Modification of Mechanical Properties of Ti–6Al–4V Using L-PBF for Anatomical Plates

Abstract

In this research, as-built Ti–6Al–4V anatomical plates were successfully fabricated using laser powder bed fusion (LPBF). This study thoroughly examines the microstructural evolution and its role in enhancing the mechanical properties of clavicle bone plates under sub-β-transus heat treatment for medical application. Scanning electron microscope (SEM) images of the as-built specimens reveal a dense formation of a hard α’ hcp martensite structure, which decomposes during annealing at 650 °C and ultimately transforms into an α + β lamellar structure at 950 °C. Additionally, coarse grains resulting from recrystallization and reduced dislocation density were observed through electron backscatter diffraction (EBSD) following heat treatment. Due to these microstructural evolutions, the desired mechanical properties of as-built Ti64 parts for surgical applications were achieved. Heat treatment of the anatomical plates at 950 °C demonstrated an excellent strength–ductility synergy under tensile deformation and the highest energy absorption capability under bending deformation, indicating sufficient durability for medical implantation applications.

1. Introduction

Surgical implantations constitute a critical domain within the medical industry [1]. For instance, the annual number of hip and knee replacement surgeries involving metal endoprosthesis implantation in the United Kingdom has reached 117,000 [2]. However, implant components manufactured using traditional methods often struggle to achieve complex geometries, potentially compromising essential properties and leading to increased surgical risks, infections, and revision procedures [3]. Morselli et al. [4] conducted a comprehensive investigation of complications of cranioplasty with titanium material, reporting a surgical revision rate of 12% in 2019.

Additive manufacturing (AM) has emerged as a prominent technology in various industries over recent decades, offering distinct advantages over conventional manufacturing methods. These benefits include enhanced precision in producing complex components, increased customization flexibility, reduced lead times, and optimized material utilization [5,6,7]. In addition, AM offers the advantage of producing roughened surfaces in conjunction with complex designs, which proves beneficial in orthopedic surgical transplantations. The roughened surface of AM-produced components generates friction against human bone, facilitating osseointegration—the process by which bone tissue integrates with the metal implant [2]. The medical and healthcare industry, in particular, has witnessed significant advancements through AM, with its potential to efficiently fabricate biomaterial components for diverse applications [8,9]. The impact of AM on the medical industry is multifaceted, encompassing the production of surgical instruments, models for preoperative planning, anatomical implants, prosthetics, and drug delivery systems [10]. Market projections indicate a substantial growth trajectory for AM in healthcare services, with an anticipated increase from 0.713 billion USD in 2016 to 3.5 billion USD by 2025. Moreover, AM has proven invaluable in medical education, providing students with enhanced opportunities to study human anatomy through high-fidelity 3D models derived from medical imaging data [11,12].

Titanium (Ti) and its alloys are predominantly utilized in biomedical applications due to their excellent biocompatibility and superior osteoconductive properties, which are further enhanced by the promotion of porous structures [13]. Vrancken et al. [14] demonstrated that additively manufactured Ti–6Al–4V (referred as Ti64) alloy exhibits exceptional tensile strength, exceeding 1 GPa compared to as-cast and as-forged conditions. Zou et al. [15] reported that laser-based powder bed fusion (LPBF), a prominent AM technology, can produce high-quality components with reduced microporosity and notable structural integrity through precise regulation of process parameters such as laser power and scanning speed.

Typically, Ti–6Al–4V fabricated via LPBF presents a columnar martensitic structure with a lower likelihood of β phase presence due to the rapid solidification of the melt pool during processing [16]. However, this microstructure may limit the required toughness and ductility of additively manufactured Ti64 components. Additionally, the higher residual stress in manufactured parts restricts optimal strength–ductility combinations, posing significant challenges for the application of the LPBF process to Ti64. Consequently, annealing treatments are essential to achieve microstructural homogeneity and mechanical stability in AM-produced Ti64 parts [17,18]. Zhang et al. [19] investigated microstructural changes between as-built and heat-treated Ti64. The material exhibited an α + β lamellar structure after annealing at 950 °C, while annealing above 1050 °C resulted in the formation of equiaxed β grains with α boundaries. Although numerous studies [14,15,16,17,20] have focused on developing the properties of additively manufactured Ti alloys, most have been oriented toward automotive and aerospace applications, which differ significantly from medical applications. While AM is rapidly expanding in the dental industry, its application in joint-specific transplant surgeries remains in its infancy due to various regulatory challenges [2,9]. This underscores the need for further research and supporting data, particularly regarding property optimization of 3D-printed medical implants.

In this study, an anatomical plate composed of Ti–6Al–4V was fabricated by laser powder bed fusion, with the input digital geometry derived from a patient-specific clavicle bone model specifically designed for orthopedic transplantation. The results suggest an optimal annealing temperature to enhance the mechanical properties of LPBF Ti64, being suitable for medical implant applications. The evolution of the microstructure was evaluated by analyzing phase evolution, dislocation density, grain morphology, and kernel average misorientation. Additionally, tensile, bending, and surface hardness tests were conducted to investigate the effects of annealing temperature on these mechanical properties, which are crucial for ensuring the longevity of implants within the human body.

2. Materials and Methods

Commercially available gas-atomized Ti–6Al–4V (referred to as Ti64) raw powder, supplied by SLM Solutions, was used as the feedstock for LPBF in this study. The average particle size is ~41 µm, with a particle size distribution of D10, D50, and D90 values of 27.7 µm, 44.0 µm, and 66.6 µm, respectively. The chemical compositions of the powder material are given in Table 1.

Parts were additively manufactured in the vertical orientation (Z-build) using an IPG fiber single-laser LPBF (SLM 280, SLM Solutions, Germany) facility (Figure 1a) in an argon atmosphere while keeping the oxygen content below 25 ppm. As shown in Figure 1b, the three types of parts—clavicle bone plate, tensile specimen, and rectangular cuboid—were fabricated using the optimized LPBF parameters (Table 2) through adjusting laser power and scan speed with individual relative density measured using Archimedes method (ASTM B962-17) and tensile testing with a typical script scanning strategy. The clavicle bone plate was subjected to 4-point bending tests to investigate its mechanical stability. Cylindrical tensile specimens with a gauge diameter and length of 6 mm and 24 mm were manufactured to investigate the tensile mechanical properties of the as-built Ti64. To analyze detailed microstructure and surface hardness, the rectangular cube specimen with a dimension of 20 × 10 × 10 mm³ was utilized. The specimens showed a relative density above 99.5% using optimized process parameters for LPBF established in this work. The energy density (E) was calculated to be 52 J/mm³ using the equation E = P/vht, where P is the laser power, v is the scan speed, h is the hatch distance, and t is the layer thickness.

Using a high-precision wire cut EDM machine (MV2400S, Mitsubishi Electric, Tokyo, Japan), the specimens were separated from the build plate. All the three types of specimens were subjected to post-heat treatment with various annealing temperatures—650 °C, 800 °C, and 950 °C for 2 h—followed by furnace cooling in the vacuum furnace (HFL160, Nabertherm, Lilienthal, Germany) with a continuous Ar gas flow. According to the annealing temperature, the heat-treated specimens were denoted by 650 HT, 800 HT, 950 HT, respectively. X-ray computed 3D tomography (XCT, XTH320, Nikon, Tokyo, Japan) was performed to identify the shape, size, and volume of the pore presence inside both as-built and heat-treated specimens. Prior to the analysis, the specimens were rotated to the Z axis with 2000 projections to follow the build direction. The operating parameters for CT analysis were set to 180 kV voltage, 300 µA current, 0.038 mm voxel size, and 267 ms exposure time. The specimen surface was then polished up to SiC paper with 1000 µm grit size to obtain a smooth oxidation-free surface for X-ray diffraction (XRD, Miniflex600, Rigaku Corporation, Tokyo, Japan) analysis of the presence of phases and their evolutions through heat treatment. For XRD analysis, a copper (Cu) radiation X-ray source, 40 mA current, 40 kV voltage with a scan range between 30°–80°, step size 0.02°, and scan speed 2°/min were followed. The surfaces were further polished up to 1 µm diamond paste suspension and etched with Kroll’s reagent for field emission scanning electron microscopy (FE-SEM, JSM7600F, Jeol Ltd., Tokyo, Japan) and energy dispersive spectroscopy (EDS, X-Max50, HORIBA Ltd., Kyoto, Japan) characterizations. For electron backscattered diffraction (EBSD) analysis, the cross-sections of the specimen were further electropolished at −30 °C using Struers LectroPol-5 equipment (a voltage of 30 V, flow rate of 15 for 60 s in 95% methyl alcohol + 5% Perchloric acid). EBSD maps, including inverse pole Figure (IPF) and kernel average misorientation (KAM), were developed using FE-SEM (SU5000, HITACHI, Tokyo, Japan) equipped with an EBSD detector (Symmetry S3, Oxford instruments, Abingdon-on-Thames, UK). A grain tolerance angle of 5° was used for grain identification.

Surface microhardness mapping was performed with a load of 1N for a dwell time of 10 s using an automated Vickers microhardness tester (AMT-X7FS series, Matsuzawa, Akita, Japan). A total of 50 points with a spatial distance of 0.5 mm on the middle of the polished surface through the build direction was considered for specimens in as-built and heat-treated conditions. A universal testing machine (DTU900-MH, Daekyung, Incheon, Republic of Korea) was employed for uniaxial tensile tests with cylindrical specimens at a 1 mm/min constant strain rate until fracture. As illustrated in Figure 1c, the 4-point bending tests were performed on the fabricated clavicle bone plates using a tensile and compression tester (TW-D2000, Taewontech, Bucheon, Republic of Korea) with a constant speed of 5 mm/min to evaluate the bending rigidity and energy absorption capability. The clavicle bone plate was placed (Figure 1c) between the 4-point bend test system so that the compressive loading acted perpendicular to the build direction. The fracture surfaces of the broken tensile and bend test specimens were then placed under SEM to investigate the failure mechanism for all the conditions.

3. Results

3.1. X-Ray Computed Tomography

Figure 2a shows the presence and distribution of pores within the as-built Ti–6Al–4V, obtained through 3D micro-XCT analysis. The pores are uniformly distributed on the XY and YZ planes, and the cube specimen contains no noticeable large disk-shaped defects. The number of pores on the XZ plane appears to be fewer than that on the other two planes. These pores primarily develop due to the partial melting of gas-atomized powder particles, the presence of moisture, and surface fluctuations [21]. Additionally, the protective gas stream could become trapped in the melt pool during LPBF, leading to the formation of gas pores [22].

Regarding the pore volume distribution (Figure 2b), there are no pores with a volume greater than 0.004 mm³. For as-built, the relative density was measured to be 99.79% by the Archimedes method; on the other hand, the pore density was measured to be 0.60% (pore volume of 10.85 mm³ with a part volume of 1792 mm³). The pore density values of the annealing conditions were 0.47%, 0.55%, and 0.65% for 650 HT, 800 HT, and 950 HT, respectively. Since it is hard, by only thermal treatment is it possible to remove or alleviate the pores inside the specimen [21,23], so the pore characteristics were unchanged after thermal annealing. The analysis result of these micro-sized pores across a considerable range, coupled with the good relative density of the as-built Ti64 alloy, further substantiates the printing quality and contributes to the microstructural characterization related to the development of mechanical properties discussed in the following sections.

3.2. X-Ray Diffraction Analysis

X-ray peaks (Figure 3a) of the as-built and heat-treated specimens illustrate the presence of various phase constituents alongside their respective ranges of diffraction angles (2θ). Magnified images of the marked locations (2θ = 34–42° and 70–80°) reveal phase evolutions after thermal annealing, along with their respective crystallographic orientation planes, as shown in Figure 3b,c. It is confirmed that only α/α’ hexagonal close-packed (hcp) martensitic phases are predominantly present in the as-built Ti64. The α/α’ phase peak intensity reached its maximum at the ($101\overline{1}$) plane at 2θ = 40°. Since α and α’ represent nearly identical crystallographic structures and lattice parameters, distinguishing them through XRD analysis is challenging; thus, they are characterized as α/α’ within the peak profiles [15]. The martensitic phase in the as-built specimen exhibited generally solely α’, which is a distorted form of the α-stable martensite phase formed due to rapid solidification during LPBF. In addition to the highest peak intensity, α/α’ phase components also appeared at the $\left(101\overline{0}\right)$, $\left(0002\right)$, $\left(112\overline{0}\right)$, $\left(112\overline{2}\right)$, and $\left(202\overline{1}\right)$ hcp planes within the 2θ range from 35° to 78°.

After thermal annealing treatment, the β (110) phases with the body-centered cubic (BCC) crystal structure appeared at 2θ = ~40° and ~75° alongside other dominant α/α’ phases. The peak broadening and separation observed in the heat-treated conditions suggest the formation of an equilibrium α phase in conjunction with a metastable α′ phase. In addition to the formation of the α phase, the emergence of β phases signifies the initiation of martensitic phase decomposition during the thermal annealing process. According to the phase transformation study conducted through in situ XRD and TEM analysis, the α′ martensite begins to decompose above 400 °C and is completely transformed into the α + β equilibrium phase at 700 °C [24]. XRD peak profiles show relatively broader peaks for the as-built specimens compared to those for the heat-treated conditions. Following thermal annealing, the peaks became narrower with sharper tips, likely due to the removal of residual stress and the decomposition of martensitic phase components, which can reduce the amount of elastic distortion within the specimens [15,18]. Since the residual stress significantly related with the dislocation density, the dislocation density was quantified from FWHM values of peaks, as shown in Figure 4. Due to the rapid cooling process, the highest dislocation density can be observed in the as-built condition [25], which decreased after heat treatment. It can be concluded that a significant reduction in internal dislocations and distortions introduced by the LPBF process can be effectively achieved through heat treatment at temperatures above 800 °C.

Furthermore, the peak intensity of the β phase was higher at 950 HT than at 800 HT, indicating an increased presence of the β phase due to substantial decomposition of α’ [8]. Thus, it is evident that the as-built Ti64 underwent significant phase changes during thermal annealing, highlighting the need for more precise microstructural characterizations and detailed analyses.

3.3. Electron Microscopy Analysis

Figure 5 presents SEM images of the as-built, 650 °C heat-treated (650 HT), 800 °C heat-treated (800 HT), and 950 °C heat-treated (950 HT) conditions along the Z-build direction. The microstructure of the as-built condition reveals a predominant α′ martensitic structure along with white nanosized particles, which are expected to be V-enriched β precipitates formed during the LPBF process due to the low solubility of vanadium (V) in the α phase within the α’ lath structure [26,27]. Given the limited resolution of X-ray scans for detecting these nanosized particles, the XRD analysis did not exhibit a peak corresponding to the β phase (Figure 3). It is observed that the number of nanosized particles decreased with increasing annealing temperature, likely due to the diffusion of V into the matrix during the annealing treatment. The formation of the α′ martensitic structure primarily resulted from the diffusionless transformation of β phases during the subsequent rapid solidification of molten Ti64 powder deposited in layers during the LPBF process [18]. Furthermore, it was found that annealing at 650 °C was insufficient for the formation of β lamellar phases, although decomposition of the α′ phase occurred.

In comparison to the as-built condition, the phase morphology underwent significant changes in the heat-treated conditions. The α′ martensitic structure of as-built Ti64 normally transforms from α′ to α phase due to the reduction in supersaturation in α′, which results from the nucleation of the β phase at the boundaries of the α′ needles. This process induces the diffusion of vanadium from α′ to the β phase during the annealing treatment. Xu et al. [28] reported that this martensite decomposition can occur at temperatures as low as 400 °C. Given that the annealing temperatures employed in this study exceeded 650 °C, it is thus expected that martensite decomposition occurred.

In the 650 HT condition, the microstructure exhibited incompletely decomposed martensite, while at higher annealing temperature of 800 °C, the α′ martensite fully decomposed into an α lamellar structure, accompanied by the emergence of β phases (Figure 5b,c). As the temperature is further increased to 950 °C, the α lamellar structure showed significant coarsening and transitions to a globular α structure compared to the finer lamellar structure observed at 800 HT, which is a phenomenon attributed to static recrystallization at higher annealing temperatures [15,29]. Additionally, substantial growth of the β phases was observed, along with rod-shaped β phase particles located near the junction of the primary α phase, as the annealing temperature approached the β-transus temperature of Ti64 [30]. According to the Hall–Petch equation, this coarsening of the α lath is known to induce a reduction in the strength of the α + β bimodal structure in the Ti64 alloy [8,19,31].

To further investigate the grain morphologies, EBSD analyses, including Inverse Pole Figure (IPF) and Kernel Average Misorientation (KAM) maps, are presented in Figure 6 and Figure 7. The locations of the EBSD analysis correspond to the SEM observations on the surface along the Z-build direction. For the as-built Ti64, the α′ grains exhibited fine columnar morphologies (with a minor axis width of 1.38 µm) and random crystallographic orientations (hcp) (Figure 6a) situated within prior β grain boundaries, as indicated by the black dashed line parallel to the Z direction.

The KAM maps (Figure 7) indicate a higher average KAM value of 0.824 in the as-built condition, reflecting an overall distorted lattice structure. The primary cause of the elevated KAM values at the as-built samples is rapid overlapping of solidified layers, which generates residual stress [18]. Following the annealing process, the average KAM value decreased from 0.823 in the 650 HT condition to 0.419 in the 950 HT condition, indicating a relief of residual stress and a reduction in dislocation density within the as-built Ti64 alloy. Notably, the KAM value was not significantly higher at the prior β grain boundaries, suggesting an independence between local distortion and the prior β grain boundary. Since the KAM value serves as an indicator of the geometrically necessary dislocation (GND) density in a localized region [32,33], the total dislocation density calculated through XRD analysis exhibited a similar trend, with dislocation density increasing as the annealing temperature rises.

It was observed that the KAM values were uniformly distributed, except for several thicker lath structures in the as-built and 650 HT conditions, which exhibited generally higher average KAM values. In contrast, in the 800 HT and 950 HT conditions, the KAM values localized at the grain boundaries rather than within the grain interiors, resulting in lower average KAM values.

The KAM value is notably higher at the low-angle grain boundaries (LAGBs) in the heat-treated specimens. The formation of LAGBs can occur due to recrystallization, leading to grain growth during the annealing treatment. This phenomenon suggests that the redistribution and reduction of dislocations may take place under heat-treated conditions. Since LAGBs can facilitate the dispersion of deformation in materials, this is likely connected to the observed improvements in ductility [34]. Consequently, it can be inferred that the sub-β-transus heat treatment plays a significant role in relieving residual thermal stress, resulting in microstructural stability and homogenization [19,21]. This process is crucial for the effective development of properties in as-built Ti64 parts for biomedical applications.

3.4. Surface Microhardness

Microhardness measurements were performed at the center of the surface along the building direction of the samples. Figure 8 exhibits the microhardness map, showing the reduction of average surface microhardness values from as-built (384.6 ± 25.3 Hv), 650 HT (378.5 ± 11.3 Hv), 800 HT (355.2 ± 3.6 Hv), and 950 HT (328.1 ± 18.5 Hv), respectively. However, the reduction in surface microhardness was under the limit of consideration (>300 Hv) for Ti64 implants and approximately similar to conventionally available Ti64 [35,36].

Surface microhardness is an essential surface property of an implant, which helps to protect the parts from wear inside the human body and requires proper optimization. The generation of thermo-residual stress inside the as-built parts impacts the surface hardness and makes it significantly harden, surpassing the required level [37,38]. Thus, the decrease in hardness values after thermal annealing treatment is mainly due to the removal of residual stress. In addition, the occurrence of soft β phases after annealing also helps optimize surface hardness and level it to the need [5].

3.5. Tensile Behavior

As shown in Figure 9, the uniaxial tensile tests for heat-treated Ti64 specimens along the Z-build direction indicate a considerable increase in ductility with losing much strength (Table 3) compared to the as-built specimen. The martensite structure and its higher dislocation density (Figure 7a) results in a higher strength at the lowest strain hardening, leading to a reduction of ductility in terms of total elongation of the as-built specimen [25,39]. After thermal annealing, strain hardening develops, improving ductility and helping to build the required level of flexibility within the additively manufactured Ti64 for biomedical applications. In addition, the ductility increased further when the annealing temperature rose from 800 °C to 950 °C.

As shown in Figure 10, the SEM fractography images of the fractured tensile specimens indicate a transition from a cleavage fracture mode, characterized by facets and river patterns in the as-built condition, to a fully ductile failure in the heat-treated specimens. The 650 HT specimen exhibited a curved surface with facets, microcracks, and tearing edges, suggesting a quasi-cleavage fracture surface. The microstructural evolution following sub-β-transus heat treatment, as discussed earlier in Section 3.3, plays a crucial role in the development of mechanical properties in the as-built Ti64 alloy. The formation of significantly finer and harder α’ distorted martensitic phase structures (Figure 5a) primarily generates a substantial local strain gradient within the as-built specimen, resulting in the highest tensile strength while reducing the necessary ductility [16]. The presence of hard phases with fewer β phase particles further contributes to a quasi-cleavage mixed mode of tensile fracture (Figure 10a,e), where most of the area exhibits cleavage facets with only a few dimples [30].

As the annealing temperature increased, dimpled fractures were observed, indicating ductile failure compared to the as-built and 650 HT conditions. The required strength–ductility balance in the additively manufactured Ti64 alloy was achieved through the formation of stable α + β phases, accompanied by a reduction in dislocation density, as shown in Figure 6 and Figure 7. This transformation resulted from the decomposition of α’ martensite and static recrystallization within the microstructure during the heat treatment [8]. Additionally, the 950 HT specimen demonstrated superior strength–ductility synergy compared to the 800 HT condition, attributed to α grain coarsening and a higher volume fraction of β phases leading to a fully lamellar structure [19]. As the softer β phases and phase particles increased with rising annealing temperatures, the 950 HT sample exhibited complete ductile fracture (Figure 10d,h) characterized by fine micro-voids and dimples. In contrast, the 800 HT condition (Figure 10c,g) still displayed some areas of facets within the fractography.

3.6. Bending Behavior

The implantation of 3D-printed components in the human body is subjected to continuous complex loading such as tensile and compressive stress after a fixation process with a bone. It is also essential that the printed components exhibit a requisite degree of flexibility, as the surgeon often manipulates them to achieve accurate fixation at the targeted joint for successful implantation through the clinical procedure known as osteosynthesis. Consequently, it is imperative to evaluate the tensile properties, as well as the bending properties of these implants through the 4-point bending test, to gain a comprehensive understanding of their flexibility prior to insertion. Figure 11 presents typical load–displacement curves obtained from the 4-point bending test. The bending curve exhibits three main stages: linear elastic deformation, a gradual increase in loading during elastic-plastic deformation, and a plastic deformation region characterized by damage progression leading to bending fracture [40,41]. Overall, the clavicle bone plate fabricated by powder bed fusion demonstrated abrupt failure upon reaching the maximum load, without exhibiting a plateau stage. The as-built specimen revealed a rapid bending failure and a limited plastic deformation region, whereas a more stable and a gradual bending behavior developed after heat treatment, particularly as the annealing temperature increases. Detailed values of bending mechanical properties, including the maximum load, rigidity, and elastic and plastic energy absorption, are provided in Table 4. The energy absorption, which can be categorized into elastic and plastic components, was calculated for each condition, including both as-built and heat treated, by integrating the area under the experimental load–displacement curve [42,43,44].

As shown in Figure 12a, a maximum load occurred in the 800 HT condition, while rigidity was the greatest in the as-built specimen. The bending rigidity, which indicates resistance to the bending deformation, correlates well with the higher tensile strength observed in as-built specimens (Table 3). This is attributed to the formation of hard phases (α’) visible in the microstructure, which makes it challenging to bend the as-built anatomical plates and limits their flexibility [45]. As the tensile strength decreased with increasing annealing temperature, a corresponding reduction in rigidity was observed. In terms of the maximum load, the values were similar across all conditions. This can be affected by improved durability, which may be beneficial for successful osteosynthesis.

Although the as-built specimens demonstrated higher rigidity, their lower energy absorption capacity promoted premature bending failure. In contrast, the energy absorption capability increased significantly after annealing, as shown in Figure 12b and Table 4. Therefore, the much higher capability of plastic energy absorption in the heat-treated specimens can delay the bending failure of the clavicle bone plate. Notably, both the elastic and plastic energy absorption increased in the 650 HT and 800 HT conditions, while only the plastic energy absorption showed a significant increase in the 950 HT case. After annealing at 950 °C, the as-built parts exhibited maximum flexibility characterized by the lowest bending rigidity of 778.38 N/mm and the highest energy absorption capability of 8830.52 mJ among all three conditions. However, considering the critical importance of elastic deformation in medical implants, the heat treatment temperature of 800 °C may represent the most optimal condition due to its superior elastic energy absorption properties in this work.

Photographic images of the broken clavicle bone plates after bending tests reveal the fracture locations (Figure 13a), which changed for the as-built to annealed conditions. The as-built clavicle bone plates exhibited fractures originating from the surface characterized by blunt edges, which indicate an inevitable failure with minimal deflection due to lower durability [45]. In contrast, both heat-treated specimens fractured at the locations of the fixation holes. The introduction of holes in the clavicle bone plates leads to a reduction in the overall cross-section, resulting in stress localization under continuous compression on the top surface and tension on the bottom surface during bending tests, ultimately leading to failure [46,47].

SEM fractography images of the fractured areas show substantial variations in the dimensions of the fracture zones. The total width of the fracture zone in the as-built bone plate was measured at 2.3 mm (Figure 13b), which decreased to 2.1 mm (Figure 13c) for the 800 HT condition and 1.8 mm (Figure 13d) for the 950 HT condition. The narrower fracture zone widths in the annealed specimens indicate that both experienced maximum deflection and reached their highest flexibility before failure. The smallest measured fracture zone width following annealing treatment at 950 °C further suggests substantial plastic deformation accompanied by necking prior to failure, which is consistent with the highest plastic energy absorption observed.

Magnified SEM images captured from the marked regions in Figure 13(b1–d2) illustrate the details of the fracture morphologies of the bone plate specimens after the bending tests. In the as-built specimen, the fracture surface (Figure 13(b1)) displays two distinct regions: facet (brittle) and river-type (ductile) patterns, indicating a mixed failure mode. In contrast, both annealed conditions (Figure 13c,d) show no facets—only evidence of ductile fracture. Additionally, the appearance and characteristics of the dimples changed significantly from the as-built to the annealed conditions. The higher magnification image of the ductile failure regions in the as-built condition (Figure 13(b2)) reveals flat dimples with tearing edges (indicated by yellow arrows), resembling a tensile fracture surface. In contrast, the higher magnification fracture surface image of the 800 HT specimen (Figure 13(c2)) displays shear dimples with distinct edges (marked with red arrows), while the 950 HT specimen (Figure 13(d2)) exhibits equiaxed, dense dimples, indicating a greater degree of flexibility and ductility compared to the as-built condition. Furthermore, the failure mechanisms observed in the bending test specimens highlight the necessity of optimizing the clavicle bone plate design to enhance support in stress-localized areas (such as fixation holes), thereby improving overall flexural strength and deflection.

4. Conclusions

This research successfully demonstrated the microstructural changes associated with the development of mechanical properties in as-built Ti64 anatomical plates subjected to sub-β-transus heat treatment. The α’ martensitic structure within the as-built Ti64 began to decompose during heat treatment at 650 °C, ultimately forming a stable α + β lamellar structure as the annealing temperature reached 950 °C. The annealing treatment of the as-built specimens at 950 °C further developed coarse α + β structures and reduced dislocation density through static recrystallization resulting from prolonged exposure to high temperatures followed by slow cooling. Based on these microstructural evolutions, the optimization of mechanical properties—including strength–ductility synergy, bending strength, and flexural rigidity—was achieved to enhance suitability for surgical implantation. During the annealing process, recrystallization and the reduction in dislocation density within the microstructure, attributed to the removal of internal residual stress, resulted in a surface microhardness of approximately 328.1 Hv for the 950 HT condition. 20condition significantly increased its ductility to 49% compared to the as-built specimens. Furthermore, bending tests on the annealed clavicle bone plates demonstrated improved durability and energy absorption compared to the sudden failure observed in the as-built plates, thereby promising the desired specifications for medical implants. It is anticipated that this research will advance the utilization of customized implants in clinical practice while paving the way for a sustainable approach to additive manufacturing in the medical industry.

Author Contributions

Conceptualization, S.B., S.-H.L., and D.-H.K. (Dong-Hyun Kim 2); Formal Analysis, S.B. and S.-H.L.; Investigation, S.B., S.-H.L., J.-R.L., and D.-H.K. (Dong-Hyun Kim 1); Resources, D.-H.K. (Dong-Hyun Kim 2); Validation, J.-R.L., J.H.L., M.B., and D.-H.K. (Dong-Hyun Kim 2); Writing—original draft preparation, S.B. and S.-H.L.; Writing—review and editing, J.-R.L., J.H.L., M.B., and D.-H.K. (Dong-Hyun Kim 2); Visualization, S.B., S.-H.L., and J.-R.L.; Supervision, M.B. and D.-H.K. (Dong-Hyun Kim 2); Project administration, D.-H.K. (Dong-Hyun Kim 2); Funding acquisition, D.-H.K. (Dong-Hyun Kim 2) All authors have read and agreed to the published version of the manuscript.

Data Availability Statement

The datasets presented in this article are not readily available because the data are part of an ongoing study. Requests to access the datasets should be directed to corresponding author.

Acknowledgments

The authors acknowledge Sung-Tae Hong for providing testing facilities at the University of Ulsan, Republic of Korea.

Conflicts of Interest

The authors declare no conflicts of interest.

Footnotes

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Figures and Tables

Figure 1. Schematic representation of (a) LPBF process for Ti–6Al–4V (SEM image of raw powder particles at inset); (b) Z-built clavicle bone plate, cuboid, and tensile specimen; (c) a 4-point bending test for clavicle bone plate.

Figure 2. (a) 3D X-ray computed tomography analysis of pore characteristics within the as-built Ti-–6Al–4V cube specimen; (b) a bar chart showing the volume distribution and count of the interior pores.

Figure 3. (a) XRD analyses of as-built and heat-treated Ti64 alloy fabricated by LPBF. (b,c) Magnified images from the marked by purple and green dashed line reveal details of peak profiling and the presence of phase constituents. β-phase peak intensity increased at 950 HT and 800 HT thermal annealing conditions.

Figure 4. Dislocation density estimated by XRD peak analysis for as-built, 650 HT, 800 HT, and 950 HT.

Figure 5. SEM images through build direction (Z axis) of (a) as-built showing α’ martensitic structures; (b) 650 HT showing decomposition of martensitic α’ and stable α; (c) 800 HT showing α + β phase; and (d) 950 HT showing primary α, α + β lamella, and rod-shaped β. All the conditions present the nanosized β particles.

Figure 6. Inverse pole Figure (IPF) maps of the (a) as-built, (b) 650 HT, (c) 800 HT, and (d) 950 HT. The prior β grain boundaries are marked with black dashed lines.

Figure 7. Kernel average misorientation (KAM) maps of the (a) as-built, (b) 650 HT, (c) 800 HT, and (d) 950 HT. The prior β grain boundaries are indicated by white dashed lines.

Figure 8. Surface hardness maps sequentially for as-built, 650 HT, 800 HT, and 950 HT, respectively. Hardness decreased as the thermal annealing temperature increased to 950 HT condition.

Figure 9. Engineering stress–strain curves for as-built, 650 HT, 800 HT, and 950 HT specimens.

Figure 10. SEM fractography and magnified tensile fracture surface images for (a,e) as-built; (b,f) 650 HT; (c,g) 800 HT; (d,h) 950 HT. The fracture mode change from quasi-cleavage failure of as-built to ductile failure of 950 HT specimen.

Figure 11. Plots of 4-point bending behavior (force–displacement curve) for various annealing treatment conditions: as-built, 650 HT, 800 HT, and 950 HT.

Figure 12. 4-point bending properties of clavicle bone plate under different conditions: (a) maximum load and rigidity; (b) energy absorption capacity in the elastic and plastic regions for as-built and heat-treated specimens.

Figure 13. (a) Pictographs of the broken specimens after bending tests and corresponding SEM images reveal morphologies of fracture zone for (b–b2) as-built, (c–c2) 800 HT, (d–d2) 950 HT. The yellow and red arrows indicate the tearing edges and shear dimples, relatively.

Table 1

Chemical composition of Ti64 raw powder.

Range	Ti	Al	V	Fe	N	C	O
Min.	Balance	5.50	3.50
Max.		6.50	4.50	0.25	0.03	0.08	0.13

Table 2

LPBF process parameters for fabrication of tensile specimen, cubic, and clavicle bone plate for Ti–6Al–4V.

Laser Power(W)	Laser Scan Speed(mm/s)	Layer Thickness(mm)	Hatching Distance(mm)	Rotation Angle(Degree)
275	1450	0.03	0.12	67°

Table 3

Tensile properties for as-built, 650 HT, 800 HT, and 950 HT specimens.

Condition(Ti64)	Yield Stress(MPa)	Tensile Strength(MPa)	Elongation(%El)
As-built	1042 ± 15	1228 ± 6	9.0 ± 0.4
650 HT	1125 ± 2	1191 ± 6	10.2 ± 0.2
800 HT	1008 ± 6	1056 ± 5	15.0 ± 1.5
950 HT	923 ± 6	978 ± 3	18.8 ± 1.2

Table 4

Bending properties of clavicle bone plates subjected to the various annealing treatment conditions: as-built, 650 HT, 800 HT, and 950 HT.

Condition (Ti64)	Maximum Load (N)	Rigidity (N/mm)	Elastic Energy (mJ)	Plastic Energy (mJ)	Energy Absorption (mJ)
As-built	2407 ± 219	900 ± 61	1260 ± 208	3007 ± 337	4267 ± 301
650 HT	2426 ± 103	800 ± 8	1502 ± 57	4048 ± 315	5549 ± 372
800 HT	2463 ± 13	795 ± 20	1573 ±48	5937 ± 425	7510 ± 425
950 HT	2410 ± 63	778 ± 7	1184 ± 129	7647 ± 440	8831 ± 330

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