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Article

Microstructural Evolution of a Pre-Alloyed Duplex Stainless Steel 2205 with Boron Addition Prepared by Powder Metallurgy

by
Pedro Morita Terceiro
and
Juliano Soyama
*
School of Mechanical Engineering, University of Campinas, Rua Mendeleyev, 200, Campinas 13083-860, Brazil
*
Author to whom correspondence should be addressed.
Powders 2025, 4(3), 24; https://doi.org/10.3390/powders4030024
Submission received: 11 June 2025 / Revised: 16 July 2025 / Accepted: 30 July 2025 / Published: 22 August 2025
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Abstract

The addition of hard particles such as borides to a ductile stainless steel matrix can be very efficient for improving mechanical properties. Powder metallurgy represents a suitable route for developing these material modifications, combining high reproducibility and cost-effectiveness. The present research investigated the effect of sintering time on an atomized, pre-alloyed 2205 stainless steel with 2.5 wt.% boron, using two different powder size distributions: fine (<45 µm) and coarse (250–500 µm). Cold uniaxial compaction was conducted using a cylindrical closed die. Sintering was carried out at 1200 °C with a dwell time of 2 and 4 h in argon atmosphere. Microstructural investigation showed that borides were formed in the powder’s atomization step and presented a small size with different morphologies. The borides significantly improved the hardness and compression strength. Compared to the reference 2205 stainless steel, specimens prepared with the fine powder size distribution achieved a twofold enhancement in yield stress, while hardness increased by 26%.

1. Introduction

In duplex stainless steels, the coexistence of austenite and ferrite phases provides an interesting combination of high strength, toughness, and corrosion resistance [1,2]. Consequently, these stainless steels are widely used in the oil, gas, and petrochemical industries [1,2,3]. Nevertheless, depending on the application, hardness and wear resistance might be an issue. A possible solution to improve the mechanical properties is the introduction of hard particles, such as borides, into their microstructure. Boride formation in steel occurs exclusively during solidification through precipitation of primary and/or eutectic borides [4]. Additionally, borides have high thermal stability and very limited solubility in ferrite and austenite phases. Therefore, they are more interesting than carbides or other solution and precipitation-dependent particles. Carbides, for instance, require heat treatment for controlled precipitation and can be redissolved, depending on the application temperature.
Powder metallurgy (PM) is an attractive processing route for fabricating components of complex geometries at a low cost because it can significantly reduce the manufacturing steps and waste of raw materials compared to conventional techniques. Moreover, PM can be applied for processing a large variety of alloys, including steels, titanium, aluminum, copper, and hard materials, such as tool steel or hard metal with high mechanical properties, repeatability, and tight geometrical tolerances [5,6]. Previous reports investigated the modification of stainless steels with boron via PM, focusing mainly on adding elemental boron and mixing/blending them in austenitic [6,7,8,9] or ferritic steels [10]. Boron additions to steel create a eutectic point, which induces liquid phase formation during sintering, resulting in high densification [6]. Another technique to obtain borides is boro-sintering, a process in which sintering occurs simultaneously with a boronizing treatment [11,12]. However, boronizing is a thermo-chemical surface treatment that introduces boron only at the surface [13]. The hardening in boro-sintering is thus superficial. Consequently, the surface hardness is significantly increased to values around 1400 HV, but the bulk of the material remains unchanged [11,12].
A prior research work reported processing by press and sinter of a pre-alloyed 2205 duplex stainless steel with 2.5 wt.% of boron [14]. In that study, the effect of sintering temperature was investigated, and the best results were achieved at 1200 °C for 2 h, obtaining densifications above 85%. The sintering temperature of 1200 °C was shown to be the highest possible because of liquid phase formation. However, changing the sintering time and powder size might lead to higher densification. Therefore, the main objective of this study was to evaluate the effect of the sintering time in the pre-alloyed 2205 stainless steel with 2.5 wt.% boron using two different particle size ranges, namely <45 μm and 200–500 μm, in terms of mechanical properties and microstructural evolution.

2. Materials and Methods

The pre-alloyed metallic powders were produced using a homemade closed-couple gas atomizer under a nitrogen atmosphere. The raw materials used were commercial super duplex stainless steel grade 2507 (Metalinox, São Paulo, Brazil) Fe–B (Núcleo Ligas, Guarulhos, Brazil) to perform the boron addition, and other materials such as Fe–Mo, Fe, and Cr (Núcleo Ligas, Guarulhos, Brazil) with purity > 99.5% to adjust the compositions. Firstly, all the raw materials were melted using an induction furnace (Inductotherm, Indaiatuba, Brazil) at a temperature of 1600–1650 °C until a homogeneous molten metal was produced inside a tundish. Finally, the liquid metals were atomized with nitrogen at a pressure of 0.40 MPa. The powders were sieved using a vibratory sieve shaker (Retsch, Haan, Germany). The powders were separated into two granulometries: < 45 µm and 250–500 µm, hereafter named Fine and coarse powders, respectively. The chemical composition of the pre-alloyed powder was measured using ICP-OES (inductively coupled plasma optical emission spectrometry) (Agilent, Santa Clara, CA, USA), and chromium and nickel equivalents were calculated based on the formulas reported in the literature [15]. Table 1 shows the average composition and chromium/nickel equivalent, which indicates only the presence of ferrite and austenite phases, based on Schaeffler’s diagram [15].
The specimens were prepared by cold uniaxial pressing of pre-alloyed powders using a cylindrical closed die with a diameter of 8.2 mm. Stearic acid (Sigma-Aldrich, Darmstadt, Germany) was employed as a solid lubricant mixed with the metal powder in a fraction of 3 wt.%. Approximately 3 g of powder was used per specimen, resulting in cylinders about 12 mm in height. An average compaction force of 68 kN was applied, accounting for 1.3 GPa of compaction stress. The green density was evaluated by measuring the mass and dimensions of the specimens. The theoretical density was calculated using the rule of mixtures based on the elemental composition and their fractions. Sintering was conducted at 1200 °C with a sintering dwell time of 2 and 4 h in a tube furnace with an argon atmosphere (Grion, São Paulo, Brazil). The sintering substrate was an alumina crucible. The applied heating rate was 5 °C/min with furnace cooling after sintering. Porosity measurements were carried out by Light Optical Microscopy using a Leica DM500 (Wetzlar, Germany) and ImageJ (version 1.54a). The image analysis consisted of applying a treatment to distinguish between pores and sintered particles. Five images were taken from the cross-section to calculate the average porosity and the standard deviation of each sintered specimen. In addition, ImageJ software was used to evaluate the pre-alloyed powder size and distribution.
Differential scanning calorimetry (DSC) was carried out in a Netzsch STA 409 (Selb, Germany) with atomized pre-alloyed powders. A mass of about 100 mg was employed. A heating and cooling rate of 10 °C/min was applied with a maximum temperature of 1300 °C. Structural characterization was conducted by X-ray diffraction (XRD) using a PANalytical X’Pert Pro MPD (Malvern, England) and Cu-Kα radiation. A 2θ range of 30–90° was employed with a step of 0.026°, 40 kV of voltage, and 30 mA current. The metallographic preparation consisted of grinding and polishing a transversal cross-section of sintered specimens and atomized pre-alloyed powders. A scanning electron microscope from Zeiss (EVO MA15) (Jena, Germany) was used for phase analysis and the observation of the powder’s morphology. The Vickers hardness tests (HV) followed the ASTM B933-20, which describes microhardness tests for materials obtained via powder metallurgy. The measurements were performed using a Future-Tech FV-800 (Kawasaki, Japan) with a load of 0.3 kgf and a dwell time of 15 s [16]. Compression tests were carried out to evaluate the mechanical properties of sintered specimens. The compression tests followed the ASTM E 9-89a standard, which recommends cylindrical bodies with an L/D ratio of 2 [17]. Thus, specimens of 8.2 mm in diameter and 16.4 mm in height were prepared for the compression tests. A load limit of 270 kN with a constant crosshead speed of 0.5 mm/min was applied using an EMIC DL-30000 universal test machine (Norwood, MA, USA).

3. Results and Discussion

3.1. Powder Characterization

Figure 1a,b show the spherical morphology of the metallic powders, which are typical of the gas atomization process. However, the fine powder (Figure 1a) also contained particles with irregular morphology, while satellites were identified in the coarse powder (Figure 1b). Both powders exhibited a rough and irregular surface originating from the atomization’s fast cooling rate. Overall, particle size distribution (Figure 1c,d) followed the targeted separation through sieving. The mean particle size of the fine powder was 56 µm (Figure 1c), while that of the coarse powder was 322 µm. The particle size distribution of the fine powder showed values above the standard deviation (>80 µm). This was a consequence of irregularly shaped particles that presented a larger area compared to the spherical particles. However, their frequency was low (approximately 10%). The particle size distribution of the coarse powder indicated a small fraction of particles below the average size.
The thermograms presented in Figure 2 show that both particle sizes presented two consecutive pronounced endothermic peaks in the heating cycle at temperatures of 1231 °C and 1252 °C (fine powder) and 1227 °C and 1245 °C (coarse powder). These peaks were associated with the formation of a liquid phase and/or dissolution of eutectic borides. The cooling cycles (Figure 2b) attested that the transformations were reversible, once exothermic peaks were observed at similar temperatures. The exothermic peaks of the fine powder were observed at 1233 °C and 1183 °C, and for the coarse powder at 1217 °C and 1193 °C. These temperatures were above the eutectic temperature based on the Fe-B binary system [18], indicating an additional effect of the alloying elements that shifted the liquidus temperature. Based on the isopleth of the Fe-Cr-Ni system [19], and the fact that when sintered at 1250 °C, as reported in a prior study [14], the specimens lost their shape due to an excess of liquid phase, it was most likely that these peaks corresponded to the following transformations: δ → δ + L and δ + L → L. Although the liquidus temperature was below the reported value for the Fe-Cr-Ni system, it could be a result of boron decreasing the liquid phase formation temperatures [7]. Reported Thermo-Calc simulations in the literature using similar compositions indicated a eutectic transformation around the same temperature range [4,20].

3.2. Sintered Density

The theoretical density calculated for the boron-modified steel was 7.93 g/m3. According to Figure 3, which summarizes the densifications achieved with the different powder particle sizes, the green density (based on geometrical measurements) reached between 60% and 70% of the theoretical density. The fine powder showed a slightly lower green density than the coarse powder because smaller particles are usually harder to compress due to the higher interparticle friction. After sintering at 1200 °C with a dwell time of 4 h, the coarse powder improved only 2% in density when compared to sintering for 2 h. Nevertheless, the fine powder increased about 9%, reaching a density of >80% of the theoretical density. This behavior may be attributed to the fact that smaller powders have a larger surface area and can activate sintering mechanisms more effectively than larger powders, such as volumetric diffusion. Moreover, they can be more efficiently covered and bound by a liquid phase. As a result, the fine powder progressed more rapidly through the sintering stages compared to the coarse powder. This difference is visually evident in the microstructural images presented in Figure 3. In Figure 3b, representing the coarse powder sample, individual primary powder particles with clearly defined sintering necks remain visible, indicating that the liquid phase did not fully cover the particles. Conversely, the fine powder sample in Figure 3c exhibits no distinguishable primary particles and features smaller, more uniformly distributed pores, demonstrating a more advanced sintering stage. This enhanced behavior is attributed to the greater surface area and the more effective liquid-phase sintering mechanism enabled by the smaller particle size.

3.3. Microstructural Characterization

The SEM images of the sintered microstructures recorded with the secondary electron detector are shown in Figure 4. The fine powder presented a considerable decrease in porosity with an increase in sintering time (Figure 4a,b), especially in the radial direction, as depicted by the rectangles in Figure 4a. The large, crack-like pores in the radial direction were caused by the uniaxial compaction process that typically leads to compaction inhomogeneities and thus density gradients. The sintering evolution with the fine powder achieved a stage where the primary powder particles were indistinguishable and the porosity became isolated. This indicates that the third stage of sintering was reached [21] and that increasing the sintering time was very effective for densification. The improved densification could be explained by the formation of liquid phase during sintering. With longer sintering times, the liquid phase was able to fill the empty spaces and pores more efficiently. Moreover, most of the pores showed an irregular shape (Figure 4b, arrows), which could be an effect of the liquid phase solidification during cooling.
The coarse powder was not significantly affected by the longer sintering time. As shown in Figure 4c, with a sintering time of 2 h, the powder particles seemed to be at the first stage of sintering with only a small approximation between the particle centers (as highlighted by the markings in Figure 4c). In this case, the large empty spaces between the particles could not be filled by the liquid phase. The density increase observed in the coarse powder with a longer sintering time was a consequence of neck growth, as shown in Figure 4d. Nevertheless, powder particles with small sintering necks could still be observed (rectangle in Figure 4d).
X-Ray diffraction patterns of atomized powders and sintered specimens are shown in Figure 5. The powder diffraction peaks could be identified as α-ferrite, γ-austenite, and borides. The boride peak was small and only observable in the fine powder due to the higher counts that are naturally achieved when measuring fine powders by XRD. The diffraction patterns of sintered specimens showed additional peaks of two intermetallic phases, M7C3 carbide and the sigma phase. The source of carbon for carbide formation was the stearic acid used as a lubricant in the compaction step. During the initial heating of the sintering cycle, the organic lubricant was degraded, but some carbon likely remained in the sintered material. Moreover, with the slow cooling rate after sintering, the precipitation of the carbides and sigma phase was expected to occur. The precipitation of the sigma phase with slow cooling rates agrees with literature reports [9,22]. Although the sigma phase is typically associated with material embrittlement, its impact is less detrimental in this case due to the substantial volume fraction of borides present in the steel, as shown in Figure 6. While the presence of sigma phase and carbides in duplex steels may contribute to embrittlement, their combined effect (alongside borides) should benefit wear resistance. Carbide formation can be mitigated by using carbon-free lubricants, such as borax, or by optimizing the sintering atmosphere. Additionally, increasing the cooling rate can help prevent sigma phase precipitation by minimizing formation kinetics.
SEM images of the atomized fine powder and the sintered specimens recorded with the backscattered electron detector are shown in Figure 6. In the atomized powder, Figure 6a, small borides could be observed embedded in a fine duplex α + γ matrix. The borides were found with a faceted and elongated morphology, as well as with a circular shape, indicated by the circles. These observations are in accordance with the XRD results. Therefore, despite the fast cooling rates of the atomization process, borides were already present in the pre-alloyed powder.
The sintered specimen prepared with the coarse and fine powders (Figure 6b,c) showed the presence of irregular pores, borides, sigma phase, and carbides, in addition to the α + γ matrix, which is consistent with the indexed XRD peaks. The borides were faceted and elongated, with some particles showing a circular shape, as indicated by the rectangles. They were larger than observed in the atomized powder, probably due to growth occurring in the presence of the liquid phase. The sigma phase was indicated by an arrow and contained heavier elements that appeared with a lighter contrast in the SEM image. Carbides could be observed as roundish or angular black features, indicated by the triangles. The morphology of carbides and the sigma phase was similar to the literature reports [23,24,25]. Pores were also shown as black features; however, they were irregular and contained mostly sharp corners or were completely crack-like.
Figure 7 presents the EDS elemental maps for each sintered specimen. The sigma phase and the carbides are clearly enriched in Cr and Mo. The borides are predominantly composed of Mo and tend to follow the boundaries of the Cr-rich phases. These borides are finely dispersed and occur even in small amounts within those regions. No effect of powder particle size could be related to the boride sizes.

3.4. Mechanical Properties

Figure 8 shows the mechanical properties of the sintered specimens. Due to the higher densification achieved with the fine powder, these specimens were selected for the compression tests. Figure 8a indicates a considerable increase in compression strength with 4 h of sintering. This improvement took place due to the higher densification (lower porosity) and the elimination of most of the pores in the radial direction. Additionally, with a sintering time of 2 h, the elastic region showed some nonlinearity oscillation due to the deformation accommodation from closure of the radial pores, which was not observed in the case of 4 h. This supports the fact that the enhancement in the mechanical properties was induced by the denser and more homogeneous microstructure. The yield and fracture stress were 528 MPa and 1154 MPa, respectively. In comparison to the literature values of a pore-free reference 2205 stainless steel [26], an almost twofold increase in the yield stress was observed.
Figure 8b shows the evolution of microhardness as a function of sintering parameters and the powder particle sizes. The fine powder showed a significant increase in hardness with longer sintering times. With 2 h of sintering, the hardness was in the same range as a pore-free 2205, which was likely due to the effect of porosity that increased the dispersion and masked the hardening effect of the borides. However, with 4 h of sintering and a porosity below 20%, the hardness reached 340 ± 19 HV, which is approximately 26% higher than the reference 2205 steel. Previous studies reported hardness values ranging from 280 to 310 HV [27,28]. This variation can be attributed to the use of elemental boron blended through conventional mixing methods. In contrast, the present study employed atomized pre-alloyed powders containing boron, resulting in the formation of borides. The use of atomized powders promotes a highly homogeneous distribution of borides, effectively minimizing the formation of segregated boron-deficient regions that could compromise mechanical properties, which is a common issue observed in traditionally mixed compositions. The coarse powder showed a trend of lower hardness for longer sintering times, but the dispersion was high, and it could be the effect of porosity. It is noteworthy to mention that the sintering time > 2 h is only beneficial for fine powders. For coarse powders, alternative strategies are needed.

4. Conclusions

The present study investigated the influence of sintering time on atomized pre-alloyed 2205 stainless steel containing 2.5 wt.% boron, using two distinct particle size ranges: <45 µm (fine) and 250–500 µm (coarse). Differential scanning calorimetry (DSC) analyses indicated that liquid phase formation occurs around 1200 °C, following the transformation sequence δ → δ + L and δ + L → L. This liquid phase proved crucial for enhancing densification during sintering. The microstructural analysis showed that specimens produced with fine powder reached the final sintering stage, whereas those made with coarse powder exhibited some growth in sintering necks but limited densification, even at extended sintering times. Post-sintering analysis showed the precipitation of carbides and sigma phase, attributed to slow furnace cooling. The mechanical properties of the sintered fine powder were significantly improved for longer sintering times as a consequence of a denser and more homogeneous microstructure. In the optimal case, a yield stress of 528 MPa and a fracture stress of 1154 MPa were recorded. Additionally, hardness increased by approximately 26% compared to standard 2205 stainless steel, surpassing values reported in previous studies. This enhancement is linked to the uniform boride distribution achieved through the use of pre-alloyed atomized powders.

Author Contributions

Conceptualization, P.M.T. and J.S.; methodology, P.M.T.; software, P.M.T.; validation, P.M.T. and J.S.; formal analysis, P.M.T. and J.S.; investigation, P.M.T.; resources, J.S.; data curation, P.M.T.; writing—original draft preparation, P.M.T.; writing—review and editing, J.S.; visualization, P.M.T.; supervision, J.S.; project administration, J.S.; funding acquisition, J.S. All authors have read and agreed to the published version of the manuscript.

Funding

This research was funded by CAPES—Coordenação de Aperfeiçoamento de Pessoal de Nível Superior, finance code 001, and FAEPEX—Fundo de apoio ao ensino, pesquisa e extensão da Unicamp, grant number 3348/23.

Institutional Review Board Statement

Not applicable.

Informed Consent Statement

Not applicable.

Data Availability Statement

Data will be made available upon request from the corresponding author.

Acknowledgments

The authors wish to acknowledge the laboratory technicians at the School of Mechanical Engineering and Agronomic Engineering, University of Campinas, for their help in carrying out the experiments. The authors are also grateful for the support provided by the Laboratory of Physical Metallurgy, Rubens Caram, and the team of the Multiuser Laboratory for Materials Characterization.

Conflicts of Interest

The authors declare no conflicts of interest.

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Powders 04 00024 g001
Figure 1. Morphology and particle size distribution: (a,c) fine powder, (b,d) coarse powder.
Figure 1. Morphology and particle size distribution: (a,c) fine powder, (b,d) coarse powder.
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Figure 2. Thermograms obtained with different particle sizes. (a) Heating cycle and (b) cooling cycle.
Figure 2. Thermograms obtained with different particle sizes. (a) Heating cycle and (b) cooling cycle.
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Figure 3. (a). Fraction of theoretical density achieved with different sintering parameters and particle sizes. (b) SEM-SE images coarse powder sintering necks and (c) SEM-SE images fine powder sintering necks.
Figure 3. (a). Fraction of theoretical density achieved with different sintering parameters and particle sizes. (b) SEM-SE images coarse powder sintering necks and (c) SEM-SE images fine powder sintering necks.
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Figure 4. SEM images of the sintered microstructures. (a) Fine powder, 1200 °C/2 h; (b) fine powder, 1200 °C/4 h; (c) coarse powder, 1200 °C/2 h; and (d) coarse powder, 1200 °C/4 h. The numbers indicate crack-like pores (1), irregular pores (2), and sintering necks (3).
Figure 4. SEM images of the sintered microstructures. (a) Fine powder, 1200 °C/2 h; (b) fine powder, 1200 °C/4 h; (c) coarse powder, 1200 °C/2 h; and (d) coarse powder, 1200 °C/4 h. The numbers indicate crack-like pores (1), irregular pores (2), and sintering necks (3).
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Figure 5. XRD of the pre-alloyed atomized powders and sintered specimens at 1200 °C/4 h. (a) Full 2θ range and (b) zoomed 2θ angles showing the boride and carbide peaks.
Figure 5. XRD of the pre-alloyed atomized powders and sintered specimens at 1200 °C/4 h. (a) Full 2θ range and (b) zoomed 2θ angles showing the boride and carbide peaks.
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Figure 6. SEM images of the atomized fine powder (a) and the sintered specimens at 1200 °C/4 h; (b) coarse powder; and (c) fine powder. The rectangles and circles indicate borides, the arrows sigma phase, and the triangles carbides.
Figure 6. SEM images of the atomized fine powder (a) and the sintered specimens at 1200 °C/4 h; (b) coarse powder; and (c) fine powder. The rectangles and circles indicate borides, the arrows sigma phase, and the triangles carbides.
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Figure 7. EDS Elemental maps for sintering at 1200 °C/4 h. (a) Fine powder and (b) coarse powder.
Figure 7. EDS Elemental maps for sintering at 1200 °C/4 h. (a) Fine powder and (b) coarse powder.
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Figure 8. Mechanical properties of the sintered specimens. (a) Stress–strain curves measured with the sintered fine powders at different sintering times, and (b) microhardness and porosity of all investigated conditions.
Figure 8. Mechanical properties of the sintered specimens. (a) Stress–strain curves measured with the sintered fine powders at different sintering times, and (b) microhardness and porosity of all investigated conditions.
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Table 1. Average chemical composition of pre-alloyed powders and Cr/Ni equivalent.
Table 1. Average chemical composition of pre-alloyed powders and Cr/Ni equivalent.
Average Composition (wt.%)
CrMnFeNiMoB
23.8 ± 0.11.2 ± 0.162.0 ± 0.14.9 ± 0.15.6 ± 0.12.5 ± 0.1
Cr equivalent (wt.%)Ni equivalent (wt.%)
29.4 ± 0.17.7 ± 0.1
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MDPI and ACS Style

Morita Terceiro, P.; Soyama, J. Microstructural Evolution of a Pre-Alloyed Duplex Stainless Steel 2205 with Boron Addition Prepared by Powder Metallurgy. Powders 2025, 4, 24. https://doi.org/10.3390/powders4030024

AMA Style

Morita Terceiro P, Soyama J. Microstructural Evolution of a Pre-Alloyed Duplex Stainless Steel 2205 with Boron Addition Prepared by Powder Metallurgy. Powders. 2025; 4(3):24. https://doi.org/10.3390/powders4030024

Chicago/Turabian Style

Morita Terceiro, Pedro, and Juliano Soyama. 2025. "Microstructural Evolution of a Pre-Alloyed Duplex Stainless Steel 2205 with Boron Addition Prepared by Powder Metallurgy" Powders 4, no. 3: 24. https://doi.org/10.3390/powders4030024

APA Style

Morita Terceiro, P., & Soyama, J. (2025). Microstructural Evolution of a Pre-Alloyed Duplex Stainless Steel 2205 with Boron Addition Prepared by Powder Metallurgy. Powders, 4(3), 24. https://doi.org/10.3390/powders4030024

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