**Effect of Process Parameters and High-Temperature Preheating on Residual Stress and Relative Density of Ti6Al4V Processed by Selective Laser Melting**

**Martin Malý 1,\*, Christian Höller 2, Mateusz Skalon 3, Benjamin Meier 4, Daniel Koutný 1, Rudolf Pichler 2, Christof Sommitsch <sup>3</sup> and David Paloušek <sup>1</sup>**


Received: 26 February 2019; Accepted: 18 March 2019; Published: 20 March 2019

**Abstract:** The aim of this study is to observe the effect of process parameters on residual stresses and relative density of Ti6Al4V samples produced by Selective Laser Melting. The investigated parameters were hatch laser power, hatch laser velocity, border laser velocity, high-temperature preheating and time delay. Residual stresses were evaluated by the bridge curvature method and relative density by the optical method. The effect of the observed process parameters was estimated by the design of experiment and surface response methods. It was found that for an effective residual stress reduction, the high preheating temperature was the most significant parameter. High preheating temperature also increased the relative density but caused changes in the chemical composition of Ti6Al4V unmelted powder. Chemical analysis proved that after one build job with high preheating temperature, oxygen and hydrogen content exceeded the ASTM B348 limits for Grade 5 titanium.

**Keywords:** Selective Laser Melting; Ti6Al4V; residual stress; deformation; preheating; relative density; powder degradation

#### **1. Introduction**

One of the most popular additive manufacturing technologies is Selective Laser Melting (SLM). SLM technology allows the production of nearly full density metallic parts with mechanical properties comparable to the ones produced by conventional methods. Components are made layer-by-layer directly from powdered material, where each layer is selectively melted in an inert atmosphere by a laser beam [1–3].

Due to non-uniform spot heating and fast cooling, thermal gradients are formed in materials, which lead to the development of residual stresses [1]. Residual stress (RS) is described as stress which remains in the material when the equilibrium with surrounding environment is reached [4]. Unwanted RS in SLM can cause part failure due to distortions, delamination or cracking.

The RS measurement is possible by mechanical and diffraction methods; magnetic and electric techniques; and by the ultrasonic and piezoelectric effect. Mechanical measurements are usually based on material removal and its relaxation or measuring part distortion. Typical mechanical measurement is a drilling method with an accuracy of ±50 MPa [4]. The main advantage of this method is its

ability to measure RS to the depth of 1.2 times the diameter of the drilled hole. The Bridge Curvature Method (BCM) is often used in case of the SLM for the fast comparison of process parameters and their influence on the RS [5]. The principle is based on measuring the distortion angle, after the sample is cut off from the base plate, using the bridge-like samples. Measuring accuracy could be affected by imprecise cutting or by angle evaluating method, thus measurement of the top surface inclination was proposed as a better technique [6,7]. Value of the RS can be determined by the simulation of measured distortion in Finite Element Method (FEM) analysis [8].

Ali et al. [9] proved that higher exposure time and lower laser power with preserved energy density lowered the RS due to lower cooling rate and temperature gradient. The variation of laser power and exposure time did not cause a change in yield strength of Ti6Al4V, but elongation increased with lower laser power and higher exposure time. The cooling rate and also the RS can be affected by layer thickness. Higher layer thickness prolonged the cooling rate and RS was lower, but with higher layer thickness the relative density was lowered [9,10].

Scanning strategy has a significant effect on the RS. Ali et al. [11] observed that with longer scanning vectors the RS increased. Due to prolonged time between scanning adjacent scan tracks, higher thermal gradients were induced. The lowest RS had a stripe strategy with the ninety-degree rotation. This conclusion was also confirmed by Robinson et al. [12]. Ali et al. [11] did not observe the positive or negative effect of the scanning strategy on mechanical properties nor relative density.

Powder bed preheating can significantly reduce the amount of RS [5,13,14]. Ali et al. [15] demonstrated that for Ti6Al4V-ELI material preheating of the build platform to the temperature of 570 ◦C effectively eliminated the RS. A positive influence of preheating on microstructure and mechanical properties of H13 tool steel was observed by Mertens et al. [16]. The preheating up to 400 ◦C in his study improved mechanical properties and the parts had more homogeneous microstructure. Formation of cracks can be also affected by preheating, which was observed during printing aluminium of 2618, 7075 [17,18] and tool steel [19].

In this study, the BCM samples made of Ti6Al4V were used to evaluate the effect of process parameters on the relative density and the RS. Investigated parameters were hatch laser speed, hatch laser power border laser velocity, waiting time between adjacent layers and powder bed preheating up to 550 ◦C. The design of experiment and the surface response method were used for a comprehensive evaluation of the effects of observed process parameters. Furthermore, the influence of high-temperature preheating on powder degradation was evaluated.

#### **2. Materials and Methods**

#### *2.1. Powder Characterization*

In this study, Ti6Al4V gas atomized powder (SLM Solutions Group AG, Lübeck, Germany) was used. The chemical composition of virgin powder delivered by the manufacturer is in Table 1. The powder shape was checked by scanning electron microscopy (SEM) LEO 1450VP (Carl Zeiss AG, Oberkochen, Germany). Figure 1a shows that the powder particles have a spherical shape with a low amount of satellites. The particles size distribution was analysed by laser diffraction analyser LA-960 (Horiba, Kioto, Japan). Measured particle mean size was 43 μm and median size 40.9 μm. The particles up to 29.97 μm represented 10% of particle distribution while particles up to 58.61 μm represented 90% (Figure 1b).


**Table 1.** Chemical composition of virgin Ti6Al4V powder.

**Figure 1.** Ti6Al4V gas atomized powder characterization (**a**) shape evaluation by SEM; (**b**) particles size distribution.

The chemical composition of used powder was evaluated by the following methods. The aluminium content was checked by the inductively coupled plasma atomic emission spectroscopy. Oxygen and nitrogen contents were evaluated by hot extraction in helium by LECO TCH 600 (LECO Corporation, Saint Joseph, MO, USA). The hydrogen concentration was verified by the inert gas fusion thermal conductivity method JUWE H-Mat 2500 (JUWE Laborgeraete GmbH, Viersen, Germany). The accuracy of all methods is Al ±0.327 wt %, O ±0.008 wt % and N ±0.0025 wt %.

#### *2.2. Sample Fabrication*

The samples were manufactured on the SLM 280HL (SLM Solutions Group AG, Lübeck, Germany) 3D printer. The machine is equipped with 400 W ytterbium fibre laser YLR-400-WC-Y11 (IPG Photonics, Oxford, MS, USA) with a focus diameter of 82 μm and a Gaussian shape power distribution. Argon was used as a protective atmosphere during the process and the O2 content was kept below 0.05 %. Before each experiment, the humidity of the powder was measured by the hydro thermometer Hytelog (B + B Thermo-Technik GmbH, Donaueschingen, Germany) with an accuracy of ±2%. The powder humidity was kept under 10%. The heating platform (SLM Solutions Group AG, Lübeck, Germany) was used to preheat the powder. This device is able to preheat the build platform up to 550 ◦C, but the build area is reduced to a cylindrical shape with 90 mm in diameter and 100 mm in height. For the preheating a resistive heating element is used and the temperature is controlled by a thermocouple placed below the base plate. The temperature of a printed component may be slightly lower than the measured temperature by the thermocouple. However, the maximum height of parts printed in this study is 12 mm, thus the temperature field should be relatively homogeneous. Build data were prepared in Materialise Magics 22.03 (Materialise NV, Leuven, Belgium).

#### *2.3. Sample Geometry*

The geometry of samples was designed according to the BCM shape (Figure 2a) [5], therefore the effect of chosen process parameters on distortion and RS can be evaluated. Support structures were used for all samples to simulate the condition during the printing of real components. To restrict distortion during the SLM process (before cutting off) the 4 mm high block supports were reinforced with 1 mm block spacing, while fragmentation was switched off. Teeth top length was set to 1 mm. Support structures were added just under the pillars. Samples were rotated to 20◦ from recoating direction to ensure consistent powder spreading. Samples were cut in the middle of support structures and the evaluated parameter was top surface angle distortion α, which is the sum of α<sup>1</sup> and α<sup>2</sup> (Figure 2b).

**Figure 2.** Samples geometry: (**a**) Dimensions of the BCM sample; (**b**) Measured bridge top surface angle distortion α is the sum of α<sup>1</sup> and α2. Dimensions presented in mm.

#### *2.4. Design of Experiment*

For data evaluation Design of Experiment (DoE) and Surface Response Design (SRD) were used. Hatch laser power (H LP), hatch laser velocity (H LV), border laser velocity (B LV), delay time (DT) and preheating temperature (T) were chosen as the variable factors. The range of parameters with central points is summarized in Table 2. The DT value is waiting delay between two adjacent layers, which was set in the printing machine. The real delay (RD) value which was used for result evaluation is composed of set DT between two layers and 13 s recoating time. If the DT is zero, then the RD value is composed of 13 s recoating time and scanning time. The temperature range was set from the common preheating temperature 200 ◦C to the maximum temperature of 550 ◦C that our equipment is capable to evolve.


**Table 2.** Table of used process parameters for Design of Experiment (DoE) and Surface Response Design (SRD).

The half fraction of the SRD was built with the five continuous variable parameters. This means twenty-six samples plus four repetition central points. To minimize the number of global parameters, which has an influence on the whole build job, the face-centered design was used.

For evaluation, Minitab 17 (Minitab Inc., State College, PA, USA) was used. Data of the top surface angle distortion α were evaluated with full quadratic terms with 95% confidence level for all intervals and with backward elimination of 0.1. Relative density data were evaluated on samples 1–16 as the half fraction of factorial design. Then the data were assessed with 95% confidence level and insignificant term combinations were manually deleted.

Border laser power was set to 100 W and hatch spacing to 0.12 mm. The layer thickness of 50 μm and stripe strategy with a maximum stripe length of 10 mm and a rotation of 67◦ was used. Fill contour was turned off. Other parameters were set as standard.

#### *2.5. Distortion Evaluation*

The 3D optical scanner Atos TripleScan 8M (GOM GmbH, Braunschweig, Germany) was used for assessing distortions of the bridges. Each sample was scanned after cut-off from the base plate and after coating by TiO2 mating spray with the thickness of around 3 μm [20]. The 3D scanned surface data were evaluated in GOM Inspect 2018 (GOM GmbH, Braunschweig, Germany).

The top surface angle distortion α was measured on the top surface of the bridge as is shown in Figure 2b. First, the Computer-aided Design (CAD) data of the undeformed bridge was fitted by Gaussian best fit function on the scanned data. Then three cross sections were created parallel to the YZ plane in distance 0, 8 and −8 mm. Lines using Gaussian best fit function were fitted on the top surface in each cross section (Figure 3a). Next, points in distance 0, 3 and −3 mm in Y direction were created on those three lines (Figure 3b). Then the distance was measured between middle and side points. Finally, X and Z components from each measured distance were used for calculating angle distortions by tangent function. Left and right sides were calculated separately. Therefore, the α value was calculated as the sum of angles on both sides. The result of the top surface angle distortion α value is the mean value of three measurements of one sample.

**Figure 3.** Distortion evaluation, fitted lines are yellow, line points are red and measured distances are green: (**a**) Isometric view on scanned data; (**b**) Top view of scanned data.

#### *2.6. Relative Density Measurement*

Relative density was determined using an optical method and was calculated as the mean value of parallel to build cross sections (Figure 4a). Value of relative density was evaluated in ImageJ v. 1.52k (National Institutes of Health, Bethesda, MD, USA). First, the picture of the cross section was converted to 8-bit type. Next, an automatic threshold was applied and relative density was evaluated in the areas defined by red rectangles (Figure 4b).

**Figure 4.** Cross sections of the BCM samples: (**a**) Cross section of sample 1; (**b**) Cross section of the sample made with the lowest energy density (Sample 3), red rectangles show area for relative density evaluation.

#### **3. Results**

#### *3.1. Top Surface Distortion and Relative Density*

The experimental design matrix and results of top surface angle distortion α and measured relative density are summarized in Table 3. Samples were sorted in printing order and horizontal lines represent a group of samples which were printed together in one build job.


**Table 3.** DoE and SRD test matrix with process parameters, the value of top surface angle distortion α and relative density.

#### *3.2. Surface Response Model for Top Surface Angle Distortion α*

Minitab 17 was used to establish a regression model for prediction of the top surface angle distortion α responses to the H LP, H LV, B LV, RD and T. Equation (1) represents the SRM-based mathematical model of significant parameters, which represent the relation between observed parameters. Table 4 and Figure 5 show results from an analysis of variance (ANOVA). The correlation of the regression model for α value is confirmed by determination coefficients R<sup>2</sup> = 91.82% and adjusted R2 = 86.04%.


**Figure 5.** Main effect plot for top surface angle distortion α.


**Table 4.** ANOVA table for the top surface angle distortion α.

In order to investigate the effect of high energy and high-temperature preheating on the distortion, an additional four bridge samples were made. Those samples were built with increasing H LP according to Table 5.

**Table 5.** Parameters of samples with increasing laser power and value of α.


<sup>1</sup> Calculated as H Ed = H LP·(H LV·Lt·Hs)−1, Layer thickness (Lt) = 50 <sup>μ</sup>m, Hatch spacing (Hs) = 120 <sup>μ</sup>m.

#### *3.3. Mathematical Model for Relative Density*

Equation (2) represents a mathematical model of significant parameters with an influence on the relative density. Figure 6 shows a Pareto chart of standardized effect for evaluated relative density data. ANOVA results are shown in Figure 7 and Table 6. The correlation of the regression model for relative density value is confirmed by determination coefficients R2 = 89.38% and adjusted R2 = 82.29%.

Relative density = 136.67 − 0.1558 H LP − 0.0628 H LV − 0.00004 B LV + 0.00766 T + 0.0390 RD + 0.000232 H LP\*H LV

**Figure 6.** Pareto chart of the standardized effect to relative density.

**Figure 7.** Main effect plot for relative density.


#### *3.4. Analysis of Used Powder*

Figure 8a shows the influence of high-temperature base plate preheating (550 ◦C) on the powder. The powder significantly changed colour from silver to brown and there is a hint of particle

(2)

agglomeration. Therefore, the powder used in build job with 550 ◦C was checked by the SEM microscopy in order to investigate the particle agglomeration and their shape (Figure 8b).

**Figure 8.** The powder used in heating unit preheated to the 550 ◦C (**a**) Build job made with 550 ◦C; (**b**) SEM microscopy photo of the powder used with 550 ◦C.

Powder chemistry analysis was done after first build job with preheating to 200 ◦C and these results were compared with the powder used with 550 ◦C preheating. Results in Table 7 confirm a rise in oxygen content from 0.12 to 0.33 wt % and in the hydrogen from 0.002 to 0.0168 wt %. Aluminium content also slightly rose from 6.05 to 6.11 wt %, but nitrogen content decreased from 0.017 to 0.0149 wt %. Results are compared with the ASTM B348 Grade 5 titanium requirements and virgin Ti6Al4V powder chemical composition received by the vendor.



#### **4. Discussion**

#### *4.1. Top Surface Angle Distortion α*

The main contribution of each parameter on the distortion, and therefore the amount of residual stresses, can be derived from the ANOVA (Table 4). Further on the significance of each parameter can be evaluated by a p-value. If the value is lower than 0.05, then the parameter is significant, while the p-value of the lack-of-fit parameter should be high, which shows that the error value is not significant. The p-value of the lack-of-fit parameter 0.685 shows that the regression model for the top surface angle distortion α fits the measured data.

Table 4 and Figure 5 show that the most significant parameters for reduction distortions are T and H LP with 46.31% and 17.22% linear contribution. P-values of those parameters are 0. The RD and H LV are not that significant in comparison with the previous two parameters. Their linear contributions are 5.26% and 3.62%, while p-values are lower than 0.05, therefore parameters are significant. The B LV can be considered as an insignificant parameter with 0.59% linear contribution and the p-value greater than 0.05. Figure 5 indicates that H LP, B LV and T have linear behaviour in contrast to RD and H LV.

From the ANOVA results, it can be concluded that increasing the preheating temperature or laser power causes a reduction in deformation. In contrast, increasing H LV, B LV and RD lead to higher deformations. This can be contributed to the cooling rate, which is lower with slower laser movement, higher preheating and shorter waiting time. Therefore, the thermal gradients, residual stresses and finally distortions are lower [9,11,15,21,22].

Optimal parameters for achieving the lowest distortion can be predicted from the fitted regression model (Figure 9). For reaching the lowest distortion it is predicted to use H LP 275 W, H LV 785 mm/s, B LV 350 mm/s, RD 17 s and T 550 ◦C. Predicted α is −0.182◦. In contrast, the lowest distortion predicted with a preheating temperature of 200 ◦C is 0.139◦.

**Figure 9.** Predicted values for the lowest distortion in the full range of observed parameters.

Figure 10 shows the influence of energy density (ED) on the top surface angle distortion α. Samples used for this comparison were made with the same preheating temperature of 550 ◦C and 200 ◦C. The effect of B LV and RD was neglected. The interpolated line for 200 ◦C samples is constantly dropping with increasing ED and as was measured by Mishurova et al. [7], and this trend constantly continues. In contrast, the interpolated line for 550 ◦C samples with added high ED samples starts much lower than 200 ◦C line. The decrease of α value is gradual until 65 J/mm3 and drops rapidly with higher ED.

**Figure 10.** Effect of energy density on the top surface angle distortion α.

#### *4.2. Relative Density*

The most significant effect out of involved parameters influencing the relative density are H LP and H LV. The linear contribution is 34.66% for H LP and 23.85% for H LV. They are also significant in their two-way interaction with a contribution of 26.41%. P-values are 0 for H LP and 0.002 for H LV.

Preheating temperature has a linear contribution of 2.88% while its p-value of 0.153 is higher than 0.05. This means that this parameter in the observed range is not that significant for the model due to the high contribution of H LP and H LV. Relative density rose with higher T. Positive influence of preheating was also confirmed on stainless steel M2. It was proved by Kempen et al. [19] that with higher preheating temperature, higher laser velocities can be used while maintaining relative density.

The real delay has a linear contribution of 1.58% and a p-value of 0.277, which means the RD has minimal influence on relative density. Prolonged RD leads to an increase in relative density.

Border laser velocity with a linear contribution of 0 % and p-value 0.993 means that this parameter was not significant for the model, which can be due to the place where porosity was measured.

From the ANOVA (Figure 7) it can be deduced that a sample will have maximum relative density when values of H LP, RD and T are set as the highest and H LV as the lowest. This means the highest energy density is in the hatch.

#### *4.3. Powder Degradation*

There is clear evidence that the chemical composition of the powder significantly changed due to oxidation. Titanium alloys suffer high chemical affinity to oxygen leading to form a thin oxide layer even on air room temperature. Exposing titanium to an oxygen-containing atmosphere at elevated temperatures around 550 ◦C increase diffusion rates through thin oxide layers, and allows penetration of oxygen in the material [23]. After experiments with 550 ◦C preheating, oxygen content increased against 200 ◦C preheating from 0.12 to 0.33 wt % which is 0.13% higher than the ASTM B348 requirement for Grade 5 titanium.

Increased oxygen content in the Ti6Al4V causes an increase in yield and ultimate tensile strength, whilst ductility up to 0.19 wt % of the oxygen content remains constant [24]. Ti6Al4V additively manufactured alloy is due to rapid cooling mainly composed from α'martensitic microstructure even with preheating up to 550 ◦C [15]. Therefore, this is sensitive to oxygen content because of the concentration of oxygen higher than 0.22 wt % leads to the brittleness of the α'martensitic structure. Critical oxygen content for α and β structure is 0.4 wt % [25]. Oxygen concentration above 0.25 wt % leads to change in the typical microstructure, which causes a sharp decrease in ductility of Ti6Al4V [26].

Hydrogen content in the powder used under 550 ◦C exceeds the approved ASTM B348 limit value of 0.0125 wt % and its content rose from 0.002 to 0.0168 wt %. The diffusion rate of the hydrogen is rapidly increasing at the elevated temperatures [23]. The origin of the hydrogen element is most probably from powder moisture which was kept below 10%. Hydrogen in titanium alloys causes a phenomenon known as hydrogen embrittlement and could lead to part failure [27,28].

It was shown that the critical issues with processing the titanium alloy by SLM at high temperatures are connected with chemical composition changes in the unused powder, although the material was processed under argon atmosphere with oxygen concentration of 0.05% and powder humidity kept below 10%. The measured concentration of oxygen and hydrogen was beyond ASTM B348 requirement for Grade 5 titanium. Therefore, the used powder cannot be used for mechanical stressed parts.

#### **5. Conclusions**

Effects of hatch laser power, hatch laser velocity, border laser velocity, preheating temperature and delay time on residual stress and relative density on SLM processed Ti6Al4V samples have been investigated. In addition, the impact of preheating temperature up to 550 ◦C on Ti6Al4V powder degradation has been discussed. The main findings are the following:


**Author Contributions:** Conceptualization, M.M., D.K.; methodology, M.M., C.H and M.S.; validation, M.M., C.H., B.M., M.S., D.K., R.P., C.S. and D.P.; formal analysis, M.M.; investigation, M.M., B.M. and M.S.; resources, R.P., C.S., D.K. and D.P.; data curation, M.M.; writing—original draft preparation, M.M.; writing—review and editing, M.M., D.K. and M.S.; visualization, M.M.; supervision, M.S., C.H., D.K., R.P. and C.S.; project administration M.M.; funding acquisition, C.S., R.P, D.K. and D.P.

**Funding:** This research was funded by the ESIF, EU Operational Programme Research, Development and Education within the research project [Architectured materials designed for additive manufacturing] grant number [CZ.02.1.01/0.0/0.0/16\_025/0007304] and faculty specific research project FSI-S-17-4144.

**Acknowledgments:** The authors express their thanks to Philipp Schwemberger from the Institute of Production Engineering, Graz University of Technology, for his help with the samples preparation.

**Conflicts of Interest:** The authors declare no conflicts of interest. The funders had no role in the design of the study; in the collection, analyses, or interpretation of data; in the writing of the manuscript, or in the decision to publish the results.

#### **References**


© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

## *Article* **Phase Studies of Additively Manufactured Near Beta Titanium Alloy-Ti55511**

#### **Tuerdi Maimaitiyili 1,2,\*, Krystian Mosur 3, Tomasz Kurzynowski 3, Nicola Casati <sup>4</sup> and Helena Van Swygenhoven 2,5**


Received: 16 March 2020; Accepted: 3 April 2020; Published: 7 April 2020

**Abstract:** The effect of electron-beam melting (EBM) and selective laser melting (SLM) processes on the chemical composition, phase composition, density, microstructure, and microhardness of as-built Ti55511 blocks were evaluated and compared. The work also aimed to understand how each process setting affects the powder characteristics after processing. Experiments have shown that both methods can process Ti55511 successfully and can build parts with almost full density (>99%) without any internal cracks or delamination. It was observed that the SLM build sample can retain the phase composition of the initial powder, while EBM displayed significant phase changes. After the EBM process, a considerable amount of α Ti-phase and lamella-like microstructures were found in the EBM build sample and corresponding powder left in the build chamber. Both processes showed a similar effect on the variation of powder morphology after the process. Despite the apparent difference in alloying composition, the EBM build Ti55511 sample showed similar microhardness as EBM build Ti-6Al-4V. Measured microhardness of the EBM build sample is approximately 10% higher than the SLM build, and it measured as 348 ± 30.20 HV.

**Keywords:** titanium alloy; Ti55511; synchrotron; XRD; microscopy; SLM; EBM; EBSD; additive manufacturing; Rietveld analysis

#### **1. Introduction**

Additive manufacturing (AM) also known as "3D printing" is an advanced manufacturing technology which allow fabrication of geometrically complex and functional parts directly form the computer-aided design (CAD) model in a short time with limited tooling cost and with almost no material waste [1–6]. Hence, AM is a potential technology in areas where a high degree of customization and on-demand manufacturing is key such as aerospace and medical industry.

Different types of AM techniques have been developed for metallic materials, and they can be categorized into many different subclasses based on energy sources, feed material, and ingredient material feeding approach as described by Liu et al. [7]. Among various AM techniques, the powder bed selective laser melting (SLM) and electron beam melting (EBM) are two common AM techniques, and these systems have been described in detail by Samy et al. [1] and Maimaitiyili et al. [2], respectively. In general, AM processes involve shaping a build plate and selectively melting the raw material (e.g.,

wire/powder) to form a three-dimensional solid object using a high energy focused laser or electron beam with multi-axis motion. Even though the basic operational principle of the AM method is relatively simple, the actual metal AM process is complex, and the results depend upon different settings of the system such as beam power (current), scanning speed, preheat temperature, etc. These are collectively referred to as processing parameters. In the AM process, processing parameters determine the build environment and cooling conditions and, consequently, affect the phase composition [1,2,8,9], residual stress [2], texture [2,8,10], surface roughness [11], density [7,12,13], and mechanical properties [3,7,12–15] of the as-built part. The lack of understanding of the relationship between process parameters and microstructure hinders the prediction of the properties of the built material and service life.

Titanium-based alloys have been widely used as an engineering material in many industries because of their excellent combination of a high strength/weight ratio and good corrosion resistance [1,2,10,12,14–17]. However, extracting high purity Ti and producing usable Ti-alloy parts are difficult and expensive processes. Therefore, there is a strong interest in using AM, such as SLM and EBM techniques, to process Ti-based materials.

The mechanical properties of Ti-alloys depend on the microstructure, chemical, and phase composition [3,7,12–15]. Based on concentrations of alloying elements, Ti-alloys can be divided into three main classes: α, α + β, and (meta-stable and stable) β-alloys [17,18]. Depending on alloy composition and heating/cooling rates, Ti-alloys can also contain metastable α (hcp) and α" (orthorhombic) phases [1,17–19]. Therefore, it is of fundamental and technological importance to investigate the influence of AM processes on the phase composition and microstructure in Ti-alloys. Many AM studies have been performed on the two-phase (α + β) Ti-alloy known as Ti-6Al-4V [1,2,5,6,8–11,14,20],while other alloy compositions such as Ti–5Al–5Mo–5V–1Cr–1Fe (Ti55511) have been less addressed [13,21].

The Ti55511 is a near β-type alloy with important application in aerospace industries [13,22–24]. Compared to Ti-6Al-4V, the Ti55511 is a superior structural material, as it provides comparable or higher strength with 15%–20% less weight [24]. Most of our knowledge has been derived from conventionally manufactured materials [23,24]. Characterization of phases and microstructure after synthesis using different AM methods has not yet been performed. Here, we report and compare the build quality, microstructure, and phase composition of Ti55511 synthesized by SLM and EBM.

#### **2. Materials and Methods**

#### *2.1. Powders*

Pre-alloyed Ti–5Al–5Mo–5V–1Cr–1Fe (Ti55511) powder prepared by gas atomization with an average particle size of 43 μm for SLM and 71 μm for EBM was obtained from KAMB Import–Export Warszawa (Nr CAS: 7440-32-6). The chemical compositions (wt%) of the as-received Ti55511 powder was Al 5.17, Mo 4.95, V 4.74, Cr 0.92, Fe 1.01, balanced by Ti.

Laser diffraction particle size analyzer Partica LA-950 V2 system (Horiba, Tokyo, Japan) was used to measure the particle size distribution. For accuracy, each measurement was repeated three times. The results are plotted in Figure 1, and some important parameters are listed in Table 1.

A scanning electron microscope (SEM) equipped with an energy-dispersive X-ray detector system was used to examine the shape, size distribution, surface morphology, and external and internal defects of the powders. In addition, optical microscopy was employed to observe the internal porosity, defects, and cross-section of the powders. Selected SEM images of powders before and after processing are shown in Figure 2. Figure 3 shows the internal defect of the powders. The size distribution determined by SEM agrees well with data presented in Figure 1 and Table 1.

**Figure 1.** Particle size distribution of Ti55511 powders.

**Table 1.** Powder size distribution before and after processing measured by the laser diffraction method.


"D10", "D50", and "D90" mean the particle sizes at 10 vol%, 50 vol%, and 90 vol%, respectively.

#### *2.2. Sample Build Processes*

The Ti55511 blocks with dimension of 2.5 × 2.5 × 5 cm were fabricated using both SLM and EBM methods in 90◦ (long side of the sample was along with the build direction), 0◦ (long side of the sample was in the build plane), and 45◦ (long side of the sample was 45◦ to the build plane) orientations. All microstructure related results presented were obtained from 90◦ samples.

The SLM samples were made with ReaLizer 250 II SLM machine (ReaLizer GmbH, Borchen, Germany) equipped with a 400 W fiber laser. The laser melting process was carried under a protective argon atmosphere with O2 content less than 0.1 vol.%. The laser power (P) was 200 W, the scan speed (v) 330 mm/s, the hatch spacing (h) 0.21 mm, and layer thickness (d) 50 μm. To reduce thermal residual stress and elemental segregation, the samples were built on a Ti-6Al-4V substrate pre-heated to 250 ◦C prior to the building process.

The EBM specimens were built on an Arcam A2 machine (Arcam AB, Mölndal, Sweden) with a layer thickness of 50 μm. The processing parameters, such as spot size and scan velocity, were defined by the Arcam A2 process control algorithm. The scanning speed was 4530 mm/s, current 15mA, focus offset 3mA, and preheat temperature 650 ◦C. To investigate only the process effect, no post-treatment was applied to the specimens.

The scan strategy used in SLM was the "island" scan strategy in which each layer was divided into 3 mm × 3 mm square islands. Scan tracks in each island were exposed at the same orientation with respect to the neighboring island and rotated 90◦ among alternating layers.

In the EBM, a bidirectional scan strategy was used in which all tracks are made with alternating direction, i.e., left-to-right, then right-to-left.

The scanning strategies used are those that resulted in this material lowest residual stress and porosity which determined after cube print tests.

#### *2.3. Characterization Methods*

All samples used for microscopic studies were prepared by using standard metallographic preparation routines. Examination of microstructure was performed using a Visible Light Microscope (VLM, Leica DMRX + SpeedXT Core5, Wetzlar, Germany) and ZEISS NVision40 scanning electron microscope equipped with an energy-dispersive X-ray spectrometer (EDS) analysis system from Oxford Instruments (Oberkochen, Germany). To obtain phase and texture related information from the as-built material, electron backscatter diffraction (EBSD) investigations were performed using a field emission gun scanning electron microscope (FEG SEM) ZEISS ULTRA 55 equipped with an EDAX Hikari Camera (Oberkochen, Germany) operated at 20 kV in a high current mode with 120 μm aperture.

To identify the phase composition of powders, synchrotron X-ray powder diffraction were carried out at the Material Science (MS) beamline X04SA-MS4 of the Swiss Light Source (Paul Scherrer Institute, Villigen, Switzerland) using the MYTHEN II detector. All measurements were made at room temperature with 25.1 keV (λ = 0.4940 Å) X-ray beam and 60s exposure.

Diffraction data of as built material were acquired using a D500 X-ray diffractometer (XRD) from Bruker–Siemens (Karlsruhe, Germany) with Cu Kα radiation (λ = 0.15406 nm) operating at 40 mA and 40 kV. The step size and the acquisition time were 0.01◦ and 1 s respectively. All measurements were conducted at room temperature at the center of each test blocks cut from each sample faces (xy-, xz- and yz-planes) from the top- and bottom-half of the sample. The quantitative phase analysis was performed with a Topas-Academic software package.

Porosity was characterized using the Archimedes technique and microscopy. Cuboid samples were sectioned at different depths, ground and polished, and inspected in SEM. On average, 120 images were captured from each sectioned part and stitched with the functions in IMAGIC IMS V17Q4. These color images were then converted into 8 bit black-and-white images using ImageJ. To understand the shape of the pores, a circularity of the pores was also calculated with this program.

#### **3. Results and Discussions**

#### *3.1. Powder Characterization*

Figure 1 and Table 1 present the results of the powder size distribution (PSD) of different powder samples obtained using a Partica LA-950 V2 laser particle size analyzer (Kyoto, Japan). Figures 2 and 3 show the morphology and external/internal defect of the powders, respectively. Both SLM and EBM powder particles predominantly in spherical shape with limited quantity of non-spherical particles and spherical imperfections. Both SLM and EBM powders have a nearly normal size distribution (Figure 1). The SLM powder had a size distribution between 28 (D10) and 60 μm (D90) with mean volume diameter around 40 μm. The EBM powder had a size distribution between 51 (D10) and 99 μm (D90) with mean volume diameter around 70 μm.

**Figure 2.** SEM images showing the morphologies of ingredient powder before (**a**,**c**) and after (**b**,**d**) the EBM (**a**,**b**) and SLM (**c**,**d**) processes.

It has been reported that a spherical powder with narrow PSD has a positive effect in both mechanical properties and finishing surface for the sample [25]. Because of charging problems associated with electron beam, commonly, a larger powder is used in the EBM [2]. According to the literature [2,4,5], measured PSD of the EBM and SLM was in the suggested PSD range for respective methods. Hence, both types of powders are ideal for processing with corresponding methods and all observation reported here can be directly related to the alloy and the manufacturing methods in use.

After processing, the powder remains predominantly spherical as shown in Figure 2b,d. Table 1 shows, however, that both the SLM and EBM process caused changes in powder size distribution. The averaged sizes tended to increase which can be due to the powder agglomeration or powder recoating. Occasionally, broken powder particles were observed as shown in Figure 2d. It is, however, clear that both process routes have limited effect on the powder quality.

Figure 2 shows cross-sections of embedded powder after etching before and after usage. The initial microstructure (before usage) of the SLM and EBM powders was very similar. The powder used in the SLM process was not different, but the powder used in EBM showed a lamellae microstructure with a mixture of very fine and coarser α and β phases. Similar lamella microstructure is reported by Li et al. [26] for thermally treated conventional Ti55511 which first solution treated at 920 ◦C for 120 min and annealed at 700 ◦C for 60 min. Therefore, it is believed that the high build plate temperature used in the EBM process together with heat dissipated from the melt zone during electron beam scanning is the main cause of such significant phase transformation observed in the remaining powders left in the build chamber.

**Figure 3.** Internal microstructure of EBM powder (**a**) before and (**b**) after processing. Microstructure of SLM powder (**c**) before and (**d**) after processing.

The chemical composition of powder samples before and after the process was assessed by EDS analysis. As in Figure 3, the elemental map of powders before SLM and EBM process and as well as after SLM process do not show any distinct regions, and they all seemed homogeneous and featureless. However, there are two distinct regions in the elemental map of powder after the EBM process: Mo dense and Mo depleted region. Such a difference indicates a difference in phases. In Ti-alloy, Mo and V work as a β stabilizers, and Al is an α stabilizer [17]. Therefore, during phase formation, these elements will preferentially partition to the respective phases.

The diffraction patterns of powders before and after the process shown in Figure 4 confirm a microstructure only composed of the β Ti-phase. Remaining powder in the SLM build chamber after the process consisted of comparable phase composition as ingredient powder, however, the remaining powder after the EBM process showed the clear presence of α Ti-phase in addition to β Ti-phase.

It is well known that titanium and its alloys have a strong affinity for oxygen, and they can react to form detrimental oxides at elevated temperatures which potentially degrade the quality of the build parts [7]. However, from Figure 4, one can see that all powders are free from oxides. In addition, all powders did not show any observable color changes after the process. Therefore, it is safe to say that both EBM and SLM process are equally effective in preventing oxygen contamination, and powder degradation related to oxygen from both methods are minimal.

**Figure 4.** Diffraction pattern of powders before and after the process. These color-coded tick marks under the diffractogram correspond to the expected peak positions of α- and β-Ti phases reported in the Inorganic Crystal Structure Database (ICSD) [27] (ICSD reference number of α-phase is 191187 and β-phase is 653278).

It is commonly reported that the powder morphology [25], oxygen content [7], and PSD [25] have significant impact on the final build material quality, and for that reason these parameters are often used/discussed in the literature for evaluating the impact of an additive manufacturing process to the powder degradation behavior. However, it is not clear whether the phase composition of the ingredient powder has any influence on the porosity of the build parts. As different phases have a different crystal structure and each phase mixtures can have specific microstructures, it can be expected that the thermal/chemical properties of the powders with phase transformation can be different from the standard/initial powders. Therefore, it might be also important to include powder phase composition in the discussion of powder degradation evaluations together with other parameters.

#### *3.2. Surface Roughness*

All builds from both methods were successful, with no warping, distortion, lifting from the base and no macro/micro-scale cracking. The physical appearance of vertically built samples in the 90◦ orientation from SLM and EBM are shown in Figures 5a and 5b, respectively. Clear band-like patterns were observed in the SLM-built specimen (Figure 5a). A limited number of rough spikes sticking to the sides of the EBM-built specimen, as shown with arrows in Figure 5b, were observed. It is believed that these spikes might be caused by fallen agglomerated powders/melts which spatter out during scanning processes. As the distribution of these spikes is random, the occurrence is limited and can be removed relatively easily, they are excluded from the surface roughness evaluations. The roughness was measured with a Veeco Dektak 8 (NY, USA) profilometer and values of 12.27 μm and 38.05 μm were obtained for SLM and EBM, respectively. The surface roughness of the additively manufactured parts was mainly influenced by the powder particle size and build layer thickness [28,29]. In general, the smaller the particle size, the thinner the layer thickness and so the higher the surface quality. The size of the powder particles used in EBM was twice the size of SLM, so a higher surface roughness was expected for EBM.

**Figure 5.** Photography of (**a**) SLM and (**b**) EBM samples build on 90◦ orientation. (**c**) Density comparison of SLM (red) and EBM (blue) samples built on three different orientations obtained from the Archimedes principle.

#### *3.3. Porosity*

Figure 5c shows the Archimedes density measurement results of samples built in three different orientations. The standard deviations of density measurement were less than 0.005 g/cm3 for all measurements. As can be observed, EBM in general gave higher density than SLM irrespective of sample build orientation, but the difference between these two was less than one percent (0.39%). With respect to reported densities in literature, both methods can produce an almost fully dense structure (99.38 and 99.77% for SLM and EBM, respectively). Results of the porosity after image analysis from the xy- and xz-planes agreed well with the results presented in Figure 5c. The SLM sample had more pores than the EBM samples. The size of the pored ranged between 5 and 300 μm in SLM and 5 and 70 μm in EBM. The distribution of the porosity in the xy planes can be observed in Figure 6. The figures confirm the difference in porosity but also reveal differences in their distribution.

Generally, two types of pores exist in powder-based AM: spherical, gas-induced pores, and irregular-shaped, process-induced pores [7,30]. The first can occur due to the presence of entrapped gas in the powder particles during atomization, the latter mainly associated with non-optimal process parameters [30]. According to Figure 6 and results from the image analysis, pores in the EBM sample were mostly gas pore type, while in SLM both types were observed.

An interesting point to notice on the micrographs presented in Figure 6 were the pore patterns. In the EBM build sample, pores seemed to form randomly at both build direction and build planes, while in SLM, there was a clear tendency to form preferentially in both directions. In SLM most pores tended to form linearly in the build direction. In the build plane, pores occurred mostly at or near to the corners of a scanning "island" (Figure 6b,c). The reason for having such a pore formation pattern in SLM may be related to the scanning strategies in use. When the microstructure was observed after etching in the xy-plane of SLM build sample (Figure 6c), one can identify the used scanning strategy. The approximate size and locations of one "island" is shown in Figure 6c by a red square. As seen in Figure 6c, pores were indeed most prevalent at the "island" intersection or corner regions. One of the possible reasons for this is when the laser reaches the "island" border and starts to melt the next scan vector, a process associated pores, such as keyhole and lack of fusion might be formed. This problem can be mitigated or eliminated by (1) increasing the overlap between "islands", (2) introducing a shift and tilt between layers, (3) adding a contour scan with lower energy density in each "island" just before or after its completion, (4) changing the scanning speed when the scan vector approaches the "island" border.

**Figure 6.** VLM image of (**a**) EBM and (**b**) SLM samples before etching. (**c**) VLM image of SLM sample after etching. All images are taken from the xy-plane (or build plane) of samples.

#### *3.4. Microstructures of the Build Material*

Figure 7 illustrates typical microstructural features from three perpendicular planes of SLM (first row) and EBM (second row) processed Ti55511 alloy. The orientation of the planes is indicated, and the axes are given in Figure 5b. In the xz surface plane of the SLM sample (Figure 7a), individual scan tracks and molten pool boundaries with the typical arc-shaped configuration are observed. The tracks are about 60–110 μm in thickness and 100–200 μm in width and are produced by the Gaussian-like energy distribution of the laser.

Parallel to the building direction (Z) columnar grains are visible in both SLM and EBM samples. These columns are much larger than the layer thickness. The width of the columnar grains is smaller in the SLM built sample than in the EBM (Figure 7c,f). In the xy-plane the cross-section of the columns was equiaxed in both samples (Figure 7b,e). Despite significant alloy composition differences, in the EBM sample fine grains with lamella and Widmanstätten-like structure (Figure 7d–f) similar to EBM, processed Ti-6Al-4V can be observed within the large columns. A similar microstructure is observed in the used powder (Figure 3b). Figure 8 shows an EBSD phase map and an inverse pole figure taken from the xz-plane of the EBM sample revealing the presence of about 67 wt% α phase and 33 wt% β phase. Areas of similar orientation that were likely to have originated from the same parent β grain can be recognized.

This microstructure and phase differences between EBM and SLM build samples can be ascribed to the difference in build plate temperature and cooling rates applied. Because of faster scanning speed and a high chamber temperature of the EBM, the cooling rate will be slower than the SLM. Therefore, the average temperature of the melt pool and heat affected zone in EBM will be relatively higher than in SLM which will consequently lead to grain growth and β→α transformation.

**Figure 7.** Microstructure of SLM (three images in first/top row) and EBM (all three images in second/bottom row) build samples. (**a**,**d**) are VLM image from the xz-plane; (**b**,**e**) are from the xy-plane; (**c**,**f**) are EBSD pattern from the xz-plane.

**Figure 8.** High-resolution EBSD (**a**) phase map and (**b**) IPF of EBM build sample.

Similar to the powder analysis, the chemical composition of the as-built samples was assessed by EDS analysis. The elemental map of the as-built SLM sample does not show any distinct regions and they all seem homogeneous and featureless. In the as-built EBM sample, however, there are two distinct regions in the elemental map like what was observed in the EBM powders after processing.

The XRD analysis, performed on the xy-plane of the as-built samples, confirmed that the SLM consists predominantly of beta phase but reveals also the presence of a limited amount of α or α' phase (<3 wt%) as shown in Figure 9a. Because of limited quantity and the almost identical unit cell parameters of α and α' phase, it is not possible to distinguish between them. According to the literature, the α' phase commonly associated with the extreme temperature change is usually observed in the SLM build Ti-6Al-4V [1,8,17]. The α phase, on the other hand, is associated with an isothermal cooling condition [2,10,17,31] and is commonly observed in the EBM build Ti-6Al-4V. Therefore, it is believed that that the minority phase present in the SLM build sample might be α' phase.

**Figure 9.** (**a**) Diffraction pattern of as-built samples. These color-coded tick marks under the diffractogram correspond to the expected peak positions of reported α- and β-Ti phases. (**b**) Rietveld refinement result of XRD data from EBM build sample.

The XRD pattern of the EBM sample confirms the presence of α and β phases. Rietveld analysis (Figure 9b) was performed and the lattice constants of hexagonal close packed α phase determined to be *a* = 2.931(7) Å and *c* = 4.658(8) Å, respectively, with *c*/*a* ratio of 1.5891. For the body-centered cubic β phase, *a* = 3.230(9) Å. As shown in the figure, the phase fraction of the α phase is 64.82 wt% and the β phase is 35.18%. This result supports the observation from the SEM.

The mechanical properties of a Ti-alloys are strongly determined by their microstructure and phase composition. Measured microhardness of various samples seemed to match with literature. The mean microhardness profiles of the SLM and EBM build samples were 315 ± 3.41 HV and 348 ± 30.20 HV, respectively. Increased microhardness of the EBM build sample can be attributed to the amount of α phase present. From the values of standard deviations of microhardness, it is evident that the SLM build sample has relatively consistent microhardness while the EBM samples show a larger deviation. This variation indicates that the SLM sample has a uniform microstructure throughout while the EBM sample has an inhomogeneous microstructure. This agrees well with the microstructure presented in Figures 7 and 8. Here it is also important to point out that despite apparent differences in alloy chemical composition, the microhardness of the EBM build sample is very similar to α phase dominated EBM build Ti-6Al-4V reported by Neikter et al. [32]. Such similarity can be explained by the similar α + β microstructure observed in both types of alloys after the EBM process. The microhardness of the EBM build sample seemed to slightly lower (but same in standard deviation) than what Kurzynowski et al. [13] reported for EBM build Ti55511, previously. As hardness of the material is the resistance of the material to plastic deformation, one can expect that the EBM build sample may show higher strength than the SLM build samples. However, this needs further investigation.

#### **4. Conclusions**

The Ti55511 parts were successfully fabricated using SLM and EBM techniques. The results indicate that both methods can process Ti55511 and achieve almost full density with limited porosity. For the process parameters used in this study, the SLM process gives a slightly lower density and better surface quality. The shape of the pores in the EBM production appears to be mostly spherical while more random shapes are observed in SLM. The dominant β phase in the original powder becomes a minority phase after EBM processing, while there is almost no phase transformation in the SLM. Because of high build temperature and relatively slow cooling rate, the EBM build samples show a lamella and Widmanstätten-like structure similar to the microstructure observed in EBM processed Ti-6Al-4V despite it is alloy composition. Because of the lamella microstructure, the EBM build sample showed about 10% higher microhardness than the SLM build samples and it measured as 348 ± 30.20 HV. To achieve near-β phase composition in Ti55511 after EBM processing, the current processing route needs to be optimized or the build part needs additional post-heat treatments.

**Author Contributions:** Conceptualization, T.M.; Data curation, T.M.; Formal analysis, T.M. and H.V.S.; Funding acquisition, H.V.S.; Investigation, T.M. and K.M.; Methodology, T.M., K.M. and N.C.; Resources, K.M., T.K. and H.V.S.; Supervision, T.K. and H.V.S.; Visualization, T.M.; Writing—original draft, T.M. and H.V.S. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded by the project PREMPA, a Strategic Focus Area project of the ETH board in Switzerland.

**Acknowledgments:** Tuerdi Maimaitiyili and Van Swygenhoven Helena thank the financial support of the project PREMPA, a Strategic Focus Area project of the ETH board. Tuerdi Maimaitiyili thanks Haydous Fatima, Martin Elsener, and Miroslav Smid at the Paul Scherrer Institut (PSI) for their help with surface roughness, PSD, and EBSD measurements, respectively. Tuerdi Maimaitiyili also thanks Christof W. Schneider at PSI for the XRD measurements of as-built specimens.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**


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