The Combination of Direct Aging and Cryogenic Treatment Effectively Enhances the Mechanical Properties of 18Ni300 by Selective Laser Melting

Abstract

This study systematically explores the synergistic effects of direct aging treatment (480 °C for 6 h) combined with cryogenic treatment (−196 °C for 8 h) on the mechanical properties and microstructural evolution of 18Ni300 maraging steel fabricated via selective laser melting (SLM). Three conditions were investigated: as-built, direct aging (AT6), and direct aging plus cryogenic treatment (AT6C8). Microstructural characterization was performed using optical microscopy (OM), scanning electron microscopy (SEM), transmission electron microscopy (TEM), and X-ray diffraction (XRD), while the mechanical properties were evaluated via microhardness and tensile testing. The results show that the AT6C8 sample achieved the highest microhardness (635 HV_0.5) and tensile strength (2180 MPa), significantly exceeding the as-built (311 HV_0.5, 1182 MPa) and AT6 (580 HV_0.5, 2012 MPa) samples. Cryogenic treatment induced a notable phase transformation from retained austenite (γ phase) to martensite (α phase), with the peak relative intensity ratio ranging from 1.42 (AT6) to 1.58 (AT6C8) in the XRD results. TEM observations revealed that cryogenic treatment refined lath martensite from 0.75 μm (AT6) to 0.24 μm (AT6C8) and transformed reversed austenite into thin linear structures at the martensite boundaries. The combination of direct aging and cryogenic treatment effectively enhances the mechanical properties of SLM-fabricated 18Ni300 maraging steel through martensite transformation, microstructural refinement, and increased dislocation density. This approach addresses the challenge of balancing strength improvement and residual stress relaxation.

1. Introduction

18Ni300 maraging steel (MS), a low-carbon high-strength alloy, has been widely used in critical fields such as automobile, aerospace, and mold manufacturing, as well as for bearing gears parts [1,2], because of its exceptional toughness and ultra-high strength. With technological advancements in these fields, there is an urgent need for complex-shaped 18Ni300 components, such as lightweight structures or parts with internal conformal cooling channels, to improve system performance. Such complexities are beyond the capabilities of conventional manufacturing methods.

Selective laser melting (SLM) is an additive manufacturing technique that uses high-powered lasers to melt and fuse small metallic or insulating particles for the development of prototypes, tools, and functional parts with the aid of a three-dimensional CAD model [3]. Compared to conventional cast and forged steel techniques, maraging steel fabricated by SLM exhibits distinct advantages, including superior strength, ductility, toughness, a high manufacturing efficiency, and the ability to produce geometrically complex components with high dimensional accuracy. These enhancements are attributed to its refined cellular microstructure, ultrafine melt pool dimensions, and extremely high cooling rates during rapid solidification [4,5,6]. In addition, due to the layer-by-layer forming characteristic, an SLM-fabricated alloy enables the building of customized parts and a reduction in the waste of materials through the recycling of unmelted powder [7,8]. There has been some research on 18Ni300 in SLM-fabricated steel, such as work examining the influences of processing parameters on the relative density or the effects of heat treatment on the mechanical properties of as-built specimens in SLM fabrication.

Zhao et al. [9] investigated the effect of heat treatment on the microstructure and hydrogen embrittlement of SLMed18Ni300 maraging steel. Their results demonstrated that solution treatment can reduce microporosity and eliminate residual stress. After further aging treatment, both 18Ni300 steels showed the worst resistance to hydrogen embrittlement. Similarly, Ma et al. [10] systematically investigated the microstructural evolution and mechanical properties of SLM 18Ni300 MS under distinct heat treatment protocols. The first treatment (HT1) entailed direct aging of as-fabricated samples at 480 °C for 6 h, followed by furnace cooling. The second treatment (HT2) consisted of solution annealing of the as-built specimens at 820 °C for 45 min, air quenching, and subsequent aging at 480 °C for 6 h. Notably, their findings demonstrated that HT1 yielded a more desirable microstructure and exhibited superior mechanical performance compared to HT2. Furthermore, Mao et al. [5] conducted processing optimization and investigated the microstructure, mechanical properties, and nanoprecipitation behavior of 18Ni300 maraging steel from selective laser melting. Their study revealed that the solution-aged treatment contributed to a marked increase in UTS (1915 MPa) but resulted in a decrease in the elongation at break of the SLM-fabricated samples from 10.5% to 5%. Existing studies [9,10] have predominantly employed heat treatment methods to mitigate the detrimental effects induced by the rapid heating and cooling cycles that are inherent in the SLM process on the mechanical properties of 18Ni300 maraging steel. However, direct aging treatment (DAT) or solution aging treatment (SAT) alone cannot effectively resolve these existing challenges. While DAT can enhance mechanical properties, it introduces additional residual stress. Although SAT can nearly eliminate residual stress, this comes at the expense of mechanical performance. Consequently, developing approaches that simultaneously improve mechanical properties while mitigating residual stresses has attracted significant research attention.

Cryogenic treatment is a heat treatment process that involves exposing materials to extremely low temperatures, typically ranging from −150 °C to −196 °C, with liquid nitrogen commonly being used as the cooling medium. This treatment is employed to enhance the physical and mechanical properties of materials. In addition, the combination of cryogenic treatment with other heat treatment methods has been extensively investigated by researchers. Ding et al. [11] investigated the effect of the combined use of cryogenic and aging treatment on the mechanical and damping properties of Mn-Cu alloy based on a response surface model. Their results showed that the maximum tensile strength was 396.7 MPa when the alloy underwent cryogenic treatment (−196 °C for 30 h) followed by aging treatment (428 °C for 2 h). This value was 38 MPa higher than that of the as-cast alloy, and the internal friction value (Q⁻¹) could reach 0.074, which was 124% higher than that of the as-cast alloy. The enhancement in mechanical properties is attributed to the formation of precipitates and the phase transformation-induced increase in twinning after cryogenic treatment (−196 °C for 30 h) combined with aging (428 °C for 2 h). Moreover, Liu et al. [12] investigated the effect of deep cryogenic treatment on the microstructure and properties of an Al-Cu-Mg-Ag alloy with re-aging treatment (RRA) and found that deep cryogenic treatment can considerably increase the alloy’s hardness and wear resistance. According to the experimental findings, the alloy’s hardness was 155.9 HV following re-aging treatment combined with deep cryogenic treatment (RRA + DCT3), which was 6.93% higher than that following RRA alone. This improvement can be attributed to the cryogenic treatment inducing solute supersaturation in small-sized precipitates that are formed during tempering, which promotes the precipitation of denser phases while simultaneously suppressing the growth and coarsening of larger untempered precipitates. Consequently, a more uniform and compact precipitate distribution is achieved during subsequent re-aging, leading to enhanced overall alloy performance. Similarly, Wang et al. [13] studied the cryogenic mechanical properties of 316L stainless steel fabricated by selective laser melting (SLM) and revealed that due to microstructural refinement and an increase in dislocations caused by the cryogenic treatment, SLM-fabricated TC4 exhibits enhanced tensile properties. Compared to untreated specimens, the hardness, tensile strength, yield strength, and elongation of CT36 increased by 5.48%, 7.1%, 17.6%, and 36.7%, respectively. This performance enhancement originates from cryogenic treatment-induced microstructural refinement and dislocation multiplication.

Therefore, this study systematically investigates the synergistic effects of direct aging (480 °C for 6 h) combined with cryogenic treatment on the mechanical properties of SLM-fabricated 18Ni300 MS. The work aims to address the research gap in cryogenic treatment applications for SLM-fabricated 18Ni300 MS and enable its potential deployment in low-temperature aeroengine components (e.g., turbine cooling systems, air pre-coolers), where cryogenic strength and dimensional stability are critical performance metrics.

2. Materials and Methods

2.1. Preparation of 18Ni300 Samples

The 18Ni300 alloys were produced using the SLMed technique on a HBD E400 machine from Hanbang Co., Ltd., Tianjin, China. The powder had a spherical shape, as shown in Figure 1, with a particle size distribution of D10 = 18 μm, D50 = 34 μm, and D90 = 56 μm, measured by laser diffraction (Malvern Mastersizer 3000, wet dispersion mode). Prior to SLM, the powder was sieved through ASTM 270-mesh (53 μm) and 325-mesh (45 μm) sieves for 30 min to remove outliers [14]. Its chemical composition was determined using an inductively coupled plasma optical emission spectrometer (ICP-OES), as detailed in

Table 1. The chemical composition of the 18Ni300 powder (in wt%).

Element	Ni	Co	Mo	Ti	Al	Si	Mn	Cr	P	C	Fe
Wt%	17.69	8.59	4.59	0.70	0.12	0.086	0.014	0.018	0.014	<0.005	Bal

. To ensure printing quality, oversized particles were removed. Printing parameters included a 320 W laser power, a scanning speed of 1000 mm/s, a 0.1 mm hatch spacing, a 60 μm layer thickness, and a zigzag pattern at a 67° angle. Argon gas was used for protection.

2.2. Sample Processing Procedure

The as-fabricated SLM 18Ni300 sample, as shown in Figure 2a, was sectioned into small cubes with dimensions of 10 × 10 × 5 mm³ and 6 × 6 × 5 mm³ using wire electrical discharge machining (WEDM). All sample surfaces were sequentially ground with SiC abrasive papers ranging from 400 to 2000 grit, followed by diamond particle polishing. Metallographic etching was subsequently performed using a 4% (HNO₃, 65–68 wt% concentration) ethanol (C₂H₅OH, ≥99.7% purity) solution. The etched microstructure was immediately examined using optical microscopy (OM). Figure 2b,c present OM images of the etched cross-sections on both the horizontal plane (perpendicular to the build direction, BD) and vertical plane (parallel to BD). Based on these observations, subsequent experiments were conducted along the alloy’s building direction (XOZ plane). The post-processing of the 18Ni300 powder consisted of three steps. First, SLMed treatment was conducted. Then, direct aging treatment was performed at 480 °C for 6 h, followed by water cooling. Finally, cryogenic treatment was carried out at −196 °C for 8 h. After cryogenic treatment, samples were allowed to heat naturally to room temperature. The naming of the samples is shown in

Table 2. Sample naming and treatment methods.

Sample Name	SLMed 18Ni300	Direct Aging Treatment at 480 °C for 6 h	Direct Aging Treatment at 480 °C for 6 h
As-built	√
AT6	√	√
AT6C8	√	√	√

. To avoid the oxidation of the samples and ensure the success of the experiment, a muffle furnace filled with high-purity argon gas was used for the direct aging treatment.

2.3. Sample Characterization and Testing

The surface of each sample was polished with grinding papers from 400 to 2000 grit and then polished with diamond particles. Metallographic etching was performed using a 4 vol% nitric acid (HNO₃, 65–68 wt% concentration) in ethanol (C₂H₅OH, ≥99.7% purity) solution and observed immediately after corrosion. The 18Ni300 samples were characterized using a Zeiss Axion Observer 3 m metallographic microscope (Oberkochen, Germany), and fractographic analysis was performed on the fractured surfaces of tensile specimens using field emission (Thermo Scientific field emission scanning electron microscope, Waltham, MA, USA) SEM, with an acceleration voltage of 5 kV and working distance of 8 mm. Specimens were cleaned ultrasonically in ethanol for 5 min prior to observation to remove surface contaminants. The phase composition of the samples was characterized by X-ray diffraction (XRD, D8 ADVANCE, Bruker, Karlsruhe, Germany) using Cu Kα radiation (λ = 0.15418 nm) at 40 kV and 40 mA, with a scanning rate of 2°/min and a 2θ range of 20–100. TEM observations were conducted on a JEOL JEM-2100 (Tokyo, Japan), operated at 200 kV. For TEM observation, specimens were prepared via electrolytic twin-jet polishing. The electrolyte consisted of 10% perchloric acid and 90% glacial acetic acid, with a voltage of 40 V and a current of 80 mA being applied. The mechanical properties of the samples were characterized through hardness and tensile tests. Vickers hardness measurements were performed using an HV_0.5 S100 digital hardness tester, with a load of 0.5 kg being applied for 10 s. For each treatment condition, three hardness measurements were conducted, and the average value was calculated. The tensile properties of the samples were tested on a microcomputer-controlled electronic universal testing machine at a tensile rate of 1 mm/min. Each condition was tested three times, and the average value was determined. The dimensions of the tensile specimens are illustrated in the upper right corner of Figure 2a.

3. Results and Discussions

3.1. Mechanical Properties

Figure 3a illustrates the trend of microhardness variation of the AT6 sample over a period ranging from 0 h to 8 h. Initially, it can be observed that the microhardness shows a fluctuating trend, with values ranging approximately from 590 HV_0.5 to 605 HV_0.5. However, as the cryogenic treatment time extends from 5 to 8 h, the microhardness rises rapidly and reaches a peak of 635 HV_0.5 at 8 h. This phenomenon may be attributed to the combined effects of dislocation and α phase transformation during the cryogenic treatment [15]. During cryogenic treatment, the extremely low temperature causes a rapid contraction in the volume of the specimens, leading to changes in internal stress [16]. This stress change promotes grain refinement and an increase in dislocation density, introducing more grain boundaries, dislocations, and fixed points into the material. In contrast to the rapid change in volume shrinkage, the diffusion capacity of elements is very low at low temperatures, making the phase transformation process relatively slow [17]. Therefore, in the early stage of cryogenic treatment, the effect of the dislocation density on the hardness is more significant than that of phase transformation, thereby improving the overall microhardness of the material. However, as the cryogenic treatment time extends, the increase in dislocation density approaches saturation [18,19,20]. Phase transformation plays a dominant role in the improvement of hardness. The γ phase is a soft phase, while the α phase is a hard phase. After cryogenic treatment for 8 h, the overall hardness shows an upward trend. Figure 3b shows the microhardness of the as-built, AT6, and AT6C8 samples. When comparing the microhardness of the as-built, AT6, and AT6C8 samples, the as-built sample has the lowest microhardness with 311 HV_0.5, while the AT6C8 sample exhibits the highest value with 635 HV_0.5.

This suggests that the AT6C8 treatment significantly enhances the material’s hardness compared to the other two treatments. The stress–strain curves for the as-built, AT6, and AT6C8 samples are shown in Figure 3c. The tensile strength of the AT6C8 sample reaches the maximum value, with a peak of 2180 MPa. Next is the AT6 sample, whose tensile strength is 2012 MPa, and finally, the tensile strength of the as-built sample is 1182 MPa. This indicates that the AT6C8 treatment not only increases hardness but also improves the material’s strength. The strength is mainly provided by the martensite and the precipitated phases, which are dispersedly distributed.

3.2. XRD Analyses

Figure 4 shows the phase analyses of three samples. By comparing the three figures, it can be seen that under different processing states (SLM, AT6, AT6C8), the phase composition of the SLM-fabricated 18Ni300 steel is basically the same, but the relative intensities of the diffraction peaks for each crystal plane vary. This reflects the impact of different processing techniques on the content and distribution of martensite and austenite. The samples were composed of a γ phase with a face-centered cubic (FCC) structure (primarily austenite) and an α phase with a body-centered cubic (BCC) structure (consisting mainly of martensite) [21]. Among them, the peak intensity of (111)_γ in Figure 4b is higher than that in Figure 4a, which implies that the aging treatment of AT6 has led to the generation of reversed austenite [6]. To quantify the phase transformation, the relative intensity ratio of the (110)_α peak to the (111)_γ peak was calculated. For the AT6 sample, this ratio was 1.42, while for the AT6C8 sample, it increased to 1.58, indicating a significant phase transformation from retained austenite (γ phase) to martensite (α phase) during cryogenic treatment. This phenomenon aligns with the observed hardness enhancement, as martensite is a hard and brittle phase [22]. Meanwhile, weak Ni₃Mo phases can also be observed in both AT6 and AT6C8 samples. It is speculated that this may be due to the precipitation of nanometallic intermetallic compounds during the aging process.

3.3. OM Analysis

OM analyses of the as-built samples are presented in Figure 5a. Unlike the microstructure of 18Ni300 cast alloys, in longitudinal sections, the melt pools exhibit a fish-scale pattern. During SLM fabrication, the laser beam scans the metal powder at a predetermined angle to generate molten pools. Through alternating cross-scanning of successive layers, the powder undergoes rapid melting and solidification. The overlapping interaction of these molten pools gives rise to a semi-elliptical laser track morphology [23], which is manifested as the observed fish-scale pattern on the longitudinal cross-section. Figure 5b displays the OM characteristics of the AT6 sample. In comparison to the as-built sample, although the macroscopic features still exhibit a fish-scale pattern, the fusion line has disappeared, and the molten pool boundaries have become blurred. Figure 5c illustrates the OM characteristics of the AT6C8 sample. The macroscopic morphology is analogous to that of the AT6 sample; however, certain molten pools have become notably more indistinct, to the extent that some have even vanished [24].

3.4. Fracture Morphologies

Figure 6a shows the tensile fracture morphology of the as-built sample. An obvious plastic deformation can be seen, and the fracture includes typical fiber zones, radial zones, and shear lips. Tear ridges are present in Figure 6a. This is a characteristic formed on the fracture surface during the crack propagation process in ductile fracture. Dimples are present in all three fracture morphologies. However, brittle regions are also evident in Figure 6b and Figure 6c, respectively. Figure 6b demonstrates a mixed ductile–brittle fracture mechanism. In comparison to the fracture morphology shown in Figure 6b, the one in Figure 6c not only displays dimples and brittle regions but also exhibits quasi-cleavage features. Therefore, the fracture surface in Figure 5c is classified as a quasi-cleavage fracture [25].

3.5. TEM Observations

Figure 7 presents the TEM microstructural images of the as-built sample. In Figure 7a, it can be observed that the reversed austenite is distributed at the boundaries of the lath martensite. Meanwhile, a large number of dislocations are present in the martensite matrix. As seen from the high-magnification image in Figure 7b, the retained austenite in elliptical and short rod-like shapes is discontinuously distributed at the martensite boundaries, with a thickness of approximately 50–200 nm [26]. Furthermore, the TEM analysis in Figure 7c reveals that numerous dislocations are entangled and distributed in the martensite matrix. It is widely acknowledged that dislocations are generated due to the repetitive heating events during laser processing. In the SLM-fabricated sample, the laser beam causes local melting of the powder, and the melt pool solidifies rapidly within milliseconds. Subsequently, each successive laser scan subjects the previously solidified regions to repeated thermal cycling [27]. As shown in Figure 7d, retained austenite is distributed between the martensite laths. Figure 7e,f show the dark-field (DF) images of the retained austenite and martensite, respectively, corresponding to the bright-field (BF) image in Figure 7d.

Figure 8 presents the TEM microstructural images of the AT6 sample. As shown in Figure 8a,b, the reversed austenite structure exists at the boundaries of the cellular or columnar morphologies. Numerous nanoprecipitates are uniformly and dispersedly distributed within the martensite, as shown in Figure 8c. Their lengths are roughly between 20 and 50 nm, and the widths are roughly between 5 and 10 nm. It was reported that these nanoprecipitates are Ni₃Ti, Ni₃Mo [28]. They can significantly obstruct dislocation motion via the Orowan mechanism [29], where the presence of these fine precipitates forces dislocations to bow around them, generating dislocation loops and increasing the resistance to plastic deformation. This mechanism substantially contributes to the enhancement of the material’s strength, as the increased resistance to dislocation glide effectively restricts the onset of plastic flow and improves the overall load-bearing capacity of the material. This also explains why the strength of the AT6 sample is significantly higher than that of the as-built sample.

Figure 9 shows the TEM microstructural images of the AT6C8 sample. The morphologies of reversed austenite and martensite are clearly observable in Figure 9a,b. The lath martensite is refined from 0.75 μm to 0.24 μm, while the reversed austenite transforms from the long strip shape in AT6 to a thin linear morphology and preferentially accumulates at the boundaries of martensite laths. This further demonstrates that in the AT6C8 sample, the unique cryogenic treatment process induces a drastic temperature drop, rendering the reversed austenite in the material unstable. Consequently, the reversed austenite transforms into martensite, leading to a significant reduction in the amount of reversed austenite and an increase in martensite content, which aligns with the XRD results of the AT6C8 sample. This morphological transformation of reversed austenite may exert significant influence on the microstructural evolution and mechanical properties of the material [30]. The increase in martensite content can enhance the hardness and strength of the material. Meanwhile, after the retained austenite transforms into martensite, a specific orientation relationship and interface exist between the newly formed martensite and the original martensite [16]. These interfaces impede the growth of martensite laths, and the newly formed martensite also spreads into a fine morphology, thereby refining the overall lath martensite structure. Additionally, cryogenic treatment substantially diminishes the atomic diffusion capability. As the atomic diffusion ability decreases, the growth rate of the martensite laths is retarded. Notably, the heightened stability of the surrounding untransformed austenite at low temperatures imposes a more pronounced hindrance on martensite lath growth, restricting their expansion space and ultimately leading to the formation of a finer-lath martensite structure [31]. The refinement of martensite laths increases the interfacial area between laths, which effectively obstructs dislocation movement. During external loading, dislocations encounter more barriers during glide, requiring higher applied stress to sustain their motion. This mechanism consequently enhances the material’s strength and hardness. Figure 9c, a magnified region of Figure 9a, reveals the presence of nanoprecipitates. These nanoprecipitates remain largely unaffected by cryogenic processing, maintaining their uniform distribution within the martensitic matrix [32]. Figure 9d illustrates the distribution of dislocations around the austenite. Cryogenic treatment induces a rapid temperature decline, triggering significant internal stress relief that generates additional dislocations, thereby further enhancing strength and hardness.

4. Conclusions

This study systematically investigates the effects of integrating direct aging treatment and cryogenic treatment on the mechanical properties and microstructural evolution of 18Ni300 MS fabricated via SLM, leading to the following key conclusions:

(1)

Direct aging (480 °C/6 h) combined with cryogenic treatment (−196 °C/8 h, AT6C8) significantly improves the mechanical properties of SLM-fabricated 18Ni300 MS. Compared to the as-built sample, AT6C8 shows 104% higher microhardness (311→635 HV_0.5) and 84% higher tensile strength (1182→2180 MPa), outperforming single direct aging (AT6). This confirms that cryogenic treatment synergizes with aging to enhance strength and hardness via microstructural modification.

(2)

The strength enhancement in directly aged AT6 (480 °C/6 h) mainly results from nanoscale intermetallic precipitation. TEM shows 20–50 nm-long, 5–10 nm-wide nanoprecipitates uniformly dispersing in the martensite matrix during aging, acting as effective dislocation barriers via the Orowan mechanism.

(3)

Cryogenic treatment induces significant austenite (γ-α) transformation, with hard–brittle martensite enhancing the hardness and strength. TEM observations reveal that the lath martensite grains are refined, with an increased density of grain boundaries. This hinders dislocation slip and effectively balances the enhancement in strength with the relaxation of residual stresses.