input of high heat, solidification cracks can always be measured in the welding zone during autogenous fusion welding. In this regard, in order to reduce the occurrence of solidification cracks, heterogeneous filler materials are often used to reduce the crack sensitivity of the alloy. However, while using heterogeneous filler materials to reduce the crack sensitivity of the alloy, the concentration of the strengthening alloying substances originally present in the AA7075 aluminum alloy also decreases. As a result, not only will the strength (or hardness) of the alloy decrease significantly in the as-welded state, but the reactivity of the alloy to post-weld heat treatment will also become worse, which will lead to various defects such as liquation cracks and pores near the fusion line (e.g., near the junction of the molten zone / welding heat-affected zone). Therefore, due to the above three main obstacles accompanying fusion welding, almost no tensile tests are carried out on the AA7075 aluminum alloy in the as-welded state. In fact, until recently (2019), the precipitation-hardened AA7075 aluminum alloy was still regarded as an alloy that could not be welded using traditional fusion welding methods.

[0042] In order to further address the crack problem caused by the fusion welding of AA7075 aluminum alloy, two research reports were recently (2019) proposed by researchers. Ipekoglu and Ken Two centimeters thick AA7075-T6 alloy plates were welded using ER5356 electrode and cold metal transfer gas shielded metal arc welding (CMT-GMAW). The main features of the cold metal transfer process include controllable material deposition and low heat input, which is assumed to suppress the formation of hot cracks. However, although cracks disappeared, large pores were observed in the molten zone. In addition, complete dissolution of nanoscale GP zones and η precipitates in the molten zone and η′→η transformation of precipitates in the weld heat affected zone were observed. As a result, the yield strength, ultimate tensile strength, elongation and hardness of the alloy deteriorated severely, from 539 MPa, 597 MPa, 14.6% and 175 Vickers hardness to incalculable (i.e., no yield strength), 312 MPa, 0.03% and 65 Vickers hardness (only 37.1% of the parent material), respectively. It is generally believed that the large number of large pores in the welded parts is the cause of brittle fracture. In this way, although the low temperature metal transfer method can repair cracks caused by molten welding, this method has the problem of hole formation, which will cause a significant decline in ductility (elongation), and the serious softening problem still needs to be solved.

[0043] On the other hand, in 2019, Sokoluk et al. used a newly developed electrode for welding AA7075 alloy, which was mixed with about 1.7 volume percent titanium carbide (TiC) nanoparticles (about 40-60 nanometers in size) into AA7075 alloy (aluminum-6.4 zinc-3.2 magnesium-1.2 copper-0.15 chromium) (Al-6.4Zn-3.2Mg-1.2Cu-0.15Cr). Their report showed that the addition of titanium carbide nanoparticles in the electrode apparently led to the following very important characteristics: (1) During the solidification process after arc welding, the presence of titanium carbide nanoparticles significantly slowed down the solidification in the front, thereby reducing the growth rate of dendritic grains. This slowed growth rate of dendritic grains in turn caused the solidification microstructure in the molten zone to become fine spherical grains, rather than the directional long dendritic grains commonly seen after typical fusion welding. (2) During the solidification process, titanium carbide nanoparticles remain in the magnesium (zinc, copper, aluminum) 2The second phase (eutectic), thus, the size, shape, and distribution of the second phase present in the eutectic region are effectively improved. Since there is no directional dendritic grain growth and the improvement of the second phase in the eutectic region, the hot cracking susceptibility of AA7075 alloy can be truly eliminated. The electrode treated with titanium carbide nanoparticles seems to have solved two insurmountable challenges, namely, the problems encountered by AA7075 alloy during fusion welding for more than 70 years, hot cracking and porosity. These amazing results are allegedly a breakthrough never seen before in the fusion welding of AA7075 alloy, and it seems to make AA7075 alloy a weldable alloy. However, in this study, the tensile test results show that for the welded parts in the as-welded state, the ultimate tensile strength and elongation are only 392 MPa and 1.5% respectively. However, through solution heat treatment at 480 °C and artificial aging treatment at 120 °C for 19 hours, the post-weld heat treatment can improve the ultimate tensile strength and elongation to 551 MPa and 5.21% respectively. The elongation values obtained for the welded parts in the as-welded state and the welded parts after post-weld heat treatment are 1.5% and 5.21% respectively. Obviously, both are still far lower than the typical elongation of the base metal T6-AA7075 (~8.5 - 14.6%). In addition, it is worth noting that the yield strength of the welded parts in the as-welded state and the welded parts after post-weld heat treatment was not mentioned in the study. The fusion zone microstructure mentioned in this study clearly shows that in both the welded parts in the as-welded state and the welded parts after post-weld heat treatment, the fusion zone is composed of a vast majority of α-dendrites and some small amounts of eutectics. More importantly, there are no precipitates in the dendritic cells in the fusion zone. Based on the precipitation hardening alloy strengthening mechanism described above, the lack of precipitates in most areas of the fusion zone will inevitably lead to severe softening in this area and will significantly affect the yield strength of the welded parts. Therefore, obviously, due to common problems such as fusion zone softening, solidification cracks, liquation cracks, and pore formation, the fusion welding of precipitation hardening AA7075 aluminum alloy remains the most challenging task in this field.

[0044] However, the specific cases described above strongly suggest that electrodes with appropriate alloy design are indispensable for the welding of precipitation hardening austenitic Fe-Mn-Al-C alloys. The following examples provided by the present invention clearly show that all the outstanding problems during the fusion welding of all precipitation hardening austenitic Fe-Mn-Al-C alloys can be simultaneously solved through the appropriate design of the chemical composition of the electrode. That is to say, by using the electrode provided by the present invention, excellent hardness, strength (especially yield strength), and ductility can be obtained in the fusion zone of the welded parts in the as-welded state, without forming solidification cracks, liquation cracks, and pores in this zone.

[0045] The following references provide more detailed descriptions and discussions of the above-described properties and characteristics.

[0046]

[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.

[0047]

[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.

[0048]

[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.

[0049]

[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-0.9Clight-weight steels”, Metall. Mater. Trans. A, 48 (2017) 4500.

[0050]

[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 NDIAGVSET symposium, Aug. 7-9, 2018, Novi, Michigan.

[0051]

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

[0052]

[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.

[0053]

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

[0054]

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

[0055]

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

[0056]

[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

[0057] As described in the background art section, precipitation hardening (or age hardening) is one of the most effective means of improving the yield strength and hardness of ductile alloys. Generally, precipitation hardening (or age hardening) type alloys, such as the strongest AA7xxx series aluminum alloys (AA7075), precipitation hardening stainless steels, and precipitation hardened austenitic Fe-Mn-Al-C alloys, after solution heat treatment, quenching treatment, and optimal aging treatment, high-density nanoscale precipitates will integrally (or semi-integrally) form within the substrate, thus causing a significant improvement in the yield strength and hardness of the alloy without serious impairment of ductility. However, as described in the prior art examples in the background art section, all these previously existing and strengthening nanoscale precipitates within the substrate due to aging treatment will completely dissolve during fusion welding. After fusion welding, the typical post-weld state microstructure of the fusion zone consists of a vast majority of dendritic cells and a small amount of eutectic regions. More importantly, there are no precipitates within the dendritic cells. As such, fusion welding inevitably leads to severe softening of the fusion zone. In addition, the fusion zone is the region most susceptible to hot cracking effects (such as solidification cracking and liquation cracking), and pores will also form. Therefore, for those skilled in the art, these three major problems still remain as obstacles to be overcome when fusion welding precipitation hardening type alloys, even if not insurmountable, but also extremely difficult challenges.

[0058] Similarly, after solution heat treatment, quenching treatment, and aging treatment, nanoscale κ-carbides with a high-density strengthening effect will form in the austenitic iron-manganese-aluminum-carbon alloy's ferrite matrix during precipitation hardening. Due to the precipitation of high-density nanoscale κ-carbides in the ferrite matrix, the strength (especially the yield strength) and hardness of the alloy will increase significantly, but the ductility will not be severely impaired. However, during the fusion welding process of the precipitation-hardened austenitic iron-manganese-aluminum-carbon alloy, similar to other precipitation-hardened alloys, problems such as severe softening in the fusion zone, regions affected by hot cracking effects (such as solidification and liquefaction cracks), and pores have remained to be solved for decades. In 2017, as pointed out by Moon et al., "The high carbon content (of the precipitation-hardened austenitic iron-manganese-aluminum-carbon alloy) makes welding of automotive structures difficult, and at the same time... poses a major obstacle to the wider application of the alloy." In 2017, the research by Howell and Gerth also pointed out that to date, there is still no welding method that can be officially used for such (precipitation-hardened austenitic iron-manganese-aluminum-carbon alloy) plates. In 2018, Sebeck et al. further pointed out, "The high magnesium content of such alloys (precipitation-hardened austenitic iron-manganese-aluminum-carbon alloys) is a major challenge for the welding process... For example, the compatibility of the welding rods is limited..." In 2018, Evans et al. called for, "There is a need to develop special metal solders for precipitation-hardened austenitic iron-manganese-aluminum-carbon alloys." It seems that the above points still remain the general consensus in this field.

[0059] To overcome all these outstanding problems encountered during the fusion welding of precipitation-hardened austenitic iron-manganese-aluminum-carbon alloys, the inventors of the present invention, based on decades of practical experience in materials research, including the design and technological development of iron-manganese-aluminum-carbon alloys, conducted many experiments on electrode design and proposed the present invention.

[0060] The feature of the present invention is that when using the electrode for fusion welding of precipitation-hardened austenitic iron-manganese-aluminum-carbon alloy provided by the present invention, in the post-weld state, the microstructure of the fusion zone and the welding bead have the following characteristics, which are not only unprecedented but also can completely solve the problems of fusion zone softening and hot cracking.

[0061] (1) In the as-welded condition, high-density nanoscale κ-carbides (about 3 - 5 nm) exist in the dendritic cells, and this unique structure is the most remarkable feature disclosed by the present invention. As described in the prior art, after fusion welding, the fusion zone in the as-welded condition will have excellent hardness and strength (especially yield strength), which is comparable to or even higher than those of the Fe-Mn-Al-C alloys with optimized age hardening and a carbon content not exceeding 1.2 wt%. This result is quite different from the results mentioned in the prior art related to the fusion welding of precipitation-hardened austenitic Fe-Mn-Al-C alloys (such as AA7075 precipitation-hardened aluminum alloy, precipitation-hardened stainless steel, and precipitation-hardened Fe-Mn-Al-C alloy, etc.). In the prior art, after the fusion welding of precipitation-hardened carbon alloys, it always causes the complete dissolution of the high-density strengthening (or hardening) nanoscale precipitates originally present in the substrate of the base metal. In the as-welded condition, the microstructure of the fusion zone is mainly composed of a vast majority of dendritic cells and a small amount of eutectic regions. More obviously, no precipitates can be observed in the dendritic cells in the fusion zone in the as-welded condition, which will cause severe softening of the fusion zone.

[0062] (2) In the as-welded condition, nanoscale κ-carbides exist at the tips of the dendritic cells within the austenite grains in the fusion zone, and this seems to retard the growth of the dendritic cells during solidification. As a result, the lengths and spacings of the austenite dendrite cells formed in the columnar austenite grains are only about 20 - 30 μm and about 5 - 10 μm respectively, unlike those observed in the fusion-welded AA7075 aluminum alloy and other Fe-Mn-Al-C alloys. After fusion welding, the microstructure of the fusion zone in these alloys usually contains a large proportion of long (greater than 200 - 300 μm) primary dendrites. In addition, the amount of nanoscale (about 6 - 10 nm) κ-carbides present in the eutectic region is much larger than the amount of carbides observed in the Fe-Mn-Al-C alloys with a carbon content not exceeding 1.29 wt% and processed by fusion welding in previous studies.

[0063] (3) By adding appropriate amounts of titanium, niobium, and vanadium to the welding electrode, in the as-welded state, in addition to the high-density nanoscale κ-carbides present in the austenite dendritic cells and eutectic regions, there are also a large number of titanium-rich Ti-carbides, niobium-rich Nb-carbides, and vanadium-rich V-carbides formed in the eutectic region. These carbides all have a ductile face-centered cubic structure (FCC) like austenite. In addition, these carbides have extremely high hardness, up to about 2000 - 3500 Vickers hardness. Therefore, the hardness of the fusion zone in the as-welded state will be significantly increased, while the ductility is not significantly reduced.

[0064] (4) In the as-welded state, all phases present in the fusion zone, including the substrate (austenite dendritic cells plus eutectic region) and all types of precipitates (including κ-carbides, titanium-rich Ti-carbides, niobium-rich Nb-carbides, and vanadium-rich V-carbides), have a ductile face-centered cubic structure (FCC). In addition, the size of all types of precipitates is only about 3 - 10 nanometers. Therefore, the tensile test results of the welded object show that the fusion zone in the as-welded state has excellent ductility.

[0065] (5) By using the welding electrode provided by the present invention, the problems of hot cracking and porosity (such as solidification and liquefaction cracks) that often occur during the fusion welding process of precipitation-hardened austenitic Fe-Mn-Al-C alloys can be completely solved. These hot cracks and porosity often appear in the weld bead, in the fusion zone, and / or in the area near the interface between the fusion zone and the welding heat-affected zone.

[0066] Due to the use of the welding electrode provided by the present invention, the microstructure of the fusion zone can have the above-mentioned characteristics. After fusion welding, the fusion zone of the welded object in the as-welded state will simultaneously have high hardness, high strength (especially yield strength), and high ductility. For example, the average microhardness value range of the ductile fusion zone obtained in the present invention falls within 365 - 465 Vickers hardness, which is much larger than the average microhardness value (350 - 400 Vickers hardness) of Fe-(17.45 - 35)Mn-(7.1 - 12)Al-(0.7 - 1.2)C after hot rolling, solution heat treatment, quenching treatment, and optimized aging treatment at 550° - 600°C described in the prior art content.

[0067] In order to achieve the goal of enabling the microstructure characteristics of the fusion zone and the weld bead to have the various characteristics described above in the post-weld state, the present invention has conducted a comprehensive investigation. Through a large number of experiments and analyses, the alloy design of this innovative welding electrode is adjusted and various melting and welding parameters of the alloy are changed. According to the present invention, in the welding electrode for the fusion welding of austenitic iron-manganese-aluminum-carbon alloy for precipitation hardening, the chemical composition ranges of the individual alloying elements are as follows: manganese (23 - 34% by weight, preferably 24 - 32% by weight); aluminum (7.5 - 11.5% by weight, preferably 8.0 - 11.0% by weight); carbon (1.35 - 1.95% by weight, preferably 1.40 - 1.95% by weight); titanium (0.0 - 2.5% by weight, preferably 0.1 - 2.5% by weight); niobium (0.0 - 3.0% by weight, preferably 0.1 - 3.0% by weight); vanadium (0.0 - 2.5% by weight, preferably 0.1 - 2.5% by weight); the combined value of titanium, niobium and vanadium (0 - 3.0% by weight, preferably 0.1 - 3.0% by weight); and the balance being iron.

[0068] Furthermore, the present invention provides the following series of welding electrodes for the fusion welding of austenitic iron-manganese-aluminum-carbon alloy for precipitation hardening, by weight percentage, comprising:

[0069] (1) 23 - 34% manganese, 7.5 - 11.5% aluminum, 1.35 - 1.95% carbon, and the balance being iron.

[0070] (2) 24 - 32% manganese, 8.0 - 11.0% aluminum, 1.40 - 1.95% carbon, and the balance being iron.

[0071] (3) 23 - 34% manganese, 7.5 - 11.5% aluminum, 1.40 - 1.95% carbon, 0.1 - 2.5% titanium, and the balance being iron.

[0072] (4) 23 - 34% manganese, 7.5 - 11.5% aluminum, 1.40 - 1.95% carbon, 0.1 - 3.0% niobium, and the balance being iron.

[0073] (5) 23 - 34% manganese, 7.5 - 11.5% aluminum, 1.40 - 1.95% carbon, 0.1 - 2.5% vanadium, and the balance being iron.

[0074] (6) 23 - 34% manganese, 7.5 - 11.5% aluminum, 1.40 - 1.95% carbon, and at least two elements selected from titanium, niobium and vanadium, with the total amount being less than or equal to 3%, and the balance being iron.

[0075] As for the reasons for the novel characteristics and unprecedented features obtained in the fusion zone in the post-weld state as disclosed in the present invention above, the main reason relies on in-depth research on the individual effects of each alloy element. More relevant details are described as follows:

[0076] (1) Manganese: Manganese is a strong austenite stabilizing element and plays the most important role in maintaining the ductile face-centered cubic structure phase. It has been confirmed and verified from existing research that an Fe-Mn-Al-C alloy with a manganese content greater than 17.5 wt% has a completely austenite microstructure at room temperature after hot rolling, solution heat treatment, and quenching treatment, as described in the prior art content. However, in Example 9 of the present invention, in the post-weld state after fusion welding, the electrode with a manganese content of 20.1 wt% has a large amount of ferrite phase present in the fusion zone of the weld bead, and its structure is body-centered-cubic (BCC). In order to obtain a completely austenite microstructure in the fusion zone in the post-weld state, the manganese content in the electrode must exceed 22 wt%. Therefore, in the electrode provided by the present invention, the concentration of manganese is set in the range of 23 - 34 wt%, preferably 24 - 32 wt%.

[0077] (2) Aluminum: Aluminum forms (Fe,Mn) 3One of the main elements of the AlC carbide (κ-carbide). As described in the prior art, the κ-carbide is the most important and useful component in strengthening the properties of precipitation hardening type Fe-Mn-Al-C alloys. Therefore, the aluminum content in the welding electrode plays an important role in triggering the most unique characteristics of the present invention. That is to say, after fusion welding, in the post-weld state, there will be a high density of nano-scale κ-carbides in the austenite dendritic cells in the fusion zone. The present invention designs a series of welding electrodes with different aluminum concentrations and conducts comprehensive observation and analysis on these welding electrodes. The results show that when the aluminum concentration is lower than 7.0% by weight percentage, in the post-weld state, no κ-carbides are formed in the austenite dendritic cells in the fusion zone. When the aluminum concentration in the welding electrode increases to more than 7.5% by weight percentage, in the post-weld state, a high density of nano-scale κ-carbides can be easily observed in the austenite dendritic cells in the fusion zone. However, when the aluminum concentration in the welding electrode increases to as high as 12% by weight percentage, in addition to the high density of nano-scale κ-carbides formed in the austenite dendritic cells and eutectic regions, some aluminum-rich particles will also appear at the austenite grain boundaries in the fusion zone. It is well known that the aluminum-rich secondary phase existing at the austenite grain boundaries will cause a serious deterioration in the ductility of the Fe-Mn-Al-C alloy. In Example 8 of the present invention, it is further observed that the aluminum-rich particles formed at the austenite grain boundaries in the fusion zone will cause the generation of solidification cracks during the welding process. Obviously, the aluminum concentration of the welding electrode of the present invention should be limited within the range of 7.5-11.5% by weight percentage, preferably within the range of 8.0-11.0% by weight percentage.

[0078] (3) Carbon: Carbon is obviously the formation of Fe-Mn-Al carbides ((Fe,Mn) 3One of the main elements of AlC, κ-carbide). Similar to the aluminum described above, the carbon content in the welding electrode is also crucial for forming the most unique characteristics of the present invention. That is to say, after fusion welding, in the post-weld state, high-density nano-scale κ-carbides will be present in the austenite dendritic cells in the fusion zone. The present invention has designed a series of welding electrodes with different carbon concentrations, and at the same time, a comprehensive observation and analysis of these welding electrodes have been carried out. The results show that when the carbon concentration is lower than 1.30% by weight percentage, no κ-carbides are formed in the austenite dendritic cells in the fusion zone in the post-weld state. When the carbon concentration in the welding electrode increases to more than 1.35% by weight percentage, high-density nano-scale κ-carbides can be easily observed in the austenite dendritic cells in the fusion zone in the post-weld state. However, when the carbon concentration in the welding electrode increases to more than 2.1% by weight percentage, it is found that in addition to the high-density nano-scale κ-carbides formed in the austenite dendritic cells and eutectic regions, some coarser κ-carbides also appear at the austenite grain boundaries in the fusion zone. It is well known that the coarser κ-carbides and their associated precipitation-free zones present at the austenite grain boundaries will have some serious negative impacts on the ductility of the Fe-Mn-Al-C alloy. In Example 7 of the present invention, it is further found that for a welding electrode with a carbon content of 2.2% by weight percentage, although the parameters related to the welding procedure are deliberately adjusted and various different parameter combination values are used, serious solidification hot cracks will still occur in the weld bead of the alloy. Therefore, the carbon content of the welding electrode of the present invention should be limited to the range of 1.35 - 1.95% by weight percentage, preferably in the range of 1.40 - 1.95% by weight percentage.

[0079] (4) Titanium, niobium, and vanadium: Titanium, niobium, and vanadium are very strong carbide-forming elements. This invention simultaneously investigated the effects of adding these elements to the welding electrodes provided by this invention on the microstructure and properties of the fusion zone of the welded parts in the post-weld state. The results showed that by adding an appropriate amount of titanium, niobium, and vanadium to the welding electrodes, the microstructure of the fusion zone in the post-weld state had the following distinct characteristics: (i) In addition to the high-density nanoscale κ-carbides formed in the austenite dendritic cells and eutectic regions, a large number of nanoscale titanium-rich titanium carbides, niobium-rich niobium carbides, and vanadium-rich vanadium carbides were also formed in the eutectic regions; (ii) The size of these eutectic carbides was only about 6 - 10 nanometers; (iii) These eutectic carbides, like the nanoscale κ-carbides formed in the austenite dendritic cells and eutectic regions, all had a ductile face-centered cubic structure; (iv) Adding these elements would cause a significant reduction in the size of the austenite dendritic cells; (v) Since these titanium-rich titanium carbides, niobium-rich niobium carbides, and vanadium-rich vanadium carbides were extremely hard (about 2000 - 3500 Vickers hardness), the presence of these carbides would cause an amazing strengthening (hardening) effect in the fusion zone, and the ductility would not be significantly reduced. It should be emphasized here that the most important feature of this invention is to enable the fusion zone to have a high yield strength and hardness, and the condition for this feature to hold is that there are high-density nanoscale κ-carbides in the austenite dendritic cells in the post-weld state. As described above, in order to obtain high-density nanoscale κ-carbides formed in the austenite dendritic cells in the fusion zone in the post-weld state, the carbon content of the welding electrode must exceed 1.35% by weight. However, comprehensive experimental results showed that when about 1.0% by weight of titanium was added to a welding electrode with a carbon content of 1.35% by weight, after fusion welding, no κ-carbides were formed in the austenite dendritic cells in the post-weld state. In that case, in order to obtain high-density nanoscale κ-carbides formed in the austenite dendritic cells in the post-weld state, the carbon content in the welding electrode must exceed 1.50% by weight. More experiments also showed that in order to obtain high-density nanoscale κ-carbides formed in the austenite dendritic cells in the post-weld state, when the total content of titanium, niobium, and vanadium exceeded 3.5% by weight, the carbon content in the welding electrode must exceed 2.2% by weight. But in that case, solidification hot cracks were often observed in the weld bead. However, by using the alloy design of this invention as shown above, the fusion zone of the obtained alloy would have the expected unique microstructure and excellent combined values of microhardness, yield strength, ultimate tensile strength, and ductility.That is to say, this clearly proves that this alloy design has unprecedented advantages in promoting mechanical strength (especially in yield strength) and improving the microhardness of the fusion zone in the post-weld state, while still maintaining excellent ductility.

[0080] (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 the welding electrode on the microstructure of the fusion zone in the post-weld state. The investigation shows that adding chromium and molybdenum to the welding electrode results in the formation of micron-sized, coarse molybdenum-rich molybdenum carbides and chromium-rich chromium carbides at the grain boundaries of austenite in the post-weld fusion zone. At the same time, a clearly visible precipitate-free zone surrounds the coarse carbides in the fusion zone, which has a negative impact on the ductility of the post-weld alloy. Therefore, in the present invention, it is not recommended to add chromium and molybdenum to the welding electrode.

[0081] Finally, it should be particularly noted that the base material used in the embodiments of the present invention is a hot-rolled high-carbon content (as-hot-rolled high carbon content) (carbon greater than 1.5% by weight percentage) Fe-Mn-Al-C alloy. Before fusion welding, the microstructure of the base material consists of a completely austenite matrix and high-density nanoscale κ-carbides with an L′1 2 structure uniformly distributed within the austenite matrix. Interestingly, after fusion welding, these high-density nanoscale κ-carbides do not show any significant changes in the heat-affected zone and the base material zone of the welded part. In other words, there is no significant dissolution or coarsening of these high-density nanoscale κ-carbides. In other precipitation-hardened alloys (such as AA7075), however, dissolution or coarsening of these carbides is often observed after fusion welding. Therefore, there is no obvious softening phenomenon in the heat-affected zone of the alloy, as described in the embodiments provided by the present invention below. Description of the Drawings

[0082] Figure 1(a) is a bright-field image of a transmission electron microscope micrograph of the base material in Example 1.

[0083] Figure 1(b) is a selected area electron diffraction pattern of the transmission electron microscope micrograph of the base material in Example 1, with the zone axis being

[001] ( hkl : Austenite; hkl: κ-carbide).

[0084] Figure 1(c) is a (100) κ dark-field image of the transmission electron microscope micrograph of the same area as shown in Figure 1(a) of the base material in Example 1.

[0085] Figure 1(d) is a macroscopic image of the weld bead obtained by the gas tungsten arc welding process of the welding electrode in Example 1.

[0086] Figure 1(e) is a scanning electron microscope image of the welded part in the post-weld state in Example 1 (FZ: fusion zone; HAZ: heat-affected zone of welding; BM: base metal; the inset shows the magnified image of the marked area in the figure).

[0087] Figure 1(f) is a scanning electron microscope image of the microhardness indentation performed in the fusion zone, heat-affected zone of welding, and base metal in Example 1.

[0088] Figure 1(g) shows the corresponding microhardness values measured at each measurement point shown in Figure 1(f) in Example 1.

[0089] Figure 1(h) is a bright-field image of a transmission electron microscope taken from the fusion zone of the welded part in Example 1. The solid and hollow arrows indicate κ-carbides in the austenite dendritic crystal cells and eutectic regions, respectively.

[0090] Figure 1(i) is a selected-area electron diffraction pattern taken from the fusion zone in Example 1. The zone axis is

[001] ( hkl : austenite; hkl: κ-carbide).

[0091] Figure 1(j) shows the analysis results of a transmission electron microscope and an X-ray energy dispersive analyzer (TEM-EDS) taken from the fusion zone of the welded part in Example 1.

[0092] Figure 1(k) shows the macroscopic images of the tensile test sample in Example 1 before and after the tensile test.

[0093] Figure 2(a) is a macroscopic image of the weld bead obtained by gas-shielded metal arc welding using an electrode in Example 2.

[0094] Figure 2(b) is a scanning electron microscope image of a welded part taken in the post-weld state in Example 2 (FZ: fusion zone; HAZ: heat-affected zone of welding; BM: base metal). The inset is the magnified image of the marked area in the figure.

[0095] Figure 2(c) is a scanning electron microscope image of the microhardness indentation performed in the fusion zone, heat-affected zone of welding, and base metal in Example 2.

[0096] Figure 2(d) shows the corresponding microhardness values measured at each measurement point in Figure 2(c) in Example 2.

[0097] Figure 2(e) is a bright-field image of a transmission electron microscope taken from the fusion zone of the welded part in Example 2. The solid and hollow arrows point to κ-carbides scattered in the austenite dendritic crystal cells and eutectic regions, respectively.

[0098] Figure 2(f) shows the selected area electron diffraction pattern taken from the molten zone in Example 2, with the zone axis being

[001] ( hkl : Austenite; hkl: κ-carbide).

[0099] Figure 2(g) shows the analysis results of the transmission electron microscope and X-ray energy dispersive analyzer obtained from the molten zone of the welded part in Example 2.

[0100] Figure 2(h) shows the macroscopic images of the tensile sample in Example 2 before and after the tensile test.

[0101] Figure 3(a) shows the bright field image of the transmission electron microscope micrograph of the base material after hot rolling in Example 3.

[0102] Figure 3(b) shows the selected area electron diffraction pattern of the transmission electron microscope micrograph of the base material after hot rolling in Example 3, with the zone axis being

[001] ( hkl : Austenite; hkl: κ-carbide)

[0103] Figure 3(c) shows the (100) taken from the same area as shown in Figure 3(a) of the transmission electron microscope micrograph of the base material after hot rolling in Example 3 κ dark field image.

[0104] Figure 3(d) shows a scanning electron microscope image of the welded part in the post-weld state in Example 3 (FZ: molten zone; HAZ: heat affected zone of welding; BM: base material), and the welding uses an electrode and tungsten inert gas arc welding.

[0105] Figure 3(e) shows a scanning electron microscope image of the microhardness test indentations performed in the molten zone, heat affected zone of welding, and base material in Example 3.

[0106] Figure 3(f) shows the corresponding microhardness values measured at each measurement point in Figure (e) of Example 3.

[0107] Figure 3(g) shows the bright field image of the transmission electron microscope taken from the molten zone of the welded part in Example 3, and the solid and hollow arrows point to the κ-carbides dispersed in the austenite dendritic crystal cells and eutectic regions respectively.

[0108] Figure 3(h) shows the selected area electron diffraction pattern taken from the molten zone in Example 3, with the zone axis being

[001] ( hkl : Austenite; hkl: κ-carbide).

[0109] Figure 3(i) shows the analysis results of the transmission electron microscope and X-ray energy dispersive analyzer obtained from the molten zone of the welded part in Example 3.

[0110] Figure 4(a) is a scanning electron microscope image of the welded part in the post-weld state in Example 4 (FZ: fusion zone; HAZ: heat-affected zone of welding; BM: base metal). The welding uses electrodes and gas tungsten arc welding.

[0111] Figure 4(b) is a scanning electron microscope image of the microhardness test indentations performed in the fusion zone, heat-affected zone of welding, and base metal in Example 4.

[0112] Figure 4(c) shows the corresponding microhardness values measured at each measurement point in Figure 4(b) of Example 4.

[0113] Figure 4(d) is a bright-field image of a transmission electron microscope of the fusion zone of the welded part in Example 4.

[0114] Figure 4(e) is a selected-area electron diffraction pattern taken from the fusion zone in Example 4. The zone axis is 001 ( hkl : austenite; hkl: κ-carbide).

[0115] Figure 4(f) shows the analysis results of a transmission electron microscope and an X-ray energy dispersive analyzer obtained from the fusion zone of the welded part in Example 4.

[0116] Figure 5(a) is a scanning electron microscope image of the welded part in the post-weld state in Example 5 (FZ: fusion zone; HAZ: heat-affected zone of welding; BM: base metal). The welding uses electrodes and gas tungsten arc welding.

[0117] Figure 5(b) is a scanning electron microscope image of the microhardness test indentations performed in the fusion zone, heat-affected zone of welding, and base metal in Example 5.

[0118] Figure 5(c) shows the corresponding microhardness values measured at each measurement point in Figure 5(b) of Example 5.

[0119] Figure 5(d) is a bright-field image of a transmission electron microscope taken from the fusion zone of the welded part in Example 5.

[0120] Figure 5(e) is a selected-area electron diffraction pattern taken from the fusion zone in Example 5. The zone axis is

[001] ( hkl : austenite; hkl: κ-carbide).

[0121] Figure 5(f) shows the analysis results of a transmission electron microscope and an X-ray energy dispersive analyzer obtained from the fusion zone of the welded part in Example 5.

[0122] Figure 6(a) is a scanning electron microscope image of the welded part in the post-weld state in Example 6 (FZ: fusion zone; HAZ: heat-affected zone of welding; BM: base metal). The welding was performed using an electrode and gas tungsten arc welding. The inserted figure shows the magnified image of the marked area in the figure.

[0123] Figure 6(b) is a scanning electron microscope image of the microhardness test indentations in the fusion zone, heat-affected zone of welding, and base metal in Example 6.

[0124] Figure 6(c) shows the corresponding microhardness values measured at each measurement point in Figure 6(b) of Example 6.

[0125] Figure 6(d) is a macroscopic image of the tensile specimen before and after the tensile test in Example 6.

[0126] Figure 7(a) is a macroscopic image of the weld bead obtained by gas tungsten arc welding using an electrode in Example 7. Note the solidification cracks in the weld bead.

[0127] Figure 7(b) is a scanning electron microscope image of the welded part in the post-weld state in Example 7 (FZ: fusion zone; HAZ: heat-affected zone of welding; BM: base metal). The inserted figure shows the magnified image of the marked area in the figure. The arrow indicates the coarse κ-carbide formed at the austenite grain boundary.

[0128] Figure 7(c) is a bright-field image of a transmission electron microscope taken from the fusion zone of the welded part in Example 7.

[0129] Figure 7(d) is a selected area electron diffraction pattern of the coarse κ-carbide marked as "K" in Figure 7(c) of Example 7. The zone axis is

[001] ( hkl : austenite; hkl: κ-carbide).

[0130] Figure 7(e) shows the analysis results of a transmission electron microscope and an X-ray energy dispersive analyzer for the coarse κ-carbide in Example 7.

[0131] Figure 8(a) is a macroscopic image of the weld bead obtained by gas tungsten arc welding using an electrode in Example 8. Note the solidification cracks in the weld bead (indicated by the arrow).

[0132] Figure 8(b) is a scanning electron microscope image of the welded part in the post-weld state in Example 8 (FZ: fusion zone; HAZ: heat-affected zone of welding; BM: base metal). The arrow indicates the aluminum-rich particles (denoted by α) formed at the austenite grain boundary.

[0133] Figure 8(c) shows the analysis results of a scanning electron microscope and an X-ray energy dispersive analyzer (SEM-EDS) for the discrete aluminum-rich particles in Example 8.

[0134] Figure 9(a) is a scanning electron microscope image of the welded part in the post-weld state in Example 9 (FZ: fusion zone; HAZ: heat-affected zone of welding; BM: base metal). The welding uses electrodes and gas tungsten arc welding, and the arrow indicates the discrete ferrite phase (denoted by α) scattered in the fusion zone.

[0135] Figure 9(b) is a bright-field image of a transmission electron microscope of the fusion zone of the welded part in Example 9. The discrete ferrite phase is denoted by α, and the circles mark the coarse κ-carbides formed at the ferrite / austenite boundary.

[0136] Figure 9(c) is a selected-area electron diffraction pattern of the area circled in Figure 9(b) in Example 9 ( hkl : austenite; hkl: κ-carbide; (hkl): ferrite).

[0137] Figure 9(d) shows the analysis results of a transmission electron microscope and an X-ray energy dispersive analyzer taken from the austenite zone in Example 9.

[0138] Figure 9(e) shows the analysis results of a transmission electron microscope and an X-ray energy dispersive analyzer taken from the ferrite zone in Example 9.

[0139] Figure 10(a) is a bright-field image of a transmission electron microscope image of the base metal after hot rolling in Example 10.

[0140] Figure 10(b) is a selected-area electron diffraction pattern of a transmission electron microscope image of the base metal after hot rolling in Example 10, and the zone axis is

[001] ( hkl : austenite; hkl: κ-carbide).

[0141] Figure 10(c) is a (100) κ dark-field image of the same area as shown in Figure 10(a) in a transmission electron microscope image of the base metal after hot rolling in Example 10.

[0142] Figure 10(d) is a macroscopic image of the weld bead obtained by gas tungsten arc welding with electrodes in Example 10.

[0143] Figure 10(e) is a scanning electron microscope image of the welded part in the post-weld state in Example 10 (FZ: fusion zone; HAZ: heat-affected zone of welding; BM: base metal).

[0144] Figure 10(f) is a scanning electron microscope image of the microhardness test indentations performed in the fusion zone, heat-affected zone of welding, and base metal in Example 10.

[0145] Figure 10(g) shows the corresponding microhardness values measured at each measurement point in Figure 10(f) of Example 10.

[0146] Figure 10(h) is a bright-field image of a transmission electron microscope taken from the fusion zone of the welded part in Example 10. The arrow points to κ-carbides and Ti-rich Ti-carbides scattered in the eutectic region.

[0147] Figure 10(i) is a selected-area electron diffraction pattern taken from the circled area in Figure 10(h) of Example 10 hkl : Austenite; hkl: κ-carbide; (hkl): Titanium carbide).

[0148] Figure 10(j) shows the analysis results of a transmission electron microscope and an X-ray energy dispersive analyzer taken from the same circled area in Figure 10(h) of Example 10.

[0149] Figure 10(k) shows the macroscopic images of the tensile sample before and after the tensile test in Example 10.

[0150] Figure 11(a) shows the macroscopic image of the weld bead obtained by gas-shielded metal arc welding using an electrode in Example 11.

[0151] Figure 11(b) is a scanning electron microscope image of the welded part in the post-weld state in Example 11 (FZ: Fusion zone; HAZ: Heat-affected zone of welding; BM: Base metal).

[0152] Figure 11(c) is a scanning electron microscope image of the microhardness test indentations made in the fusion zone, the heat-affected zone of welding, and the base metal in Example 11.

[0153] Figure 11(d) shows the corresponding microhardness values measured at each measurement point in Figure 11(c) of Example 11.

[0154] Figure 11(e) is a bright-field image of a transmission electron microscope taken from the fusion zone of the welded part in Example 11. The arrow points to κ-carbides and Nb-rich Nb-carbides scattered in the eutectic region.

[0155] Figure 11(f) is a selected-area electron diffraction pattern taken from the circled area in Figure (e) of Example 11 hkl : Austenite; hkl: κ-carbide; (hkl): Niobium carbide).

[0156] Figure 11(g) shows the analysis results of a transmission electron microscope and an X-ray energy dispersive analyzer taken from the same circled area in Figure 11(e) of Example 11.

[0157] Figure 11(h) shows the macroscopic images of the tensile specimens before and after the tensile test in Example 11.

[0158] Figure 12(a) shows the macroscopic image of the weld bead obtained by gas tungsten arc welding using an electrode in Example 12.

[0159] Figure 12(b) shows the scanning electron microscope image of the welded component in the post-weld state in Example 12 (FZ: fusion zone; HAZ: heat affected zone of welding; BM: base metal).

[0160] Figure 12(c) shows the scanning electron microscope image of the microhardness test indentations performed in the fusion zone, heat affected zone of welding, and base metal in Example 12.

[0161] Figure 12(d) shows the corresponding microhardness values measured at each measurement point in Figure 12(c) in Example 12.

[0162] Figure 12(e) shows the bright-field image of a transmission electron microscope taken from the fusion zone of the welded component in Example 12. The arrow points to the κ-carbides and V-rich V-carbides dispersed in the eutectic region.

[0163] Figure 12(f) shows the selected area electron diffraction pattern taken from the circled area in Figure 12(e) in Example 12 ( hkl : austenite; hkl: κ-carbide; (hkl): V-carbide).

[0164] Figure 12(g) shows the analysis results of a transmission electron microscope and an X-ray energy dispersive analyzer taken from the same circled area in Figure 12(e) in Example 12.

[0165] Figure 12(h) shows the macroscopic images of the tensile specimens before and after the tensile test in Example 12.

[0166] Figure 13(a) shows the macroscopic image of the weld bead obtained by gas tungsten arc welding using an electrode in Example 13.

[0167] Figure 13(b) shows the scanning electron microscope image of the welded component after welding in Example 13 (FZ: fusion zone; HAZ: heat affected zone of welding; BM: base metal).

[0168] Figure 13(c) shows the scanning electron microscope image of the microhardness test indentations performed in the fusion zone, heat affected zone of welding, and base metal in Example 13.

[0169] Figure 13(d) shows the corresponding microhardness values measured at each measurement point in Figure 13(c) in Example 13.

[0170] Figure 13(e) is a bright-field image of a transmission electron microscope taken from the fusion zone of the welded part in Example 13. The arrow points to κ-carbides, titanium-rich titanium carbides, and niobium-rich niobium carbides scattered in the eutectic region.

[0171] Figure 13(f) is the analysis result of a transmission electron microscope and an X-ray energy dispersive analyzer taken from the same circled area in Figure 13(e) in Example 13.

[0172] Figure 13(g) is a macroscopic image of the tensile sample in Example 13 before and after the tensile test.

[0173] Figure 14(a) is a scanning electron microscope image taken from the welded part after welding in Example 14 (FZ: fusion zone; HAZ: heat-affected zone of welding; BM: base metal). The welding uses a welding rod and gas tungsten arc welding. Note that the base metal is hot-rolled, solution heat-treated, quenched, and then aged at 550 °C for 12 hours.

[0174] Figure 14(b) is a scanning electron microscope image of the microhardness test indentations performed in the fusion zone, heat-affected zone of welding, and base metal in Example 14.

[0175] Figure 14(c) is the corresponding microhardness value measured at each measurement point in Figure 14(b) in Example 14.

[0176] Figure 14(d) is a macroscopic image of the tensile sample in Example 14 before and after the tensile test.

[0177] Figure 15(a) is a scanning electron microscope image taken from the welded part after welding in Example 15 (FZ: fusion zone; HAZ: heat-affected zone of welding; BM: base metal). The welding uses a welding rod and gas tungsten arc welding. The arrow points to thick Mo-rich Mo-carbides formed at the austenite grain boundaries.

[0178] Figure 15(b) is the analysis result of a scanning electron microscope and an X-ray energy dispersive analyzer of the thick Mo-rich Mo-carbides formed at the austenite grain boundaries in Example 15.

[0179] Figure 16(a) is a scanning electron microscope image taken from the welded part after welding in Example 16 (FZ: fusion zone; HAZ: heat-affected zone of welding; BM: base metal). The welding uses a welding rod and gas tungsten arc welding. The arrow points to thick Cr-rich Cr-carbides formed at the austenite grain boundaries.

[0180] Figure 16(b) shows the results of scanning electron microscopy and X-ray energy dispersive analyzer analysis of the coarse chromium-rich chromium carbides formed at the austenite grain boundaries in Example 16. Detailed implementation mode

[0181] The present invention will be further described below in conjunction with specific embodiments, and the advantages and features of the present invention will become clearer as the description progresses. However, these embodiments are only exemplary and do not constitute any limitation to the scope of the present invention. Those skilled in the art should understand that the details and forms of the technical solutions of the present invention can be modified or replaced without departing from the spirit and scope of the present invention, but such modifications and replacements all fall within the protection scope of the present invention.

[0182] In the following text, unless otherwise stated, the Fe-Mn-Al-C base materials with different compositions investigated in the present invention were prepared in an atmospheric induction furnace. After being homogenized at 1150 °C for 2 hours in a protective argon atmosphere, the metal ingots were hot-rolled from 80 mm thick to 8 - 12 mm thick plates, and then quenched in water to room temperature. The hot-rolled plates were then machined to form a single V-shaped joint, which was subsequently used as a welding joint with a groove angle of 60 degrees. The electrodes were prepared according to the designed alloy compositions in an atmospheric induction furnace. The molten metal was cast into a steel mold of 80 mm × 80 mm × 1500 mm. Then the metal ingots were heated to 1150 °C for 2 hours, and then hot-rolled into coils with a diameter of about 5 mm, and then pickled and cold-drawn to become electrodes with diameters of about 1.2, 2.4, and 3.2 mm respectively. Fusion welding was carried out by the gas metal arc welding (GMAW) or gas tungsten arc welding (GTAW) process, using pure argon or a mixed gas of argon and helium (He) as the shielding gas. For gas tungsten arc welding and gas metal arc welding, the typical flow rates of the shielding gas are about 10 and 18 liters per minute respectively. In the gas metal arc welding process, the electrode feed rate is about 200 mm per minute, and the voltages and currents used are 22 - 26 volts (V) and 140 - 170 amperes (A) respectively. For the gas tungsten arc welding process, the voltages and currents used are 15 - 20 volts and 125 - 200 amperes, depending on the diameter of the electrode used. In all the examples, the root opening was maintained at about 1 mm during fusion welding. The microstructure of the welded parts was analyzed using an optical microscope (OM), a scanning electron microscope (SEM), and a transmission electron microscope (TEM) respectively. The preparation of the scanning electron microscope specimens was preceded by mechanical polishing before etching with 5% nitric acid. For transmission electron microscope observation, the thin film specimens were prepared by a twin jet electro-polisher, and the composition of the electrolyte used was 60% acetic acid, 30% ethanol, and 10% perchloric acid.When the user needs it, an Energy Dispersive Spectrometer (EDS) attached to a scanning electron microscope or a transmission electron microscope is used to analyze the composition of an alloy or a specimen. It should be noted that the results of quantitative analysis of elements with an atomic number less than 10 (such as carbon) using an X-ray energy dispersive analyzer are unreliable, so the results are often ignored. The specimens for the tensile test are prepared according to the ASTM E8 standard (Standard Test Method for Tensile Testing of Metallic Materials) of the American Society for Testing and Materials (ASTM). The gauge length of the specimen is 25 mm, the gauge width is 6.25 mm, and the thickness is 4 mm. The tensile test is carried out at room temperature using an Instron 8501 tensile testing machine at a strain rate of 6.7x10. -4 s -1 The tensile test is carried out at a strain rate of 6.7x10. The tensile strength is measured at an offset strain of 0.2%. The hardness of different regions of the welded joint, such as the fusion zone, the heat-affected zone of the weld, and the base metal, is measured using a Vickers microhardness tester to summarize its characteristics. When performing the Vickers hardness test, a load of 100 grams-force (gf) is applied, and the load duration is 15 seconds.

[0183] Example 1

[0184] Figure 1(a) is a bright field (BF) image of a hot-rolled Fe-28.5Mn-9.0Al-1.55C base metal observed using a transmission electron microscope, showing that high-density nanoscale (about 3-5 nm) precipitates are evenly distributed in the matrix. As indicated by the arrow in Figure 1(a), it should also be noted that the precipitates present at the dislocations are slightly larger in size (about 5-8 nm), indicating that these precipitates are formed during the hot-rolling process. Figure 1(b) is the selected-area diffraction pattern (SADP) of the hot-rolled base metal, clearly showing that these high-density nanoscale precipitates are iron-manganese-aluminum carbides ((Fe,Mn) 2 AlC, κ-carbides) with an L′1 3 structure. Figure 1(c) is the (100) of the same region as shown in Figure 1(a). κDarkfield (DF) images showing the austenite substrate and the κ-carbide at the dislocations. This figure clearly shows that the microstructure of the base material used in this example after hot rolling is single-phase austenite with a high density of nanoscale κ-carbides uniformly dispersed therein. The tensile test results show that the yield strength, ultimate tensile strength and elongation of the hot-rolled base material are 952 MPa, 1100 MPa and 56%, respectively.

[0185] Two hot-rolled base material plates with dimensions of 80 mm x 80 mm x 8 mm were machined to form a single V-groove butt weld. In this example, the nominal composition of the welding electrode is Fe-27.8Mn-9.1Al-1.86C, and the diameter (φ) is about 2.4 mm. Fusion welding was carried out by tungsten inert gas arc welding process, using pure argon as the shielding gas. Figure 1(d) is a macroscale image of the weld bead, showing no visible solidification cracks or pores in the fusion zone of the post-weld sample.

[0186] Figure 1(e) shows a typical scanning electron microscope image of a region within the post-weld sample that encompasses the fusion zone, the weld heat-affected zone, and the base metal. Several characteristics that are distinct from those of existing fusion-welded precipitation-hardened Fe-Mn-Al-C alloys (carbon content not greater than 1.29 weight percent, C≤1.29 wt.%) are immediately noticeable in this figure: (i) It is clearly visible that within the fusion zone, there are typical columnar austenite grains, which consist of austenite dendritic cells and eutectic regions. More importantly, within the austenite dendritic cells and eutectic regions, there are high-density nanoscale κ-carbides. This is one of the many remarkable features first disclosed by the present invention. It has been pointed out in the prior art that in fusion-welded Fe-Mn-Al-C alloys (carbon content not greater than 1.29 weight percent), no precipitates have been observed either in the primary or secondary austenite dendrites, except for a small amount of eutectic κ-carbides present in the eutectic regions. Additionally, in this example, the amount of nanoscale κ-carbides present in the eutectic regions is much greater than the amount of carbides observed in alloys with a carbon content not greater than 1.29 weight percent in previous studies. (ii) From Figure 1(e), it is obvious that there are no significantly long primary dendrites in the columnar austenite grains. The lengths and spacings of the austenite dendritic cells formed within the columnar austenite grains are approximately 20 - 30 microns and about 5 - 10 microns, respectively. Clearly, the austenite dendritic cells in this example are significantly refined. This is quite different from what can be observed everywhere in existing or traditional fusion welding. Among them, the microstructure of the post-weld bead always consists of visibly long primary dendrites, and the length of the dendrites extends up to hundreds of microns (e.g., greater than 300 microns). The numerous reasons for these differences are explained as follows. As shown in the magnified image of the marked area in the upper right corner of Figure 1(e), it can be clearly seen that in addition to the fine eutectic κ-carbides formed between the austenite dendrites, there are precipitates at the tip of each austenite dendritic cell. During solidification after fusion welding, the stability of the solid-liquid interface depends on the thermal state and constitutional supercooling state present near the interface, and the size of the dendrites within the bead mainly depends on the cooling rate and the alloy composition. In this case, the precipitates that appear at the tips of the austenite dendritic cells due to constitutional supercooling seem to slow down the solidification front and effectively hinder the growth of the dendritic cells, thus leading to much shorter dendritic cells as shown in Figure 1(e). Therefore, the composition of the welding electrode seems to be a more important reason for the short and narrow dendritic structure observed in this example.(iii) The molten zone in this embodiment is a completely austenitic phase. However, in the prior art, after autogenous fusion welding, the solidified microstructure often contains a certain amount of ferrite phase, even if the base material is completely austenitic before welding. Additionally, Fig. 1(e) also shows that no significant microstructural changes were observed in the weld heat-affected zone and the base material zone. Moreover, for the κ-carbides in the molten zone, the weld heat-affected zone, and the base material zone in this embodiment, there were no perceivable differences in terms of size and density. Therefore, there is no softening phenomenon in the weld heat-affected zone. In addition, it should be noted that no trace of microcracks or porosity was observed in the molten zone or near the interface between the molten zone and the weld heat-affected zone, indicating that the use of the molten electrode of the present invention can significantly eliminate solidification and liquefaction cracks, both of which are common problems during the fusion process of precipitation-hardening alloys. Fig. 1(f) shows the scanning electron microscope images of the microhardness test indentations carried out in the molten zone, the weld heat-affected zone, and the base material zone. The corresponding values measured at each measurement point are shown in Fig. 1(g). As shown in Fig. 1(g), the average microhardness of the molten zone (about 425 Vickers hardness) is actually slightly greater than the average microhardness of the weld heat-affected zone (about 417 Vickers hardness) and the base material zone (about 412 Vickers hardness). This is in strong contrast to the so-called common sense that the molten zone is always the weakest area after fusion welding. It is worth mentioning here that the microstructure of the base material zone used in this embodiment in the hot-rolled state is a completely austenitic phase with a high density of nanoscale κ-carbides uniformly dispersed in the austenite matrix, and the microhardness of the base material zone is significantly greater than the hardness (350 - 400 Vickers hardness) of an Fe-Mn-Al-C alloy with optimized aging treatment and a carbon content not exceeding 1.2 wt%, as described in the prior art.

[0187] In order to fully understand the main reasons leading to these unprecedented results, a more detailed analysis of the microstructure of the molten zone must be carried out. Transmission electron microscope inspection shows that in this embodiment, the substrate of the molten zone is a completely austenite phase, and both the nano-scale precipitates in the austenite dendritic cells and the eutectic region are κ-carbides. As shown in Fig. 1(h), which is a bright-field electron micrograph (BF electron micrograph) taken from the molten zone. This figure clearly shows the κ-carbides in the austenite dendritic cells (as indicated by the solid arrows) and the κ-carbides in the related eutectic region (as indicated by the hollow arrows), with their sizes belonging to the 3-5 nm and 6-10 nm grades respectively. The length of the austenite dendritic cells is approximately in the 100-300 nm grade. This shows that the alloy design of the welding rod of the present invention can effectively inhibit the growth rate of the front of the dendritic cells during the solidification process after fusion welding. As shown in Fig. 1(i), it is confirmed by the selected area electron diffraction pattern that only the austenite phase and κ-carbides exist in the molten zone of the welded part. In addition, as shown in Fig. 1(j) by the analysis results of the transmission electron microscope and X-ray energy dispersive analyzer (TEM-EDS) obtained from the molten zone, the contents of iron, manganese and aluminum in the molten zone are 64.3%, 26.7% and 9.0% respectively.

[0188] The tensile test shows that the welded parts obtained using the welding rod provided in this embodiment exhibit excellent ductile deformation behavior, with a yield strength, ultimate tensile strength, and elongation rate of approximately 946 MPa, 1086 MPa, and 44%, respectively. Figure 1(k) shows the macroscopic images of the sample before and after the tensile test. It is obvious that the welded parts as a whole exhibit typical characteristics of ductile plastic deformation. In particular, the fusion zone also shows typical characteristics of plastic deformation, with a corrugated surface and a considerable elongation. Since the tensile sample contains regions with various mechanical properties (i.e., the fusion zone and the base metal zone), it is usually difficult to directly infer the strengths of the fusion zone and the base metal zone from the measured yield strength. Nevertheless, by judging from the width changes in the various regions of the sample broken by the tensile test, it may be possible to distinguish the relative values of the yield strengths between the fusion zone and the base metal zone. For example, if the base metal zone has a higher yield strength, when the tensile stress exceeds the yield strength of the fusion zone but is still lower than the yield strength of the base metal zone, then the fusion zone will be the first to exhibit plastic deformation (such as becoming longer and thinner), and there will be a width reduction phenomenon, while the base metal zone remains within the elastic regime, so its width remains unchanged, and vice versa. Therefore, when there is a significant difference in the yield strengths of the fusion zone and the base metal zone, there will be a certain degree of width change in the various regions of the sample broken by the tensile test, and the edges of the sample will not be parallel. In contrast, if the widths of the various regions in the sample broken by the tensile test are still approximately the same, and the edges of the sample are basically parallel, it can be reasonably inferred that the fusion zone and the base metal zone have approximately the same yield strength. The fact is that the widths across the regions including the base metal zone, the welding heat-affected zone, and the fusion zone (labeled as BM, INT, and FZ in Figure 1(k), respectively) are basically equal, and the edges of the sample across this zone remain parallel, indicating that the fusion zone and the base metal zone of this tensile test sample have approximately the same yield strength. It is particularly worth emphasizing that the extraordinary yield strength present in the fusion zone in this embodiment is completely expected. It is well known that the presence of high-density nanoscale κ-carbides will significantly improve the yield strength of the material without a significant decrease in ductility, which is also the main reason why most precipitation-hardened alloys have excellent combinations of yield strength and elongation rate. This is exactly the phenomenon seen in Figure 1(h). On the other hand, in Figure 1(k), it is clearly visible that necking and final fracture (indicated by the arrow) occur in the region on the other side of the base metal zone, indicating that the ultimate tensile strength of the fusion zone is greater than that of the base metal zone, which may also be consistent with the slightly lower hardness of the base metal zone described in Figure 1(g).

[0189] In fact, the high-density nanoscale κ-carbides present in the austenite dendritic cells should be the main reason for the high hardness and high strength (especially the yield strength) obtained in the molten zone of this example, but still maintain high ductility. The microstructure of the molten zone described in this example is undoubtedly the most unique and unprecedented characteristic of the present invention, and this characteristic is derived from the composition design of the welding rod provided by the present invention.

[0190] Example 2

[0191] The purpose of this example is to attempt to clarify the influence of changing the carbon content in the welding rod on the characteristics of the molten zone of the sample in the post-weld state. The base material used in this example is the same as that used in Example 1. However, in this example, the gas shielded metal arc welding process is applied. The dimensions of the base material plate are 80 mm × 80 mm × 12 mm, and a single V-groove butt joint is formed after machining. The gas shielded metal arc welding process is carried out, 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 about 18 liters per minute, and the welding rod feeding speed is about 200 mm per minute. The standard composition of the welding rod used in this example is iron-28.2 manganese-9.2 aluminum-1.72 carbon (Fe-28.2Mn-9.2Al-1.72C), and the diameter is about 1.2 mm.

[0192] Figure 2(a) shows the macroscopic image of the weld bead obtained by gas shielded metal arc welding. As can be seen from Figure 2(a), the weld bead has a very smooth morphology, and no visually visible macroscopic cracks and holes appear throughout the weld bead. Figure 2(b) shows the scanning electron microscope image of a region in the welded sample including the molten zone, the welding heat affected zone, and the base material zone. Similar to that shown in Example 1, the molten zone is composed of typical columnar austenite grains, and at the same time, there are high-density nanoscale κ-carbides in the austenite dendritic cells and the eutectic region. As shown by the arrow in Figure 2(b) and the enlarged 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 also κ-carbides at the tip of each austenite dendritic cell. This shows that the microstructure of this molten zone indeed has very similar characteristics to those observed in Example 1. In addition, it is worth noting that no traces of microcracks and holes are observed in the molten zone and near the interface between the molten zone and the welding heat affected zone. Similar to Example 1, Figure 2(b) also shows that the fusion welding process does not seem to cause the dissolution and significant coarsening of the κ-carbides in the welding heat affected zone of this example.

[0193] Figure 2(c) shows a scanning electron microscope image of the Vickers microhardness test indentations made across the fusion zone, the weld heat-affected zone, and the base metal zone. The corresponding microhardness values measured at each measurement point are shown in Figure 2(d). The average microhardness values of the fusion zone, the weld heat-affected zone, and the base metal zone are approximately 410 Vickers hardness, 412 Vickers hardness, and 413 Vickers hardness, respectively. Obviously, the microhardness of the fusion zone and the weld heat-affected zone is almost the same as that of the base metal zone, indicating that no significant softening has occurred in the fusion zone and the weld heat-affected zone, which is also consistent with the unique microstructure shown in Figure 2(b). The main strengthening component, namely the high-density nanoscale κ-carbide, remains basically intact in these three zones after fusion welding.

[0194] Examination of the fusion zone using a transmission electron microscope revealed a large number of κ-carbides formed within the austenite dendritic cells and eutectic regions, as shown in Figure 2(e). In this bright-field image of the transmission electron microscope, high-density nanoscale κ-carbides (about 3-5 nm) can be clearly seen dispersed within the austenite dendritic cells (indicated by solid arrows) and eutectic regions (indicated by hollow arrows). Compared with the results described in Example 1 (Figure 1(h)), the density of the κ-carbides formed within the austenite dendritic cells and eutectic regions decreased slightly. Figure 2(f) is a selected area electron diffraction pattern taken from Figure 2(e), which confirmed that the base of the fusion zone is a single austenite phase, and the precipitates present within both the austenite dendritic cells and eutectic regions are κ-carbides, which have a regular L′1 2 structure. Figure 2(g) shows the analysis results of the transmission electron microscope and X-ray energy dispersive analyzer obtained from the fusion zone, which clearly shows that the weight percentage concentrations of iron, manganese, and aluminum components in the fusion zone are approximately 63.4%, 27.5%, and 9.1%, respectively.

[0195] The tensile test shows that the welded parts obtained by welding with the welding electrode of this embodiment also have excellent ductile deformation characteristics, and their yield strength, ultimate tensile strength and elongation are approximately 942 MPa, 1075 MPa and 36% respectively. Figure 2(h) shows the macroscopic images of the sample before and after the tensile test. Obviously, the whole welded part has the typical characteristics of ductile plastic deformation. In particular, a zig-zag weaving fracture pattern and a corrugated fracture surface appear in the fusion zone, which are the ductile deformation characteristics. In addition, interestingly, as described in Example 1, it can be observed that the widths of the regions including the base metal zone, the fusion zone and the interface between them (denoted as BM, FZ and INT respectively in Figure 2(h)) are basically the same, and the edges of the sample across this region remain parallel, which indicates that the fusion zone and the base metal zone of this tensile test sample have approximately the same yield strength. The main reason for the relatively high yield strength in the fusion zone is the same as that described in Example 1. However, in this case, different from the phenomenon observed in Example 1, there is slight necking and fracture in the fusion zone rather than in the base metal zone in this embodiment, which indicates that the fusion zone has a slightly lower ultimate tensile strength compared with the base metal zone. In fact, this is consistent with another phenomenon, that is, the microhardness value in the fusion zone is slightly lower (about 410 Vickers hardness) compared with the base metal zone (the microhardness value is about 410 Vickers hardness). Based on the results shown in Example 1 and Example 2, it is obvious that the post-weld samples obtained by using the welding electrode with the composition design provided by the present invention still obviously have excellent ductility, extraordinary yield strength and microhardness even after remelting, alloy element redistribution and solidification during fusion welding.

[0196] Example 3

[0197] In this embodiment, we investigated what kind of influence a slightly lower carbon content would have on the characteristics of the molten zone in the sample in the post-weld state. In this embodiment, a base material that had been hot-rolled and had a standard composition of iron-28.8 manganese-8.9 aluminum-1.62 carbon (Fe-28.8Mn-8.9Al-1.62C) was used. Fig. 3(a) is a bright-field image taken by a transmission electron microscope of the base material after hot rolling. It is obvious in Fig. 3(a) that the microstructure of the hot-rolled base material used in this embodiment is quite similar to that used in Embodiment 1 and Embodiment 2. That is to say, there are high-density nanoscale precipitates (about 3-5 nm) uniformly distributed in the matrix and at the dislocations (as indicated by the arrows). Fig. 3(b) is the selected area electron diffraction pattern of the hot-rolled base material, clearly confirming that the matrix of the base material is a completely austenite phase, and these high-density nanoscale precipitates are κ-carbides with a regular L′1 2 structure. Fig. 3(c) is a (100) κ dark-field image taken from the same area as shown in Fig. 3(a), showing the appearance of high-density nanoscale κ-carbides. The investigation using a transmission electron microscope showed that the microstructure of the hot-rolled base material used in this embodiment is a single-phase austenite with uniformly distributed high-density nanoscale κ-carbides. And the tensile test showed that the yield strength, ultimate tensile strength and elongation of the hot-rolled base material are 975 MPa, 1152 MPa and 50% respectively.

[0198] The dimensions of the two hot-rolled base metal plates are 80 mm x 80 mm x 8 mm, and a single V-groove butt weld is formed after machining. The nominal composition of the welding rod in this example is Fe-28.2Mn-10.6Al-1.62C, with a diameter of approximately 2.4 mm. All the parameters of gas tungsten arc welding are similar to those used in Example 1. Figure 3(d) shows a typical scanning electron microscope image of the post-weld sample including regions such as the fusion zone, the heat-affected zone of the weld, and the base metal zone. Similar to what was seen in Example 1 and Example 2, typical columnar austenite grains exist in the fusion zone, which are composed of austenite dendritic cells and the eutectic region. At the same time, the true characteristics of the high-density nanoscale κ-carbides distributed in the austenite dendritic cells and the eutectic region can be observed. However, compared with Example 1 (Figure 1(e)) and Example 2 (Figure 2(b)), the amount of nanoscale κ-carbides in this example seems to be slightly reduced, presumably due to the lower carbon content of the welding rod used in this example. In addition, it is worth noting that no traces of microcracks or pores were observed in the fusion zone and near the interface between the fusion zone and the heat-affected zone of the weld. At the same time, the density of the nanoscale κ-carbides originally present in the heat-affected zone of the weld and the base metal zone is obviously not affected by the welding process.

[0199] Figure 3(e) shows a scanning electron microscope image of the Vickers microhardness test indentations taken across the fusion zone, the heat-affected zone of the weld, and the base metal zone. The corresponding microhardness values measured at each measurement point are shown in Figure 3(f), which shows that the average microhardness of the fusion zone, the heat-affected zone of the weld, and the base metal zone is approximately 402 Vickers hardness, 417 Vickers hardness, and 420 Vickers hardness, respectively. Obviously, there is a negligible softening phenomenon in the fusion zone. It is worth noting that in this example, compared with the microhardness (350 - 400 Vickers hardness) of the precipitation-hardened austenitic Fe-Mn-Al-C alloy after ideal aging treatment and with a carbon content not exceeding 1.2% by weight percentage (as described in the prior art), the microhardness of the fusion zone in this example is still quite high.

[0200] Figure 3(g) is a bright-field image of a transmission electron microscope taken from the fusion zone. Compared with that shown in Example 1 (Figure 1(h)) and Example 2 (Figure 2(e)), the density of nanoscale κ-carbide in this figure slightly decreases. The selected area electron diffraction pattern shown in Figure 3(h) confirms that the microstructure of the fusion zone is composed of austenite and κ-carbide. Figure 3(i) is the analysis result of a transmission electron microscope and an X-ray energy dispersive analyzer taken from the fusion zone, showing that iron, manganese, and aluminum in the fusion zone are approximately 61.9%, 27.6%, and 10.5% respectively. The tensile test of this as-welded state sample shows that there is a large amount of plastic deformation in the fusion zone in the sample fractured by the tensile test, and necking and fracture mainly occur in the fusion zone. On the other hand, the base metal zone has a relatively small amount of deformation. This may be due to the fact that the microhardness value of the fusion zone (about 402 Vickers hardness) is quite low compared with that of the base metal zone (about 420 Vickers hardness). As a result, although there is obvious ductile deformation in the fusion zone, the total elongation rate of this example slightly decreases compared with the above-mentioned examples. The yield strength, ultimate tensile strength, and elongation rate of the welded part of this as-welded state sample are approximately 916 MPa, 1035 MPa, and 33% respectively.

[0201] From the above results, it is obvious that the microstructure of the welded part of this example and the resulting mechanical properties are very similar to those seen in Example 1 and Example 2.

[0202] Example 4

[0203] In this example, the base metal used is the same as that used in Example 3. The base metal plate used in this example has a size of 80 mm × 80 mm × 8 mm and is machined to form a single V-groove butt joint. The standard composition of the welding electrode is iron-29.5 manganese-9.8 aluminum-1.48 carbon (Fe-29.5Mn-9.8Al-1.48C), and the diameter is about 2.4 mm. Pure argon is used as the shielding gas, and tungsten inert gas arc welding is used for fusion welding. Its welding parameters are similar to those used in Example 1.

[0204] Figure 4(a) shows a typical scanning electron microscope image of the post-weld sample including the fusion zone, the weld heat-affected zone, and the base metal zone. Obviously, the microstructure of the fusion zone is very similar to that seen in the previous Examples 1 to 3, that is, typical columnar austenite grains, which are composed of austenite dendritic cells and eutectic regions. This feature that nano-sized κ-carbides are densely distributed in the austenite dendritic cells and eutectic regions can still be observed. However, due to the further reduction of the carbon content of the welding electrode used in this example, the amount of nano-sized κ-carbides scattered in the austenite dendritic cells and eutectic regions seems to have slightly decreased. In addition, it is worth noting that no traces of microcracks and pores were observed in the fusion zone and near the interface between the fusion zone and the weld heat-affected zone. At the same time, the density of the nano-sized κ-carbides originally present in the weld heat-affected zone and the base metal zone was obviously not affected during the welding process. Figure 4(b) shows a scanning electron microscope image of the Vickers microhardness test indentations made across the fusion zone, the weld heat-affected zone, and the base metal zone. 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 fusion zone, the weld heat-affected zone, and the base metal zone is approximately 385 Vickers hardness, 414 Vickers hardness, and 419 Vickers hardness, respectively. The decrease in the microhardness of the fusion zone is presumably due to the reduction of the carbon content of the welding electrode used in this example. It is worth noting that, as seen in the previous examples, there is no obvious softening phenomenon caused by fusion welding.

[0205] Figure 4(d) is a bright-field image taken by a transmission electron microscope from the fusion zone, which shows the nano-scale κ-carbides present in the austenite dendritic cells and eutectic regions. However, it is clearly visible that the density of the nano-scale κ-carbides present in the austenite dendritic cells and eutectic regions has slightly decreased compared to the nano-scale κ-carbides observed in the aforementioned Examples 1 to 3. The selected area electron diffraction pattern taken from the fusion zone shown in Figure 4(e) confirms that the microstructure of the fusion zone is indeed composed of austenite and κ-carbides. Figure 4(f) is the analysis result taken from the transmission electron microscope and X-ray energy dispersive analyzer from the fusion zone, showing that the iron, manganese, and aluminum in the fusion zone are approximately 61.7%, 28.8%, and 9.5% respectively. The tensile test shows that the yield strength, ultimate tensile strength, and elongation of the welded parts in this example are approximately 875 MPa, 1012 MPa, and 29.8% respectively. After careful inspection, it is found that the samples fractured by the tensile test have typical characteristics of extended plastic deformation in the fusion zone. However, due to the significant microhardness difference between the fusion zone (about 385 Vickers hardness) and the base metal zone (about 419 Vickers hardness) (see Figure 4(c)), the deformation almost all occurs in the fusion zone, while there is a relatively small amount of deformation in the weld heat-affected zone and the base metal zone, thus causing a general decline in the overall elongation. Nevertheless, after the tensile test, the severe necking, non-wavy surface, and meandering fracture pattern observed in the fusion zone indicate that the weld bead has good ductility.

[0206] It should be noted that the microhardness of the fusion zone in this example (about 385 Vickers hardness) is similar to the microhardness (350 - 400 Vickers hardness) of the precipitation-hardened austenitic Fe-Mn-Al-C alloy with an ideal aging treatment and a carbon content not greater than 1.2% by weight in the prior art. From the above scanning electron microscope and transmission electron microscope results, it can be seen that the electrode with the designed composition used in this example still retains most of the microstructure characteristics in the fusion zone. In other words, this characteristic refers to the high-density nano-scale κ-carbides present in the austenite dendritic cells and eutectic regions. Therefore, it can be reasonably expected that the fusion zone should have an excellent combination value of yield strength and ductility, which is due to the use of the electrode provided by the present invention. In fact, the yield strength (about 875 MPa) and elongation (about 29.8%) measured for the welded parts in this example are quite similar to the values measured for the precipitation-hardened austenitic Fe-Mn-Al-C alloy with an ideal aging treatment and a carbon content not greater than 1.2% by weight (yield strength: about 680 - 990 MPa, elongation: 55 - 26%).

[0207] Example 5

[0208] In this embodiment, the base material used is the same as that used in Example 1. The size of the base material plate used in this embodiment is 80 mm × 80 mm × 8 mm. Two base material plates are machined to form a single V-groove butt joint. The standard composition of the welding rod used in this embodiment is Fe-33.8Mn-10.6Al-1.38C, with a diameter of about 2.4 mm. 75% argon plus 25% helium is used as the shielding gas, and fusion welding is carried out by gas tungsten arc welding. Its welding parameters are similar to those used in Example 1.

[0209] Figure 5(a) is a scanning electron microscope image, which shows the microstructure of the post-weld sample including the fusion zone, the welding heat-affected zone, and the base material zone. Similar to the previous Examples 1 to 4, the fusion zone is composed of typical columnar austenite grains, and the feature that a large number of nanoscale κ-carbides are evenly distributed in the austenite dendritic cells and eutectic regions can still be observed. However, the density of the nanoscale κ-carbides has been greatly reduced. On the other hand, the density of the nanoscale κ-carbides originally present in the welding heat-affected zone and the base material zone is apparently not affected during the welding process. In addition, it is worth noting that no traces of microcracks and pores are observed in the fusion zone and near the interface between the fusion zone and the welding heat-affected zone. Figure 5(b) shows a scanning electron microscope image of the Vickers microhardness test indentations made across the fusion zone, the welding heat-affected zone, and the base material zone. The corresponding microhardness values measured at each measurement point are shown in Figure 5(c). Figure 5(c) shows that the average microhardness of the fusion zone, the welding heat-affected zone, and the base material zone is approximately 369 Vickers hardness, 385 Vickers hardness, and 412 Vickers hardness, respectively. The decrease in microhardness in the fusion zone is presumably due to the lower carbon content of the welding rod used in this embodiment, and this lower carbon content obviously causes a further decrease in the density of the nanoscale κ-carbides distributed in the austenite dendritic cells and eutectic regions, making it even lower.

[0210] Figure 5(d) is a bright-field image of a transmission electron microscope taken from the fusion zone, showing the nanoscale κ-carbides present in the austenite dendritic cells and eutectic regions. At the same time, it can be found that, compared with the nanoscale κ-carbides observed in the previous Examples 1-4, the density of the nanoscale κ-carbides present in the austenite dendritic cells and eutectic regions in this embodiment has been greatly reduced. The selected area electron diffraction pattern shown in Figure 5(e) taken from the fusion zone reconfirms that the microstructure of the fusion zone is indeed composed of austenite and κ-carbides. Figure 5(f) is the analysis result of a transmission electron microscope and an X-ray energy dispersive analyzer taken from the fusion zone, showing that the iron, manganese, and aluminum in the fusion zone are approximately 56.6%, 33.1%, and 10.3%, respectively.

[0211] The tensile test shows that the yield strength, ultimate tensile strength and elongation of the welded parts in this embodiment are approximately 725 MPa, 1005 MPa and 23.8% respectively. After careful inspection, it is found that the samples fractured by the tensile test have obvious plastic deformation in the fusion zone. However, due to the significant difference in microhardness between the fusion zone (about 369 Vickers hardness) and the base metal zone (about 412 Vickers hardness) (see Fig. 5(c)), in the samples fractured by the tensile test, the deformation almost all occurs in the fusion zone, while only slight deformation occurs in the weld heat-affected zone and the base metal zone, resulting in a significant drop in the overall elongation.

[0212] From the above results, it is obvious that the welding rod using the design composition in this embodiment still retains most of the microstructural characteristics in the fusion zone disclosed in the present invention. In addition, the microhardness obtained in the fusion zone (about 369 Vickers hardness) is comparable to that of the precipitation-hardened austenitic Fe-Mn-Al-C alloys that have undergone ideal aging treatment and have a carbon content not greater than 1.2% by weight (about 350 - 400 MPa). As described in Example 4, due to the high similarity of the microstructure and the resulting microhardness, it can be expected that the fusion zone of this embodiment also has a satisfactory combination value of yield strength and elongation. It should be emphasized here that the characteristics of the microstructure of the base metal used in Examples 1 to 5 of the present invention. That is, in the hot-rolled state, the microstructure of the base metal is composed of fully austenite and high-density nano-sized κ-carbides uniformly distributed in the austenite matrix. The density of the nano-sized κ-carbides is even much higher than that in the precipitation-hardened austenitic Fe-Mn-Al-C alloys that have undergone ideal aging treatment and have a carbon content not greater than 1.2% by weight. Therefore, the microhardness of the base metal used in Examples 1 to 5 (about 412 - 420 Vickers hardness) is much higher than that of the precipitation-hardened austenitic Fe-Mn-Al-C alloys that have undergone ideal aging treatment and have a carbon content not greater than 1.2% by weight (about 350 - 400 Vickers hardness).

[0213] Example 6

[0214] The objective of this embodiment is to attempt to clarify what kind of impact would be caused to the characteristics of the fusion zone in the post-weld state sample if the carbon content in the welding electrode is further reduced. The base material used in this embodiment is the same as that used in Embodiment 1. The size of the base material plate is 80 mm × 80 mm × 8 mm, and a single V-groove butt joint is formed after machining. Gas tungsten arc welding is carried out, and its welding parameters are similar to those used in Embodiment 1. The standard composition of the welding electrode used in this embodiment is Fe-31.8Mn-9.2Al-1.25C, with a diameter of about 2.4 mm.

[0215] Figure 6(a) is a typical scanning electron microscope image of a region in the post-weld sample that includes the fusion zone, the welding heat-affected zone, and the base material zone. The microstructure of the fusion zone clearly shows several characteristics that are completely different from those described in the above-mentioned Embodiments 1 to 5. First of all, the base is still fully austenite containing austenite dendritic cells and eutectic regions. However, there are no κ-carbides scattered in the austenite dendritic cells, and the amount of κ-carbides in the eutectic region is significantly reduced. To understand this characteristic more clearly, please refer to the enlarged image of the marked area (indicated by the arrow) in the upper right corner of the inserted Figure 6(a). This characteristic is completely different from the microstructures described in the above-mentioned Embodiments 1-5. In fact, this characteristic is very similar to what was observed by Chou and Lee in the autogenous welding of Fe-Mn-Al-C alloys with a carbon content not exceeding 1.29 wt%, that is, as described in the prior art, there are no κ-carbides scattered in the austenite dendritic cells.

[0216] Figure 6(b) shows a scanning electron microscope image of the Vickers microhardness test indentations made across the fusion zone, the welding heat-affected zone, and the base material zone. The corresponding microhardness values measured at each measurement point are shown in Figure 6(c). It can be clearly seen from these two figures that the microhardness of the welding heat-affected zone and the base material zone is about 412 Vickers hardness. However, the microhardness of the fusion zone quickly drops to only 242 Vickers hardness. This Vickers hardness value is very close to the Vickers hardness value of 210-230 Vickers hardness of the precipitation-hardened austenitic Fe-Mn-Al-C alloy with a carbon content not exceeding 1.2% in the solution heat treatment and quenched state as described in the prior art. It is generally believed that the absence of κ-carbides in the austenite dendritic cells and the significant reduction of κ-carbides in the eutectic region are the main reasons for the significant drop in the microhardness of the fusion zone.

[0217] Figure 6(d) shows the macroscopic images of the sample before and after the tensile test. The results show that the yield strength, ultimate tensile strength and elongation of the welded joint are approximately 462 MPa, 938 MPa and 14% respectively. Obviously, for the welded joint obtained using this embodiment, its yield strength and elongation are hardly satisfactory. The main reason for the significant decrease in elongation may be that the microhardness of the fusion zone (about 242 Vickers hardness) is much smaller than that of the heat-affected zone and the base metal zone of the weld (about 412 Vickers hardness). Thus, during the tensile test, the deformation of the entire welded joint is basically limited to the fusion zone, while the heat-affected zone and the base metal zone of the weld hardly deform and remain unchanged, resulting in a significant decline in elongation. At the same time, it can be observed in Figure 6(d) that the fracturing crack seems to start from the interface between the fusion zone and the heat-affected zone of the weld (as indicated by the arrow), and then penetrates the fusion zone horizontally, forming a zigzag fracture line. The elongation within the fusion zone is estimated to be about 27%. In fact, its microhardness is only about 242 Vickers hardness, and the deformation, the onset of necking, and subsequent cracks all occur within the fusion zone, indicating that the strength of the fusion zone is much lower than that of the heat-affected zone and the base metal zone of the weld. The main reason may be that there is no κ-carbide dispersed in the austenite dendritic cells, and the amount of κ-carbide in the eutectic zone is significantly reduced. The welding electrode with a carbon content of 1.25 wt.% seems to be insufficient to promote the formation of nano-scale κ-carbide in the austenite dendritic cells and the eutectic zone during solidification, and the formation of the above nano-scale κ-carbide is also the most unique characteristic disclosed in the present invention.

[0218] Example 7

[0219] The purpose of this example is to investigate the impact of a higher carbon content in the welding electrode on the characteristics of the fusion zone in the post-weld sample. The base metal used in this example is the same as that used in Example 3. The size of the base metal plate is 80 mm × 80 mm × 10 mm, and a single V-groove butt joint is formed after machining. Gas tungsten arc welding is used for fusion welding, and its welding parameters are similar to those used in Example 1. The standard composition of the welding electrode used in this example is iron-30.5 manganese-8.9 aluminum-2.2 carbon (Fe-30.5Mn-8.9Al-2.2C), with a diameter of about 2.4 mm.

[0220] Figure 7(a) is a macroscopic image of the weld bead, clearly showing a large macrocrack penetrating the weld bead. Obviously, the relatively high carbon content in the welding rod results in severe solidification hot cracking. Figure 7(b) is a scanning electron microscope image of a region within the post-weld state sample that includes the fusion zone, the weld heat-affected zone, and the base metal zone. This figure clearly shows that the fusion zone is mainly composed of typical columnar austenite grains, and at the same time, has a high density of nanoscale κ-carbides dispersed in the austenite dendritic cells and eutectic regions. This characteristic is similar to that observed in Examples 1 to 5. However, some micron-sized coarse carbides can be observed at the grain boundaries (as indicated by the arrow in Figure 7(b)). Please refer to the upper right corner of the inserted Figure 7(b), an enlarged image of the marked area (indicated by the arrow), which clearly shows a precipitate-free zone near the coarse carbides formed at the grain boundaries. It is well known that both grain boundary precipitates and the associated precipitate-free zones have a negative impact on mechanical properties, especially elongation. In addition, Figure 7(b) also shows that during the fusion welding process, the density of the nanoscale κ-carbides originally present in the weld heat-affected zone and the base metal zone is apparently not affected.

[0221] Figure 7(c) is a bright-field image of a transmission electron microscope taken from the fusion zone, clearly showing that in addition to the high density of nanoscale κ-carbides, there are also some micron-sized coarse precipitates (labeled as K) present at the austenite grain boundaries. In addition, there is an obvious precipitate-free zone near the coarse precipitates. Figure 7(d) is a selected area electron diffraction pattern taken from the coarse grain boundary precipitate (labeled as K), showing that these coarse grain boundary precipitates are κ-carbides with an L′1 2 structure, which is the same as the structure of the nanoscale κ-carbides dispersed in the austenite dendritic cells and eutectic regions. Figure 7(e) is a typical analysis result of a transmission electron microscope and an X-ray energy dispersive analyzer for a coarse grain boundary κ-carbide, showing that the iron, manganese, and aluminum contents of the composition of the coarse carbides are approximately 56.7%, 34.2%, and 9.1% respectively.

[0222] Based on the above various observations, it can be clearly seen that although the relatively high carbon content (i.e., 2.2% by weight) in the welding rod results in a high density of nanoscale κ-carbides in the austenite dendritic cells and eutectic regions, it also causes some coarse κ-carbides to appear at the grain boundaries. Therefore, it can be expected that these coarse grain boundary κ-carbides and their associated precipitate-free zones will not only have some negative impacts on the welded parts, such as ductility, but also lead to severe solidification hot cracking in the weld bead, as shown in Figure 7(a).

[0223] Example 8

[0224] In this embodiment, the effects of a relatively high aluminum content in the welding electrode on the characteristics of the fusion zone in the post-weld sample will be described. The base material used in this embodiment is the same as that used in Example 1. The two base material plates have dimensions of 80 mm × 80 mm × 8 mm and are machined to form a single V-groove butt joint. The welding parameters used during gas tungsten arc welding are similar to those used in Example 1. The standard composition of the welding electrode is iron-30.5 manganese-12.5 aluminum-1.55 carbon (Fe-30.5Mn-12.5Al-1.55C), with a diameter of approximately 2.4 mm.

[0225] A preliminary inspection of the welded component revealed solidification hotcracking in some of the weld beads, as indicated by the arrow in Figure 8(a). The scanning electron microscope image shown in Figure 8(b) depicts the microstructure of the fusion zone, which mainly consists of typical columnar austenite grains composed of austenite dendritic cells and eutectic regions. Similar to what was observed in the previous Examples 1 to 5, high-density nanoscale κ-carbides are present in the austenite dendritic cells and eutectic regions. However, in addition to these characteristics described above, some discrete particles can be observed at the austenite grain boundaries, as indicated by the arrow and marked with the symbol α in Figure 8(b). Figure 8(c) shows the transmission electron microscope and X-ray energy dispersive spectrometer analysis (TEM-EDS) taken from the discrete particles, indicating that the iron, manganese, and aluminum contents of their composition are approximately 52.3%, 34.2%, and 13.5% respectively. Apparently, the aluminum content concentration in the discrete particles is as high as 13.5%, which means that aluminum-rich particles will form at the austenite grain boundaries during the welding process. It has been determined that the presence of the aluminum-rich secondary phase at the austenite grain boundaries will significantly deteriorate the ductility of the iron-manganese-aluminum alloy. Therefore, an excessive aluminum content in the welding electrode will not only lead to the formation of solidification cracks during the welding process but also cause the formation of an aluminum-rich phase with negative effects at the austenite grain boundaries.

[0226] Example 9

[0227] The main objective of this embodiment is to examine the effects of a relatively high manganese content in the welding electrode on the characteristics of the fusion zone in the post-weld state sample. The base material used is the same as that used in Example 1. The base material plate has dimensions of 80 mm × 80 mm × 8 mm and is machined to form a single V-groove butt joint. The welding parameters used during gas tungsten arc welding are similar to those used in Example 1. The standard composition of the welding electrode used in this embodiment 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.

[0228] Figure 9(a) is a scanning electron microscope image of a region within the post-weld sample that includes a fusion zone, a weld heat-affected zone, and a base metal zone. The microstructure of the fusion zone is clearly very different from that observed in Examples 1 to 5 described above. Although a high density of nanoscale κ-carbides can still be observed within the austenite dendrite cells and eutectic regions, there are also a large number of discrete ferrite phases (indicated by the arrow and labeled with the symbol α in Figure 9(a)) scattered throughout the fusion zone. It is estimated that the volume fraction of the ferrite phase is approximately 18%, which is a value calculated using the point-counting technique.

[0229] Figure 9(b) is a bright-field image taken by a transmission electron microscope from the fusion zone, showing nanoscale κ-carbides scattered within the austenite dendrite cells and eutectic regions. In addition, it is clearly visible that there are also some micron-sized coarse κ-carbides formed at the austenite / ferrite grain boundaries. Figure 9(c) is a selected area electron diffraction pattern taken from the circled region in Figure 9(b), clearly showing the presence of three phases: the austenite phase, the ferrite phase, and the κ-carbide. The crystal orientation relationships between the austenite and ferrite phases are respectively and Figures 9(d) and 9(e) respectively show the typical analysis results of a transmission electron microscope and an X-ray energy dispersive analyzer taken from the austenite dendrite cell region and the ferrite region. The quantitative analysis results within the austenite dendrite cells and ferrite show that their chemical compositions are Fe-20.6%Mn-8.2%Al and Fe-15.8%Mn-8.8%Al respectively.

[0230] Based on the above-described results, an interesting and notable point in this embodiment is that for a Fe-Mn-Al-C alloy with a composition of Fe-(17.5 to 35)Mn-(7.1 to 12)Al-(0.8 to 1.2)C, after hot rolling, solution heat treatment, and quenching, the microstructure of the alloy always exhibits single-phase austenite, as described in the prior art. Therefore, this embodiment further demonstrates that even for materials with the same chemical composition, after fusion welding, their microstructures may be very different from those obtained by general heat-treatment processes. Obviously, in order to make the microstructure of the fusion zone ferrite-free full austenite with high-density nanoscale κ-carbides dispersed in the austenite dendritic cells and eutectic regions, the manganese content in the welding rod used in this embodiment should basically be higher than about 20% (by weight).

[0231] Example 10

[0232] In this embodiment, the effect of adding titanium to the welding rod on the characteristics of the fusion zone in the post-weld sample was investigated. The base material used in this embodiment is a hot-rolled base material with a standard composition of Fe-29.3Mn-9.3Al-1.75C. Figure 10(a) is a bright-field image of a transmission electron microscope of the hot-rolled base material, clearly showing that during the hot-rolling process, high-density nanoscale precipitates (about 3 - 5 nm) are uniformly dispersed in the matrix and formed at the dislocations (indicated by the arrow). Figure 10(b) is a selected-area electron diffraction pattern, clearly confirming that the matrix of the base material is a completely austenite phase, and the high-density nanoscale precipitates are regular κ-carbides with an L′12 structure. Figure 10(c) is a (100)κ dark-field image taken from the same region as shown in Figure 10(a), showing the high-density nanoscale κ-carbides present in the austenite matrix and located at the dislocations. Microstructural analysis by transmission electron microscope confirmed that the microstructure of the hot-rolled base material used in this embodiment is single-phase austenite with uniformly dispersed high-density nanoscale κ-carbides. Tensile tests showed that the yield strength, ultimate tensile strength, and elongation of the hot-rolled base material were approximately 1020 MPa, 1198 MPa, and 45%, respectively.

[0233] The dimensions of the two hot-rolled base metal plates are 80 mm × 80 mm × 12 mm, and a single V-groove butt joint is formed after machining. The welding parameters used during tungsten inert gas arc welding are similar to those used in Example 2. The standard composition of the welding electrode used in this example is iron-28.5 manganese-9.3 aluminum-1.82 carbon-1.6 titanium (Fe-28.5Mn-9.3Al-1.82C-1.6Ti), with a diameter of approximately 1.2 mm. Figure 10(d) shows the macroscopic image of the post-weld sample, showing that the entire weld bead appears to have a very smooth morphology, without any visible large cracks or holes. Figure 10(e) is a scanning electron microscope image of a region of the post-weld sample including the fusion zone, the heat-affected zone of welding, and the base metal zone. Several significant characteristics caused by the addition of titanium can be seen from Figure 10(e). First, the size of the austenite grains in the fusion zone is significantly smaller than those seen in the previous Examples 1 to 5, and even smaller than those in the heat-affected zone of welding and the base metal zone. Second, the morphology of the austenite grains becomes more spherical and less directional. However, it is obvious that there are still high-density nanoscale κ-carbides evenly distributed in the austenite dendritic cells and the eutectic region. In addition, although it can be seen that the austenite grains in the heat-affected zone of welding have slightly grown, the high-density nanoscale κ-carbides originally present in the base metal zone remain unchanged. Finally, it is worth noting that no traces of microcracks or holes are observed near the interface between the fusion zone and the heat-affected zone of welding, indicating that the use of the welding electrode provided by the present invention can significantly eliminate the solidification cracks and liquefaction cracks commonly found in fusion-welded precipitation-hardened alloys. Figure 10(f) shows a scanning electron microscope image of the Vickers microhardness test indentations made across the fusion zone, the heat-affected zone of welding, and the base metal zone. The corresponding microhardness values measured at each measurement point are shown in Figure 10(g), which shows that the average microhardness values of the fusion zone, the heat-affected zone of welding, and the base metal zone of the post-weld sample of the present invention are approximately 451 Vickers hardness, 435 Vickers hardness, and 436 Vickers hardness, respectively. Obviously, no softening phenomenon, which is most common in most fusion-welded precipitation-hardened alloys, appears in the fusion zone. This is mainly because in the post-fusion-welded state (Figure 10(e)), high-density nanoscale κ-carbides are present in the austenite dendritic cells and the eutectic region. In addition, in addition to the nanoscale κ-carbides, there are also a large number of high-hardness (2200 - 3500 Vickers hardness) nanoscale titanium-rich titanium carbides formed in the eutectic region (Figure 10(h)). As a result, not only does no softening phenomenon occur in the fusion zone, but it also has a higher microhardness than the base metal zone.

[0234] The thin foil in the molten zone was examined using a transmission electron microscope. The results showed that with the addition of titanium, a large number of nano-sized titanium carbides with a regular face-centered cubic structure were formed in the eutectic zone. Figure 10(h) is a typical example. In the bright field image of the transmission electron microscope, it can be clearly seen that the length of the austenite dendrite cells is only about 100-150 nanometers, and in the austenite dendrite cells and the eutectic zone, there are high-density nano-sized κ-carbides (about 3-5 nanometers) evenly distributed. In addition, there are also some slightly larger nano-sized titanium carbides (about 6-10 nanometers) in the eutectic zone, as shown by the arrows in Figure 10(h). Figure 10(i) is a selected area electron diffraction pattern taken from the circled area in Figure 10(h), showing the coexistence of austenite substrate, κ-carbides and titanium carbides. The results of the selected area electron diffraction pattern analysis indicate that the crystal orientation relationship between the three phases is cubic-to-cubic. FIG10(j) is the results of TEM and EDXA analysis taken from the same area, which proves that titanium carbides rich in titanium do appear in the eutectic region. Since the standard composition of the welding rods used in Example 1 and Example 10 has approximately the same aluminum and carbon content, except that about 1.6% by weight of titanium is added in Example 10. Therefore, the molten zone microstructures shown in FIG1(h) and FIG10(h) should be compared with each other. In FIG1(h), although the size of the austenite dendrite cell is slightly larger than that of the austenite dendrite cell in FIG10(h), the density of nano-scale κ-carbides in the austenite dendrite cell in FIG1(h) seems to be slightly higher than that in FIG10(h) (the sample with titanium addition). In addition, both have a large number of high-density nanoscale precipitates in the eutectic region (κ-carbide in Figure 1(h) and titanium-rich titanium carbide in Figure 10(h)). However, the microhardness of the molten zone in Example 10 (about 451 Vickers hardness) is much greater than that of Example 1 (about 425 Vickers hardness). The main reason for this phenomenon may be that the hardness of titanium-rich titanium carbide is about 2200-3500 Vickers hardness, which is lower than that of (Fe, Mn) 3 The hardness of AlC carbide (about 630-670 Vickers hardness) is too high.

[0235] The tensile test shows that the yield strength, ultimate tensile strength, and elongation of the welded part are 1006 MPa, 1145 MPa, and 40%, respectively. Figure 10(k) shows the macroscopic images of the sample before and after the tensile test. The corrugated deformation surface in the fusion zone clearly confirms the ductility characteristics of the weld bead. In addition, it can be seen that the edges of the tensile-fractured sample still remain parallel, spanning the fusion zone and the base metal zone, and are basically of the same width, as shown by FZ, BM, and INT in Figure 10(k), where INT represents the width of the interface between the fusion zone and the base metal zone. As described in the details discussed in Example 1 and Example 2, such special deformation characteristics strongly imply that the yield strengths of the fusion zone and the base metal zone are approximately at the same level. As mentioned above, the amazing yield strength and microhardness obtained in the fusion zone can be attributed to the presence of high-density nanoscale κ-carbides and nanoscale titanium-rich titanium carbides in the fusion zone. At the same time, it is worth noting the slight necking phenomenon that occurs in the base metal zone, which ultimately causes fracture, further showing that the tensile strength of the fusion zone is higher than that of the base metal zone, and this phenomenon is consistent with the phenomenon that the fusion zone shown in Figure 10(g) has a larger microhardness.

[0236] Based on the above analysis, it is obvious that adding titanium to the welding rod will cause the following advantageous effects. First, the austenite grains in the fusion zone will change from columnar to globular, and at the same time, the grain size is significantly refined. Second, in the as-welded state, high-density nanoscale κ-carbides exist in the greatly refined austenite dendritic cells and eutectic regions. In addition, a large number of high-hardness nanoscale titanium-rich titanium carbides are formed in the eutectic regions. Third, the austenite dendritic cells, nanoscale κ-carbides, and nanoscale titanium-rich titanium carbides that make up the fusion zone all have a ductile face-centered cubic structure. In addition, the sizes of the κ-carbides and titanium-rich titanium carbides in the fusion zone are only 3-10 nanometers, so they can very effectively improve the mechanical properties, especially the yield strength, without causing a significant decline in ductility. Therefore, after using the welding rod provided in the present invention, the fusion zone of the as-welded sample can simultaneously have a combined value of excellent hardness, yield strength, ultimate tensile strength, and ductility.

[0237] Example 11

[0238] In this embodiment, an investigation was made of what effect the addition of niobium to the welding electrode would have on the characteristics of the fusion zone in the post-weld sample. The base materials used were the same as those used in Example 10. The dimensions of the two base material plates were 80 mm × 80 mm × 8 mm, and after machining, a single V-groove butt joint was formed. The welding parameters used during gas tungsten arc welding were similar to those used in Example 1. The standard composition of the welding electrode used was Fe-29.5Mn-8.2Al-1.56C-1.0Nb, with a diameter of approximately 3.2 mm. Figure 11(a) shows the macroscopic image of the welded state sample, showing that the entire weld bead has a very smooth morphology without any visually visible large cracks or holes. Figure 11(b) is a scanning electron microscope image of a region within the welded state sample that includes the fusion zone, the heat-affected zone of welding, and the base material zone. Different from what was seen in the previously titanium-added sample (Example 10), the morphology of the austenite grains in the fusion zone of this embodiment is a mixture of spherical ("globular") and columnar-like. However, similar to the phenomenon observed in Example 10, a high density of nanoscale κ-carbides can be observed to be uniformly dispersed in the austenite dendritic cells and eutectic regions. Similarly, the austenite grains in the heat-affected zone of welding have slightly grown, and the high density of nanoscale κ-carbides originally present in the hot-rolled base material zone basically remains unchanged during the fusion welding. In addition, it is worth noting that no traces of microcracks or holes were observed in the fusion zone and near the interface between the fusion zone and the heat-affected zone of welding, indicating that the use of the welding electrode provided by the present invention can significantly eliminate solidification cracks and liquefaction cracks. Figure 11(c) shows a scanning electron microscope image of the Vickers microhardness test indentations made across the fusion zone, the heat-affected zone of welding, and the base material zone. The corresponding microhardness values measured at each measurement point are shown in Figure 11(d), which shows that the microhardness of the fusion zone, the heat-affected zone of welding, and the base material zone of the welded state sample of the present invention is approximately 431 Vickers hardness, 435 Vickers hardness, and 437 Vickers hardness, respectively. Obviously, no softening phenomenon, which is most common in most fusion-welded precipitation-hardening alloys, occurred in the fusion zone. Similar to the phenomenon observed in Example 10, the main reason is the presence of a high density of nanoscale κ-carbides in the austenite dendritic cells and eutectic regions, and the presence of a large number of high-hardness (2000 - 3200 Vickers hardness) nanoscale niobium-rich Nb-carbides in the eutectic region (Figure 11(e)). Therefore, the microhardness of the fusion zone is approximately at the same level as that of the heat-affected zone of welding and the base material zone.

[0239] Similar to the phenomenon observed in the titanium-added sample (Example 10), transmission electron microscopy analysis shows that a large number of nanoscale niobium-rich niobium carbides with a regular face-centered cubic structure are formed in the eutectic region. Figure 11(e) is a typical transmission electron microscopy bright-field image taken from the fusion zone of the welded part, which clearly shows that the morphology of the austenite dendritic cells and the eutectic region in this example is quite different from that observed in Example 10. The dendritic cells and the surrounding eutectic region seem to be more roundish. However, the characteristic of having high-density nanoscale (about 3-5 nm) κ-carbides scattered in the austenite dendritic cells and the eutectic region remains unchanged. In the eutectic region, in addition to the eutectic κ-carbides, there are also a large number of slightly larger nanoscale precipitates (about 6-10 nm), as indicated by the arrows in Figure 11(e). Figure 11(f) is a selected area electron diffraction pattern taken from the area circled in Figure 11(e), which shows the coexistence of the austenite substrate, κ-carbides, and niobium-rich niobium carbides. According to the analysis results of the selected area electron diffraction pattern, the crystal orientation relationship between these three phases is cubic to cubic. Figure 11(g) is the analysis result of transmission electron microscopy and X-ray energy dispersive analyzer taken from the same area, which confirms the existence of niobium-rich niobium carbides. Since the hardness of the nanoscale niobium-rich niobium carbides is about 2000-3200 Vickers hardness, it can be expected that the precipitation of nanoscale niobium-rich niobium carbides will have a similar effect on the characteristics of the fusion zone, which also explains the microhardness results shown in Figure 11(d).

[0240] Tensile tests show that the yield strength, ultimate tensile strength, and elongation of the welded part are 980 MPa, 1108 MPa, and 46%, respectively. Figure 11(h) shows the macroscopic images of the sample before and after the tensile test. In the tensile-fractured sample, it is clearly observable that a wavy deformed surface and a tortuous meandering fracture pattern can be seen in the fusion zone, indirectly indicating the ductile characteristics of the fusion zone. In addition, a region including the base metal zone, the fusion zone, and the interface between them (denoted as BM, FZ, and INT) can be seen simultaneously, and its width is basically the same, and the edges of the sample across this region remain parallel, as shown in Figure 11(h). This special property implies that the yield strength of the fusion zone and the base metal zone is approximately at the same level, and the deformation under tensile stress also continues uniformly until necking and final fracture occur in the fusion zone. In fact, the slight necking and final fracture that occur in the fusion zone show that the ultimate tensile strength in the fusion zone is less than that in the base metal zone.

[0241] Based on the above-described results, it is clearly visible that adding niobium to the welding rod has a similar effect to adding titanium (Example 10). That is to say, in addition to retaining the high-density nanoscale κ-carbides in the austenite dendritic crystal cells and eutectic regions, the nanoscale niobium-rich niobium carbides formed in the eutectic region obviously also contribute to strengthening the mechanical strength and microhardness of the molten zone in the welded sample, while still maintaining excellent ductility. As described in Example 10, since all the phases constituting the microstructure of the molten zone are face-centered cubic structures with ductility, and the strengthening nanoscale κ-carbides and niobium-rich niobium carbides are only about 3-10 nanometers in size, they are very effective in increasing the yield strength. Therefore, by using the welding rod provided by the present invention, the molten zone in the welded sample simultaneously has a combined value of excellent hardness, yield strength, ultimate tensile strength, and ductility.

[0242] Example 12

[0243] In this embodiment, an investigation was made of what effect the addition of vanadium to the welding electrode would have on the characteristics of the molten zone in the welding state sample. The base material used in this embodiment was the same as that used in Example 10. The dimensions of the two base material plates were 80 mm × 80 mm × 10 mm, and a single V-groove butt joint was formed after machining. The welding parameters used during tungsten inert gas arc welding were similar to those used in Example 1. The standard composition of the welding electrode used was 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 the macroscopic image of the welding state sample, showing that the entire weld bead has a very smooth shape without any visually visible large cracks or holes. Figure 12(b) is a scanning electron microscope image of a region in the post-weld sample that includes the molten zone, the welding heat affected zone, and the base material zone. Similar to the previously niobium-added sample (Example 11), the image shows that the morphology of the austenite grains in the molten zone is a mixture of spherical and columnar-like. In addition, as observed in Examples 10 and 11, a high density of nanoscale κ-carbides is uniformly distributed in the austenite dendritic cells and the eutectic zone. Similarly, the austenite grains in the welding heat affected zone have slightly grown, and the high density of nanoscale κ-carbides originally present in the hot-rolled base material zone remains essentially unchanged during the fusion welding. Finally, it is worth noting that no traces of microcracks or holes were observed in the molten zone and near the interface between the molten zone and the welding heat affected zone. Figure 12(c) shows a scanning electron microscope image of the Vickers microhardness test indentations made across the molten zone, the welding heat affected zone, and the base material zone. The corresponding microhardness values measured at each measurement point are shown in Figure 12(d), which shows that the microhardness of the molten zone, the welding heat affected zone, and the base material zone of the welding state sample of the present invention is approximately 438 Vickers hardness, 434 Vickers hardness, and 436 Vickers hardness, respectively. Obviously, no softening phenomenon occurs in the molten zone and the welding heat affected zone, which is consistent with the unique microstructure shown in Figure 12(b), where the main strengthening component, the high density of nanoscale κ-carbides, remains essentially unchanged in all three zones. In addition, the addition of vanadium causes a large number of high-hardness (2200 - 3000 Vickers hardness) nanoscale vanadium-rich V-carbides to form in the eutectic zone (Figure 12(e)), resulting in the microhardness of the molten zone being approximately at the same level as that of the welding heat affected zone and the base material zone, which is similar to the phenomenon observed in Example 11.

[0244] Similar to the phenomena observed in the titanium-added sample (Example 10) and niobium-added sample (Example 11), transmission electron microscopy analysis shows that adding vanadium has a similar effect. That is, a large number of nano-scale vanadium-rich vanadium carbides with a regular face-centered cubic structure are formed in the eutectic region. Figure 12(e) is a transmission electron microscopy bright-field image taken from the molten region, which clearly shows that the length of the austenite dendritic cells is only about 20 - 60 nanometers, and its morphology is different from that observed in the titanium-added sample (Figure 10(h)) and niobium-added sample (Figure 11(e)). However, the characteristic of high-density nano-scale κ-carbides (about 3 - 5 nanometers) scattered in the austenite dendritic cells and the eutectic region can still be observed. In addition, in the eutectic region, there are obviously a large number of precipitates with slightly larger sizes (about 6 - 10 nanometers), as indicated by the arrow in Figure 12(e). Figure 12(f) is taken from the area circled in Figure 12(e), and its selected-area electron diffraction pattern shows the coexistence of austenite substrate, κ-carbide, and vanadium-rich vanadium carbide. According to the analysis results of the selected-area electron diffraction pattern, the crystal orientation relationship between these three phases is cubic to cubic. Figure 12(g) is the analysis result of transmission electron microscopy and X-ray energy dispersive analyzer taken from the same area, which confirms the existence of vanadium-rich vanadium carbide. Since the hardness of vanadium-rich vanadium carbide is similar to that of niobium-rich niobium carbide and titanium-rich titanium carbide, it can be expected that vanadium-rich vanadium carbide will have a similar impact on the molten region characteristics, which also explains the microhardness results shown in Figure 12(d).

[0245] The tensile test shows that the yield strength, ultimate tensile strength, and elongation of the welded joint are 998 MPa, 1167 MPa, and 44% respectively. Figure 12(h) shows the macroscopic images of the sample before and after the tensile test. Interestingly, note that the edges of the tensile fractured sample remain parallel, and the width across the region including the fusion zone and the base metal zone is basically the same. In Figure 12(h), each region is represented by BM, FZ, and INT respectively (where IN represents the interface width between the base metal zone and the fusion zone). As discussed in Examples 1 and 2, and Examples 10 and 11, this special property implies that the yield strengths of the fusion zone and the base metal zone are approximately at the same level. A more careful examination of the tensile fractured sample shows that the wavy deformed surface (an indicator of ductility) in the weld bead fusion zone is particularly obvious. It is estimated that the elongation rate in the fusion zone can be as high as about 28%. In addition, when the applied stress exceeds the yield strength, the plastic deformation of the entire sample, including the fusion zone, the weld heat affected zone, and the base metal zone, seems to have a very uniform deformation until near the interface between the fusion zone and the weld heat affected zone, as indicated by the arrow in Figure 12(h), where slight necking or even rupture occurs. As a result, even after rupture, the overall edge remains nearly parallel. In fact, the slight necking and final rupture occurring near the interface between the fusion zone and the base metal zone indicate that the ultimate tensile strengths in each of the fusion zone, the weld heat affected zone, and the base metal zone are very close to each other.

[0246] Based on the above results, it is obvious that adding vanadium to the welding rod produces similar effects to adding titanium (Example 10) and adding niobium (Example 11). Anyway, in addition to retaining the high-density nano-sized κ-carbides present in the austenite dendritic cells and the eutectic region, the nano-sized vanadium-rich vanadium carbides formed in the eutectic region obviously also contribute to strengthening the mechanical strength and microhardness of the fusion zone in the welded sample.

[0247] Furthermore, since the nano-sized κ-carbides (3 - 10 nm), titanium-rich titanium carbides, niobium-rich niobium carbides, and vanadium-rich vanadium carbides all have the same ductile face-centered cubic structure as the austenite dendrites, an amazing strengthening effect can be achieved without causing a significant reduction in ductility. The above various detailed analyses clearly indicate that by using the welding rod provided by the present invention, the fusion zone obtained will have an excellent combination value of microhardness, yield strength, ultimate tensile strength, and ductility.

[0248] Example 13

[0249] In this embodiment, further investigation was carried out on what kind of influence the addition of titanium and niobium in the welding electrode would have on the characteristics of the molten zone in the post-weld state sample. The base materials used in this embodiment were the same as those used in Example 10. The dimensions of the two base material plates were 80 mm × 80 mm × 12 mm, and a single V-groove butt joint was formed after machining. The welding parameters used during tungsten inert gas arc welding were similar to those used in Example 1. The standard composition of the welding electrode used was Fe-30.6Mn-9.2Al-1.88C-0.8Ti-1.0Nb, with a diameter of approximately 3.2 mm. Figure 13(a) shows the macroscopic image of this welded state sample, showing that the entire weld bead has a very smooth morphology without any visually visible large cracks and holes. Figure 13(b) is a scanning electron microscope image of a region within this welded sample that includes the molten zone, the heat-affected zone of welding, and the base material zone. In this figure, it is obvious that the morphology of all austenite grains in the molten zone has become spherical, and the size of these spherical austenite grains has been significantly refined, even much smaller than the austenite grains in the heat-affected zone of welding and the base material zone. As described in Examples 10 to 12, it can be observed that there are high-density nano-scale κ-carbides evenly distributed in the austenite dendritic cells and eutectic regions. Moreover, the high-density nano-scale κ-carbides originally present in the hot-rolled base material zone basically remained unchanged during the fusion welding process. In addition, it is worth noting that no traces of microcracks and holes were observed in the molten zone and near the interface between the molten zone and the heat-affected zone of welding. Figure 13(c) shows a scanning electron microscope image of the Vickers microhardness test indentations carried out across the molten zone, the heat-affected zone of welding, and the base material zone. The corresponding microhardness values measured at each measurement point are shown in Figure 13(d), which shows that the average microhardness of the molten zone, the heat-affected zone of welding, and the base material zone of this welded state sample is approximately 467 Vickers hardness, 444 Vickers hardness, and 435 Vickers hardness, respectively. Not only did no softening phenomenon occur in the molten zone, but it also had a much higher microhardness than the heat-affected zone of welding and the base material zone. Similar to Examples 10 to 12, this is mainly because, in the post-weld state (Figure 13(b)), there are high-density nano-scale κ-carbides present in the austenite dendritic cells and eutectic regions. In addition, in the eutectic region, in addition to nano-scale κ-carbides, there are a large number of high-hardness nano-scale titanium-rich titanium carbides and niobium-rich niobium carbides (Figure 13(e)). Therefore, the microhardness in the molten zone is much greater than that in the heat-affected zone of welding and the base material zone, and no sign of softening appears. In addition, the addition of approximately 0.8% titanium and 1.0% niobium seems to result in a higher microhardness in the molten zone than that in the previous Examples 10 to 12, which may be due to the increase in the total amount of high-hardness nano-scale titanium-rich titanium carbides and niobium-rich niobium carbides in this embodiment.

[0250] Figure 13(e) is a bright-field transmission electron microscope image taken from the fusion zone, which clearly shows that the morphology of the austenite dendritic crystal cells and eutectic regions in this example is quite different from that observed in Examples 10 to 12. The dendritic crystal cells and their surrounding eutectic regions seem to be more refined and more intertwined with each other. However, the characteristic of having high-density nanoscale (about 3 - 5 nm) κ-carbides dispersed in the austenite dendritic crystal cells and eutectic regions remains unchanged. In addition, within the eutectic region, κ-carbides and titanium-rich titanium carbides and niobium-rich niobium carbides with slightly larger sizes (about 6 - 10 nm) seem to increase significantly, as indicated by the arrows in Figure 13(e). Figure 13(f) is a selected area electron diffraction pattern taken from the circled area in Figure 13(e), which confirms the presence of titanium-rich titanium carbides and niobium-rich niobium carbides.

[0251] The tensile test shows that the yield strength, ultimate tensile strength, and elongation of the welded part are 1015 MPa, 1168 MPa, and 38% respectively. Figure 13(g) shows the macroscopic images of the sample before and after the tensile test. The wavy deformed surface observed in the fusion zone indirectly shows the ductility of the fusion zone. In addition, it can be seen that the edges of the tensile-fractured sample still remain parallel, and the distances on both sides from the base metal zone, fusion zone to the interface range between the base metal zone and the fusion zone are basically the same, as shown by BM, FZ, and INT in Figure 13(g). This special property implies that the yield strengths of the fusion zone and the base metal zone are approximately at the same level, and the deformation under tensile stress continues uniformly until necking and final fracture occur in the base metal zone, as indicated by the arrows in Figure 13(g). In fact, the slight necking and final fracture occurring in the base metal zone show that the ultimate tensile strength in the heat-affected zone is greater than that in the base metal zone, and this phenomenon is consistent with the fact that the microhardness obtained in the heat-affected zone shown in Figure 13(d) is much greater than that in the base metal zone.

[0252] Based on the above results, it is obvious that the addition of titanium and niobium to the welding electrode will have a more obvious effect on the morphological change and the refinement of the austenite grain size in the molten zone. In addition, this addition will also result in a significant increase in the amount of high-density nanoscale titanium-rich titanium carbides and niobium-rich niobium carbides in the eutectic zone compared to the amount of carbides observed in Examples 10 and 11. More importantly, the characteristics of the high-density nanoscale κ-carbides present in the austenite dendritic cells and the eutectic zone are still largely retained. As described above, since the nanoscale κ-carbides (about 3-10 nanometers), titanium-rich titanium carbides, niobium-rich niobium carbides, vanadium-rich vanadium carbides, and austenite dendritic cells all have the same ductile face-centered cubic structure, an amazing strengthening effect can be achieved without a significant decrease in ductility. The detailed analysis in the above examples clearly indicates that by using the welding electrode provided by the present invention, the obtained molten zones will all have excellent combined values of microhardness, yield strength, ultimate tensile strength, and ductility.

[0253] Example 14

[0254] In the development of high-strength, especially high tensile strength and high ductility, precipitation-hardened austenitic Fe-Mn-Al-C alloys, significant progress has mainly relied on heat treatment, including solution heat treatment, quenching treatment, and ideal aging treatment, to obtain a high density of nanoscale κ-carbides within the austenite matrix. However, as described in the prior art, despite a large number of related studies conducted by many people in the past few decades, most of the problems related to the weldability of these alloys remain unresolved. Recently, in order to confirm the key role played by the high density of nanoscale κ-carbides within the austenite matrix in retaining the strength (especially tensile strength) and ductility of welded parts, Jeong et al. attempted to use Gleeble simulation in 2019 to reveal the microstructure in the weld heat-affected zone during welding and its influence on the mechanical properties of the weld heat-affected zone of a lightweight precipitation-hardened austenitic Fe-31.4Mn-11.4Al-0.9C alloy. The alloy they used was hot-rolled to a thickness of 13 mm at 1200 °C, water quenched, solution heat-treated at 1050 °C for 2 hours, then water quenched again, and then aged at 550 °C for 100 minutes to obtain a high density of nanoscale κ-carbides within the austenite matrix. The aging treatment obviously increased the hardness of the alloy, increasing it from 298 Vickers hardness in the solution heat-treated state to 349 Vickers hardness after aging at 550 °C for 100 minutes. The simulation test was carried out using a Gleeble simulation testing machine. The results showed that when the simulated test specimens cooled to room temperature, softening occurred in each of the specimens that had been age-hardened, and for the specimens with a peak temperature of 1150 °C, it was observed that the κ-carbides completely dissolved. From these observed phenomena, it can be inferred that there is definitely a relationship between the κ-carbide precipitates and the transformation of the mechanical properties in the weld heat-affected zone. However, it should be noted that in these Gleeble simulation test studies, no direct fusion welding was carried out, so there was no remelting and solidification process, and thus the actual microstructure and property values of the fusion zone could not be directly obtained. However, in their Gleeble simulation test studies, after the aged specimens were heated to a peak temperature of 750 °C to 1150 °C and then cooled to room temperature, dissolution of the previously existing strengthening κ-carbides and severe softening were observed in the weld heat-affected zone. In fact, this kind of phenomenon is very similar to what happens during fusion welding of 7xxx-series precipitation-hardened aluminum alloys, as described in the prior art.

[0255] Regarding the above-described features and characteristics, more details and related discussions can be found in the following published literature.

[0256]

[32] S. Jeong, G. Park, B. Kim, J. Moon, S. J. 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.

[0257] In this embodiment, the effect of fusion welding on the characteristics of a welded joint using an age-hardened FeMnAlC alloy as the base material was examined. The standard composition of the base material was Fe - 31.5Mn - 8.5Al - 1.25C. Before fusion welding, the hot-rolled base material was solution heat-treated at 1050 °C for 1 hour, water quenched, and then aged at 550 °C for 12 hours. Two aged base material plates with dimensions of 80 mm × 80 mm × 8 mm were machined to form a single V-groove butt joint. The welding parameters used during tungsten inert gas arc welding were similar to those used in Example 1. The welding electrode used in this embodiment was the same as that used in Example 3, with a diameter of approximately 2.4 mm.

[0258] Figure 14(a) is a scanning electron microscope image that shows the microstructure of a region within the post-weld sample that includes a fusion zone, a weld heat-affected zone, and a base metal zone. In Figure 14(a), it is clearly seen that the microstructure of the fusion zone is very similar to that seen in Figure 3(d), with the typical microstructure of columnar austenite grains and a high density of nanoscale κ-carbides distributed in the austenite dendritic cells and eutectic regions. However, in this embodiment, it can be clearly observed that the amount of nanoscale κ-carbides located within the weld heat-affected zone is significantly reduced, indicating that during fusion welding, most of the nanoscale κ-carbides originally present in the age-hardened base metal zone have dissolved. Based on the investigations of the present invention, this seems to be because of the insufficient carbon content in the base metal used in this embodiment. This result is the same as the research report by Jeong et al. in 2019, where they used Gleeble simulation tests to investigate the effect of fusion welding on the weld heat-affected zone of an age-hardened Fe-31.4Mn-11.4Al-0.9C alloy, and found the dissolution of κ-carbides and significant softening in the weld heat-affected zone. In fact, this is also quite similar to the phenomenon observed during fusion welding of precipitation-hardened AA7075 aluminum alloy, as described in the prior art.

[0259] Figure 14(b) shows a scanning electron microscope image of the Vickers microhardness test indentations made across the fusion zone, the weld heat-affected zone, and the base metal zone. The corresponding microhardness values measured at each measurement point are shown in Figure 14(c), which shows that the average microhardness of the fusion zone, the weld heat-affected zone, and the base metal zone of this post-weld sample is approximately 406 Vickers hardness, 282 Vickers hardness, and 380 Vickers hardness, respectively. Clearly, the dissolution of the strengthening nanoscale κ-carbides in the weld heat-affected zone results in severe softening, as described in most of the prior art cited above. Figure 14(d) shows the macroscopic images of the sample before and after the tensile test. It can be clearly seen from the tensile fracture sample that due to the severe softening in the weld heat-affected zone, the fracturing path propagates along the interface between the weld heat-affected zone and the base metal zone, as indicated by the arrow.

[0260] Based on the above-described results, it is clearly understood that in order to retain the strengthening nanoscale κ-carbides in the fusion zone and the weld heat-affected zone, appropriate alloy design is basically important for both the welding electrode and the base metal.

[0261] Example 15

[0262] The purpose of this embodiment is to attempt to clarify what kind of influence adding molybdenum to the welding electrode will have on the characteristics of the fusion zone in the post-weld sample. The base material used in this embodiment is the same as that used in Example 1. The size of the two base material plates is 80 mm × 80 mm × 10 mm, and a single V-groove butt joint is formed after machining. The welding parameters used during tungsten inert gas arc welding are similar to those used in Example 1. The standard composition of the welding electrode 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 about 3.2 mm.

[0263] Figure 15(a) is a scanning electron microscope image, which shows the fusion zone, the heat-affected zone of the weld, and the base material zone in the welded sample. It is obvious that the microstructure of the fusion zone mainly contains typical columnar austenite grains. At the same time, a high density of nano-scale κ-carbides can be observed in the austenite dendritic cells and eutectic regions, which is similar to the phenomena observed in Examples 1 to 5. However, at the same time, a large number of micron-sized coarse particles and wide precipitate-free zones are also observed at the austenite grain boundaries (as indicated by the arrows). Figure 15(b) shows the analysis results of a scanning electron microscope and an X-ray energy dispersive analyzer for a coarse particle, showing that these coarse particles are Mo-rich Mo-carbides. Many studies have confirmed that the coarse particles existing at the grain boundaries and the accompanying precipitate-free zones have a serious destructive effect on the ductility and strength of the alloy. Therefore, it is not recommended to add molybdenum alloying elements that can strongly form carbides to the welding electrode provided by the present invention.

[0264] Example 16

[0265] The purpose of this embodiment is to attempt to clarify what kind of influence adding chromium to the welding electrode will have on the characteristics of the fusion zone in the post-weld sample. The base material used in this embodiment is the same as that used in Example 1. The size of the two base material plates is 80 mm × 80 mm × 10 mm, and a single V-groove butt joint is formed after machining. The welding parameters used during tungsten inert gas arc welding are similar to those used in Example 1. The standard composition of the welding electrode used is iron-29.3 manganese-8.8 aluminum-1.58 carbon-1.5 chromium (Fe-29.3Mn-8.8Al-1.58C-1.5Cr), with a diameter of about 3.2 mm.

[0266] Figure 16(a) is a typical scanning electron microscope image of a region within the post-weld sample that includes a fusion zone, a weld heat-affected zone, and a base metal zone. The microstructure of the fusion zone mainly includes typical columnar austenite grains. At the same time, high-density nanoscale κ-carbides can be observed within the austenite dendritic cells and eutectic regions, which is similar to the phenomena observed in Examples 1 to 5. However, at the same time, a large number of micron-sized coarse particles are also observed at the austenite grain boundaries (as indicated by the arrows), along with a precipitate-free zone surrounding the coarse particles. Figure 16(b) shows the scanning electron microscope and X-ray energy dispersive analyzer analysis results of a coarse particle, indicating that these coarse particles are Cr-rich Cr-carbides. As described in Example 15 above, the presence of coarse particles at the grain boundaries and the accompanying precipitate-free zone can cause highly detrimental effects on the ductility and strength of the alloy. Therefore, it is not recommended to add chromium alloying elements that strongly form carbides to the welding electrodes provided by the present invention.

[0267] In the present invention, interestingly and notably, although titanium, niobium, vanadium, molybdenum, and chromium are all alloying elements that strongly form carbides, in the above-described examples (Examples 10 to 13, Example 15, and Example 16), it has been confirmed that these metals have very different effects on the microstructure of the fusion zone of the welded samples. For titanium, niobium, and vanadium, the present invention clearly reveals that during fusion welding, these metals can form nanoscale (6 - 10 nm) titanium-rich titanium carbides, niobium-rich niobium carbides, and vanadium-rich vanadium carbides within the eutectic region. These metal carbides formed in the eutectic region during fusion welding can not only significantly improve the microhardness and yield strength but also retain excellent ductility within the fusion zone of the welded alloy. In contrast, our tests have also shown that adding molybdenum and chromium to the welding electrode results in the formation of micron-sized coarse molybdenum-rich molybdenum carbides and chromium-rich chromium carbides at the austenite grain boundaries, along with an obvious precipitate-free zone surrounding the coarse carbides within the fusion zone, and these two can have a negative impact on the ductility of the alloy in the post-weld state.

Claims

1. An electrode for fusion welding of austenitic iron-manganese-aluminum-carbon alloy for precipitation hardening, characterized in that, by weight percentage, it contains: 23 - 34% manganese, 7.5 - 11.5% aluminum, 1.48 - 1.95% carbon, and the balance is iron.

2. An electrode for fusion welding of austenitic iron-manganese-aluminum-carbon alloy for precipitation hardening, characterized in that, by weight percentage, it contains: 24 - 32% manganese, 8.0 - 11.0% aluminum, 1.48 - 1.95% carbon, and the balance is iron.

3. An electrode for fusion welding of austenitic iron-manganese-aluminum-carbon alloy for precipitation hardening, characterized in that, by weight percentage, it contains: 23 - 34% manganese, 7.5 - 11.5% aluminum, 1.48 - 1.95% carbon, 0.1 - 2.5% titanium, and the balance is iron.

4. An electrode for fusion welding of austenitic iron-manganese-aluminum-carbon alloy for precipitation hardening, characterized in that, by weight percentage, it contains: 23 - 34% manganese, 7.5 - 11.5% aluminum, 1.48 - 1.95% carbon, 0.1 - 3.0% niobium, and the balance is iron.

5. An electrode for fusion welding of austenitic iron-manganese-aluminum-carbon alloy for precipitation hardening, characterized in that, by weight percentage, it contains: 23 - 34% manganese, 7.5 - 11.5% aluminum, 1.48 - 1.95% carbon, 0.1 - 2.5% vanadium, and the balance is iron.

6. An electrode for fusion welding of austenitic iron-manganese-aluminum-carbon alloy for precipitation hardening, characterized in that, by weight percentage, it contains: 23 - 34% manganese, 7.5 - 11.5% aluminum, 1.48 - 1.95% carbon, and at least two elements selected from titanium, niobium, and vanadium, the total amount of which is less than or equal to 3%, and the balance is iron.

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