# The Dependence of Properties on Structural Features in Macro/Micro Levels for PBF-LB Fabricated Porous NiTi Alloy

Wen-Hao Zeng, Yang Yang, Xing-Rong Chu, Yang Zou, Ren-Zhou Deng, Ke Huang, Philippe Castany, Thierry Gloriant

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Wen-Hao Zeng, Yang Yang, Xing-Rong Chu, Yang Zou, Ren-Zhou Deng, et al.. The Dependence of Properties on Structural Features in Macro/Micro Levels for PBF-LB Fabricated Porous NiTi Alloy. Rare Metals, 2025, pp.e70053. 10.1002/rar2.70053. hal-05461611

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# The Dependence of Properties on Structural Features in Macro/Micro Levels for PBF-LB Fabricated Porous NiTi Alloy

Wen-Hao Zeng<sup>1,2</sup> | Yang Yang<sup>3</sup> | Xing-Rong Chu<sup>1,2</sup> | Yang Zou<sup>3</sup> | Ren-Zhou Deng<sup>1,2</sup> | Ke Huang<sup>4</sup> | Philippe Castany<sup>5</sup> |
Thierry Gloriant<sup>5</sup>

<sup>1</sup>School of Airspace Science and Engineering, Shandong University, Weihai, China | <sup>2</sup>Preparation and Application of Aerospace High-Performance Composite Materials, Future Industry Laboratory of Higher Education Institutions in Shandong Province, Shandong University, Weihai, China | <sup>3</sup>State Key Laboratory for High Performance Tools, School of Electromechanical Engineering, Guangdong University of Technology, Guangzhou, China | <sup>4</sup>State Key Laboratory of Mechanical Manufacturing System, School of Mechanical Engineering, Xi'an Jiaotong University, Xi'an, China | <sup>5</sup>INSA Rennes, CNRS, ISCR-UMR 6226, University of Rennes, Rennes, France

**Correspondence:** Yang Yang (yy@alum.imr.ac.cn; yy@gdut.edu.cn) | Xing-Rong Chu (xrchu@sdu.edu.cn) | Philippe Castany (Philippe.castany@insa-rennes.fr)

**Received:** 19 July 2025 | **Revised:** 13 October 2025 | **Accepted:** 13 October 2025

**Keywords:** mechanical property | PBF-LB | porous NiTi alloy | recovery property | structural characterization

## ABSTRACT

Porous NiTi alloys are widely used in biomedical and engineering fields. Using the PBF-LB technique, the TPMS and Voronoi porous parts with a porosity of ~67% and a minimum feature size of 0.15 ± 0.02 mm were fabricated. Although the bulk sample fabricated with the same parameters is defect-free, μ-CT reconstruction on the porous sample revealed inevitable defects with a volume fraction of 0.23%–0.43% and sizes ranging from 15 to 65 μm. For the two structures, a significant increase in thickness of the horizontal walls or struts was observed, with obvious fluctuations at the bottom contours. The limited epitaxial growth of B2 grains due to heat conduction constraints and weakened crystallographic texture were observed. Considering the optimized low-temperature aging (450°C/1 h) that promotes the formation of a high-volume fraction of R phase leading to superior elongation of 18.8% in bulk sample, this treatment was applied to the porous structures to prevent premature fracture during compression. Based on compressive cyclic tests, the mechanical properties as well as the functional properties were investigated. Results revealed that the designed porosity indeed affects the ultimate strength and Young's modulus, but the spatial configuration as TPMS or Voronoi gives a limited influence on the recovery property for the porous NiTi alloy. The mechanical recoverable strain presently peaks at 4.2% when the applied compressive strain reached at 21.3%. This study provides theoretical evidence for tailoring structures at both macro and micro levels to achieve the desired mechanical and recovery properties for porous NiTi alloys.

## 1 | Introduction

Porous NiTi alloy refers to a nickel-titanium alloy that contains intentionally designed macro/micro voids or pores within its structure [1]. The inherent advantages of NiTi alloys, such as low Young's modulus, superior ductility (elongation > 50%) [2], superelasticity (mechanical recoverable strain > 10%) [3], along

with the additional benefits introduced by porosity, make porous NiTi alloys valuable in biomedical [4–6] and engineering fields [7].

The performance of porous NiTi alloys is significantly influenced by porosity. Porosity leads to a reduction in the overall density of the material and alters its related mechanical

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© 2025 The Author(s). *Rare Metals* published by John Wiley & Sons Australia, Ltd on behalf of Youke Publishing Co., Ltd.
properties, especially at higher levels of porosity, where the
decrease in elastic modulus and compressive strength becomes
very noticeable [8]. On the other hand, the introduction of pores
also affects the functional properties which can be evaluated by
the recoverable strain originating from the martensitic phase
transformation [9]. For instance, the recoverable strain may be
decreased with increased porosity as the pores enhances uneven
stress distribution and in turn locally suppresses phase trans-
formation in regions near the pores [10, 11]. Therefore, the
choice of porous structures with weak stress concentration is of
great importance in order to achieve a more uniform deforma-
tion across the porous structure. Among all types of porous
structures, triply periodic minimal surface (TPMS) structures
are a class of geometrically defined periodic designs that feature
continuous surfaces and minimize mean curvature [12]. Their
advantage in avoiding anisotropy in mechanical behavior makes
them an excellent structure for reducing stress concentration.
Voronoi structures are another type of geometric design based
on Voronoi tessellation [13]. Theoretically, Voronoi structures
are potential to exhibit isotropic mechanical behavior once seed
points are randomly and uniformly distributed. This similarity
between TPMS and Voronoi structures provides the possibility
of achieving the desired mechanical and functional properties in
porous NiTi alloys compared to other porous structures. How-
ever, there is currently no comparative study on TPMS and
Voronoi porous structures in NiTi alloys on the topic.

In order to fabricate the porous NiTi alloy, powder bed fusion of
metals with a laser beam (PBF-LB) as an important method due
to its higher manufacturing precision compared to other addi-
tive manufacturing methods attracts attention. However, using
PBF-LB to fabricate porous NiTi alloys still faces several chal-
lenges. Firstly, the laser beam size used in the PBF-LB process is
typically 50–100 µm [14], which seems unfavorable for pro-
ducing pore structures smaller than 100 µm. Secondly, improper
selection of laser parameters can lead to too deep or too wide
molten pools, and the excessive heat conduction from the
molten pools to the surrounding powders will inevitably result
in powder adhesion, thereby reducing obtained porosity [15].
Furthermore, the fine features of porous structures, such as
single struts, thin walls, and narrow designed pores [16], may
vary in different directions in space. Inappropriate laser pa-
rameters can further increase the size deviations of the formed
struts and walls and cause the designed structure to deviate
from the actual structure [17]. Therefore, the laser energy input
for PBF-LB processing of porous structures needs to be precisely
controlled to eliminate unacceptable structural and dimensional
deviations.

Based on the successful fabrication of bulk NiTi alloy by present
authors [18–21], the lower elongation showing the brittle
feature at room temperature is regarded to be a crucial problem
which inhibits the further industrial application of both bulk
and porous NiTi alloys. Additionally, it was found that although
the mechanical recovery rate of bulk NiTi samples reached
98.7%, the elongation was only 7.2% [22]. Although superior
elongation was claimed to be obtained for the PBF-LB fabri-
cated NiTi alloy through optimization of laser scanning strat-
egy, the elongation at 15.6% seems still remains far lower than
the conventional values [23, 24]. In another study, when PBF-
LB fabricated NiTi alloys achieved an elongation of 11%, it

was regarded that a ductile fracture mechanism had been
developed [25]. In order to improve the ductility, researchers
used both optimization of laser scanning strategies and post
heat treatment to induce the transformation from brittle to
ductile for NiTi alloy [26–36]. For instance, Finazzi et al. [30]
employed the heat treatment at 550°C for 15 min, resulting in
an elongation approximately at 7% with a slight increase
compared with the PBF-LB fabricated alloy. On the other hand,
some researchers resorted to laser scanning strategies as the
post heat treatment was regarded to link with a lower elonga-
tion. For example, Zhang et al. [31] obtained the elongation less
than 10% on NiTi alloy bulk sample fabricated with PBF-LB and
the further decrease of elongation on the heat-treated samples
at 300°C–1000°C. Zhong et al. [32] revealed a total elongation of
8.7%–14.2% originating from the hierarchical microstructure. It
should be admitted that the brittle characteristics of PBF-LB
fabricated NiTi alloys remain far from being resolved. It has
to be noted that the ductility for alloy fabricated with conven-
tional casting and forging methods depends mainly on the
microstructure as there are few defects inside, so the ductility
evaluation can be evaluated via either or compressive tests to
large deformation strain. However, the potential defects like
non-spherical pores and undetectable micro-cracks existing in
the PBF-LB fabricated alloys would cause premature fracture
during tensile test and finally decrease the ductility. That is why
the tensile tests on bulk samples are irreplaceable for opti-
mizing the ductility of PBF-LB fabricated NiTi alloys presently.

In this present study, a specific TPMS structure and a Voronoi
structure with similar designed porosity will be designed for
fabrication the porous NiTi alloy with the PBF-LB technique.
An optimal low-temperature heat treatment (LTA) confirmed
on bulk sample via tensile tests will be directly used to tailor
the microstructure of the porous alloy to avoid potential pre-
mature fracture. The structural features of porous NiTi alloy
from macro to micro levels will be revealed with the techniques
of micro-computer tomography (μ-CT), optical microstructural
observation, differential scanning calorimetry (DSC), X-ray
diffraction (XRD), electron backscattered diffraction (EBSD),
scanning electron microscope (SEM), and transmission electron
microscope (TEM). The cyclic compressive tests will be per-
formed to evaluate the mechanical properties and mechanical
recoverable strains for porous samples. The dependence of
properties relating to macro/micro-structures on porous NiTi
alloy will be finally discussed.

2 | Experimental

The chemical composition of the near equal-atomic NiTi pre-
alloyed powders used in the present study is shown in Support-
ing Information S1: Table S1. Figure 1A,B presents two types
porous structures which are TPMS and Voronoi. The dimensional
features of the two structures are indicated inside, with the
minimum hole size and minimum wall thickness being approx-
imately 0.15 ± 0.02 mm which were obtained on basis of at least
five measurement on with software of “Materialise Magics.” The
average wall thickness for TPMS and average strut thickness for
Voronoi are measured to be 0.20 ± 0.02 and 0.34 ± 0.3 mm,
respectively. The bulk samples with 3 mm × 20 mm × 1.5 mm
<figure><img src="image_2.png" /><figcaption>**FIGURE 1** | The designed structures and raw powders used in the present study: (A) TPMS porous structure, (B) Voronoi porous structure; (C) SEM Morphology of the powders, (D) DSC curves of the powders, (E) size distribution, (F) schematic diagram of linear rotation scanning strategy, (G) NiTi alloy samples on the substrate.</figcaption></figure>

gauge dimensions and tensile direction parallel to the horizontal direction were fabricated with PBF-LB method for evaluating tensile properties. The porous samples with diameter of 5 mm and height of 10 mm were fabricated with PBF-LB method for evaluating compressive properties. The powders morphology is observed to be spherical as shown in Figure 1C. DSC curve tested with NiTi powders was shown in Figure 1D. The results indicate the initial NiTi powders contains nearly fully B2 austenitic phase at room temperature. The size distribution with Gaussian profile in Figure 1E demonstrates the powders are suitable to spread homogeneously and compactly onto substrate.

The additive manufacturing process was performed on a PBF-LB equipment (HBD 100). The IPG YLR-200-AC continuous ytterbium-doped fiber laser which produces continuous Gaussian laser with maximum laser power of 200 W is installed on the machine. The laser wavelength and laser diameter at focal plane is 1064 nm and 100 µm, respectively. Three key processing parameters which are laser power (P), laser scanning

speed (v), and hatch distance (h) were emphasized for PBF-LB process optimization. Actually, the layer thickness (t) was kept at a constant value of 30 µm. The laser power is set between 50 and 150 W for guarantee the power stability. According to the equations of $E_1 = P/v$ and $E_v = P/(v \times t \times h)$ [37, 38], the linear energy density ($E_i$) and volumetric energy density ($E_v$) can be calculated. Argon with purity higher than 99.9% was used as protective gas and the oxygen content can be controlled lower than 100 ppm. In addition, the laser scanning strategy is schematically displayed in Figure 1F. A linear scanning strategy with the initial rotation angle of 0° was used. The interlayer laser rotation angle increment is set to be 67° for decreasing the thermal stress between layers. The laser scanning distance between adjacent tracks are kept at a constant value of 80 µm. All PBF-LB processing parameters were concluded in Supporting Information S1: Table S2. According to Figure 1F,G, it is shown that the original point of coordinating system is marked by “O.” The substrate is located on the X-Y plane, whereas the Z direction is parallel to the build direction (BD).
The samples fabricated directly by PBF-LB without post thermal treatment was designated as “PBF-LB” state. In order to optimize ductility of LPBF NiTi samples, post heat treatment should be generally employed [26, 39]. However, according to previous investigation [19, 31], high amount of oxygen element existing as brittle Ti₂O or Ti₄Ni₂O in the PBF-LB fabricated NiTi alloy are inevitable to form high-temperature heat treatment which results in brittleness in the NiTi alloy. In contrast in another public paper [40], LAT was suggested to possibly configure the properties while simultaneously avoid the unfavorable brittle phase. On basis, the LTA with aging temperature ranging from 350°C to 450°C for 1, 3 and 18 h followed by water quenching (WQ) to room temperature were used to tailor the microstructure and enhance the ductility for PBF-LB fabricated NiTi alloy. According to the exact conditions, LTA samples were designated as “350°C/1 h, 350°C/3 h, 350°C/18 h and 450°C/1 h.”

Microstructures were all observed in detail with optical microscopy, SEM (Hitachi, SU8220), and TEM (FEI, Talos F2000S) on mechanical polished and ion milled samples. Phase identification and crystallographic texture were analyzed with XRD (D8 Advance, Bruker) and EBSD techniques with step sizes of 0.5 µm. The potential phases including the B2 phase, B19' phase, and R phase lattice parameters and group number listed in Supporting Information S1: Table S3 [41] were all used for data analysis. The thermal phase transformation was characterized on DSC equipment (Mettler Toledo, TGA/DSC3+). Relevant heating and cooling rates for DSC were set at 10°C min⁻¹ with liquid nitrogen. Internal metallic defects were examined with the non-destructive testing µ-CT (nanoVoxel-3000). The voltage and current of X-ray source were set at 150 kV and 40 µA, while the exposure time was 0.5 s. To ensure the quality of CT scanning results, the resolution and rotation steps were set at 4 µm and 0.5°, respectively. All obtained digital signals were used to reconfigure the three-dimension structure with software of Avizo. In order to guarantee the reliability of properties, the three replicate samples were prepared and tested, but only one curve for each condition was displayed in detail. The supersonic vibration was used to remove the loosely adhering powders on the porous samples before mechanical testing. The mechanical properties and recovery property were evaluated by cyclic tensile tests on a LISHI instrument (LD24.304) with an extensometer and cyclic compressive tests on MTS instrument (MTS/E45.105).

# 3 | Results and Analysis

## 3.1 | Optimization of PBF-LB Parameters

Because the single molten pool is regarded to be the unit for PBF-LB fabrication process, the continuity and morphology of single molten track are of great importance for assurance of high quality. The single molten tracks were initially deposited with various laser processing parameters for preliminary optimization. Relevant results of the observations on the top surface and optical microstructure in the cross section are shown in Figure 2A,B, respectively. According to the surface morphologies shown in Figure 2A, it is known that the laser was scanned from left to right.

All the tracks in Figure 2 were generally classified into three groups “a,” “b,” and “c” marked inside.

For group “a,” the linear energy density ($E_i$) ranges from 0.38 to 0.75 J mm⁻¹. Excessive input of laser energy for this group causes the formation of keyholes highlighted by red circles in Figure 2B. As heat affected zones exist once high energy input was used for PBF-LB, this probably results to more powders attached around the contours and is not suitable for fabrication of tiny porous structure with high quality. In contrast for the c group, the value of $E_i$ calculated between 0.04 and 0.11 J mm⁻¹ is not sufficient. Under such circumstance, the NiTi powders cannot be melted completely. The single tracks also display discontinuity and spherical droplets as shown in Figure 2A by yellow circles. Attentions can be paid on group “b” for which single molten tracks display good continuity and well stability. Its linear energy density ranges between 0.15 and 0.25 J mm⁻¹ which seems quite appropriate for PBF-LB fabrication of NiTi alloys.

Based on morphology of tracks in Figure 2A,B, the width and depth of a single melt pool are measured and presented in Figure 2C,D in relation to the scanning speed and laser power. As shown, either the increase of scanning speed or decrease of laser power would lead to a decrease of width and depth for the molten pools. It is also worth noting that even when the $E_i$ value remains constant, for example, 0.25 J mm⁻¹ for the samples 50 W–200 mm s⁻¹, 100 W–400 mm s⁻¹, and 150 W–600 mm s⁻¹, the width and depth of the molten pools still vary greatly, as indicated by red arrows in Figure 2C,D. This is because the importance of laser power and scanning speed is different. The morphology of single molten pool actually depends on the action of the laser spot and also the heat diffusion which links more closely with the scanning speed [42]. Indeed, the combination of lower laser power of 50 W and lower scanning speed of 200 mm s⁻¹ results in a wider but shallower molten pool, whereas the combination of higher laser power of 150 W and higher scanning speed of 600 mm s⁻¹ creates a narrower but deeper molten pool. The superiority of a specific combination of laser power and scanning speed also depends on how well the resulting melt pool dimensions match the precise structure of the fabricated product.

Figure S1 displays additionally the relationship between porosity on PBF-LB parameters for bulk samples fabricated with constant layer thickness of 30 µm and hatch distance of 80 µm. When the laser power is kept constant at 50 W, the number of black defects gradually increases with the scanning speed, whereas at the laser power of 150 W, the number of black defects decreases as the scanning speed increases. However, the morphologies of defects for these two groups are different. The defects in the former group exhibit irregular shapes and are referred to as lack of fusion (LOF) defects [43], suggesting that the metallic powders may not have fully melted. In contrast, the spherical defects shown in Supporting Information S1: Figure S1 are typical keyhole defects [44]. In Figure S1H, when the laser power and scanning speed are kept at 100 W and 600 mm s⁻¹, respectively, the number of defects seems to be minimal.

It can be found that for bulk samples, the setting for two other parameters, hatch distance, and layer thickness should be
<figure><img src="image_3.png" /><figcaption>**FIGURE 2** | Observations on the continuity and morphology of single molten tracks fabricated with varied PBF-LB processing parameters: (A) top surface morphology observed with SEM, (B) optical microstructure in the cross section, (C) variation of width for single molten pool and (D) variation of depth for single molten pool.</figcaption></figure>

further considered. Bur for porous samples, attention on the match between the precise structure feature and sizes of single molten pools should also be paid. For the present study, the porous structure shown in Figure 1A,B is characterized by the minimum hole size and thin walls approximately in the range from 0.15 to 0.2 ± 0.1 mm. Hence, the processing parameters relating to depth and width of molten pools higher than 200 µm are absolutely not suitable for present fabrication even if fewer defects could be obtained in bulk samples. In contrast, a smaller depth and width lower than 200 µm are proposed here. Therefore, a general approach for optimization of the PBF-LB processing parameters for fabrication of the porous metallic materials is obtained. The laser power and scanning speed, as indicators of the spatial and temporal distribution of laser energy, are definitely critical for the continuity and stability of laser molten tracks. The hatch distance and layer thickness should be considered secondarily to further match with the laser power and scanning speed to ultimately determine a processing

parameters package which can be compatible with the precise structural features. Based on this approach, the optimized parameters here are 100 W–600 mm s⁻¹ with a fixed layer thickness of 30 µm and hatch distance of 80 µm for fabrication the porous NiTi alloy.

## 3.2 | Characterization of Macro/Micro-Structures

Based on the above optimized PBF-LB processing parameters, NiTi samples with TPMS and Voronoi porous structures were fabricated. Relevant macro/micro structural characterizations are presented in Figures 3–5. The structural features were firstly visualized by 3D reconstruction of µ-CT data in Figure 3. It can be noted that the gray images in the first column display the external morphology of TPMS and Voronoi structures. The blue images in the second column show the distribution of interior
<figure><img src="image_4.png" /><figcaption>**FIGURE 3** | 3D reconstruction of micro-CT data to reveal macro/micro structural features: (A) TPMS porous structure, (B) Voronoi porous structure, (C) comparison of statistical frequency for defects with varied diameters. (a and b: 2D cross-sections of different planes from the 3D CT images in the first column).</figcaption></figure>

defects with diameters greater than 10 µm in both structures. Statistical analysis performed with Avizo software reveals that the volume fraction of all defects is 0.23% ± 0.08% and 0.43% ± 0.15% for TPMS and Voronoi structures, respectively. The images in the third and fourth columns provide 2D cross-sections of planes “a” and “b” from the 3D CT images in the first column, where tiny spherical defects can be observed inside the samples. Further statistical analysis of the defects detected is shown in Figure 3C in which the normalized frequency of defects with diameters ranging from 15 to 65 µm could be observed. Actually, the diameter of the defects follows a Gaussian distribution. Among them, approximately 26% of defects detected have sizes between 35 and 40 µm, and these results are in good agreement with the optical microstructure observations, further enhancing the credibility of the CT results in Figure 3.

Macro/micro-structures were further observed with SEM and optical microscopy to reveal the surface morphology and internal metallurgical microstructure as shown in Figure 4. It can be observed that partial powders previously metallurgically bonded to surface had been removed away from the TPMS and Voronoi samples in Figure 4A,B after chemical cleaning for few

seconds with acid solution composed of 50% HNO₃ and 50% HF in volume fraction. The morphology of designed porous structures can be recognized well. However, as shown in Figure 4B, some powders are still adhered to the walls of deep holes. These metallic powders adhered to the surface of PBF-LB metallic alloys are difficult to remove, which could become a weakness and stress concentration area during the subsequent mechanical deformation process.

The optical microstructural observations in the middle and right columns of Figure 4A,B provide more detailed information. Firstly, based on the coordinates shown by the red arrows in the bottom left corner (BD//Z direction), the horizontal and vertical structures can be distinguished in the middle columns of Figure 4. It is obvious that the horizontal structure is thicker than the vertical structure. To be specific, this interesting fact is highlighted by yellow arrows and dashed lines. The thickness of the vertical strut and horizontal strut was measured to be 281.8 ± 6.5 and 484.8 ± 7.5 µm, respectively. Such difference is because of the laser beam energy being transferred into the surrounding powder bed, which enhances the undesired epitaxial growth or sintered powders of the horizontal strut and therefore thicken the horizontal strut in vertical direction. For
A

<figure><img src="image_5.png" /></figure>

B

<figure><img src="image_6.png" /><figcaption>**FIGURE 4** | Macro/micro-structures observed with SEM and optical microscopy to reveal the surface morphology and internal metallurgical microstructure: (A) TPMS porous structure, (B) Voronoi porous structure.</figcaption></figure>

A

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

B

<figure><img src="image_8.png" /></figure>

C

<figure><img src="image_9.png" /><figcaption>**FIGURE 5** | EBSD IPFs-mappings, PFs, and KAM maps on different PBF-LB fabricated samples: (A) bulk sample, (B) TPMS porous sample, and (C) Voronoi porous sample.</figcaption></figure>
the vertical structure, the laser beam energy transferring into
the preceding solidified layers below the present layer could
enhance the epitaxial growth but cannot lead to the thickening
of the vertical structure in horizontal direction. Additionally,
some black dots of defects indicated by green arrows, were
observed in the optical microstructure of both TPMS and Vor-
onoi structures which is in consistence with the µ-CT
reconstructions.

Figure 5A displays EBSD inverse pole figure (IPF) mapping, pole figures (PFs) and kernel average misorientation (KAM) maps for the bulk sample, whereas Figure 5B,C provides IPFs for porous samples. Since potential precipitates such as B19', Ni₄Ti3, and R phases feature with small grain size, only the lattice information of the B2 phase listed in Supporting Information S1: Table S3 was used for Kikuchi indexation. It can be observed that columnar B2 grains are clearly formed along the Z direction due to epitaxial growth, which is parallel to the build direction. There is a noticeable difference in the microstructure between the porous and bulk samples. Specifically, the columnar grains in Figure 5B,C for porous NiTi samples contain smaller grains at the boundaries of the columnar B2 grains.

Further, the misorientation between adjacent layers in the
molten pools for porous samples is significantly greater
compared to the bulk samples. It can be proved by the KAM
images which is actually superimposed with the grain bound-
aries misorientations in the second row of each sample. It can be
observed that the length proportion for high-angle grain
boundaries with misorientation higher than 10° is 39.6%,

whereas the proportion increased to 53.8% and 57.1% for TPMS
and Voronoi samples, respectively. Indeed, the local strain
accumulation generally exists around the molten pool bound-
aries. Similar phenomenon on the strain accumulation at the
bottom of molten pools can be found in the reference [36].
Below the IPF mapping in Figure 5C, the corresponding {100}-
B2 pole figures (PFs) are shown. It is noted that the three
strongest {100}-B2 poles, marked by black hollow square sym-
bols, correspond to a maximum value of 6.48 according to the
color bar. For the TPMS and Voronoi structure samples shown
in Figure 5B,C, the maximum values for the {100}-B2 poles are
2.76 and 5.79, respectively, which are lower than the corre-
sponding values of 6.48 for the bulk sample.

3.3 | Phase Identification

As shown in Figure 6A, XRD was firstly used to preliminarily identify the phases in the NiTi alloy. However, due to the overlapping of the main diffraction peaks for the B2, B19', and R phases, it was not possible to accurately index the diffraction peaks in XRD spectrum. Subsequently, DSC and TEM experiments were conducted on all PBF-LB and heat-treated samples. For the PBF-LB sample, only a single broadened endothermic peak and exothermic peak corresponding to the thermally induced reversible phase transformation between martensitic B19' phase and austenitic B2 phase were observed on the heating and cooling curves in Figure 6B, highlighted in blue. It should be noted that the exothermic peak in the cooling curve for the PBF-LB sample appears in the temperature range lower

<figure><img src="image_10.png" /><figcaption>FIGURE 6 | Characterization of thermal-induced phase transformation: (A) XRD profiles, (B) DSC curves, (C) volume fraction measurement, and (D) XRD profiles for different samples.</figcaption></figure>
than room temperature (RT), whereas in heat-treated samples such as "350°C/18 h" and "450°C/1 h," the main exothermic peak on the cooling curve almost appears above room temperature. This suggests that initial phase constitution for the PBF-LB sample which underwent a full thermal cycle and cooling down to room temperature, likely contain nearly 100% of the B2 phase. In contrast, for the heat-treated samples "350°C/18 h" and "450°C/1 h," the B2 phase has almost transformed into R phase. The phase transformation types corresponding to each endothermic and exothermic peaks on the DSC curves have been clearly identified in Figure 6B.

It is noted that the DSC curves for the heat-treated samples in Figure 6B greatly differ which also depend on the specific temperature and duration. On cooling curve highlighted in orange for the sample "350°C/1 h," two distinct exothermic peaks can be observed. For heat-treated samples "350°C/3 h," "350°C/18 h," and "450°C/1 h" highlighted in violet, pink, and blackish green, both the double exothermic and endothermic peaks can be observed. The double phase transformation peaks indicate the involvement of the R phase which suppresses the formation of the B19' phase to some extent [40, 45] during the thermally induced phase transformation process. Under such condition, the thermal-induced phase transformation sequence of B2 ↔ B19' in PBF-LB sample changes to a two-step transformation of B2 ↔ R ↔ B19' in heat treated samples [46]. This double-phase transformation sequence is marked in Figure 6B for these heat-treated samples. Furthermore, it is also worth noting that for heat-treated samples "350°C/18 h" and "450°C/1 h," the exothermic peaks on cooling curve corresponding to the B2 → R occur at temperatures higher than RT. This indicates the B2 → R phase transformation is nearly completed at room temperature, and thus the phase constitution for the two heat-treated samples should be a mixture of a few B2 phase and more R phase. Taking the sample "450°C/1 h" in Figure 6C as an example, the volume fractions for B2 phase and R phase are calculated to be 18% and 82%, respectively, via integrating the orange and green areas of the exothermic peak corresponding to the B2 → R phase transformation. The volume fractions for all samples were analyzed using the method shown in Figure 6C and relevant results were displayed in Figure 6D. It can be seen that the heat-treated sample "450°C/1 h" contains the highest amount of R phase.

The TEM observations on PBF-LB bulk samples are shown in Figure 7A–L for further phase identification, whereas the TEM observations on PBF-LB porous samples were comparatively shown in Figure 7M–P. Supporting Information S1: Figure S2 displays TEM observations on all the heat-treated samples. Attention can be first on the comparison between bright field images (BFIs) for bulk and porous samples. It can be noticed the typical BFIs in Figure 7M–P is similar to the BFIs in Figure 7A–D. The, the observation on "450°C/1 h" sample shown in Figure 7J–L would be emphasized.

Typical morphology displays two features for PBF-LB NiTi alloy including the equiaxed grains in Figure 7A and elongated grains surrounded by much smaller grains shown in Figure 7B. The magnified observation in Figure 7C reveals the presence of precipitates sized between 10 and 20 nm. The same magnified area was tilted to another diffraction condition for precise identification, as shown in Figure 7D. The inserted diffraction pattern in

the top right corner proves that only the B2 phase can be indexed with its zone axis as < 111 > B2. The dark field image in STEM mode is shown in Figure 7E, along with the corresponding EDX results. It can be observed that these precipitates are rich in oxygen. EDX results in Figure 7F–H confirm the lack of nickel element. The high-resolution imaging of the precipitated phase in Figure 7I with inserted fast Fourier transform pattern provides explicit evidence to index the precipitate as TiO₂ phase which could be easily formed in NiTi alloy [31, 47]. It should be noticed that precipitates of Ni₄Ti₃ and Ti₄Ni₂O [18, 19, 33] found in the NiTi alloy samples fabricated by present authors with another different PBF-LB machine cannot be detected here. Based on the both results on DSC and TEM, it is known that phase constitution is nearly 100% of B2 phase with slight amount of nano-sized Ti₂O phase for PBF-LB fabricated NiTi alloy.

In Figure S2, the typical microstructures in the upper row display with BFIs, whereas the corresponding selected area electron diffraction patterns (SAEDPs) in the middle row corresponds to the area selected by yellow circles. In the bottom row, dark field images (DFIs) were highlighted with the diffracted spots highlighted by red circles for each sample. Phase identifications are affirmed through the indexation of SAEDPs shown in Supporting Information S1: Figure S2. It should be noticed that the Figure S2D were repetitively displayed in Figure 7J–L for better readability. It is clear that R phase exists in all the heat-treated sample with dispersed nano-sized morphology inside the grains and also in the grain boundary. For sample heat treated at 350°C for 1 h shown in Supporting Information S1: Figure S2A, both B2 and R phase were indexed, with very faint diffraction spots indicating that only a small amount of R phase exists within the B2 matrix. For sample heat treated at 350°C for 3 h, the amount of R phase is increased, some of them even exists along the grain boundaries as shown in DFI of Supporting Information S1: Figure S2B. For sample heat treated at 350°C for 18 h, slight amount of martensitic B19' phase can be observed via faint diffraction spots. The corresponding DFI in Supporting Information S1: Figure S2C also reveals the coarsened and lamellar morphology of R phase. Therefore, it can be concluded that the heat treatment at 350°C–450°C leads to the dissolution of nano-sized Ti₂O phase which was formed in previous PBF-LB fabricated process back into the NiTi matrix and favors to the formation of R phase. Once the heat-treated temperature is stabilized at 350°C, increasing the duration leads to an increase in both the quantity and size of R phase. For sample heat-treated at 450°C for 1 h, R phase is also coarsened as proven by the DFIs in Figure 7J–L. It should also be noted the phase identification obtained via DSC and TEM corresponds well with each other.

## 3.4 | Mechanical and Recovery Properties

The ductility of PBF-LB fabricated NiTi alloy was evaluated via tensile testing for bulk sample, relevant results were shown in Figure S3. Special attention was first paid on the elongation at rupture. As shown in Supporting Information S1: Figure S3A, the initial elongation for PBF-LB bulk sample is 12.2% ± 0.3%, which was increased after heat treatment. For example, the elongation for sample "450°C/1 h" is 18.8% ± 0.5% and for sample
<figure><img src="image_11.png" /><figcaption>**FIGURE 7** | (A–L) TEM observation on bulk NiTi samples and (M–P) porous samples: (A) typical BFI showing equiaxed grains, (B) typical BFI showing elongated grains separated by dislocation nets, (C) magnified area with precipitates, (D) the same magnified area tilted to diffracted conditions with diffraction pattern indexed as B2 phase on the inserted figure, (E) DFI in STEM mode with EDX results shown on the inserted figure on right corner, (F–H) EDS mapping for different chemical elements, (I) high resolution imaging of the precipitated phase with inserted fast Fourier, (J) another typical BFI, (K) corresponding SADP for area highlighted by yellow circle in panel (J), (L) DFI highlighted by diffraction spot circled by red in panel (K), (M–P) typical BFI showing similar morphology with the bulk samples.</figcaption></figure>

"350°C/18 h" is 16.7% ± 0.5%. It can be seen that the heat-treated sample "450°C/1 h" exhibits the maximum elongation at fracture. It is known that the ductility could be definitively optimized via employment of heat treatment of "450°C/1 h."

The mechanical compressive properties were then comparatively measured on samples "TPMS-PBF-LB," "TPMS-450°C/1 h," "Voronoi-PBF-LB," and "Voronoi-450°C/1 h" as shown in Figure 8A–D, respectively. The dependence of critical stress including the first yield stress (first σ<sub>c</sub>), the second yield stress (second σ<sub>c</sub>), and ultimate stress at fracture on state of the sample can be explored. The methods with the intersection of slopes used to measure the critical stress were shown in Figure 8A. All the measured critical stresses and Young's modulus were listed either in the right bottom corner of each figure in Figure 8 and compared in Figure 9A. For the Young's modulus, it indeed

depends on the chemical composition of intrinsic materials as well as precise structural porosity, and thus nearly keeps constant at 22–23 GPa for bulk sample but varies greatly among 1.5–4.6 GPa for porous samples as shown in Figure 9A.

The mechanical recoverable strain (ε<sub>rev</sub>) as a typical indicator to evaluate the properties for porous NiTi alloy was measured using the method shown in Supporting Information S1: Figure S3B. The relationships between the mechanical recoverable strain and applied strain for porous samples were plotted in Figure 9B. Further, the recovery rate calculated by the mechanical recoverable strain divided by the applied strain were plotted with respect to the applied strain as shown in Figure 9C. It is observed that the mechanical recoverable strain increases linearly with the applied strain, then stabilizes more or less at a certain level although differs among samples and increase
<figure><img src="image_12.png" /><figcaption>**FIGURE 8** | Compressive cyclic curves for porous NiTi samples: (A) TPMS-PBF-LB, (B) TPMS-450°C/1 h, (C) Voronoi- PBF-LB, (D) Voronoi-450°C/1 h.</figcaption></figure>

linearly again in the end. The maximum mechanical recover-
able strain is approximate 4.2% measured on TPMS-PBF-LB
sample. According to Figure 9B,C, the recovery rate for NiTi
alloy more or less decreases and finally stabilizes at approximate
20% in the final deformation stage for porous samples. What is
more, the variation curves for mechanical recoverable strain
with applied strain nearly coincide for all three 450°C/1 h
samples of porous or bulk types.

4 | Discussion

4.1 | Structural Features of Porous NiTi Alloy

Based on the morphologies of single molten tracks observed in Figure 2, the optimized laser parameters for porous NiTi alloy are suggested to be laser power of 100 W, scanning speed of 600 mm s⁻¹ with constant layer thickness of 30 µm and hatch distance of 80 µm. The morphology of selected single molten pool is featured with the depth of 126 µm and the width of 175 µm. It has to be noted that no apparent defects can be detected in bulk samples with such optimized parameters as shown in Supporting Information S1: Figure S1H. However, the CT scanning results in Figure 3 revealed a number of defects existing in porous samples in micro level sized in 15–65 µm with a total volume fraction of 0.23% ± 0.08% for TPMS and 0.43% ± 0.15% Voronoi structures. Here, it can be concluded the main difference between the porous sample and bulk sample fabricated with the same processing parameters lies on the internal defects. In order to explain the formation mechanism for such difference, the laser scanning strategies were schematically drawn in Figure 10A,B.

For porous samples, due to the limited scanning length which
varies frequently among different areas in one layer, the laser
has to turn frequently resulting to molten pool less stable
compared to the bulk samples. The unstable feature during
PBF-LB fabrication would lead to the uneven heat distribution
and more spatters which eventually results to more internal
defects. Although previous investigation has categorized spat-
ters into five types [48], it has not been distinguished because
the dependence between spatter and defects are not presently
focused. In fact, the spatters monitored with high-speed cam-
era has demonstrated the dependence on scanning length [49].
The finding that more defects and spatters were observed in
specimens fabricated with smaller scanning length revealed by
in reference [50] proved the rationality. Statistical investigation
would be conducted with high-speed camera in the near
future.

Another structural feature worthy of note is the intentionally
designed hollow porosity, which is comparatively listed in
Table 1. According to the structural features shown in Figure 1,
the theoretical porosity values of the hollow-designed TPMS and
Voronoi structures are 67.8% and 66.8%, respectively. However,
the real porosities obtained from the CT scanning results for the
two porous structures are only 47.4% ± 0.9% and 44.8% ± 1.1%,
respectively. Meanwhile, as shown in Figure 4, the decrease in
porosity is accompanied by an increase in wall or strut thick-
ness. Such increase in thickness closely related to the mea-
surement direction also shows anisotropy. Taking the TPMS
structure as an example, the designed wall thickness is
approximately 200 µm. However, based on the measurements
highlighted in Figure 4A, the average wall thickness measured
on the XZ plane is 396.9 ± 50.3 µm, and the average wall
<figure><img src="image_13.png" /></figure>

<figure><img src="image_14.png" /></figure>

<figure><img src="image_15.png" /><figcaption>**FIGURE 9** | Mechanical properties and functional properties for PBF-LB fabricated and heat treated porous NiTi alloys: (A) comparison on measured critical stresses, (B) variation of mechanical recoverable strain with increase of applied strain, (C) variation of mechanical recovery rate with increase of applied strain.</figcaption></figure>

thickness measured on the XY plane is 275.8 ± 9.5 µm. Therefore, for the TPMS structure, the wall thickness measured on the XZ plane is greater than that on the XY plane. A schematic diagram is drawn in Figure 10C,D to illustrate the formation mechanism of this phenomenon.

As shown in Figure 10C, the wall thickness difference between the designed contour (highlighted in red) and the

actual formed contour (highlighted in green) on the XY plane mainly originates from the deviation between the heat-affected zone and the outermost molten track. However, the wall thickness on the XZ plane shown in Figure 10D increases significantly at the bottom of wall and presents fluctuation. The first reason is the continuous heat conduction input from the upper region, which inevitably leads to epitaxial grain growth in the BD//Z direction and the sintered powder formed via full or semi-metallurgical bonding at the bottom of the wall [51]. Another reason is the gravity, which causes the molten metal at the bottom of thin walls or struts to sink into the powder bed [52, 53], leading to asymmetric morphology and non-uniformity of the contour as shown in green in Figure 10D. These factors result in the anisotropy of the porous structural features in three-dimensional space, presenting a structure significantly different from that of bulk material in relative macro-level.

With regard to the structural features in grain level, smaller B2 grains with limited growth and weakened crystallographic texture can be observed in porous NiTi alloy according to EBSD mapping in Figure 5. These results indicate that the macro-structure, whether bulk or porous, indeed affects the microstructure and despite using the same laser parameters. The main reasons can be attributed to differences in cooling rate, temperature gradient and thermal history influenced by the dimensional characteristics [54]. Less powders are melted during PBF-LB for the porous samples compared with the bulk samples, which leads to less heat accumulation. Additionally, the shorter scanning length in the PBF-LB process for porous structures disrupts the stabilization of the molten pool, ultimately affecting the solidification orientation of the grains and leading to the formation of defects. The heat conduction in porous samples is additional different as powders are filled in the gaps between solidified struts, therefore, heat could transfer to several directions in space. However, for the bulk sample, the molten tracks are stacked and heat conduction continuously transfers to adjacent areas, primarily from the upper region to the bottom solidified regions. These factors may all influence the crystallographic texture and grain size.

The presence of metallurgical defects could lead to local stress concentration and failure in advance. On this basis, the reduction in porosity could result in an increase in strength and Young's modulus compared to the theoretical values. However, the specific ultimate strength is significantly related to the porosity which requires detailed analysis in future with specific cases. Taking into consideration that number of potential defects and sintered powders of borders [54] would be increased with higher laser power and lower scanning speed due to larger heat affected zone, a lower laser power coupling moderate scanning speed should be a better choice for assurance of the dimensional accuracy for tiny or porous structure metals. Laser pre-sintering (LPS) [55] is probably beneficial for fabrication of metallic materials with tiny structure. All these scanning parameters and strategies will be tried in the future.
<figure><img src="image_16.png" /><figcaption>B Bulk structure</figcaption></figure>

<figure><img src="image_17.png" /><figcaption>C XY plane</figcaption></figure>

<figure><img src="image_18.png" /><figcaption>D XZ plane</figcaption></figure>

<figure><img src="image_19.png" /></figure>

**FIGURE 10** | Schematically illustration on mechanisms for the structural defects for (A) porous structure and (B) bulk structure; schematically illustration on increased wall or strut thickness due to heat conduction leading to epitaxial growth and undesired sintered powders on (C) XY plane and (D) on XZ plane.

<table><caption>**TABLE 1** | Comparison between designed structural features and obtained counterparts.</caption><thead><tr><th rowspan="3">Structure type</th><th rowspan="3">Defects ratio (%)</th><th colspan="2">Hollow porosity (%)</th><th colspan="3">Wall or strut thickness (µm)</th></tr><tr><th rowspan="2">Designed</th><th rowspan="2">Obtained from CT</th><th rowspan="2">Designed</th><th colspan="2">Obtained from OM</th></tr><tr><th>On XY plane</th><th>On XZ plane</th></tr></thead><tbody><tr><td>TPMS</td><td>0.23 ± 0.08</td><td>67.8</td><td>47.4 ± 0.9</td><td>200 ± 20</td><td>275.8 ± 9.5</td><td>396.9 ± 50.3</td></tr><tr><td>Voronoi</td><td>0.43 ± 0.15</td><td>66.8</td><td>44.8 ± 1.1</td><td>340 ± 30</td><td>355.7 ± 120</td><td>351.7 ± 60</td></tr></tbody></table>

## 4.2 | The Dependence of Properties on Microstructure for Porous NiTi Alloy

The relationship between the microstructure and properties of porous NiTi alloys can be explored. Initially, the intrinsic properties of bulk NiTi alloys are analyzed with a focus on ductility and critical stress determined by the microstructure. Subsequently, the impact of designed porosity on the mechanical recoverable strain of the porous sample is discussed to reveal the main factors affecting the mechanical recoverable performance of shape memory alloys.

The heat-treated sample “450°C/1 h” contains a higher amount of R phase, whereas the PBF-LB sample contains more B2 phase. Among all bulk samples, the former exhibits the highest elongation at 18.8% ± 0.5% and the latter shows the lowest at 12.2% ± 0.3%. Additionally, based on Supporting Information S1: Figure S3, it can be seen that the former sample “450°C/1 h” demonstrates the lowest first critical yield stress at 104.7 ± 3.2 MPa, but the highest ultimate tensile strength (UTS)

at 900.5 ± 72.0 MPa, resulting in the maximum stress difference on the corresponding stress-stress curve. The bulk heat-treated sample “450°C/1 h” hence shows a lower first critical stress and distinct strain-hardening phenomenon before fracture which indicates a characteristic of ductile deformation. In contrast, the PBF-LB fabricated bulk sample shows a higher first critical stress and a weaker strain hardening. These facts remind us to hypothesize that adjusting the microstructure to decrease the first critical stress and obtain the continuous strain-hardening over a wider stress range could contribute the superior ductility.

The hypothesis above is supported by several studies on both PBF-LB manufactured NiTi alloys [24] and conventionally manufactured NiTi alloys [56]. Xue et al. [24] reported that when the first critical yield stress decreased from 202 to 158 MPa, the tensile elongation increased from 9.7% to 15%. Lu et al. [22] found that when the first critical yield stress is lowered at 112 MPa, the elongation peaks at 7.4%. These results are consistent with the present study. For conventionally fabricated NiTi alloys [56], research also shows that when the first critical
yield stress is as low as 160 MPa, the maximum elongation can reach up to 60%, whereas when the first critical yield stress is around 350 MPa, the elongation is only 5%. Indeed, the first critical yield stress in the stress-strain curve corresponds to the elastic limit and is related to the activation of the stress-induced phase transformation including various types, which is linked to the initial microstructure and phase constitution [57]. For shape memory alloys, including β-type Ti-based alloys and NiTi alloys [58, 59], a higher volume fraction of the parental phase leads to a higher first critical yield stress. Therefore, adjusting the microstructure to B19' or R phase is necessary to enhance the ductility of bulk NiTi alloys. However, although reports show that a lower first critical stress and exceptional strain hardening are associated with multiple deformation mechanisms and hierarchical microstructures [20], the formation of Ti₄Ni₂O phase leads to an elongation of only 7.0% which seems unsatisfactory. This is because the superior ductility of shape memory alloys is also related to crystallographic texture, defects, and residual stress, which should all be considered to further suppress the involvement of plastic deformation during the stress-induced phase transformation stage.

With respect to the functional property which was specified to be the mechanical recoverable strain and recovery rate as shown in Figure 9B,C, the two values were also found to be lowest on the heat-treated sample "450°C/1 h" but the highest for PBF-LB fabricated samples whatever bulk or porous. This fact means that the PBF-LB sample mechanically recovers more once unloading from a certain applied strain compared with heat-treated "450°C/1 h" samples. Reasons can be attributed to the phase transformation, the stress-induced phase transformation from B2 phase to B19' phase could partially reverse back to B2 phase after releasing stress for the PBF-LB sample which contains more B2 phase. Further, it is interestingly to find that the variation curves in Figure 9B for samples undergoing the same heat treatment nearly overlap with each other, regardless of whether their macro structures are porous or bulk. Further, it can still be found that the variation curves in Figure 9B for TPMS-PBF-LB and Voronoi-PBF-LB samples highlighted in blue marked by different symbols can also coincide with each other. These findings imply that the most important factor affecting mechanical recovery property for porous NiTi alloy is the material itself which is deeply linked with the microstructure rather than the designed porosity or defects inside. The designed porosity cannot contribute to additional mechanical recoverable strain for the porous NiTi alloy, while the defects existing inside the porous alloy is not harmful to the recoverable properties.

From another viewpoint, the recovery rate shown in Figure 9C finally stabilizes at approximate 20% in the final deformation stage for porous samples. It can be assumed that our previous publication [59], the recovery rate for Ti-6Al-4V alloy fabricated with PBF-LB are stabilized at 22%–25%. This comparison between the different alloys thus proves the superelasticity was not obtained for the present NiTi alloy fabricated with PBF-LB method. Nevertheless, the recoverable strain can be compared with other reference which claimed the observation of superelasticity manifested with recoverable strain of 3% at the applied strain of 5% [30, 33]. The reason why the superelasticity is not

obtained probably due to the involvement of plasticity during the SIM transformation stages and will be further confirmed. To be concluded, for the bulk NiTi alloy, the properties in terms of elongation, the first yield critical stress and recovery rate all depend on the initial microstructure and subsequent stress-induced phase transformation. For the porous NiTi alloy, the introduction of the designed porosity for NiTi alloy cannot enhance the recovery property which nearly keeps the same with the value measured in bulk samples.

## 5 | Conclusions

Because of the instability of molten pools and the difficulty in precisely controlling the heating input, metallurgical defects and sintered powders are inevitable for PBF-LB fabricated porous NiTi alloy. In both TPMS and Voronoi structures, a significant increase in the thickness of the horizontal walls or struts was observed with obvious fluctuations at the bottom contours. The limited epitaxial growth of grains due to heat conduction constrains and weakened crystallographic texture were observed in porous NiTi alloys compared with the bulk samples. These structural features account for the varied ultimate compressive strength and Young's modulus for porous NiTi alloy. TEM microstructure observation on porous and bulk samples show similarity. The post heat treatment of LTA-450°C/1 h followed by water quenching enhance the tensile elongation to a maximum value of 18.8% which is linked with the highest volume fraction of R phase. This is consistent with the lowest critical yield stress and minimum mechanical recoverable strain compared with other LTA samples. The comparison between the TPMS and Voronoi samples revealed that the spatial configuration gives a limited influence on the recovery property for the porous NiTi alloy. Thus, the introduction of designed porosity for porous NiTi alloy cannot additionally enhance the mechanical recoverable strain. A continuous increase of mechanical recoverable strain is observed to peak at 4.2% with increasing the compressive applied strain till 21.3%.

## Author Contributions

**Wen-Hao Zeng:** data curation, formal analysis, methodology, validation, writing – original draft. **Yang Yang:** methodology, project administration, supervision, writing – review and editing. **Xing-Rong Chu:** methodology, software, supervision. **Yang Zou:** data curation. **Ren-Zhou Deng:** methodology, validation. **Ke Huang:** data curation, formal analysis. **Philippe Castany:** supervision, writing – review and editing. **Thierry Gloriant:** data curation, formal analysis, writing – review and editing.

## Acknowledgments

The work was financially supported by the Special Project for Research and Development in Key areas of Guangdong Province (Grant No. 2023B0909020004), General Program of National Natural Science Foundation of China (Grant No. 52475328) and project of Innovation Research Team in Zhongshan (Grant No. CXTD2023006).

## Conflicts of Interest

The authors declare no conflicts of interest.
# Data Availability Statement

The data that support the findings of this study are available from the
corresponding author upon reasonable request.

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**Supporting Information**

Additional supporting information can be found online in the Supporting Information section.

**Supporting Information S1:** rar270053-sup-0001-suppl-data.doc.

**Figure S1:** Observation of defects on mechanically polished bulk samples with varied laser power and scanning speed at constant layer thickness of 30 µm and hatch distance of 80 µm: (A–E) 50 W, (F–J) 100 W, (K–O) 150 W. **Figure S2:** TEM microstructure observation on LPBF fabricated NiTi samples followed by different heat treatments presenting typical microstructure with BFI, SAEDP and DFI: (A) 350°C/1 h, (B) 350°C/3 h, (C) 350°C/18 h, (D) 450°C/1 h. **Figure S3:** Tensile test curves for bulk samples: (A) PBF-LB, (B) 350°C/1 h, (C) 350°C/3 h, (D) 350°C/18 h, (E) 450°C/1 h, (F) comparison on the measured critical stress.