PSI - Issue 69

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Procedia Structural Integrity 69 (2025) 1

13th European Symposium on Martensitic Transformation 2024 (ESOMAT 2024) Preface Ausonio Tuissi a , Carmine Maletta b a CNR-ICMATE, National Research Council of Italy – Institute of Condensed Matter Chemistry and Technologies for Energy, Via Previati 1E, 23900, Lecco (LC) Italy. b University of Calabria, Dept of Mechanical, Energy and Management Engineering - Via P. Bucci 44C, 87036 - Rende (CS) - Italy The European Symposium on Martensitic Transformations (ESOMAT) has, since its first edition in 1989, established itself as one of the most significant international events for scientists studying martensitic transformation-based phenomena. This includes research on shape memory alloys, magnetic shape memory alloys, and martensitic transformations in steels and ceramics. The 13th edition of the symposium, ESOMAT 2024 , was held in Lecco, Italy, from August 26 to 30, 2024. The event was co organized by the National Research Council of Italy (CNR-ICMATE) and the University of Calabria, with the valued support of the Italian Association of Metallurgy (AIM). The symposium brought together more than 200 participants, including scientists and students from approximately 30 countries. Over the course of five days, the program featured around 170 scientific contributions, including five plenary lectures, 22 oral sessions, and one poster session. The welcoming and collaborative atmosphere of the conference provided an ideal setting for the exchange of ideas and the strengthening of international scientific networks. We would like to express our sincere gratitude to the ESOMAT International Advisory Committee, the National Organizing Committee, AIM for organizational support, the Politecnico di Milano, the sponsors, invited speakers, contributing authors, and session chairs. The success of ESOMAT 2024 was made possible only through their commitment and generous support. This special issue of Procedia Structural Integrity includes a selection of 16 peer-reviewed papers presented during ESOMAT 2024. Together, they reflect the richness and diversity of current research in the field and the high scientific level of the symposium. We hope that this collection will serve as a valuable reference for researchers and practitioners working in this exciting and evolving area of materials science.

Ausonio Tuissi & Carmine Maletta ESOMAT 2024 Chairs and Special Issue Guest Editors

2452-3216 © 2025 The Authors. Published by ELSEVIER B.V. This is an open access article under the CC BY-NC-ND license (https://creativecommons.org/licenses/by-nc-nd/4.0) Peer-review under responsibility of Guest Editors 10.1016/j.prostr.2025.07.001

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Procedia Structural Integrity 69 (2025) 69–75

13th European Symposium on Martensitic Transformation 2024 (ESOMAT 2024) Carbon Redistribution in a Martensitic Medium Mn Steel During Heating to Intercritical Region: An In-Situ Synchrotron XRD Study R. Surki Aliabad a, * , S. Sadeghpour a , P. Karjalainen a , J.Kömi a , H. Singh b , V. Javaheri a a Materials and Mechanical Engineering, Centre for Advanced Steels Research, University of Oulu, Oulu 90570, Finland b Nano and Molecular Systems Research Unit, University of Oulu, Oulu 90570, Finland Abstract This study investigates the microstructural evolution and carbon redistribution during the heating stage prior to intercritical annealing treatment (IAT) in a medium manganese steel (MMnS) with the nominal composition of Fe-0.40C-6Mn-2Al-1Si-0.05Nb (wt.%). The material was characterized using high-energy X-ray diffraction, scanning electron microscopy, and transmission electron microscopy. The initial microstructure primarily consisted of tempered martensite containing nano-sized plate-like η carbides and 7 vol.% retained austenite (RA) with thicknesses of 10–20 nm and 300 nm in average, respectively. During heating, carbon partitioning caused an increase in carbon content within the RA up to 530 °C, rising from 0.4 wt.% to 1 wt.%. η-carbides initially coarsened and subsequently transformed into cementite with an average diameter of ~ 20 nm. Above 530 °C, RA began to decompose, resulting in the formation of a pearlite-type microstructure. Concurrently, the carbon content in the remaining RA decreased, facilitating further growth of cementite formed in the earlier stages. The microstructure at the onset of IAT at 640 °C consisted of tempered martensite with nano-sized spherical cementite, 9 vol.% RA with >1 wt.% carbon and a small fraction of pearlite-type decomposed RA. The study highlights the complex interplay between carbon redistribution, carbide formation, and RA stability during the heating stage of MMnS and emphasizes the importance of accurately characterizing the initial microstructure to tailor the properties of these advanced high-strength steels. © 2025 The Authors. Published by ELSEVIER B.V. This is an open access article under the CC BY-NC-ND license (https://creativecommons.org/licenses/by-nc-nd/4.0) Peer-review under responsibility of Guest Editors

Keywords: Medium Mn Steel; Synchrotron XRD, Carbon Partitioning, Auto-Tempered Martensite

* Corresponding author. Tel.: +358-505649693. E-mail address: Roohallah.surkialiabad@oulu.fi

2452-3216 © 2025 The Authors. Published by ELSEVIER B.V. This is an open access article under the CC BY-NC-ND license (https://creativecommons.org/licenses/by-nc-nd/4.0) Peer-review under responsibility of Guest Editors 10.1016/j.prostr.2025.07.010

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1. Introduction Medium manganese steels (MMnS), containing 4-12 wt% manganese, have emerged as promising candidates for the third generation of advanced high-strength steels (AHSS). These alloys offer an excellent combination of high strength, large ductility, and cost-effectiveness, making them suitable for diverse industrial applications [1]. Among these, MMnS with 3-10 wt% Mn and 0.05-0.3 wt% C tend to exhibit almost fully martensitic microstructures with a small fraction of retained austenite (RA) after hot rolling, owing to their remarkable hardenability [2]. By carefully controlling thermomechanical processing, these steels can be widely tailored to develop a multi-phase microstructure, ensuring a desired RA distribution and stability. This flexibility facilitates the customization of mechanical properties across a broad spectrum, as highlighted by the recent review article [3]. However, tailoring MMnS microstructures to achieve desired mechanical properties poses significant challenges. A comprehensive understanding of the thermodynamic and kinetic factors influencing phase transformations, particularly during heating stage and intercritical annealing treatment (IAT), is critical. For example, manganese segregation to lattice defects and carbide precipitation significantly affect the thermodynamic driving force for austenite nucleation [4]. The competition between carbide formation and martensite-to-austenite reversion adds further complexity to the microstructural evolution [5]. However, this competition is not always observed. For instance, after hot rolling, MMnS may undergo auto-tempering during slow cooling resulted in (partially) tempered martensitic microstructure [2]. Similarly, tempering prior to IAT [6] results in a tempered martensitic structure, where carbides, particularly cementite, serve as the primary nucleation sites for new austenite formation during IAT. Recent multi-stage processing of MMnS has leveraged Mn-rich cementite particles to improve mechanical properties by serving as nucleation sites for austenite reversion [7]. The resulting austenite is finer and more Mn enriched than that found in conventionally treated alloys [8]. Moreover, Enomoto and Hayashi found that austenite is more likely to nucleate on cementite at prior austenite grain boundaries and martensite packet boundaries than at inter lath boundaries [9]. Additionally, larger cementite particles would be more effective nucleation sites, as noted by Zhang et al. [10]. An additional complexity in this intricate microstructure is the redistribution of carbon during carbides formation. Zhao et. al [11] observed that tempering at 500°C for 60 minutes applied to an intercritically annealed microstructure led to reduction in RA volume fraction from 18.5% to 12% due to RA decomposition, accompanied by an increase in carbon content from 0.1% to 0.8% by weight in remaining RA. Recent advancements have incorporated in-situ high-energy X-ray diffraction (HEXRD) techniques to gain deeper insights into microstructural changes during heating and IAT of MMnS. For instance, Muller et al. investigated the evolution of microstructure during heating and holding of IAT, starting from a martensitic microstructure with 2 vol.% RA. Their study revealed that new film-like austenite forms rather than the preservation of existing RA [12]. Similarly, Hu et al. [13] observed RA decomposition and cementite precipitation during heating, followed by an increase in austenite fraction during isothermal holding. Mehrabi et al. [14] monitored austenite fraction and lattice parameters throughout the annealing process under varying temperature and time conditions. Despite extensive research on phase transformations during the isothermal holding stage of IAT and the resulting final microstructures [15–18], limited attention has been given to the heating stage prior to reaching the intercritical annealing temperature. This critical stage involves possible complex microstructural changes, including carbon redistribution, carbide formation, and the decomposition or preservation of existing austenite, which significantly influence the initial conditions for subsequent phase transformations. Therefore, this work focuses on examining the microstructural evolution and carbon redistribution occurring during the heating stage prior to IAT. This is crucial for accurately characterizing the initial microstructure at the onset of IAT, including carbide formation, carbon partitioning between RA and the surrounding matrix, and the preservation of existing austenite as nucleation sites. 2. Materials and method The steel material used in this study, with the nominal composition of Fe-0.40C-1Si-6Mn-2Al-0.05Nb (wt.%), was produced in a vacuum induction furnace. After casting, the material was subjected to austenitization at 1200 °C for 2 hours, followed by hot rolling with a 90% thickness reduction to achieve a final thickness of 4 mm. The finish rolling was conducted at 950 °C, and the material was subsequently cooled in ambient air.

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In-situ HEXRD was employed to explore the microstructural evolution of the hot-rolled (HR) sample upon heating. The sample was heated at a rate of 1°C/s to 700°C, slightly above the Ac1 temperature (640 °C). This HEXRD analysis was performed at the Brockhouse High Energy Wiggler Beamline, located at the Canadian Light Source (CLS). The HR sample, with dimensions of 1 × 1 × 20 mm³, was housed in a quartz capillary while it was heated in a high temperature capillary furnace. A 70 keV monochromatic X-ray beam (wavelength: 0.1779 Å) was employed, with the sample-to-detector distance set at 1116 mm. Calibration was done at room temperature using a nickel standard. The resulting two-dimensional diffraction data were collected using an area detector, Varex XRD 4343CT detector, and radially integrated and analyzed using the Rietveld refinement method in GSAS-II software. Field emission scanning electron microscope (FESEM) couples with electron backscatter diffraction (EBSD) detector was utilized to analyze the microstructure. Samples were prepared through several steps, including sectioning, grinding, polishing with a 1 μm diamond paste, and chemical etching in a 2% Nital solution (2% nitric acid in ethanol) for about 10 seconds at room temperature. The analyses were performed using an InLens and EBSD detectors at 5 kV and 15 kV accelerating voltage, with a working distance of 4 mm and 15 mm, respectively. For detailed microstructural characterization, Scanning Transmission Electron Microscopy (STEM) was employed using a JEOL JEM-2200FS Energy Filtered Scanning Transmission Electron Microscope (EFTEM/STEM) at 200 kV. Thin lamella specimens measuring approximately 8 × 10 μm were prepared using a focused ion beam-field emission scanning electron microscope (FIB-FESEM) equipped with a Ga ion source. Thermodynamic and kinetic simulation was carried out using Thermo-Calc 2024a software. The phase equilibrium was modeled utilizing the TCFE13 thermodynamic database. 3. Results and Discussion The microstructure of the hot-rolled and subsequently air-cooled steel was analyzed using TEM and EBSD methods, as illustrated in Fig. 1. The microstructure consisted of mainly a lath martensitic matrix with plate-like carbides heterogeneously precipitated within the laths. These carbides, measuring 10–20 nm in thickness and less than 100 nm in length, are clearly visible obtained on a wide block of tempered martensite (Fig. 1a). In addition to the carbides, HEXRD results at room temperature detected approximately 7 vol.% RA. EBSD phase map (Fig. 1b) revealed RA with an average thickness of 300 nm.

Fig. 1. (a) TEM BF image of plate-like carbides within the martensite matrix; (b) EBSD phase and band contrast maps highlighting RA morphology and size; (c) SAED patterns of carbides and martensite; (d) key diagram for diffraction pattern, beam // [111] ! // [110] " ; (e) TEM DF image of -carbides with zone axis 2 -2 0 taken from the area shown in (a). In this study, carbides precipitated in the martensite matrix (Fig. 1a) were analyzed using TEM microscopy. The diffraction pattern revealed a structure consistent with the η-Fe₂C carbide, as shown in Fig. 1c, with the [111] ! //

72 R. Surki Aliabad et al. / Procedia Structural Integrity 69 (2025) 69–75 [110] " orientation relationship with the matrix martensite phase [19]. Fig. 1e displays a dark field (DF) image using the 2 2' 0 η reflection from the selected area electron diffraction (SAED) pattern shown in Fig. 1d. The heterogeneous distribution of carbides observed in Fig. 1a is attributed to martensite formation occurring at varying temperatures. Auto-tempering of martensite was active at temperatures just below the martensitic start (Ms) temperature, gradually slowing as the temperature decreased further [2]. The presence of these plate-like carbides in the studied temperature range has also been documented for other MMnS [6,20,21].

Fig. 2. (a) Carbide peaks evolution tracked by HEXRD at different temperatures; (b) The volume fraction and the lattice parameters of RA across different heating temperatures.

Fig. 2a illustrates HEXRD patterns of the sample at different temperatures. Before reaching 375°C, the evolution of carbide peaks intensity only shows a slight increase in the background, adaptable with 1 2 1 and 2 2 0 reflections of η-Fe 2 C carbide. Up to 375°C, only a faint bump appears in the background, likely due to an increase in the tempering level. Comparing the observed XRD pattern at 375°C with Fe 2 C, Fe 2.5 C, Fe 5 C 2 and Fe 3 C patterns reveals that this bump could be more related to η-Fe 2 C carbide that eventually gives way to cementite at higher temperatures as cementite peaks become distinctly visible in the patterns above 530°C. Fig. 2b displays both the volume fraction and the lattice parameter of RA across different heating temperatures. Additionally, it shows the various phenomena occurring in the different regions. The HEXRD results reveal that the initial RA volume fraction increases from 7.5% in the as-rolled specimen to 10.5% in the specimen heated to 375°C. This increase is likely related to the movement of the γ/α′ interface caused by carbon partitioning [22]. The carbon content of RA, calculated using Equation 1 [23], initially increases from 0.4 wt.% in the as-annealed specimen to 1 wt.% between ~ 375°C to 530°C. Beyond this temperature range, RA volume fraction exhibits only a slight fluctuation while its lattice parameter changed linearly with temperature; showing the end of C partitioning at ~ 375°C. On the other hand, Fig. 2a indicates that between 375°C and 530°C, η carbide transitions to cementite. However, the kinetics of this transition appear to be very slow, possibly due to the nucleation stage of cementite. a (Å) = 3.572 + 0.033%C + 0.0012%Mn + 0.00157%Si + 0.0056%Al (1) Where a is austenite lattice parameter in Å, and the elements composition in mass weight.

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Fig. 3. SEM micrographs of the heated samples step-quenched from different temperatures (denoted in Fig. 2b by black dotes); (a and d) 250°C; (b and e) 450°C; (c and f) 680°C.

The microstructural evolution of the studied steel at three different heating temperatures is shown in Fig. 3. The general microstructure, including tempered martensite, remains largely consistent with the initial structure across all temperatures. At 250°C, some degree of carbide coarsening becomes evident. As the temperature increases to 450°C, significant coarsening of carbide particles is observed. At higher temperature, 680°C, carbide transitions dominate, resulting in the formation of numerous nano-sized spherical carbides distributed both within and on the martensitic blocks. Figures 3d and 3e illustrate this coarsening behavior at 250°C and 450°C, respectively. HEXRD data reveals clear carbon enrichment in retained austenite (RA) between 165°C and 375°C. Following this enrichment, the estimated carbon content in the BCC matrix phase, attributed to cementite precipitation, can be estimated by mass balance calculation as shown in Equation 2: C γ .f γ +C α .f α =0.4 →1×0.1+C α ×0.9=0.4 → C α =0.33wt.% (2) Where C and f are carbon content in mass weight and phase fraction of each phase, respectively. Using Thermo-Calc software, the maximum amount of cementite that could precipitate, based on the nominal composition of the material with a carbon content of 0.33 wt.%, was calculated to be 4.8 vol% at 540°C. Although SEM observations suggest the potential formation of significant amounts of cementite, HEXRD failed to detect carbide peaks with high intensity. One key factor contributing to this is the size of the carbides. When the particle size falls below a critical threshold, the corresponding diffraction peaks become broadened, which prevents their detection with the current experimental setup [24]. Between 530°C and 640°C, the volume fraction of retained austenite (RA) decreased to 9.0%. At 680°C, SEM micrographs (Fig. 3f) revealed the formation of a pearlite-type microstructure, resulting from the decomposition of RA into colonies of large cementite lamellae. This transformation is consistent with previous studies. For instance,

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Zhao et al. [11] observed the transformation of RA into pearlite during tempering in a 0.1C-5Mn steel, while Ribamar et al. [25] detected a pearlite-like constituent after tempering 0.99C-1.0Mn-1.5Cr-1.82Si steel at 500°C. Similarly, Sadeghpour et al. [7] identified the same decomposition in a 4Mn–0.31C–2Ni–0.5Al–0.2Mo steel. Moreover, in this region, the carbon content of RA decreases, a trend also previously reported by Zhao et al. [11]. This decline could be attributed to the depletion of carbon in the martensitic matrix, which results from several carbon consuming phenomena occurring in the earlier stages. At this point, for further growth of cementite, the only available carbon source at high temperatures could be RA. This applies to cementite located near the RA or farther away, as the BCC matrix could act as a pathway for carbon diffusion [26]. 4. Conclusions This study investigated the microstructural evolution and carbon redistribution in MMnS with a nominal composition of Fe-0.40C-1Si-6Mn-2Al-0.05Nb (wt.%) during the heating stage prior to IAT. The initial microstructure of the hot-rolled and air-cooled steel consisted of tempered martensite containing nano-sized plate-like η-carbides and 7 vol.% RA. Through HEXRD and SEM observations, the overlapping phenomena over various temperature ranges were deconvoluted as follows: 1. Low temperatures (165°C - 375°C): Carbon enrichment in RA from 0.4 to 1 wt.% and coarsening of carbides occurred simultaneously. Despite the high fraction of carbides in the microstructure, they were undetectable by HEXRD, likely due to their nano-sized nature. 2. Mid-range temperatures (375°C - 530°C): Cementite nucleation occurred. 3. Higher temperatures (530°C - 640°C): Decomposition of RA resulted in a pearlite-type microstructure and at the same time spherical carbides nucleated from transitional carbides in the matrix grew. This process reduced the carbon content of the remaining RA. 4. Above 640°C: An increase in the austenite volume fraction was observed, identifying this temperature as the Ac1 point for the studied steel. Acknowledgements The authors would like to thank Jane and Aatos Erkko (J&AE) Foundation and Tiina and Antti Herlin (TAH) Foundation for their financial supports on Advanced Steels for Green Planet (AS4G) project. References [1] Y.K. Lee, J. Han, Current opinion in medium manganese steel, Materials Science and Technology (United Kingdom) 31 (2015) 843– 856. https://doi.org/10.1179/1743284714Y.0000000722. [2] J.H. Nam, S.H. Yu, Y.K. Lee, Effect of auto-tempering on the cold roll-ability of medium-Mn steel, Materials Science and Technology (United Kingdom) 35 (2019) 2069–2075. https://doi.org/10.1080/02670836.2018.1547474. [3] D. Kumar, I. Sen, T.K. Bandyopadhyay, A Systematic Review of Medium-Mn Steels with an Assessment of Fatigue Behavior, Steel Res Int 95 (2024). https://doi.org/10.1002/srin.202300375. [4] D. Raabe, S. Sandlöbes, J. Millán, D. Ponge, H. Assadi, M. Herbig, P.P. Choi, Segregation engineering enables nanoscale martensite to austenite phase transformation at grain boundaries: A pathway to ductile martensite, Acta Mater 61 (2013) 6132–6152. https://doi.org/10.1016/j.actamat.2013.06.055. [5] A. Kwiatkowski da Silva, G. Inden, A. Kumar, D. Ponge, B. Gault, D. Raabe, Competition between formation of carbides and reversed austenite during tempering of a medium-manganese steel studied by thermodynamic-kinetic simulations and atom probe tomography, Acta Mater 147 (2018) 165–175. https://doi.org/10.1016/j.actamat.2018.01.022. [6] S. Lee, S.H. Kang, J.H. Nam, S.M. Lee, J.B. Seol, Y.K. Lee, Effect of Tempering on the Microstructure and Tensile Properties of a Martensitic Medium-Mn Lightweight Steel, Metall Mater Trans A Phys Metall Mater Sci 50 (2019) 2655–2664. https://doi.org/10.1007/s11661-019-05190-4. [7] S. Sadeghpour, M.C. Somani, J. Kömi, L.P. Karjalainen, A new combinatorial processing route to achieve an ultrafine-grained, multiphase microstructure in a medium Mn steel, Journal of Materials Research and Technology 15 (2021) 3426–3446. https://doi.org/10.1016/j.jmrt.2021.09.152. [8] Q. Ye, H. Dong, Q. Guo, Y. Yu, L. Qiao, Y. Yan, Tailoring the austenite characteristics via dual nanoparticles to synergistically optimize the strength-ductility in cold rolled medium Mn steel, J Mater Sci Technol 169 (2024) 158–171. https://doi.org/10.1016/j.jmst.2023.07.002. [9] M. Enomoto, K. Hayashi, Simulation of Austenite Formation During Continuous Heating from Low Carbon Martensite with Poly dispersed Cementite, Metall Mater Trans A Phys Metall Mater Sci 51 (2020) 618–630. https://doi.org/10.1007/s11661-019-05569-3.

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Procedia Structural Integrity 69 (2025) 80–88

13th European Symposium on Martensitic Transformation 2024 (ESOMAT 2024) Characterization of Vibration Damping in Nitinol Samples Manufactured through Selective Laser Melting Diego Scaccabarozzi a , Carlo Alberto Biffi b,* , Jacopo Fiocchi b , Abdelrahman Mohamed Ragab M. Ahmed a , Marco Giovanni Corti a , Ausonio Tuissi b , Bortolino Saggin a,c a Politecnico di Milano, Polo Territoriale di Lecco, Lecco b CNR, ICMATE, Lecco c Università degli Studi di Padova, Padova Abstract This research explores the potential of 3d-printed Nitinol samples to mitigate transmitted vibration levels, aiming at the development of interface dampers for space applications. The study investigates the correlation between the measured damping behaviour of Nitinol samples, measuring the loss factor with the half-power bandwidth approach in free-constrained conditions to avoid unwanted and unpredictaqw02ble dissipation given by the constraints. Among the influencing parameters, sample printing directions and post-additive heat treatments have been investigated. Results showed that by carefully tailoring the heat treatment process, it is possible to modulate the damping properties of Nitinol parts. This opens possibilities for engineers to design and manufacture lightweight aerospace components with optimised structural properties, minimising the overall mass and enhancing vibration characteristics. Moreover, this study paves the way for further research to optimise the Selective Laser Melting (SLM) process and post-treatment procedures for Nitinol alloys. © 2025 The Authors. Published by ELSEVIER B.V. This is an open access article under the CC BY-NC-ND license (https://creativecommons.org/licenses/by-nc-nd/4.0) Peer-review under responsibility of Guest Editors

Keywords: NiTiNol, damping characterisation, interface dampers, Additive Manufacturing.

* Corresponding author. Tel.: E-mail address: diego.scaccabarozzi@polimi.it

2452-3216 © 2025 The Authors. Published by ELSEVIER B.V. This is an open access article under the CC BY-NC-ND license (https://creativecommons.org/licenses/by-nc-nd/4.0) Peer-review under responsibility of Guest Editors 10.1016/j.prostr.2025.07.012

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1. Introduction Nowadays, the demand for advanced materials in aerospace and space applications has significantly increased. For instance, there is an urgent need for lightweight structures with superior mechanical properties and vibration-damping capabilities, especially in such applications. In dynamic environments, vibration control is critical to ensure structural integrity and functionality. Shape memory alloys (SMAs), such as Nitinol, are promising for vibration-damping applications, as their exceptional superelasticity grants considerable damping capacity. These properties allow Nitinol to be the ideal choice for interface dampers that reduce transmitted vibrations without adding significant weight. Additive manufacturing (AM), particularly laser powder bed fusion (LPBF), has transformed the production of Nitinol components, enabling the fabrication of complex geometries with precise control over the microstructure and relative density of the material (Sinha et al., 2023). However, challenges such as residual stresses, porosity, and microstructural inconsistencies remain significant and can negatively impact the damping properties of LPBFed parts (Yan et al., 2024). Overcoming these challenges is essential to fully exploit Nitinol’s potential in vibration-damping systems. Several studies have demonstrated that the damping properties of metallic alloys can be enhanced through LPBF design and post-processing treatments. For example, Fiocchi et al. (2020) showed that trabecular structures produced in Ti6Al4V alloys via LPBF enhanced energy dissipation without compromising mechanical strength. Similarly, Colombo et al. (2020) demonstrated that stress-relieving thermal treatments improved the damping behaviour of LPBFed AlSi10Mg alloys. These findings emphasise the role of both structural design and thermal processing in optimising damping capacity. The latter characteristic is of paramount importance for space applications and payloads (Saggin et al., 2022; Scaccabarozzi et al., 2024a, 2024b). The LPBF production of Nitinol components has been widely explored in recent years; nevertheless, this task has been shown to face considerable challenges, as careful control of the material composition, stress state and microstructure is needed for obtaining satisfactory mechanical properties (Biffi et al., 2024). Nevertheless, the possibility of joining the design freedom granted by LPBF and the superelastic / shape-memory behaviour of Nitinol-based alloys is extremely interesting in several fields, as it may allow the production of architected structures with integrated complex functionalities (Mehrpouya et al., 2024). This research investigates the influence of post-processing heat treatments on the damping properties of LPBFed Nitinol. By understanding this relationship, this study aims to add more knowledge towards optimising the treatment processes for lightweight, high-performance interface dampers. The findings offer significant potential for vibration control systems in aerospace applications, which could allow for the design of structures that minimise mass while enhancing vibration resistance. 2. Materials and methods Sample Preparation Rectangular test samples, sized 33 × 4 × 2 mm, are fabricated by LPBF, using a Renishaw AM400 system, which employs a 400 W pulsed-wave laser and a reduced build volume. The feedstock material was gas-atomised, Ni-rich Nitinol powder with a nominal composition of Ni54Ti46 (wt.%). This composition was chosen for its excellent superelasticity and suitability for additive manufacturing. The processing parameters, summarised in Table 1, were optimised to achieve a relative density exceeding 99.5%. A schematic of the principal process parameters is depicted in Figure 1.

Table 1. Printing processing settings.

Value 150 75 Meander

Setting/Parameter

Power, P

Exposure time, t exp Scanning strategy

Argon 30 50

Atmosphere

Layer thickness Hatch distance, d h

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120 65 30 ° < 20

Point distance, d p Laser spot size

Platform temperature

Oxygen level

Fig. 1. Schematic of the principal process parameters in the LPBF process: a) exposure time cycle, b) laser spot and related hatch and point distances.

Samples were oriented at a 45° angle relative to the building platform during fabrication. This orientation improved residual stress distribution and maximised the build volume efficiency. After the printing, the produced samples were subjected to two ageing treatments aimed at improving their functional properties. The treatments were carried out at 500 °C for 5 (T500-5’) and 10 minutes (T500-10’) in a muffle furnace, followed by water quenching. Martensitic Transformation Analysis The martensitic transformation was analysed using differential scanning calorimetry (DSC, model Q25 from TA Instruments, USA). Measurements were conducted on small samples (15-20 mg) with a heating/cooling rate of 10 °C/min, in a temperature range from -80°C to 100°C. Measurement method description The range for the acoustic excitation to vibrate the sample was numerically determined through a modal analysis in free conditions, conducted using commercial software. The material properties used in the simulation were density (6450 kg/m³), modulus of elasticity (44.1 GPa), and Poisson's ratio (0.30). These parameters were derived from testing of the manufactured samples by mechanical testing. Fig.2 shows the first mode of vibration and the theoretical position of the nodal lines. The location of these lines is critical for conducting the vibration tests since additional fictitious damping can be introduced in the measured frequency response function if the component is not vibrating on the nodal lines. Thus, the theoretical position of nodal lines was used as a reference to mount the samples on the testing setup.

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Fig. 2. Modal analysis – first vibration mode in free constraint condition.

The numerical simulation revealed a first bending natural frequency at 5112 Hz, thus providing an excitation range between 5 kHz and 6 kHz. The measurement setup and chain are illustrated in Fig. 3. A Laser Doppler Vibrometer (LDV – Polytec OFV 505) was used for a noncontact vibration measurement. Additionally, a BSWA MA 201 microphone with a sensitivity of 51.79 mV/Pa measured the sound pressure generated from the loudspeaker. Furthermore, to monitor ambient temperature and ensure consistency during tests, a thermocouple measuring the ambient temperature was added. All signals were acquired using an NI 9234 data acquisition board. The frequency response function was computed by averaging the data from at least ten frequency sweeps. The samples were mounted on elastic cords positioned at the nodal lines.

Fig. 3. The experimental setup. (Left) testing setup scheme, (right) view of the sample over the suspension and LDV spot.

Each sample underwent at least ten subsequent sweep sine cycles, each lasting 50 seconds. To further confirm the accuracy of the results and consider the effect of the positioning on the elastic cords, the tests were repeated three

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times for every sample, replacing each time the samples. This methodology provided a robust dataset to assess the reproducibility of the damping measurement. The loss factor ( ) was calculated using the half-power bandwidth method. This method determines the damping capacity based on the resonance curve (Lee et al., 2008). Physically, the loss factor reflects how quickly the system loses energy during an oscillation. Having a higher loss factor means the system dissipates more energy (stronger damping), which leads to a broader resonance peak. Furthermore, a low loss factor indicates less energy loss and a sharper resonance. This makes mandatory in evaluating the damping characteristics of materials or mechanical structures under dynamic loading. The loss factor can be obtained by: = ! ! "! " ! # (1) where: • # and $ are the frequencies at 3 dB below the resonance peak amplitude. • % is the resonance frequency at the peak amplitude. Previous frequencies shall be obtained from the measured Frequency Response Function (FRF), derived from the acquired stimulus and response of the tested specimens. 3. Results Differential scanning calorimetry Fig. 4 shows the evolution of the martensitic transformation in the explored thermal treatment conditions and in the as-built condition. The as-built sample exhibits a single-stage transformation, both upon heating and cooling, whereas the heat-treated ones present a double stage upon cooling (from austenite to R phase and from R phase to martensite). It may be observed that both heat treatments caused a slight shift of transformations to higher temperatures (A f is about 18 °C in as-built condition and 28 °C in T500-5’ condition). In details, it is well known that LPBF process promotes residual stresses, due to the fast cooling during the rapid solidification; such residual stresses could alter the martensitic transformation, as it can occurs in plastic deformation, where the phase transformation can be tuned at different temperatures and also completely suppressed under high levels of deformation degrees [Miller, 2001], [Biffi, 2019]. The measured transformation temperatures of the Nitinol samples in the investigated conditions are listed in Table 2. DSC scans show that the predominant phase present at room temperature is austenite in all the studied conditions, because the peak of the phase transformation lies below room temperature. Therefore, under limited loading conditions, it can be affirmed that further damping evaluation would be carried out for the austenitic phase. Conversely, the increase of treatment time did not induce a considerable peak shift, but rather produced a slight sharpening of said peaks.

Fig. 4. DSC scan of the Nitinol samples in as-built and heat-treated conditions.

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Table 2: Transformation temperatures detected from the DSC scans

Af [°C] Ms [°C] Mf [°C] [° ] Rf [°C]

As [°C]

As built

-27.5

15.9

3.7

-61.9

-

-

After (T500-5’)

-0.8

20.0

-31.3

-58.5

13.2

-17.6

After (T500-10’)

3.5

18.3

-27.9

-53.3

9.6

-0.6

Damping Properties In Fig. 5 acquired vibrometer waveform and related FRF are shown.

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Fig. 5. (Top) Time domain signal acquired from the LDV – sample #3 (bottom) FRF amplitude, phase, and coherence.

Table 3 summarises the measured natural frequencies, loss factors, and their variations across the samples, which reflect the reproducibility of measurements. The measured repeatability from a single test replication, i.e. comprising 10 sweeps, showed 1σ worst-case variability of 2%, whereas averaging on the three test repetitions aiming to assess the reproducibility of the measured data, as shown in Table 3, the repeatability of the measured loss factors was found to vary between 1% and 10%. The measurements were performed at 25.8°C ± 0.6°C. Fig. 6 highlights the changes in loss factor before and after heat treatments. All samples in heat-treated condition exhibited loss factors, which are higher than the ones obtained in as-built condition by a factor of 2 or 3.5, a result that reflects a considerably improved damping capacity. This is consistent with the material microstructure before ageing, where the absence of optimised phase distribution limits energy dissipation, depending on the phase exactly present at room temperature (see DSC scans in Figure 3). Moreover, it can also be mentioned that the residual stresses in the as-built condition may reduce the damping capacity of the Nitinol samples.

Table 3. Measured natural frequency and loss factor for the three samples before and after heat treatments.

As built

After (T500-5)

After (T500-10’)

!

!

!

η

η

η

Average [Hz]

1σ [%]

Average [10 -3 ]

1σ [%]

Average [Hz]

Average [10 -3 ]

Average [Hz]

1σ [%]

Average [10 -3 ]

1σ [%]

1σ [%]

1σ [%]

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Sample 1 5355.9 0.03

2.56

3.6 5302.1 0.13

5.56

11

5311.2 0.12

5.30

9.7

Sample 2 5624.3 0.02

2.44

5.8 5381.9 0.10

6.71

7.2

5395.8 0.05

6.07

0.77

Sample 3 5525.8 0.01

2.61

7.0 5291.8 0.11

7.07

3.2

5288.1 0.05

7.35

9.3

0 0.001 0.002 0.003 0.004 0.005 0.006 0.007 0.008 0.009

Loss factor (ζ)

1

2

3

Samples

Before Treatment

A8er (T500-5’)

A8er (T500-10’)

Fig. 6. Trend of loss factor vs heat treatment. 1σ bands are visible.

The first heat treatment (T500-5’) significantly enhanced the damping capacity across all samples. The loss factor increased to values ranging from 0.6% to 0.7%. This increase indicates that thermal treatment improves internal friction and energy dissipation mechanisms. Concurrently, the natural frequencies decreased, which aligns with the expected microstructural softening reported in the literature. Following the second heat treatment (T500-10’), the loss factors stabilised with minimal further changes compared to the (T500-5’) condition, suggesting that the microstructural modifications introduced by the second treatment were incremental. This consolidates the major role of the initial heat treatment. Moreover, the natural frequencies remained relatively constant, which indicates no substantial alterations in the material stiffness. 4. Conclusion The study demonstrates the effect of heat treatments to optimise the damping characteristics of LPBFed Nitonol components. The results showed that the heat treatment effectively enhances the damping performance up to three times with respect to the as-built condition. These findings are pivotal for developing lightweight, vibration-resistant structures in aerospace engineering, suggesting avenues for further research to refine heat treatment techniques for maximum performance consistency and efficiency. In fact, additional improvement of the damping capacity is expected in the case that AMed structures would undergo larger deformation, as expected in the vibrational environment of space payloads and structures. References [Biffi, 2019] C.A. Biffi, A. Tuissi, Laser shape setting of superelastic NiTi wire: effects of laser beam power and axial pre-load, 2019 Smart Materials and Structures. https://doi.org/10.1088/1361-665X/ab1e86.

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[Biffi, 2024]Biffi, C.A., Lemke, J.N., Bassani, P., Fiocchi, J., Bregoli, C., Tuissi, A,, 2024. Characterization and postprocessing of additively manufactured shape memory alloys. In “Additive Manufacturing of Shape Memory Materials - Techniques, Characterization, Modeling, and Applications”, Elsevier. [Colombo, 2020] Colombo, C., Biffi, C.A., Fiocchi, J., Scaccabarozzi, D., Saggin, B., Tuissi, A., Vergani, L.M., 2020. Modulating the damping capacity of SLMed AlSi10Mg trough stress-relieving thermal treatments. Theoretical and Applied Fracture Mechanics 107, 102537. https://doi.org/10.1016/j.tafmec.2020.102537 [Fiocchi, 2002] Fiocchi, J., Biffi, C.A., Scaccabarozzi, D., Saggin, B., Tuissi, A., 2020. Enhancement of the Damping Behavior of Ti 6 Al 4 V Alloy through the Use of Trabecular Structure Produced by Selective Laser Melting. Adv Eng Mater 22. https://doi.org/10.1002/adem.201900722 [Lee, 2008] Lee, D.-G., Lee, S., Lee, Y., 2008. Effect of precipitates on damping capacity and mechanical properties of Ti–6Al–4V alloy. Materials Science and Engineering: A 486, 19–26. https://doi.org/10.1016/j.msea.2007.08.053 [Mehrpouya, 2024] Mehrpouya, M., Biffi, C.A., Lemke, J.N., Bregoli, C., Fiocchi, J., Mohajerani, S., Tuissi, A., Elahinia, M., 2024. Additive manufacturing of architected shape memory alloys: a review. Virtual and physical prototyping, 19, https://doi.org/10.1080/17452759.2024.2414395. [Miller, 2001] Miller D.A., Lagoudas D.C., Influence of cold work and heat treatment on the shape memory effect and plastic strain development of NiTi, Materials Science and and Engineering A 308 (2001) 161-175. [Saggin, 2022] Saggin, B., Scaccabarozzi, D., Appiani, A., Rusconi, F., Naon, M.G., Bellucci, G., 2022. RIIFS spectrometer optical bench design, in: 2022 IEEE 9th International Workshop on Metrology for AeroSpace (MetroAeroSpace). IEEE, pp. 107–112. https://doi.org/10.1109/MetroAeroSpace54187.2022.9856085 [Scaccabarozzi, 2024a] Scaccabarozzi, D., Ahmed, A.M.R.M., Saggin, B., Esposito, F., Porto, C., Mongelluzzo, G., Franzese, G., Silvestro, S., Donnarumma, I., Turchi, A., Cortesi, U., D’Amico, F., Gai, M., Argan, A., 2024a. Feasibility Design of LD GRIDS, a Dust Analyzer for the Moon, in: 2024 11th International Workshop on [Scaccabarozzi, 2024b] Scaccabarozzi, D., Potemkin, K., Saggin, B., Vieira, E., Corti, M.G., Martina, C., Appiani, A., Ortega, A.M., Arruego, I., Jimenez Martin, J.J., Fernandez, L.M.G., Palomino, M.S., Garcia-Ibarrola, D.G., Moreno, A.G., Rodriguez, M.F., Rebate, N.M., Braukhane, A., Quantius, D., 2024b. Feasibility Design of MiLi, a Miniaturized Lidar for Mars Observation, in: 2024 11th International Workshop on Metrology for AeroSpace (MetroAeroSpace). IEEE, pp. 432–436. https://doi.org/10.1109/MetroAeroSpace61015.2024.10591546 [Sinha, 2023] Sinha, A., Kumar Rajak, D., Shaik, N.B., Mohapatra, R.K., Saxena, K.K., Singh, R., Xu, J., Behera, A., 2023. A review on 4D printing of Nickel-Titanium smart alloy processing, the effect of major parameters and their biomedical applications. Proceedings of the Institution of Mechanical Engineers, Part E: Journal of Process Mechanical Engineering. https://doi.org/10.1177/09544089231154416 [Yan, 2024]Yan, Z., Wu, K., Xiao, Z., Hui, J., Lv, J., 2024. The Effect of Scanning Strategy on the Thermal Behavior and Residual Stress Distribution of Damping Alloys during Selective Laser Melting. Materials 17, 2912. https://doi.org/10.3390/ma17122912 Metrology for AeroSpace (MetroAeroSpace). IEEE, pp. 427–431. https://doi.org/10.1109/MetroAeroSpace61015.2024.10591536

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Procedia Structural Integrity 69 (2025) 76–79

13th European Symposium on Martensitic Transformation 2024 (ESOMAT 2024) Deformation Behavior of NiTi Shape Memory Alloy Subjected to Severe Torsion Deformation Victor Komarov a, b, c *, Roman Karelin a, b , Vladimir Cherkasov b , Ivan Postnikov c , Grzegorz Korpala c , Irina Khmelevskaya a, b , Vladimir Yusupov a , Rudolf Kawalla c , Ulrich Prahl c , Sergey Prokoshkin b Abstract In the present work, the possibility of application severe torsion deformation (STD) to a bulk near-equiatomic NiTi shape memory alloy in order to accumulate high strain was studied. STD was performed using the multidirectional test system “BÄHR MDS 830” in a temperature range of 300-600 °C (the dynamic polygonization range) in 10 turns with accumulated true strain of e =3. The results obtained show the prospects of application STD to bulk NiTi samples in terms of large strains accumulation. The STD at 350-600 °C leads to the formation of steady state flow stress stage. The polygonized dislocation substructure can be formed in this temperature range at e = 0.6 to 3. The obtained flow curves can be used for mathematical modeling and optimization of deformation processes of NiTi alloys. © 2025 The Authors. Published by ELSEVIER B.V. This is an open access article under the CC BY-NC-ND license (https://creativecommons.org/licenses/by-nc-nd/4.0) Peer-review under responsibility of Guest Editors a Baikov Institute of Merallurgy and Materials Science RAS, Leninskiy pr. 49, 119334 Moscow, Russia b National University of Science and Technology MISIS, Leninskiy pr. 4, 119049 Moscow, Russia c TU Bergakademie Freiberg, Bernhard-von-Cotta-Straße 4, 09599 Freiberg, Germany

Keywords: shape memory alloys; NiTi; severe plastic deformation; torsional deformation; deformation behavior

1. Introduction NiTi shape memory alloys (SMA) are attractive functional materials that are widely used in various fields of engineering and medicine for production of shape-memory devices [1-5]. The severe plastic deformation (SPD) is one

* Corresponding author. E-mail address: vickomarov@gmail.com

2452-3216 © 2025 The Authors. Published by ELSEVIER B.V. This is an open access article under the CC BY-NC-ND license (https://creativecommons.org/licenses/by-nc-nd/4.0) Peer-review under responsibility of Guest Editors 10.1016/j.prostr.2025.07.011