PSI - Issue 69

Haofei Zhu et al. / Procedia Structural Integrity 69 (2025) 113–120

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1. Introduction Failures of aerospace structural Ultra-high strength steels (UHSSs) are predominantly attributed to crack propagation under cyclic fatigue loading [1, 2]. Consequently, the principal design objective for UHSSs is to enhance damage tolerance, requiring the simultaneous optimization of both strength and toughness. However, the commonly used 300M steel in aircraft landing gears exhibits limited toughness, with a K IC /YS value of approximately 60 MPa m 1/2 /1570 MPa, which is insufficient to meet the evolving performance standards for aerospace structural components [3, 4]. Driven by the development goal of improving the toughness of UHSSs without sacrificing strength, several high Co-Ni secondary hardening steels such as AerMet100 [5] and M54 [6] have been successfully developed. The K IC /YS values for AerMet100 and M54 are approximately 126 (MPa m 1/2 )/1724 MPa and 115 (MPa m 1/2 )/1723 MPa, respectively, making them excellent candidates to replace 300M steel. However, both commercial steels have their limitations; AerMet100, with its high cobalt content (13.4 wt%), results in raw material costs that are ten times higher than 300M steel. Although M54 steel reduces the cobalt content from 13.4 wt.% to 7 wt.% compared with Aermet100 steel [7], leading to a 16.6% reduction in raw material costs, the introduction of 1.3 wt.% tungsten induces the formation of primary M 6 C carbides. This results in an increase in the solution temperature from 885 °C to 1060 °C, consequently elevating fuel costs. To address the high cost of secondary hardening steels, a new type of low-Cobalt M 2 C and NiAl co-precipitated secondary hardening steel has recently been developed. During aging, NiAl precipitates coherently in the martensitic matrix due to low lattice mismatch, which inhibits dislocation recovery, promotes M 2 C precipitation, and strengthens the material in compensation for the reduced M 2 C strengthening due to the lower cobalt content [8]. Moreover, due to its simple heat treatment process, low-Cobalt M 2 C and NiAl co-precipitated secondary hardening UHSS holds significant potential for broader applications [9, 10]. Optimizing the heat treatment process is the most effective way to increase strength without sacrificing the toughness of UHSS. Solid-solution treatment, the first step in steel heat treatment, is crucial as it involves several microstructural features that directly and significantly affect the mechanical properties of the steel. These microstructural features include: (1) the prior austenite grain size (PAGs), (2) the type, number density, and volume fraction of primary carbides, and (3) the amount of retained austenite [1, 11]. Zhang et al. [11] found that increasing the solid-solution treatment temperature increases the retained austenite content, promotes the dissolution of primary phases, and maximizes toughness. Furthermore, the dissolution of primary carbides has a more significant effect on strength than the growth of grain size. Wang et al. [1] also demonstrated that the dissolution of primary carbides benefits the improvement of tensile strength and toughness, while the coarsening of the PAG is associated with a decrease in both strength and toughness. Therefore, determining the appropriate solid-solution temperature range is critical for the strength and toughness of newly developed M 2 C and NiAl co-precipitated secondary hardening UHSS. This study aims to investigate the effect of solid-solution temperature on the mechanical properties of the new M 2 C and NiAl co-precipitated secondary hardening steels and to find a balance between undissolved primary carbides and optimal PAGs. 2. Experimental details The steel ingot, with a nominal composition of Fe-0.23C-5Co-11Ni-3Cr-1.3Mo-1Al (wt. %), was produced via vacuum induction melting (VIM) followed by electroslag remelting (ESR). The ingot was first homogenised at 1200 °C for 24 h and then forged into a square with an interface of 100×100 mm 2 . Fig. 1 illustrates the variation in equilibrium phase fractions with temperature, as calculated using Thermo-Calc software. It is evident that the dissolution temperature of M 7 C 3 carbide is the highest, at 862 °C. Therefore, to ensure the dissolution of M 7 C 3 carbide, a solid-solution temperature exceeding 862 °C is required. The forged samples underwent solid-solution treatment at temperatures ranging from 870 to 960 °C for 1 h, followed by oil quenching to room temperature. The quenched samples were then cryogenically treated at -73 °C for 1 h, after which they were allowed to return to room temperature in air. Finally, the samples were aged at 482 °C for 5 h and air-cooled to room temperature.