While classical high-strength steels (HSSs) such as the precipitation-hardened stainless steel 17–4 PH (Fe-17Cr-4Ni-4Cu, wt.%) and the maraging steel 18Ni-300 (Fe-18Ni-9Co-4Mo, wt.%) have been extensively investigated and adopted by additive manufacturing (AM), the state-of-the-art HSS grades are not that well accepted [1]. One major reason for the difference in adoption rates is the occurrence of cracking during rapid solidification, like many other alloy systems [2,3]. Yet, different from most crack-susceptible materials such as Al and Ni alloys, steels often experience multiple phase transformations during fabrication, making the task of crack alleviation within them ever more arduous.
The Custom 465® (C465, Fe-11Cr-11Ni-1.5Ti-1.0Mo, wt.%) HSS has been attracting growing attention from academia and industry since 2010 [4] due to its superior corrosion resistance, compared to conventional HSSs like 18Ni-300. Since it can operate without the need for protective coatings (e.g., cadmium and chrome that are toxic), it offers a more sustainable environmental footprint [5]. Within the group of stainless HSSs, C465 is reported to have one of the highest strengths, i.e., about 50 % higher than 17–4 PH (1751 vs. 1172 MPa), and at the same time possessing comparable corrosion resistance and toughness [4]. If these attributes are combined with the design freedom of AM, it is possible to envision several new applications and functionalities, e.g., impact absorbing foam on high performance cars, complex drill bits with internal cooling channels, and intricate high-thrust ship propellers, etc.
The conventional manufacturing (CM) route of C465 involves three processing steps after casting. (1) The alloy is first solution annealed (SA) around 900 °C for 1 h, to obtain a fully austenitic face-centered cubic (FCC) phase. (2) The material is then deeply quenched to −73 °C and held for 8 h to assist the martensitic transformation. (3) A final aging process at temperatures between 480 and 650 °C is then carried out, to fine-tune the characteristics of the hexagonal-close-packed (HCP) η-Ni3Ti strengthening precipitate and the reverted austenite [6]. In general, the volume fraction of the reverted austenite increases with the aging temperature, with a maximum amount of about 14 to 19 vol.% at the peak condition that occurs around 650 °C [7,8]. These reverted austenite grains preferentially nucleate at the prior-austenite or lath boundaries of the parent martensite phase, generally having a lower dislocation density compared to their neighboring martensite [7]. The η-Ni3Ti precipitate mainly nucleates within the martensite obeying an orientation relationship of and [6]. These precipitates have a rod-like morphology, and their length can range from ∼5 to ∼180 nm depending on the aging treatment employed [9]. The variation in the content of reverted austenite and the morphology of precipitate can yield a wide range of tensile, corrosion, and hydrogen resistance properties [8], [9], [10].
In this study, we attempt to fabricate C465 using the laser powder bed fusion (LPBF) technique. Our first results show that this alloy is highly susceptible to hot cracking during LPBF. The presence of hot cracks (also known as hot tearing) in AM alloys is often ascribed to the partition-induced liquid film near the end of solidification, which has lower solidus temperatures than its surrounding materials [2]. Several hot-crack mitigation approaches have been formulated in prior AM works. Kontis et al. [11] prevented hot cracking in nickel-base superalloy by lowering the volumetric heat input, hence reducing elemental partition and impeding low-solidus-temperature liquid film formation. Sun et al. [12] introduced secondary precipitates along grain boundaries during solidification to facilitate dendrite bridging, hence avoiding hot cracks. Sun et al. [3] also eliminated the hot cracks in nickel-base superalloy by controlling the competitive solute partition with the assistance of thermodynamic computations. Opprecht et al. [13] resolved the hot cracks in aluminum alloy via grain refinement.
Despite these earlier successes, as will be shown later, a direct adoption of the aforementioned hot-crack mitigation methods remains problematic for hot-crack-susceptible steels. This is because modern HSSs (mostly of the maraging grade) experience multiple phase transformations (unlike Al or Ni alloys which always have an FCC matrix), which could include primary and secondary phase formations during solidification, martensitic transformation during cooling, austenite reversion upon subsequent heating, and possible deformation-induced martensitic transformation under loading. As a result, retrieval of the elemental partition information right after solidification via experimental means is not straightforward. This, in turn, significantly hampers efforts through alloy design for eliminating hot cracking during AM.
Through the experimental work reported in this paper, we highlight the intricate nature and important aspects one needs to consider, when attempting to resolve hot cracking in high strength maraging steels, by using C465 alloy as an example. We will first present the severity of hot cracks in the as-built material, and the limitations of several existing hot-crack elimination methods. Crack-free specimens were obtained after introducing TiN particles into the precursor steel feedstock powders. The effect of TiN on the microstructure and tensile properties were also investigated and discussed in detail. We believe this work will not only assist in the adoption of high-strength stainless maraging steels to AM, but could also serve as a guide for any alloy that experiences phase transformations during production involving rapid solidification.
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