Laser additive manufacturing of TI-6AL-4V alloy at a high deposition-rate using a rectangular laser spot Open Access

Ti-6Al-4V is the most widely used titanium alloy, accounting for more than half of the titanium alloy market.¹ As an (α + β) alloy, it exhibits high strength, moderate density, and excellent corrosion resistance, making it highly suitable for applications in the aerospace, automotive, medical, chemical, and power generation industries.² This lightweight and strong alloy is particularly advantageous for highly stressed structures, where weight reduction is critical to improve efficiency.

Machining the titanium alloy Ti-6Al-4V, especially through milling or turning, is notably difficult due to its material characteristics.^1–3 To address this challenge, near-net-shape manufacturing offers a clear advantage by minimizing the material that needs to be removed to reach the final dimensions.^4,5 While additive manufacturing (AM) can achieve near-net shapes,^6,7 its slow deposition rate limits its practical use in industry. Meiners et al. emphasize that traditional manufacturing processes for titanium components, such as forging, generate significant material waste and require costly postprocessing. AM complements forging by facilitating near-net-shape production, which reduces both material use and tooling costs. AM also offers greater design flexibility, especially for complex geometries that forging cannot easily achieve, thus improving resource efficiency and minimizing production steps.

AM of titanium alloys, particularly in the aerospace industry, has attracted significant interest due to its potential for new design possibilities as well as its utility in repair and hybrid manufacturing.³ Components such as turbine blades, blade-integrated disks, and combustors are parts typically produced using AM. However, the standard deposition rate of Ti-6Al-4V using laser material deposition (LMD) is less than 1 kg/h,⁴ which is not cost effective for building complete components and is, thus, mainly used for repair purposes.

Near-net-shape technologies such as AM can improve resource efficiency compared to traditional manufacturing methods. However, AM production costs escalate as part size increases.⁵ One way to mitigate this is to increase the deposition rate to several kilograms per hour.

AM alone often fails to deliver the mechanical properties required for high-performance applications. Combining AM with hot forming presents a promising solution. This hybrid approach enhances near-net shape preforms through deformation, improving both strength and durability while reducing forming forces. As a result, it makes the production of high-performance components more efficient and cost effective.^3,6 The mechanical properties, such as strength, creep resistance, and ductility, are significantly influenced by the macro- and microstructure, particularly the morphology and distribution of the α-β phases.^1,7 The microstructure of AM samples produced at high deposition rates using a rectangular laser spot remains largely unexplored. This study also examines how laser spot size and cooling rates influence the formation of micro- and macrostructure, with a focus on β-grain size and α-lamellae width.

In the domain of laser-based AM applications, laser beam shape profiles typically exhibit variability. For instance, in laser-based cladding and coating, rectangular beams with top-hat or flat-hat distributions are frequently favored over Gaussian beams, owing to their uniform energy distribution.^8–14 In addition, experiments are underway to explore the use of diverse beam profiles in a range of AM techniques, including doughnut-shaped,¹⁵ convex¹⁶ distributions, and others. LMD has already demonstrated the utilization of beam distributions that deviate from Gaussian profiles.^17,18 In laser material processing (e.g., laser ablation), beams with Gaussian profiles are often less efficient than those with top-hat profiles. This inefficiency arises because Gaussian beams have a larger area where excess energy exceeds the threshold, while the energy in the outer regions falls below the necessary threshold (Fig. 1).¹⁹ 

A key factor influencing bonding defects in single tracks and layers is the geometry and intensity distribution of the laser spot. Figure 2(a) shows different intensity distributions across the spot width, and Fig. 2(b) shows the laser–material interaction time as a function of spot geometry. The circular shape of the laser spot results in shorter irradiation times at the edges compared to the center, resulting in less energy input at the boundaries. This problem is exacerbated by a Gaussian intensity distribution, where the power intensity decreases toward the edges of the spot [Fig. 2(a) (bottom)].

As energy input falls, the frequency of pores and bond defects increases.²⁰ In contrast, rectangular laser spots with a quasihomogeneous intensity distribution [Fig. 2(a) (top)] provide adjusted energy input over the entire spot area. In addition, the contact angle θ influences the bonding defects (Fig. 3).²¹ As θ increases, more material accumulates at the track edges, requiring additional energy to remelt overlapping tracks when creating layers or bodies.²¹ Gaussian intensity distributions and round laser spots exacerbate this problem since they provide insufficient energy for melting.

By using a rectangular laser spot with a quasihomogeneous intensity distribution, we can achieve a more adapted energy input (Fig. 4), reduce defects, and improve the quality of additively manufactured Ti-6Al-4V parts. This approach is expected to improve the overall efficiency and effectiveness of the LMD process and make it easier to produce high-quality, defect-free devices.

The objective of this work is to investigate high deposition-rate LMD produced Ti-6Al-4V samples and develop parameters for a rectangular laser spot with a quasihomogeneous intensity distribution. These parameters were used to produce single tracks, layers, and solid bodies. Requirements for the samples included the absence of cracks and bonding defects as well as minimal porosity. A parameter study was conducted to investigate the influence of spot size on microstructure development and correlating microstructure evolution with the cooling rate. A hot forming test on additively manufactured material was conducted using a deformation dilatometer to examine flow behavior under different strain rates and temperatures.

In laser material deposition using powder, metal powder is delivered by an inert gas, typically argon, into a focused laser beam that generates a melt pool on the surface. The carrier gas, in conjunction with an additional gas stream traversing the beam path, establishes an inert atmosphere surrounding the melt pool. As the laser beam progresses, the melt pool undergoes rapid cooling and solidification.

The inert gas supplied through the powder feed nozzle serves to shield the melt pool; however, the remainder of the part is exposed to the ambient atmosphere, which results in surface oxidation. To circumvent this issue, a global inert gas atmosphere was established utilizing an inert gas chamber filled with argon, which housed the processing head (Fig. 5). A 12 kW diode laser (LDF 12000-100, Laserline GmbH) was utilized as the beam source. A motorized zoom optic (OTZ, Laserline GmbH) was used to form a rectangular beam shape. A six-jet discreet powder feeding nozzle with a 23 mm stand-off was used. This specialized powder nozzle (HighNo 20-6, HD Sonderoptiken GmbH) was necessary to create a stabile powder-gas stream and withstand the back radiation and heat.

The Ti-6Al-4V powder material utilized in the experiments was provided by AP&C, a GE Additive Company. The powder material, produced via the electron induction-melting gas atomization process, exhibits spherical particle with a particle size distribution ranging between 45 and 63 μm. A carrier gas pressure of 4 bar was employed to facilitate powder mass flow rates more about 90 g/min. The Ti-6Al-4V substrate utilized in the LMD process, provided by ATI Specialty Materials, and had dimensions of 250 × 100 × 10 mm³. The chemical composition of the powder and the substrate is provided in Table I for reference. Prior to utilization, the substrate plates were subjected to sandblasting, which served to clean the surface and enhance laser radiation absorption.

The parameters developed in this study are shown in Table II.

Several Zeiss scanning electron microscopes (SEMs 1540XB, ULTRA55, and LEO1550) equipped with Gemini columns were used to achieve nanometer resolution and to perform semiquantitative elemental (EDX) and texture (EBSD) analyses. These SEMs were coupled with Oxford Instruments EDX detectors (UltimMax170, X-MaxN150, and INCA X-act), Oxford Tetra BSE detectors, and Nordlys EBSD cameras.

The porosity of the cross sections of all deposited tracks was measured using quantitative optical metallography, with an Olympus optical metallography microscope and Stream Motion. The porosity was determined using the ratio of the areas of all pores present in the cross section.

The ImageIR^© 7350 infrared camera system (IR camera) from InfraTec GmbH was employed to capture thermal images of the LMD process. The temperature field of a single deposited track was obtained using the IR camera, which facilitated the analysis of the molten pool state and the measurement of the sample temperature. The IR camera was set at an inclination angle of 30° and operated at a frequency of 100 Hz. A temperature curve was generated over the entire track length to evaluate the recordings. The cooling rate was determined from these data in conjunction with the feed rate. The thermographically calibrated profile was designed to accommodate temperatures within the range of 400–1700 °C, with measurements outside this range deemed to be of limited utility.

Compression tests were performed to determine the hot deformation behavior of the AM material at different temperature. Cylindrical samples with a diameter of 5 mm and a length of 8 mm isothermally deformed at temperatures of 850, 900, and 950 °C and at a strain rates of 0.001, 0.01, 0.1, and 1 per second.

Single layers were fabricated using the following process parameters, with cross sections exhibiting no bonding defects: The deposition rate was found to be greater than 5 kg/h, the dilution was below 30%, and the aspect ratio (track height to track width) was about 0.25 (Fig. 6).

The empirically determined factor b was found to be 0.94. Single layers were produced using a meandering strategy, with each track offset and a 180° rotation after each pass. The average layer height was 2.37 mm, with an increase from 2.23 to 2.51 mm across the width. This resulted in a part density of 99.99% ± 0.0025 without defects (Fig. 8).

The average cooling rates are differentiated according to the melting temperature (T_m) and the β-transus temperature (T_β), which is defined as the temperature at which the material transitions from the β-phase to the α-phase. The average cooling rate is then calculated for temperatures between T_β and 600 °C, and the overall cooling rate is determined. Figure 9 illustrates the average cooling rates of single tracks with spot sizes of 5 × 5, 6 × 6, and 7 × 7 mm².

For all spot sizes, the cooling rates above T_β are higher than those below T_β. An increase in spot size results in a reduction in the overall cooling rate. A spot size of 5 × 5 mm² exhibits a cooling rate of 466.14 K/s, while a 7 × 7 mm² spot size exhibits a cooling rate of 210.93 K/s.

As evidenced by studies conducted by Yu et al., smaller spot sizes are associated with higher cooling rates reaching up to 1066 K/s.²² Reduced cooling rates with larger laser spots can be attributed to the increased melt pool size and prolonged laser–material interaction time. The introduction of larger spots entails the incorporation of a greater quantity of material at a reduced feed rate, which consequently expands the melt pool volume and heat content. Due to the low thermal conductivity of Ti-6Al-4V (7 W/mK), the heat dissipates (conduction), and the material solidifies more slowly.

The process parameters from Table II were employed to ascertain a density of 99.98% ± 0.0195%. Cross sections of single tracks (spot sizes 5 × 5 to 7 × 7 mm²) exhibited a density of 99.99% ± 0.0115%, while single layers an average density of 99.99% ± 0.0025%. For volumetric builds, the average density was 99.99% ± 0.0057%. Despite the high densities observed, porosity and bonding defects can still occur when powder accumulates at the start and end of the deposition paths and sticks to the surface (Fig. 10).

A rectangular laser spot with a quasihomogeneous intensity distribution likely reduced bonding defects. As Tammas et al. have observed, a uniform energy distribution across the track width could reduce the likelihood of bonding defects.²³ Isolated round pores were identified as gas pores introduced by shielding or carrier gas that could not escape the melt pool, possibly due to rapid solidification and Marangoni flow effects.² 

All samples exhibited a macrostructure of columnar β-grains growing epitaxially from the substrate surface. These grains spanned the entire cross section, with lengths exceeding 30 mm (Fig. 11).

The average grain width in single tracks was determined to be 334.67 ± 143.56 μm². The average grain width in single layers was found to be 428.33 ± 182.55 μm², with a maximum of 1028.78 μm. Increasing the laser spot size resulted in larger β-grains. For a 5 × 5 mm² spot size, the β-grain width was 273.72 μm, while a 7 × 7 mm² spot size increased it to 468.93 μm (Fig. 12). In solid bodies, the average grain width was 1748.34 ± 1169.40 μm², with a maximum of 5420.3 μm.

The larger spot sizes resulted in lower cooling rates due to the increased melt pool volume and longer laser–material interaction time. For instance, the cooling rates decreased from 483.47 K/s at 5 × 5 mm² to 223 K/s at 7 × 7 mm² (Fig. 9). This allowed for a greater period of time for grain growth, resulting in larger grains in both single tracks and solid bodies. In solid bodies, the cooling rates in the center dropped to as low as 4 K/s, whereas single tracks exhibited cooling rates between 200 and 400 K/s. The epitaxial growth of β-grains continued through multiple layers, with heat dissipation primarily through the substrate, promoting columnar growth. The β-grain structure, reconstructed using EBSD measurements, demonstrated the presence of the $⟨ 001 ⟩$ orientation in all samples, corresponding to the build direction N_z (Fig. 13).

The β-grains exhibit a preferred orientation parallel to the build direction, indicating a significant degree of anisotropy. Additionally, the β-grains exhibit considerable width, which contributes to the overall structural characteristics observed in the build. This anisotropic grain structure is more pronounced in single layers than in single tracks. Grain growth occurs during the solidification of previously deposited AM layers and reflects the crystallographic nature of the AM structure. It is driven by partial or complete remelting of the underlying layer. The direction of growth at the solid-liquid interface is determined by the direction of maximum heat flow.

At temperatures up to 980 °C (β-transus temperature), the microstructure transitions completely to the body-centered cubic β-phase, where atoms are relatively mobile, allowing a uniform distribution of alloying elements. The size of the β-grains depends on the cooling rate: slower cooling allows more time for grain growth.

The microstructural analysis indicates the presence of α-lamellae within primary β-grain boundaries, with the α-lamellae and grain boundary α exhibiting no preferred orientation. The α-lamellae in solid bodies are coarser than those in single layers and single tracks. Also using a bigger laser spot size influences the α-lamellae (Fig. 14). During slow cooling, the formation of larger α-lamellae within the β-grains was favored, while fast cooling resulted in the production of finer α-lamellae (fine Wittmannstätten structures) due to limited atomic diffusion (Fig. 15). It is observed that a phase of discontinuous grain boundary α also occurs at the boundaries of prior β-grains. The cooling rate is contingent upon several variables, including the geometry of the part, its position within the part, the flow rate of the process gas, the power of the laser, the speed of the laser scanning, and the strategy employed for the deposition of layers.

In Fig. 16, the results of the deformation test at different temperature and strain rates are presented and compared to conventionally wrought materials. As deformation speed increases and temperature decreases, the flow stress of the material rises. However, it is observed that material produced through AM with high deposition rates exhibits comparatively lower flow stresses.

The reduced flow stresses observed in such materials are advantageous for the forging process, resulting in reduced tool load and enabling lower forging temperatures for the same degree of deformation and deformation speed. Additionally, the influence of anisotropy due to deformation is reduced.

A possible reason for the lower flow stress in AM materials is the presence of fewer grain boundaries and thinner α-lamellae compared to conventional materials.

The thermomechanical treatment leads to the formation of globularized α-phases. As the forming temperature and degree of forming increase, the globularized phase becomes increasingly coarse (Fig. 17).

The objective of this study was to develop a process capable of achieving high deposition rates up to 5 kg/h using a rectangular laser spot. The resulting material exhibited a low porosity of less than 0.05% and no cracks. The microstructure was characterized by the presence of large β-grains and a directed hexagonal α-phase aligned in the build direction. The formation of large β-grains in build direction (⁠ $⟨ 001 ⟩$ orientation) resulted in strong anisotropy. Moreover, as the laser spot size increased (from 5 × 5 to 7 × 7 mm²), the cooling rate declined (from 466.1 to 210.93 K/s), resulting in the formation of larger β-grains (maximum width of 5420.3 μm) and coarser α-lamellae.

At elevated temperatures and slow deformation speeds, the material exhibited reduced flow stresses. The forging process and subsequent heat treatment led to the formation of globularized α-phases, resulting in quasi-isotropic material behavior regardless of the build direction.

The authors gratefully acknowledge financial support by the Federal Ministry for Economic Affairs and Energy (BMWi) for the LuFo project SAMT64 “Forging and additive manufacturing as a process combination for the resource-efficient production of aerospace structural components made of Ti-6Al-4V on flexible production scales.” Special thanks to Norbert Pirch for conducting the simulation.

The authors have no conflicts to disclose.

Rebar Hama-Saleh: Conceptualization (lead); Data curation (equal); Formal analysis (lead); Funding acquisition (equal); Investigation (lead); Methodology (lead); Validation (lead); Visualization (lead); Writing – original draft (lead); Writing – review & editing (lead). Andreas Weisheit: Supervision (equal). Susanne Hemes: Formal analysis (equal). Constantin Leon Haefner: Supervision (equal).