Manufacturing oxide-dispersion-strengthened steels using the advanced directed energy deposition process of high-speed laser cladding

The used metallic powder materials Fe20Cr4.5Al0.5Ti (wt.%) were received in an Ar-atomized state (Nanoval GmbH, Berlin, Germany) in a particle size distribution of 15–45 µm with a d₅₀ value of 32.1 µm (Fig. 1a) according to the manufacturer. Y₂O₃ powder (abcr GmbH, Karlsruhe, Germany) was received in a particle fraction of 20–40 nm. Due to massive agglomeration by electrostatic forces, nanometer-sized powder material cannot be conveyed via conventional powder feeding systems, which makes the manufacturing of a powder compound consisting of micrometer-sized metallic powder and nanometer-sized powder necessary. Manufacturing of a powder compound was performed via mechanical alloying in the planetary mill Pulverisette 4 classic line (Fritsch GmbH, Idar-Oberstein, Germany) using milling containers and grinding balls manufactured from yttrium stabilized zirconia (YSZ). The result was a powder compound consisting of metallic particles covered with nanometer-scaled Y₂O₃ particles on the surface. Only minor deformation of the metallic particles by the milling process was observed; hence, the powder compound was considered suitable for the used powder feed system (Fig. 1b).

EDS measurements on the surface of the powder particles shown in Table 1 confirm the presence of yttrium and oxygen on the surface of the mechanically alloyed powder material.

The HSLC process was performed using a 4-axis gantry system equipped with a fiber-coupled diode laser system LDF 2000-30 (Laserline GmbH, Mülheim-Kärlich, Germany) emitting laser radiation with wavelengths of 1025 nm and 1064 nm simultaneously. The laser beam was guided by an optical fiber and shaped via collimation (f_c = 200 mm) and focusing lens (f_f = 182 mm) to generate a focal spot with a diameter of approx. 0.6 mm and a top-hat beam profile. The powder material was fed by conventional disk-based powder feeding units PF 2/2 (GTV Verschleißschutz GmbH, Luckenbach, Germany) via Argon carrier gas. Before entering the powder nozzle, the powder gas stream entered a diffusor unit dividing the powder gas stream into three independent streams to homogenously feed the powder material into the continuous coaxial powder nozzle D40 (Fraunhofer ILT, Aachen), injecting the powder material into the melt pool formed by the laser beam on the substrate surface. A local argon shielding gas atmosphere was created by the coaxial powder nozzle to protect the melt pool from atmospheric oxidation (Fig. 2).

As a substrate material, steel tubes (AISI 304) with a wall thickness of 5 mm and an outer diameter of 50 mm were used, which were mounted on a lathe machine Twister FU-180 (Neureiter Maschinen GmbH, Kuchl bei Salzburg, Austria). The substrates were sand-blasted and cleaned with ethanol before mounting to avoid surface contaminations. The process parameters for the HSLC process used are shown in Table 2.

The determination of the temperature distribution and the solidification kinetics during the HSLC process was performed via simulations. The simulation is based on a mathematical model operating with the integration of the time-dependent heat equations in conjuction with calculations of the track geometry using nonlinear Young–Laplace equations and solved with the finite element method (FEM) [27]. Temperature-dependent thermophysical data for heat capacity, thermal conductivity, and thermal expansion were taken from Kanthal APMT FeCrAl-based steel [41]. A density of 7.25 g/cm³ [42] is used for modeling. Due to their small volume fraction, the influence of the dispersed nanoparticles was considered negligible for the FEM melt pool simulations.

Figure 3 provides LOM images of cross sections consisting of two subsequently manufactured layers for both material types. The non-reinforced coating, manufactured with Fe20Cr4.5Al0.5Ti (wt.%) material, (Fig. 3a) features full metallurgical bonding to the substrate material without any sign of delamination or defects in the fusion zone. The full double-layered coatings are found to be almost defect-free, exhibiting no cracks and low porosity, resulting in a relative density of 99.94%. The curvature of the fusion zone proves that substantial melting of substrate material occurred, resulting in an effective layer height of approx. 100 µm. The fact that the upper layer exhibits a larger layer height indicates significant remelting of the previously applied layer.

A strong texture in building direction (BD) is observed, indicating that the main heat diffusion direction is perpendicular to the building direction. The inclination of the grain growth direction of grains starting at the substrate–cladding interface is caused by the curvature of the melt pool [43, 44]. The curvature of the melt pool is directly connected with the melt pool geometry and describes the solid–liquid interface from the bottom to the surface of the melt pool during the deposition process. [45]

The grain growth direction is not altered by the layer boundary (indicated by the red line), and epitaxial grain growth (indicated by purple arrows) is observed regularly. However, the growth direction changes significantly parallel to the traverse direction in the top area of the cladding (indicated by blue arrows). The abrupt change of the grain growth orientation can be explained by the curvature of the melt pool boundary and the energetically preferred growth direction, which is the smallest angle to the temperature gradient [44]. Because all these features can also be observed in the cladding manufactured with Fe20Cr4.5Al0.5Ti + 0.5% Y₂O₃ (wt.%) (Fig. 3b), it can be concluded that no apparent change between reinforced and non-reinforced cladding could be observed on a scale that can be resolved by LOM.

The microstructure of the non-reinforced material (Fe20Cr4.5Al0.5Ti) was examined using SEM images shown in Fig. 4a and is characterized by grain boundaries (marked with yellow arrows) in the build direction and a predominantly dendritic solidification structure. Within the microstructure, no precipitations or dispersoids can be detected (Fig. 4b). The sporadically observed flakes (green circles) could be identified as remnants of the etching process and were therefore discarded in the microstructure analysis.

SEM images of the reinforced material reveal the presence of sub-micrometer scaled oxide particles homogeneously distributed inside the grains, appearing as bright white spots throughout the microstructure (Fig. 4c). The microstructure evolution is not affected by the presence of the nanometer-scaled oxide particles, which are not acting as a site for heterogeneous nucleation of the ferritic matrix. This may be attributed to the low wettability of Y₂O₃ by iron–chromium melts [46]. The size of the nanoparticles (marked with red arrows) ranges between 50 and 150 nm (Fig. 4d). However, smaller nanoparticles may also be present, which cannot be resolved with the used SEM system. The nanoparticles are found to be significantly smaller than in PM2000 material processed via conventional DED (100–200 nm) [25] but larger than nanoparticles found in L-PBF processed with approx. 60 nm [21] and conventionally manufactured PM2000 with approx. 60 nm [47].

This process-related refinement of dispersed oxide particles can be further examined by modeling of the respective temperature distribution in the melt pool, the cooling rate, and the solidification rate encountered in the high-speed deposition process with the process parameters used (Table 2).

Comparing the calculated process-related values for HSLC with DED and L-PBF processes processing the same alloy, indicate that the considered values of maximum melt pool temperatures, solidification rates and cooling rates in HSLC are shifted from DED toward L-PBF (Table 3).

The maximum melt pool temperature directly influences the melt viscosity and may enhance nano-particle mobility in the melt pool. Therefore, low melt pool temperatures are considered favorable for preventing nano-particle agglomeration. In Fig. 5c, high-temperature gradients in HSLC-formed melt pools are observed, which are believed to cause severe melt pool dynamics induced by maragoni convection, effectively preventing particle agglomeration. We hypothesize that even disruption of occasionally formed agglomerates by maragoni-induced shear forces might occur. Another factor for successful dispersion of nanoparticles is high cooling and solidification rates limiting the existence time of the melt pool significantly inhibiting diffusion controlled agglomeration mechanisms [48]. Nanoparticles dispersed in the melt pool are instantaneously captured by the advancing solidification front and integrated in the mushy zone inhibiting the nanoparticles’ mobility and therefore preventing agglomeration tendencies. Low solidification rates allow nanoparticles to agglomerate in the melt pool impeding capture by the mushy zone and are consequently being pushed in front of the solidification front [49]. Additionally, advanced particle agglomeration enforces buoyancy-induced flotation effects of nanoparticles leading to slag formation on the clads’ surface and therefore reduced dispersion strengthening. [50] However, a direct comparison between the powder-blown deposition techniques of DED and HSLC, and the powder bed manufacturing process of L-PBF is difficult because of the strongly differing beam diameters (DED and HSLC: 0.66 mm [25], L-PBF: 0.08 mm [27]) leading to formation of increased melt pool dimensions in DED and HSLC, altering the solidification conditions significantly. The direct comparison of DED and HSLC with identical beam diameters used proves the ability to increase cooling and solidification rates leading to improved dispersion of nanoparticles in the stainless steel matrix reflected by smaller oxide particles in HSLC (50–150 nm) compared to DED (100–200 nm) [25].

The chemical composition of the dispersoids in the ODS specimen could be evaluated using EDS analysis on rarely observable agglomerates, revealing enrichments of yttrium, oxygen and titanium. This indicates the formation of Y₂Ti₂O₇ or Y₂TiO₅ compounds typically found in titanium-alloyed Fe–Cr-based ODS steels [51]. The formation of Y–Al–O particles, found in Al-alloyed Fe–Cr-based ODS steels [52, 53], could not be observed. These findings are in agreement with DED-manufactured material. [25] However, a detailed study using high-resolution electron microscopy is necessary to confirm these results.

The hardness profiles of both studied materials show a higher hardness of the reinforced material with hardnesses locally exceeding 300 HV0.1 (Fig. 6) proving the impact of nanoparticles on the hardness of the material. Please note that the nominal nano-particle mass loadings are different between the conventionally DED-manufactured sample (0.3 wt.%) [25, 54] and the HSLC (0.5 wt.%) processed material from this study.

However, the HSLC processed non-reinforced stainless steel (blue) exhibits a hardness of approx. 250 HV0.1, which is higher compared to its DED-manufactured specimen (approx. 230 HV0.1). Grain refinement enabled by higher solidification kinetics in the HSLC processed specimen may be a reason for this behavior. It can be concluded that nano-scaled dispersed oxides increase the hardness presumably by the Orowan mechanism leading to pronounced dispersoid hardening.