# Microstructural evolution in high-temperature direct aging on PBF-LB 15-5PH stainless steel

L.F. Kultz Unti, L.S. Aota, E.S.N. Lopes, G.G. Ribamar, N. Schell, J.P.

Oliveira, Baptiste Gault, J.A. Avila, A.L. Jardini, K.D. Zilnyk

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## To cite this version:

L.F. Kultz Unti, L.S. Aota, E.S.N. Lopes, G.G. Ribamar, N. Schell, et al., Microstructural evolution in high-temperature direct aging on PBF-LB 15-5PH stainless steel. Acta Materialia, 2026, 303, pp.121747. 10.1016/j.actamat.2025.121747. hal-05526330

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# Microstructural evolution in high-temperature direct aging on PBF-LB 15-5PH stainless steel

L.F. Kultz Unti<sup>a, 1*</sup>; L.S. Aota<sup>b,c</sup>; E.S.N. Lopes<sup>d</sup>; G.G. Ribamar<sup>ef</sup>; N. Schell<sup>f</sup>; J.P Oliveira<sup>e</sup>; B. Gault<sup>bg</sup>; J.A. Avila<sup>h</sup>; A.L. Jardini<sup>i</sup>; K.D. Zilnyk<sup>c</sup>

a) Departamento de Materiais e Processos, Instituto Tecnológico de Aeronáutica. Pç. Mal. Eduardo Gomes, 50, 12228-900. São José dos Campos-SP, Brazil

b) Department of Microstructure Physics and Alloy Design, Max Planck Institute for Sustainable Materials, Max-Planck-Straße 1, 40237. Düsseldorf, Germany.

c) Centro de Competência em Manufatura, Instituto Tecnológico de Aeronáutica (ITA). Pç. Mal. Eduardo Gomes, 50, 12228-900. São José dos Campos-SP, Brazil

d) Departamento de Engenharia de Manufatura e Materiais, Faculdade de Engenharia Mecânica, Universidade Estadual de Campinas (UNICAMP). R. Mendeleyev, 200 - Cidade Universitária, 13083-860. Campinas-SP, Brazil.

e) CENIMAT/I3N, Department of Materials Science, NOVA School of Science and Technology, Universidade NOVA de Lisboa, Caparica, 2829-516, Portugal

f) Helmholtz-Zentrum Hereon, Institute of Materials Physics, Max-Planck-Str. 1, Geesthacht, 21502, Germany

g) Univ Rouen Normandie, CNRS, INSA Rouen Normandie, Groupe de Physique des Matériaux, UMR 6634, F-76000 Rouen, France

h) Universidade Estadual Paulista (UNESP), Campus São João da Boa Vista, Av. Profª Isette Corrêa Fontão, 505, Jardim das Flores, 13876-750, São João da Boa Vista-SP, Brazil

i) Instituto de Biofabricação (BIOFABRIS), Universidade Estadual de Campinas. Cidade Universitária Zeferino Vaz, 13083-970. Campinas-SP, Brazil.

¹ Present address: Laboratório de Manufatura Avançada, Universidade Estadual de Campinas (UNICAMP). R. Mendeleyev, 200 - Cidade Universitária, 13083-860. Campinas-SP, Brazil.

* Corresponding author: luizunti@unicamp.br

## Abstract

High solidification rates and *in situ* heat treatments are commonly found in additive manufacturing (AM) of steels, resulting in a complex and far-from-equilibrium microstructure. Therefore, standard post-processing heat treatments commonly applied to wrought steels can favor the occurrence of different phenomena and can change the phase transformation sequence, due to the unique microstructure obtained by powder bed fusion – laser beam (PBF-LB). This work reports the microstructural evolution of 15-5 precipitation hardening (PH) stainless steel manufactured by PBF-LB during direct aging heat treatments at 621 °C (AMS H1150 standard condition), a route used to increase fracture toughness due to the martensite reversion and precipitates coarsening. The reversion of martensite into a Ni-rich austenite, predicted by kinetic calculations, was confirmed by high-energy X-ray diffraction (HE-XRD), being preferentially nucleated close to the copper-rich precipitates (CRPs), which can act as a preferential nucleation site. CRPs presented an oval shape, as confirmed by electronic microscopy (SEM and TEM) and atom probe tomography (APT). Fast Fourier transform (FFT) analysis of high-resolution TEM (HR-TEM) images suggests CRPs still present the metastable untwined 3R-type structure after 8h, rather than the most stable FCC structure. The presence of retained austenite, inherent to PBF-LB-processed PH steels, affects the CRPs evolution in different phases, and the CRPs themselves act as nucleation sites for Nb(C,N) secondary precipitation. These findings emphasize the necessity of microstructure-oriented heat treatment routes to unlock the full potential of additively manufactured PH stainless steels.

Keywords: Precipitation-hardening; Martensite reversion; Stainless steels; Additive Manufacturing; Phase transformation; CALPHAD.
# 1. Introduction

Laser Powder-Bed Fusion (PBF-LB) is an additive manufacturing method commonly used for processing metallic materials, which involves melting fine powder particles (~ 45 µm diameter) using a laser as the heat source. It has gained significant attention in the past decade owing to its capability to fabricate intricate and customized components with high-dimensional precision [1]. Successive layers of powder are added and consolidated above the previous ones, until the desired part is completed. Metals processed by PBF-LB method typically have a unique hierarchical and non-equilibrium microstructure that varies depending on the material and process [2], [3]. The melted material experiences high solidification rates, leading to fine (~ 500 nm diameter) columnar cells spanning across melt pools with a “fish scale” morphology due to epitaxial growth [4]. Another phenomenon during additive manufacturing is the appearance of heat-affected zones (HAZs) [3], [5], [6]. Due to the high laser power used in the equipment, the melt pool reaches elevated temperatures, and the surrounding areas are thermally affected. This iterative process acts as an in-situ heat treatment upon consolidation [7]. The intricate microstructure emerging from the PBF-LB manufacturing processes has a profound influence on the properties of the resultant materials [3], [8].

The 15-5PH stainless steel is a precipitation hardening stainless steel with high corrosion resistance in most environments. It presents a martensitic microstructure reinforced by the nanoprecipitation of copper-rich particles (CRPs). This steel grade is suitable for processing by PBF-LB, but the resulting microstructure presents particularities, such as a high fraction of retained austenite and the presence of crystalline nano-oxides [9]. The optimal aging temperature range for this steel, as recommended by the AMS 2759/3E standard [10] and supported by other authors [11], [12], [13], [14], is approximately 482 °C / 1 h (H900 condition, i.e. 900 °F / 1 h) for maximum mechanical strength (1170 MPa of UTS), and 621 °C / 4 h (H1150 condition, i.e. 1150 °F / 4 h) for improved fracture toughness (KIC up to 149 MPa.m¹/²) [15]. Variations in the size, shape, and crystal structure of the metallic Cu precipitates primarily cause the difference in the mechanical properties [16], [17], [18]. The aging process begins with forming small coherent spherical clusters within the martensitic matrix, as they possess a similar BCC crystal structure, reducing the energetic barrier for such nucleation. It evolves into a semi-coherent hexagonal 9R-type precipitate when it reaches a size between 2.5 and 4 nm [17]. With further heat treatment, these spherical particles preferentially grow in one direction, assuming an ellipsoidal shape, until they ultimately transform into a face-centered cubic (FCC) structure incoherent with the martensitic matrix (overaged state). Generally, this transformation occurs when the precipitate size exceeds 10 nm. It is also known that Ni enrichment can be commonly found near the precipitates, forming a core-shell structure [19].

In general, when it comes to the heat treatments of PH stainless steels manufactured by AM, and particularly for 15-5PH steel, there is still a wide discrepancy between the microstructures reported in the literature [20], especially regarding phase quantification and microstructural evolution during heat treatments. Although the AMS 2759/3E standard establishes 621°C/4h as the overaging condition (H1150) for conventional steels, the unique microstructure of PBF-LB-processed materials necessitates a critical re-evaluation of these parameters, particularly regarding phase stability and precipitation behavior. Within the
conditions selected for this study, hardness showed little variation among the 1, 2, and 4 h samples, while the 8 h specimen exhibited significant softening due to overaging. This work aims to elucidate the microstructural evolution mechanisms after standard heat treatment that generally lead to improved toughness, in PBF-LB-processed alloys. By clarifying these mechanisms, the study provides a basis for the development of optimized and process-specific heat treatment strategies for PBF-LB materials. The far-from-equilibrium microstructure can favor the occurrence of multiple different phase transformations simultaneously. Therefore, another interesting research topic pertains to the effectiveness of standard heat treatments on materials processed by AM, which must be adapted to non-equilibrium microstructures [21]. In a previous study [9], direct aging at 482°C (H900 condition) was investigated, resulting in an optimized maximum hardness of 505 HV after 2 h.

In the present work, a combined approach using experimental data and CALPHAD were used to thoroughly analyze the phenomena associated with direct aging during different holding times up to 8 h at 621 °C (H1150), elucidating the sequence of phase transformations during aging of an additively manufactured PH stainless steel grade. Furthermore, it was found that the formation of Nb(C, N) carbides follow a different path compared to wrought 15-5PH, due to the different thermal history compared to PBF-LB. The temperature variation during the AM process affects even the martensite lath size, resulting in a material with increased hardness even in an overaged state, compared to the conventional steel grade. By employing a combination of characterization techniques and thermodynamic/kinetic simulations, we unveil the interplay between precipitation and coarsening of CRPs, martensite reversion, and carbonitrides precipitation responsible for the microstructure evolution upon aging. These results show the necessity to adapt standard heat treatments, developed for conventional materials, to materials processed by PBF-LB.

## 2. Materials and Methodology

The used raw material was 15-5PH stainless steel (Stainless Steel PH1, supplied by EOS GmbH, Germany) nitrogen-atomized virgin powder with PSD (particle size distribution) between 15 and 45 µm [22]. Cubic-shaped samples with 10 mm of edge length were manufactured on an additive manufacturing machine (PBF-LB) EOS DMLS M280, using the standard process parameters recommended by the equipment manufacturer [9]. The heat source's scan strategy used an alternating stripe approach with a 67° laser scan rotation between layers. Samples were aged at 621°C directly from the as-built condition in a furnace under an argon flow, followed by air cooling. The holding times at the aging temperature were 1, 2, 4, and 8 h. For microstructural and hardness comparison, a wrought 15-5PH grade, supplied by Villares Metals (Brazil) in the annealed condition, was also characterized after 1 and 8h of aging.

The CALPHAD (Calculation of Phase Diagrams) methodology was employed using ThermoCalc 2025a software equipped with the TCFE12 database to predict the produced phase under thermodynamic equilibrium at 621 °C. Diffusional calculations were carried out using the diffusion module (DICTRA) equipped with the MOBFE8 database.
*In situ* high-energy synchrotron X-ray diffraction (HE-XRD) analysis was performed to simulate aging heat treatment at 621°C for up to 2h of holding time. This heat cycle was performed in a modified Bähl quench dilatometer installed at the P07 High Energy Materials Science Beamline (HEMS), at the Deutsches Elektronen-Synchrotron (DESY), PETRA III. Complete Debye-Scherrer rings were collected during the heat cycle at a rate of approximately 4 s per image using a monochromatic 87.1 keV X-ray beam (λ = 0.14235 Å) in transmission diffraction geometry. LaB₆ powder was used as a calibrant. HighScore Plus software was used for phase identification and data refinement [23], [24], while the phase quantification was performed using the Nelson-Riley function.

Samples were ground with SiC paper (up to #2400 mesh) and mechanically polished with diamond paste and colloidal silica (up to 0.02 µm) for SEM. A 390 x 360 µm² EBSD mapping was performed on the direct-aged AM samples (at 621°C for 1 and 8h) using a Zeiss Merlin SEM, with a step size of 0.1µm. Electron channeling contrast imaging (ECCI) micrographs [25] were acquired using a Zeiss Merlin field emission gun scanning electron microscope. Thin foils were prepared through grinding, followed by twin-jet electrolytic thinning and perforation using a 10% aqueous solution of HClO₄. STEM and energy dispersive spectroscopy (EDS) analysis were performed in a JEOL JEM-2100F Transmission Electron Microscope operated at 200 kV to identify precipitates and nanoscale particles on the 621°C/4h and 621°C/8h AM samples. In contrast, HRTEM analysis was performed in a JEOL JEM-2100 TEM.

Atom probe specimens were prepared from the alloys aged at 621°C/4 and 8 h following the lift-out procedure outlined in Ref. [26] with a FEI Helios Nanolab 600, Ga focused ion beam – scanning electron microscope (FIB-SEM). These specimens were analyzed by atom probe tomography (APT), using a Cameca LEAP (local electrode atom probe) 5000 XS equipment to probe the 3D chemical composition on a near-atomic scale[27]. The measurements were performed at 60 K while employing a laser pulse energy and frequency of 50-65 pJ and 200 kHz, respectively. The target ion detection rate was 1 in each 100 laser pulses. The acquired data was reconstructed based on the voltage curve [28] and analyzed with the CAMECA APSuite 6.3 software.

Vickers hardness tests by micro indentation were employed to monitor the effect of the heat treatments and determine the overaged conditions. The tests were conducted in a DuraScan 50 benchtop hardness tester. Ten measurements (0.1 kgf load and 10 s of dwell time) were taken randomly on a section parallel to the build direction for each sample.

# 3. Results and Discussion

## 3.1. Martensite reversion and microstructural evolution

*In-situ* high-energy X-ray diffractometry was used to investigate the martensite reversion during an aging heat treatment for up to 2h at 621 °C, with a heating and cooling rate of 10 °C/s. Figure 1 shows the results obtained during the isothermal stage: the FCC volume
fraction evolution (in grey), and the lattice parameters for both BCC and FCC phases. A high austenite (FCC phase) volume fraction was present in the as-built condition (near 21.5%, Supplementary Figure 1). Still, the further increase in FCC volume percentage could be evidence of a phase transformation on the CRPs, since copper in its most stable condition for the Fe-Cu binary system, known as ε-Cu, also has an FCC structure [29]. In fact, CRPs evolution presents a singular path with intermediary structures between BCC and FCC [29], [30], [31], which will be discussed later in the manuscript (Section 3.2.); yet the low fraction of CRPs (> 3%) and size (> 15-20 nm) decreases the detectability of these structures. The lattice parameters for both FCC and BCC phases presented a slight modification during the heat treatment, suggesting chemical composition changes in these phases. The austenite volume fraction achieved during the heat treatment did not transform into martensite during cooling to room temperature, reaching 27.1 % (Supplementary Figure 1).

Thermodynamic and atomic diffusion calculations were conducted to simulate the increase of the FCC phase volume fraction (martensite reversion), showing that the austenite stabilization at room temperature is due to compositional variation resulting from the long-range diffusion during the phase transformation. The following approach analyzed the martensite reversion on the whole system: a linear cell of 1 µm with phase fractions according to the as-built condition (78.5% Fe-α' and 21.5% Fe-γ) was used, as determined by HE-XRD. The martensite phase was inserted with Cu-depletion to simulate the martensite reversion into austenite on the interface after aging, i.e., after CRPs formation, while the Fe-FCC phase remained with the nominal composition since little precipitation occurs in the austenite [19]. Since the simulations only consider the diffusional nucleation and growth of austenite from martensite, the reversion reaction is mostly controlled by a diffusional mechanism, without reversion via a displacive mechanism [32].

<figure><img src="image_3.png" /><figcaption>Figure 1 – Austenite (FCC) volume fraction (in gray) and lattice parameter evolution (red and blue) obtained during *in situ* HE-XRD (high energy – X-ray diffraction) showing isothermal stage at 621 °C.</figcaption></figure>
Figure 2a-b shows the variation of the chemical composition near the BCC/FCC interface, while Supplementary Figure 2 indicates the effect of Cr and Ni variations in the Ms temperature. The reverted austenite should have a chemical composition different from that of the matrix martensite and the pre-existing retained austenite in the material (Figure 2a), being richer in Ni (12 at.%) and with a slightly lower Cr content than the nominal composition (13 at.%). Ni-rich reverted austenite is indeed found after annealing at 621°C for 1 h due to the formation of CRPs (shown in Section 3.2). As nickel has a smaller atomic radius (0.124 nm) than iron (0.126 nm) [33], this enrichment is related to the slight decrease in lattice parameter of austenite in Figure 1, considering the experimental lattice parameter as a global value resulting from both reverted and retained austenite. Since Ni is an austenitizing element, this reverted austenite remains stable at room temperature, given the reduced Ms temperature to values near 0 °C (Supplementary Figure 2), a behavior similar to the experimental observation described in the first paragraph (higher value of austenite at room temperature after direct aging) and reported in the literature for similar conditions [34]. As described by Zhou et al. [35] and it will be discussed later in Section 3.2 and Section 3.3, reverted austenite forms through a process in which Cu precipitates at martensite/retained austenite interfaces or martensite lath boundaries, followed by the partitioning of austenite-stabilizing elements, such as Ni, toward the interfaces of the Cu precipitates. It is also worth noting that the martensite composition exhibits much less variation, agreeing with what has been reported for wrought 15-5PH steel aged at 505 °C [19].

<figure><img src="image_4.png" /><figcaption>Figure 2 – Variation of the a) Cr and b) Ni composition over different holding times at 621°C, considering an interface between a copper-depleted (aged) martensite and retained austenite.</figcaption></figure>
A microstructural comparison was conducted between AM 15-5PH stainless steel subjected to an aging process at 621°C for 1 and 8 hours, employing EBSD analysis (Figure 3), showing inverse figure pole (IPF) and phase distribution maps for each condition. Selected areas, i.e., the melt pool (MP, the localized area of molten material) and the heat-affected zone (HAZ, the surrounding area affected by the thermal energy in the MP), were analyzed to compare morphology and the occurrence of austenite reversion. Each selected area was defined based on the grain size distribution, and their exact location can be found in Supplementary Figure 3. The selected aging conditions could be suitable for understanding the phase evolution in the whole isothermal aging at the selected temperature. It has already been reported that heat treatments above 650 °C can produce metastable austenite [36]. Figure 4 shows the grain size distribution for the martensitic (BCC) phase for both conditions, in each selected area mentioned above. Full lines represent the particle density function (PDF), while dashed lines represent the total fraction of martensitic grains.

After 1 hour at 621 °C, distinct variations in the grain size persist in some areas of the melt pool (Figure 4), and the 'fish-scale' morphology remains intact. The austenite phase was barely discernible because of its fine size, and it exhibited nearly twice as much FCC phase in more refined HAZ areas (Figure 3b and c) when compared to the coarse grains within the melt pool. The average lath size (for the whole area shown in Figure 3a) was approximately 2 µm, which closely resembled the findings for the modified H900 condition (482 °C / 4h) [9]. The total FCC phase fraction remained lower than previously observed by HE-XRD diffraction due to artifacts from strain-induced martensite transformation during metallographic preparation, driven by austenite stress relief during mechanical grinding and polishing, causing its transformation if it is not stable enough [24]. Also, as the austenite concentrates preferentially in the refined area (HAZ), a higher probability of non-indexed pixels occurs due to the overlap of the Kikuchi diffraction pattern from neighboring grains [37].
<figure><img src="image_5.png" /><figcaption>Figure 3 – EBSD (electron backscatter diffraction) results for direct-aged samples from as-built condition, at 621°C. The first column refers to IPF (inverse pole figure) maps for a) 1h and d) 8h, while the second column shows phase maps in the heat-affected zone (HAZ) for b) 1h) and e) 8h, and the melt pool (MP) in c) 1h and f) 8h. The building direction (BD) is also indicated, and the band contrast is used as background for all the maps.</figcaption></figure>

Subsequently, samples subjected to an 8-hour aging process yielded a slightly different morphology, as shown in Figure 3d: the resemblance to previous conditions faded as the distinctive 'fish-scale' morphology became more difficult to distinguish. The microstructure exhibited sections with larger grains enveloped by smaller ones, as shown in Figure 3e and f,
with a similar FCC fraction in both areas. However, significantly differing grain sizes were still apparent between MP and HAZ areas, as shown in Figure 4. This highlights a crucial aspect of AMed 15-5PH: even under heat treatments that generally yield an overaged condition with particles coarsening [38], the refined martensitic microstructure endured. Despite the finer microstructure, martensite grain size does not coarsen (Figure 4), similar to the behavior observed for a coarser martensite in the wrought counterpart [39]. The diffusion-controlled reverted austenite forms mostly at lath boundaries, block or packet boundaries due to enhanced solute diffusion at such interfaces, whereas these boundaries can also provide heterogeneous nucleation sites.

<figure><img src="image_6.png" /><figcaption>Figure 4 – Grain size distribution for martensite after aging at 621°C for 1h and 8h, in both the heat-affected zone (HAZ) and the melt pool (MP). Dashed lines represent the total fraction of grains from 1 to 10 µm. Average Grain Size for each area is shown in gray.</figcaption></figure>

## 3.2. Cu-rich precipitates (CRPs): nucleation and coarsening

CRP formation during heat treatment is a critical point of interest for understanding the mechanical performance of this material, since CRPs are responsible for the high mechanical resistance of the material, as well as acting as a preferential nucleation site for the austenite reversion [19]. Figure 5 shows EDS and ECCI images obtained for aged samples at 621°C for 1 and 8 hours, respectively. The dashed yellow lines in Figure 5c-d represent the Ni-enriched areas found in Figure 5a-b. The CRPs are visible as light-gray particles dispersed throughout the entire microstructure without the Ni-enrichment, in the images obtained by ECCI (Figure 5a-c), but due to their small size, are not identifiable by EDS in this magnification; the dark, spherical particles are cristobalite, identified in a previous work [9]. These oxides are Si-rich
(35.8 at. % Si, 61.2 at. %; Supplementary Table 1), with a content of other elements (Cr, Al, Mn, Nb, and C) smaller than 1.2 at. %. However, the quantity of precipitates visible in the Ni-rich areas seems to be lower. The same behavior can be observed after 8h of aging, as depicted in Figure 5b-d.

<figure><img src="image_7.png" /></figure>

Figure 5 – EDS maps for Ni and ECCI (electron channeling contrast imaging) images from aged samples at 621°C for 1h, a) and c), and 8h, b) and d). In c) and d), CRPs (Cu-rich nanoparticles) are visible as bright particles and oxide inclusions as dark particles. Dashed yellow lines in c) and d) show the Ni-rich areas in a) and b), respectively.

Figure 6a shows a STEM image that provides a magnified view of a similar region presented in Figure 5d, for the 8h aged sample. Figure 6a-b show the CRPs in the martensitic matrix, with a spherical and/or elongated shape. Another interesting characteristic, shown in Figure 6a-c, demonstrates an elongated Ni-rich area, a morphology already reported for reverted austenite according to the results obtained by Xiao et al. [40] and Sun et al. [41] for the wrought 15-5PH grade after aging at 500 °C. A line scan was also conducted in this region, shown in the yellow line in Figure 6a. It reveals the transition from Ni enrichment in the reverted austenite to Cu enrichment in certain particles (Figure 6d-e).

Despite not being the expected equilibrium phase, the BCC CRPs cluster nucleation in the BCC matrix requires a lower undercooling to trigger a burst of nucleation. It results from the lower activation energy associated with the creation of a coherent interface upon nucleation, due to the similarity of crystal structure and lattice parameter between phases when compared
to semi-coherent or incoherent interfaces. The coarsening of CRPs is initially slow, since the
total reduction of a low-energy coherent interface area drives it. A metastability cascade is
expected, since these phases exhibit gradually lower free energy as they approach equilibrium
and the previous intermediate phase decreases the nucleation energy for the next phase. The
interplay between volume free energy, which tends towards the phase in thermodynamic
equilibrium, and interface or coherent strain energy dictates the CRP evolution in size and
crystal structure throughout the aging treatment. The coherent elastic stresses at the interface
scale with the precipitate size, providing an additional driving force for the formation of
intermediate phases (BCC → 9R → 3R, with twinned and untwinned structures [11], [16]) and
resulting in the formation of semi-coherent interfaces. In the present work, even though it was
not possible to successfully index the crystal structure in the diffraction pattern (obtained by
applying the FFT in HRTEM images), there is evidence that the CRPs present an intermediate
twined 9R-structure after 4h aging, which then transforms into an untwined 3R precipitate after
8h of aging at 621 °C, as will be shown and discussed later in this section, in Figure 7 and
Figure 8. The formation of a semi-coherent interface, with dislocations accommodating the
lattice misfit between phases, reduces the free energy of the system. Further growth eventually
triggers the transformation into an FCC phase with incoherent interfaces due to the combination
of increased coherent elastic stresses and the presence of intermediate phases, reducing the
nucleation energy for the FCC phase.
<figure><img src="image_8.png" /><figcaption>Figure 6 – STEM image of an aged sample for 8h, showing a), a high density of precipitates (α' + CRPs) and areas with apparently no particles (γ). EDS analysis was performed as the b) Cu and c) Ni enrichment. d) EDS line scan (yellow line in a) of the aged sample for 8h, showing Ni-enrichment at the elongated grains near the center of the image (reverted austenite) and several peaks for CRPs present on almost the entire section, including Ni-rich areas.</figcaption></figure>
CRPs can also nucleate in the austenite. However, in this case they tend to nucleate directly as FCC precipitates to reduce the nucleation barrier. Since retained austenite is present in the as-built condition, the nucleation will be energetically favored as the lattice parameter is closer to the CRP in equilibrium. As described by Xiao et al. [40], since there is a coherence between the nucleated particle and the FCC phase, the driving force for particle coarsening is low, resulting in finer particles in the retained austenite. However, for Ni-rich reverted austenite, CRPs can be up to four times the size of those in retained austenite [40]. This larger size originates from their nucleation within the BCC martensite via the complex pathway previously described. Once formed, however, Ni enrichment and the transformation to reverted austenite result in slower coarsening kinetics. Nonetheless, both are still smaller than the precipitates in the BCC phase, which exhibit growth along the [110]Cu direction, due to the higher driving force and mobility for coarsening from incoherent interfaces [42], [43]. To summarize, there is evidence in literature that CRPs can be found in martensite, reverted Ni-rich austenite and retained austenite, with a decreasing size respectively due to interfacial energy, which is low, resulting in a reduction in the driving force for coarsening. As will be shown in the following paragraphs, we were able to identify the larger CRPs which were nucleated and coarsened in the martensitic matrix.

HRTEM and APT analysis were performed to investigate the areas near the CRPs and their evolution. Figure 7 summarizes the findings for the 4h aged sample, while Figure 8 represents the 8h-aging. HRTEM image (Figure 7a) depicting a CRP after 4 hours of aging exhibits a spherical shape with an intermediate size (near 20 nm in diameter). FFT were applied to selected areas shown in Figure 7a to analyze the CRPs atomic structure. It shows the presence of a herring-bone pattern, implying the formation of an intermediate twined 9R-structure preceding precipitate growth into the 3R structure, which also occurs for the AM-processed steel. These data are corroborated by several studies on copper precipitation on Fe-systems, showing a step-phase transformation [29], [30], [31]. Upon size comparison, those observed after 4 hours are nearly half the size of those identified after an 8-hour heat treatment. As mentioned before, this growth is expected to occur in a preferential direction, and higher aging temperatures provide more mobility to copper diffusion throughout the microstructure, thus increasing the CRP size. The precipitates contain approximately 80% atomic copper near their core (Figure 7e), a characteristic also observed in the wrought counterpart [19].

The findings after 8h of aging, shown in Figure 8a, depict a CRP with an oval shape and size (32 nm) typical of overaged conditions, but exhibiting a structure similar to that of CRPs discovered after 4 hours (Figure 7a). The larger axis is almost double the size of the smaller one, suggesting that the growth of precipitates for PBF-LB-processed materials occurs similarly for the wrought [11], [16], [40], [41] and DED-fabricated [44], [45], [46] processes. The horizontal lines show another characteristic, the fringe pattern within the precipitate known as the Moiré fringes. This morphology has already been observed for the wrought steel [41], and it is evidence of an untwined 3R precipitate [29], characteristic of another intermediate state between the BCC and FCC structure. The ellipsoidal shape found in this structure occurs due to a preferential growth parallel to the direction with the lowest interfacial energy [29]. The CRP shown in Figure 8a is also smaller than 40 nm in its longer dimension, a characteristic
described in this structure [29]. Even though it was impossible to completely identify the
structure after applying the FFT to the CRP image, the interplanar distances of 0.213 and 0.210
nm were found in different regions of the precipitate, as shown in Figure 8b. An average angle
of about 70° between the fringes was found by analyzing an area of the FFT pattern from the
bigger CRP, as shown in Figure 8c (yellow box). A plane spacing of 0.215 and 0.209 nm from
two different regions of the (111)Cu plane and a value of 68° were already reported in the
literature for a partially transformed 3R-structure [30] between FCC and FCT structures. As
described in the study, those two variants represent different microstructures inside the CRPs:
FCC (0.209 nm) and FCT (0.215 nm) copper crystals. Therefore, those findings suggest that
even after 8h at 621 °C, it is not enough for the CRPs to achieve the final and most stable FCC
structure (ε-Cu) entirely [29]. A fully martensitic microstructure can promote a more
homogeneous precipitation [47], since the BCC structure is on a super-saturated solid solution
condition, resulting in a higher volume fraction and size of the CRPs, which is expected for the
wrought counterpart in comparison to the AMed alloy.

<figure><img src="image_9.png" /></figure>

Figure 7 – Aged sample at 621°C for 4h. a) HRTEM image of a CRP, with the inverse FFT applied to the diffraction pattern, resulting in b) and c), showing two different morphologies within the precipitate. APT reconstructed maps on d) showing CRPs and Nb(C, N) particles as orange and green regions, respectively. The chemical composition of CRP and Nb(C, N) are obtained from 40 at.% Cu and 3 ionic% NbN⁺ concentration isosurfaces, in e) and f), respectively.
### 3.3.Phase transformation sequence during direct aging at 621 °C

The CRPs contain approximately 80% atomic copper near their core after 621°C/4 h aging (Figure 7d), a characteristic also observed in wrought steel [19], [41]. In this case, residual nitrogen enrichment is observed in these precipitates. The lower Cu content could be caused by a reduction in the nucleation barrier of these precipitates with a Cu content below the expected from equilibrium thermodynamics (Supplementary Figure 4). At this stage, they are already surrounded by Nb(C, N) precipitates (Figure 7d), which are absent in the as-built condition. Nb-rich precipitates are commonly observed in wrought and aged 15-5 PH [19], [41] and 17-4 PH [48] stainless steel grades. Adding niobium to the steel serves the purpose of acting as a carbon/nitrogen trap, which helps to prevent the formation of Cr-rich $M_{23}C_{6}$ carbides. In the sequence of precipitation, CRPs should nucleate before the Nb(C, N) precipitates. The diffusion coefficient [49] of Cu in α-Fe (5-6 x 10$^{-22}$ m$^2$/s) is higher than that of Nb (5.4 x 10$^{-23}$ m$^2$/s), which would kinetically favor the CRPs nucleation and growth. Due to the low Nb solubility in Cu [50] (Supplementary Figure 4), there is partitioning of Nb ahead of the advancing reaction interface during CRP growth, causing local Nb accumulation. In addition, the nitrogen within the CRPs should also be partitioned away from the precipitate upon growth to reach the equilibrium composition. The partitioning of nitrogen in the reaction front can further increase the driving force for local Nb accumulation due to their strong affinity, resulting in the formation of Nb(C, N) precipitates. Some Nb(C, N) precipitates are seen isolated from any possible Cu precipitate. Still, in this case, they should have originated from Cu precipitates which dissolved due to Ostwald ripening during the 4 h aging at 621°C. It must be noted that a Cr content lower than the matrix is also present in these carbonitrides, while the Cr content in the matrix remains homogeneous. These results suggest that the intermediate Cu-rich structure formation and Nb(N,C) precipitation occur during the initial stages of the 621°C heat treatment.
<figure><img src="image_10.png" /><figcaption>Figure 8 – Structural and chemical analysis performed for 621°C/8h aged condition. a) HRTEM image of a CRP showing the Moiré fringes in the microstructure. b) FFT patterns from b) red and c) yellow highlighted regions shown in a). APT data reconstruction showing d) the chemical segregation near the CRPs, Nb(C, N) precipitates, and Si-rich oxides, and e) Ni enrichment around the CRPs, as indicated by the concentration isosurfaces of 10 at.% Cu and 3.5 at.% Ni.</figcaption></figure>

Figure 9 depicts the APT data reconstruction for the same regions shown in Figure 8e, using different combinations of elemental iso-surfaces. Figure 9c shows the proxigram [51] for the regions of interest (ROI) indicated by black arrows in Figure 9b, while Figure 9e shows the proxigram for the ROI highlighted with a blue cylinder in Figure 9d. During aging for 8 h, Cr preferentially segregates to linear features resembling dislocations, reaching up to 33 at.% Cr in a single dislocation (as shown in Figure 9a-b, indicated by the black arrows). A similar observation was made by Couturier et al. [19], although at a lower applied temperature (505 °C), which triggers a spinodal decomposition due to the presence of a miscibility gap under these conditions. Considering the quasi-chemical model, a positive enthalpy of mixing (endothermic) indicates a tendency of atoms to separate into solute- and solvent-rich regions [52], such as Fe-rich and Cr-rich in α-BCC, as seen in Supplementary Figure 5. Solutes forming an endothermic solution tend to segregate to sites with a lower coordination number, such as lattice defects [53] to reduce the likelihood of having a solvent neighbor, such as Fe. Thus, the
free energy of the solution is reduced by equilibrium segregation of Cr to dislocations, due to a lower local coordination number, even if the employed composition and temperature combination is outside the miscibility gap. Cr is also found near Nb(C, N), reaching up to 25 at.% Nb (Figure 9c). Most of the Cr-rich regions correspond to Nb(C, N) precipitates, even those with a linear shape. Even though these phases were found near each other, no mixing between them is expected, since the equilibrium composition for Nb(C, N) exhibits up to 3 at.% Cr (Supplementary Figure 6). Carbonitrides also exhibit an increased Nb, N, and C contents and Fe depletion after aging from 4 to 8 h, compared to Figure 7f and Supplementary Figure 7, indicating an evolution of these carbonitrides towards the equilibrium composition.

Several studies [19], [40], [41], [54] have pointed out a distinctive Ni-enrichment near the ends of the rod-like precipitates, as shown in Figure 8d and Figure 9e-f. Ni enrichment is caused by solute ejection from the growing CRPs, and the resulting solute accumulation in front of the reaction interface, as indicated by the high Cu content near the Ni-rich regions in Supplementary Figure 8. The Ni-rich region sometimes coincides with the Nb-, C-, and N-rich region, whereas sometimes it only exhibits a local Ni enrichment near a CRP. Local Ni enrichment facilitates austenite nucleation due to an increased driving force. Thus, the martensite reversion mediated by the diffusion mechanism is very likely facilitated by the Ni-enrichment near the CRPs, creating nano-austenite grains in these regions.

Silicon-based oxide nanoparticles are distributed throughout the microstructure in the as-built condition of alloys produced by PBF-LB. In previous work [9], these oxides were identified as cristobalite. Here, the chemical composition in both the as built condition and after aging at 621°C for 8 h is shown in Supplementary Table 1. Following aging, a slight increase in Si, Cr, Al, Mn, Ni, and N is observed. However, in both cases they are mostly Si-O nanoparticles with some minor impurities, often related to elements with strong oxygen affinity.

Even though a Ni-Si-Mn segregation appears to be found in this sample (Figure 8d, near the Si-rich oxide), and which could be a precursor for the formation of G-phase, it presents a slow formation kinetics. As shown by Xiao et al. [40], this phase can be present after several hours of aging (above 100h/500°C). However, holding times longer than 8 h were not explored in this work to support this hypothesis and, therefore, this phase is not expected to be found in the conditions analyzed here.

In summary, a complex phase transformation sequence occurs during the aging of 15-5 PH stainless steel processed by PBF-LB. The precipitation of CRPs occurs via the formation of metastable phases with increased Cu content. Upon nucleation and growth of CRPs, the rejection of Nb, N, Ni, and Cr trigger the nucleation of reverted austenite and Nb(C, N) near the CRPs. With longer aging treatments, Cr segregates to dislocations, further facilitating kinetically and thermodynamically the local nucleation near the Nb(C, N).
<figure><img src="image_11.png" /><figcaption>Figure 9 – a-b) APT data reconstruction with different iso-surfaces for Cu, NbN⁺, Cr, and C, with (c) showing the proximity histogram from 20 at.% Cr concentration isosurfaces, indicating the composition near the Cr segregated regions highlighted in b) with the black arrows. Cr-rich dislocations reach up to 33 at.% Cr, with local Nb, C, and N enrichment. d-e) shows different iso-surfaces for Cu, NbN⁺, and Ni, near the CRPs. f) chemical composition of the region of interest highlighted with a blue cylinder in e). The region in f) corresponds to a region enriched in Ni, Nb, N, C, and Cr.</figcaption></figure>

### 3.4. PBF-LB processing vs. wrought counterparts: microstructural evolution

The wrought counterpart usually exhibits both negligible retained austenite content and a homogeneous prior austenite grain size distribution before aging. In contrast, the PBF-LB steel in the as-built condition exhibits a bimodal grain size distribution in the HAZ and MP (Figure 4), with a higher content of reversed austenite in the fine-grained region within the HAZ (Figure 3b). Upon aging, the content of austenite becomes more uniform within different regions in the PBF-LB steel (Figure 3e-f), but with a distribution of different grain sizes across the microstructure (Supplementary Figure 9). The presence of retained austenite and martensite
in the PBF-LB material creates a heterogeneity in terms of CRPs' size and structure, due to the local difference in chemical composition and crystal structure. In our case, we observe that martensite reversion is at least partly mediated by the growth of CRPs, meaning that no CRPs will be nucleated in the reverted austenite. Yet, the nucleation and growth sequence in both martensite and retained austenite are different, resulting in a more heterogeneous microstructure when compared to the wrought counterpart.

The Nb added to this steel grade acts as a carbon trap, to avoid formation of fragile Cr-rich $M_{23}C_{6}$ carbides, which would reduce both the pitting corrosion resistance and ductility. In conventional processing routes (forging + solubilizing/homogenization), the carbonitrides (Nb(C, N)) typically form before or during the solution annealing treatment preceding aging [19], [34], [55]. On the other hand, the inherent rapid solidification rates in the PBF-LB process can significantly suppress this precipitation, resulting in the supersaturation of the matrix with carbonitride-forming elements (Nb, C, N). As discussed previously, the complex local thermal history generates a heterogeneous microstructure, with different sizes of CRPs. Due to the segregation near these precipitates, they can act as a preferential site for carbonitrides nucleation, promoting a finer and more uniform distribution of Nb(C, N) during dedicated post-build heat treatments (Figure 7d and Figure 9a). Therefore, the potential for achieving a finer average size and higher number density in the PBF-LB material offers an opportunity to control these precipitates in contrast to the larger and dispersed carbonitrides found in conventionally processed wrought steel.

To fully leverage the PBF-LB microstructure (i.e., the austenite grain size heterogeneity for TRIP activation and the potential for controlled fine precipitation), tailored thermal processing routes are essential. Figure 10 shows the hardness Vickers variation due to the influence of direct aging at 621 °C for up to 8 h. Other conditions of PBF-LB and wrought material are also shown for comparison. The aged PBF-LB steel achieves hardness levels comparable to, or even exceeding, those of the aged wrought material despite its drastically different initial microstructure. The H1150 condition evaluated in this study presents only a small difference (~20-30 HV) to wrought H900 steel, and a considerable gap (near 100 HV) to wrought H1150. This is a remarkable feature, since the H900 condition is the standard aging for obtaining maximum mechanical strength/hardness, while the H1150 condition is recommended for improved ductility and fracture toughness.

Despite the co-precipitation of carbonitrides around CRPs, the softening behavior during aging is similar, indicating a similar CRP coarsening behavior when comparing PBF-LB and wrought 15-5 PH. However, the hardness in the overaged state is still superior for the PBF-LB sample due to the maintenance of a finer martensite inherited from the additive process. Supplementary Figure 10 shows the grain size and phase fraction EBSD maps of wrought 15-5PH stainless steel aged at 621 °C for 1h and 8h. When comparing these results with Figure 3 and Figure 4, it is possible to note larger martensite laths (average grain size for wrought is at least 4 times bigger than MP grain in PBF-LB) can be seen after the same aging routes applied to the PBF-LB steel were used; however, only a small fraction of martensite reverted into austenite even after 8h, different from what happened to AMed 15-5PH. Wrought 15-5PH is generally subjected to homogenization routes before aging, leading to a more homogenous microstructure, with larger grains and no retained austenite.
<figure><img src="image_12.png" /><figcaption>Figure 10 – Hardness variation of 15-5PH after aging at 621 °C up to 8 hours in black. * Hardness levels for PBF-LB as-built and H900 conditions [9], and wrought H900 (482 °C/1h) and H1150 (621 °C/4h) conditions [56] are shown for comparison. Data for wrought H1150 shown in green is adapted from [34], while experimental data is shown in purple.</figcaption></figure>

The combined impact of the martensitic/austenitic microstructure, the bimodal grain size distribution, the presence of metastable austenitic phases, and the distributions of both CRPs and carbonitrides, can play as synergistic strengthening mechanisms and improve mechanical properties, such as hardness. The CRP heterogeneity itself may contribute to more stable hardening during deformation, mitigating early localization, while the refined carbonitrides can provide additional strengthening and microstructural stabilization. Those findings corroborate the results obtained for the H900 condition [9], showing the necessity of modifying heat treatment routes specifically for additively manufactured steels.

4. Conclusion

The present discussion focused on the evolution of the far-from-equilibrium microstructure after direct aging of 15-5 PH stainless steel at 621 °C. Future investigations should evaluate how the phase transformations reported here influence the fracture toughness, ductility and mechanical strength of this grade. The martensite reversion into a nickel-rich austenite and the formation of CRPs and Nb(C, N) are intricately interconnected processes.
Furthermore, identifying notable microstructural features such as Moiré fringes in the CRPs and the distinct presence of Ni-enrichment in precipitates in the same way they appear for wrought steel broadens the understanding of these complex processes in far-from-equilibrium microstructures such as those found for AM-processed alloys. After a comprehensive analysis of the microstructural evolution as well as various post-processing heat treatments, valuable insights were found:

* The reverted austenite during aging at 621°C in the AM-fabricated steel is Ni-rich, as predicted by thermodynamic and diffusional simulations. It nucleates and coarsens near the CRPs due to Ni rejection from these precipitates, similar to what occurs in wrought 15-5PH steel.

* There is evidence that the CRPs do not yet have an FCC structure but rather a metastable structure, most likely an untwined 3R-type, after 8h of direct aging at 621 °C.

* CRP formation leads to the expulsion of Nb, creating Nb-rich regions, acting as preferential site for nucleation of Nb(C, N) carbonitrides. Cr-rich regions are formed at dislocations between 4 and 8 h aging at 621°C, and are located close to the Nb(C, N).

* Even after an aging heat treatment that promotes a lower hardness (H1150), the PBF-LB processed steel presents hardness levels comparable to the H900 (peak hardness) condition in the wrought material due to the fine microstructure inherited from PBF-LB, mainly martensite grain size, which is stable up to 8 h at 621°C.

## Acknowledgments

This research used the facilities of the Brazilian Nanotechnology National Laboratory (LNNano), part of the Brazilian Centre for Research in Energy and Materials (CNPEM), a private non-profit organization under the supervision of the Brazilian Ministry for Science, Technology, and Innovations (MCTI). The SEM-FIB staff is acknowledged for the assistance during the experiments (proposal number 20210563). This study was financed, in part, by the São Paulo Research Foundation (FAPESP), Brazil (Process numbers 2019/06679-1 and 2019/00691-0). LFKU is thankful to CNPq (National Council for Scientific and Technological Development) for the financial support under grant 156279/2019-3. J.A. Avila is a fellow from CNPq under grant 306960/2021-4 and acknowledges his former fellowship to the Serra Hunter program. JPO acknowledges the funding of CENIMAT/i3N by national funds through the FCT-Fundação para a Ciência e a Tecnologia, I.P., within the scope of LA/P/0037/2020, UIDP/50025/2020, and UIDB/50025/2020 of the Associate Laboratory Institute of Nanostructures, Nanomodelling and Nanofabrication - i3N. The present study was developed in the scope of the Project "Agenda ILLIANCE" [C644919832-00000035 | Project nº 46], financed by PRR - Plano de Recuperação e Resiliência under the Next Generation EU from the European Union. We acknowledge DESY (Hamburg, Germany), a member of the Helmholtz Association HGF, for the provision of experimental facilities. Parts of this research were carried out at PETRA III. Data was collected using P07 beamline operated by Helmholtz-Zentrum Hereon.
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SUPPLEMENTARY FIGURES

<figure><img src="image_13.png" /><figcaption>Supplementary Figure 1 – Austenite (FCC) volume fraction and lattice parameter evolution obtained during in situ HE-XRD upon heating and cooling stages (Q̇ = 10 °C/min). HE-XRD: high energy – X-ray diffraction.</figcaption></figure>
<figure><img src="image_14.png" /><figcaption>Supplementary Figure 2 – M<sub>s</sub> temperature as a function of chromium and nickel composition. M<sub>s</sub>: martensite start temperature, or the temperature where martensitic transformation starts upon cooling.</figcaption></figure>
<figure><img src="image_15.png" /><figcaption>Supplementary Figure 3 – Grain size distribution for the EBSD maps shown in Figure 3, with the yellow dashed boxes indicating the selected areas for analysis of the Melt Pool and the Heat Affected Zone.</figcaption></figure>
<figure><img src="image_16.png" /><figcaption>Supplementary Figure 4 – Chemical composition of a CRP (FCC_A1#2) in the equilibrium condition, near the aging temperature. CRP: Cu-rich nanoparticles.</figcaption></figure>
<figure><img src="image_17.png" /><figcaption>Supplementary Figure 5 – Enthalpy of mixing in the BCC phase as function of the Cr content.</figcaption></figure>
<figure><img src="image_18.png" /><figcaption>Supplementary Figure 6 – Chemical composition of Nb(C,N) in the equilibrium (FCC_A1#3), near the aging temperature.</figcaption></figure>
<figure><img src="image_19.png" /><figcaption>Supplementary Figure 7 - Proximity histogram from a 3 ionic% NbN⁺ concentration isosurface after 8 h ageing at 621 °C. Note the presence of Cr and Ni enrichment near the Nb-, C- and N-rich regions. Cu-rich regions are observed near the (Nb,Cr)(C,N) particles, as indicated by the Cu content at -1 to -2 nm from the concentration isosurface.</figcaption></figure>

<figure><img src="image_20.png" /><figcaption>Supplementary Figure 8 - Proximity histogram from a 3at.% Ni concentration isosurface after ageing at 621°C for 8 h. The Ni-rich regions are formed at the vicinity of CRP. CRP: Cu-rich nanoparticles.</figcaption></figure>
<figure><img src="image_21.png" /><figcaption>Supplementary Figure 9 – Grain size distribution for austenite (FCC) after aging at 621°C for 1h and 8h, in both heat-affected zone (HAZ) and the melt pool (MP). Dashed lines represent the total fraction of grains from 0.1 to 1 µm. Note the grain growth during ageing in both regions, although a grain size difference remains, with finer grains at the HAZ.</figcaption></figure>

<figure><img src="image_22.png" /><figcaption>Supplementary Figure 10 – Phase fraction and grain size distribution obtained from EBSD data for wrought 15-5 PH stainless steel after aging at 621°C for 1h and 8h. The band contrast is used as background for all the maps.</figcaption></figure>
<table><caption>Supplementary Table 1 – Composition of the oxide nanoparticles in the as-built condition and after ageing at 621°C-8 h.</caption><thead><tr><td>Concentration (at.%)</td><td>As-built</td><td>621°C-8 h</td></tr></thead><tbody><tr><td>O</td><td>61.263</td><td>55.669</td></tr><tr><td>Si</td><td>35.810</td><td>39.123</td></tr><tr><td>Cr</td><td>1.164</td><td>2.470</td></tr><tr><td>Al</td><td>0.374</td><td>0.919</td></tr><tr><td>Ni</td><td>-</td><td>0.490</td></tr><tr><td>Mn</td><td>0.521</td><td>0.420</td></tr><tr><td>Nb</td><td>0.391</td><td>0.376</td></tr><tr><td>C</td><td>0.477</td><td>0.281</td></tr><tr><td>N</td><td>-</td><td>0.199</td></tr></tbody></table>