*Article* **Simulation of Ti-6Al-4V Additive Manufacturing Using Coupled Physically Based Flow Stress and Metallurgical Model**

#### **Bijish Babu 1,***∗***, Andreas Lundbäck <sup>2</sup> and Lars-Erik Lindgren <sup>2</sup>**


Received: 3 October 2019; Accepted: 18 November 2019; Published: 21 November 2019

**Abstract:** Simulating the additive manufacturing process of Ti-6Al-4V is very complex due to the microstructural changes and allotropic transformation occurring during its thermomechanical processing. The *α*-phase with a hexagonal close pack structure is present in three different forms—Widmanstatten, grain boundary and Martensite. A metallurgical model that computes the formation and dissolution of each of these phases was used here. Furthermore, a physically based flow-stress model coupled with the metallurgical model was applied in the simulation of an additive manufacturing case using the directed energy-deposition method. The result from the metallurgical model explicitly affects the mechanical properties in the flow-stress model. Validation of the thermal and mechanical model was performed by comparing the simulation results with measurements available in the literature, which showed good agreement.

**Keywords:** dislocation density; vacancy concentration; Ti-6Al-4V; additive manufacturing; directed energy deposition

#### **1. Introduction**

Powder Bed Fusion (PBF) is the technique of building thin layer over layer by melting the fine metal powder. Directed energy deposition (DED), on the other hand, is usually used for building features on large existing parts as well as for repairing damaged ones. PBF typically adds layers that are thinner than DED and can therefore create high-resolution structures, whereas DED produces components at a higher built rate. The primary challenge of DED is that the higher energy input from the heat source may lead to substantial distortion and higher residual stresses.

DED additive manufacturing (AM) can be considered as computer numerically controlled (CNC) multipass welding with progressive weldments made on a substrate to create free-form structures. The added metals can be in either powder or wire form and the heat source a laser or electron beam. The deposition path is generated from computer-aided design (CAD) geometry and is preprogrammed in a CNC machine, which makes the process very flexible and suitable for low volume production, eliminating the need for tooling and dies.This also enables the production of complicated geometries that are traditionally difficult to produce with conventional manufacturing processes. Additively manufactured parts of Ti-6Al-4V are traditionally found in human implants [1] and aerospace components because of the criticality of their applications. However, AM has also been used to repair aerospace components [2] that have developed defects during operation or production.

A few researchers have performed AM simulations or similar processes for Ti alloys using thermomechanical–microstructural (TMM) coupled material models. In Baykasoglu et al. [3], a thermomicrostructural model for Ti6Al-4V was presented and applied to a DED process. Salsi et al. [4]

presented a similar model and applied it on a PBF process, while Vastola et al. [5] compared the results when modelling electron-beam melting (EBM) and PBF processes. Song et al. [6] performed a welding simulation by using a TA15 alloy employing a TMM model. A similar model was utilized for performing a quenching simulation by Teixeira et al. [7] for alloy Til7. Cao et al. [8] showed an AM simulation using electron-beam melting without including microstructural coupling. A TMM material model was employed by Ahn et al. [9] for welding simulation ignoring strain-rate dependence.

In this work, a material model combining metallurgical and flow-stress models described by Babu et al. [10] is used. This model works for arbitrary phase composition and is an improved version of that of Babu and Lindgren [11]. The AM process involves cyclic heating and cooling, resulting in nonequilibrium phase evolution, which can be addressed with this model. The metallurgical model used in this work was also utilized in the simulation of the AM case described by Charles Murgau et al. [12], which is included in the current special issue.

#### **2. Physically Based Flow-Stress Model**

An incompressible von Mises model was used here with the assumption of isotropic plasticity. Flow stress was split into two parts [11,13–15]:

$$
\sigma\_{\mathcal{Y}} = \sigma\_{\mathcal{G}} + \sigma^\*.\tag{1}
$$

Here, *σ<sup>G</sup>* is a thermal stress contribution from long-range interactions of the dislocation substructure. The other term, *σ*∗, is the required friction stress to move dislocations within the lattice and to cross short-range barriers. Thermal vibrations can assist dislocations to overcome these barriers. Conrad [16] proposed a similar formulation after analyzing titanium systems.

#### *2.1. Long-Range Stress Component*

The long-range term from Equation (1) is derived from Seeger [13] as

$$
\sigma\_{\mathbb{G}} = \text{maG} \sqrt{\rho\_i}.\tag{2}
$$

Here, *m* is the Taylor factor that translates the resolved shear stress in various slip systems to effective stress, *b* is the magnitude of Burgers vector, *G*(*T*) is the temperature-dependent shear modulus, *ρ<sup>i</sup>* is the immobile dislocation density and *α*(*T*) is a calibrated proportionality factor.

#### *2.2. Short-Range Stress Component:*

The strain-rate-dependent part of the yield stress from Equation (1) can be derived according to the Kocks–Mecking formulation [17,18] as

$$
\sigma^\* = \tau\_0 G \left[ 1 - \left[ \frac{kT}{\Delta f\_0 G b^3} \ln \left( \frac{\dot{\varepsilon}^{ref}}{\xi^p} \right) \right]^{1/q} \right]^{1/p} . \tag{3}
$$

Here, shear strength in the absence of thermal energy is denoted by *τ*0*G*, and the activation energy required to overcome lattice resistance is denoted by Δ*f*0*Gb*3. Parameters *p* and *q* define the shape of the obstacle barrier for dislocation motion. Further, *k* is the Boltzmann constant, *T* is the temperature in kelvin and (*ε*˙ *ref*) and (*ε*˙ *<sup>p</sup>*) are the reference and plastic strain rates, respectively.

#### *2.3. Evolution of Immobile Dislocation Density*

The evolution of *ρ<sup>i</sup>* in Equation (2) is modelled as having two components, hardening and restoration.

$$
\rho\_i = \rho\_i^{(+)} - \rho\_i^{(-)}.\tag{4}
$$

#### 2.3.1. Hardening Process

The average distance moved by dislocations before they are annihilated or immobilized is called mean free path Λ. The Orowan equation shows that the density of dislocations and their average velocity are proportional to the plastic strain rate. Assuming that the immobile dislocation density also follows the same relation leads to

$$
\rho\_i^{(+)} = \frac{m}{b} \frac{1}{\Lambda} \mathbb{1}^p. \tag{5}
$$

The mean free path is computed from grain size (*g*) and dislocation subcell or subgrain diameter (*s*) as

$$\frac{1}{\Lambda} = \frac{1}{\mathcal{g}} + \frac{1}{\mathcal{s}}.\tag{6}$$

The subcell formation and evolution are modelled using a relation proposed by Holt [19].

$$s = K\_c \frac{1}{\sqrt{\rho\_i}}.\tag{7}$$

#### 2.3.2. Restoration Processes

Vacancy motion is relevant to the recovery of dislocations. Restoration of the lattice commonly occurs at elevated temperatures and is therefore a thermally activated restructuring process. Creation of vacancy requires energy and increases entropy. With increasing temperature and deformation, vacancy concentration also increases. High stacking fault materials usually exhibit constant flow stress because of the balance between hardening and recovery. The current model assumes that the mechanisms of restoration are dislocation glide, dislocation climb and globularization.

$$
\rho\_i^{(-)} = \rho\_i^{(glide)} + \rho\_i^{(climb)} + \rho\_i^{(glolulariization)}.\tag{8}
$$

The model for recovery by glide can be written on the basis of the formulation by Bergström [20] as

$$
\rho\_i^{(gldc)} = \Omega \rho\_i \pounds^p,\tag{9}
$$

where Ω is a function dependent on temperature.

Militzer et al. [21] proposed a model for dislocation climb on the basis of Sandström and Lagneborg [22] and Mecking and Estrin [23]. With a modification of diffusivity according to Reference [11], the model can be written as

$$\rho\_i^{(climb)} = 2c\_\gamma D\_{app} \frac{Gb^3}{kT} \left(\rho\_i^2 - \rho\_{c\eta}^2\right),\tag{10}$$

where *c<sup>γ</sup>* is a material coefficient and *ρeq* is the equilibrium value of the dislocation density. Here, *Dapp* is the apparent diffusivity that includes the diffusivity of the *α* − *β* phases weighted by their fractions *X<sup>α</sup>* and *Xβ*, pipe diffusion *Dp*, as well as effects of excess vacancy concentration *cv*.

Babu and Lindgren [11] proposed a model for the evolution of dislocation density during globularization where the effect of grain growth on the reduction of flow stress is only included when dislocation density is above a critical value *ρcr*.

$$\begin{array}{ll}\text{if} & \rho\_{\text{i}} \ge \rho\_{\text{c7}}\\ & \dot{\rho}\_{\text{i}}^{\left(\text{globalization}\right)} = \psi \dot{X}\_{\text{\mathcal{S}}} \left(\rho\_{\text{i}} - \rho\_{\text{eq}}\right); \text{ until } \rho\_{\text{i}} \le \rho\_{\text{eq}} \end{array} \tag{11}$$

else

$$
\phi\_i^{\,(glubularization)} = 0.\tag{12}
$$

Here, *ρeq* is the equilibrium value of dislocation density, *X*˙ *<sup>g</sup>* is the globularization rate and *ψ* is a calibration constant. Thomas and Semiatin [24] modelled the two-stage process of dynamic and static recrystallization. Owing to the similarities between globularization and recrystallization, this model can be adapted.

$$X\_{\mathfrak{F}\_{\mathfrak{F}\_{\mathfrak{F}}}} = \; \; X\_d + (1 - X\_d) \; X\_{\mathfrak{k}}.\tag{13}$$

Here, volume fractions *Xg*, *Xd*, and *Xs* denote total globularization, its dynamic component and the static component, respectively.

Assuming that grain growth and static recrystallization have the same driving force, the static globularization rate can be modelled as [25,26]

$$
\dot{X}\_{\\$} \quad = \quad M \frac{\dot{\mathcal{G}}}{\mathcal{G}'} \tag{14}
$$

where, *M* is a material parameter. The rate of dynamic globularization was computed on the basis of a model by Thomas and Semiatin [24] as,

$$\mathcal{X}\_d \quad = \frac{Bk\_p \mathcal{E}\_p^\*}{\varepsilon\_p^{r} \boldsymbol{1}^{1-k\_p} \boldsymbol{e}^{\operatorname{B} \mathcal{E}\_p^{k\_p}}} \prime \tag{15}$$

where, *B* and *kp* are material parameters.

#### *2.4. Evolution of Excess Vacancy Concentration*

The formation and evolution of excess vacancy concentration was modelled by Militzer et al. [21]. In the current work, Militzer's model was extended by adding the effect of temperature changes. Further, assuming that only long-range stress contributes to vacancy formation, the model can be rewritten as

$$\mathbf{c}\_{v}^{\prime cx} = \left[ \chi \frac{maGb^2 \sqrt{\rho\_i}}{Q\_{vf}} + \zeta \frac{c\_j}{4b^2} \right] \frac{\Omega\_0}{b} \pounds - D\_{vm} \left[ \frac{1}{s^2} + \frac{1}{g^2} \right] \left( c\_v - c\_v^{eq} \right)$$

$$+ c\_v^{\alpha \eta} \left( \frac{Q\_{vf}}{kT^2} \right) \mathcal{T}. \tag{16}$$

Here, *χ* = 0.1 is the fraction of mechanical energy spent on vacancy generation, Ω<sup>0</sup> is the atomic volume and *ζ* is the neutralization effect by vacancy emitting and absorbing jogs. The concentration of jogs (*cj*) and *Dvm* and the diffusivity of vacancy are given in Babu and Lindgren [11]. Additionally, *Qv f* is the activation energy of vacancy formation.

#### **3. Phase-Evolution Model**

A simplified model [27] for the transition between the liquid and solid state was implemented to take care of temperatures above melting temperature *Tmelt*. If the temperature is above *Tmelt*, the volume fraction of the solid phases was set to zero. In the solid state, the Ti–6Al–4V microstructure comprises the high-temperature stable *β*-phase and the lower-temperature stable *α*-phase. Depending on temperature and heating/cooling rates, a variety of *α*/*β* morphologies can form that gives varying mechanical properties. The complex relationship between thermomechanical-processing, microstructure and mechanical properties was investigated by References [28,29]. On the basis of the literature [30–33], few microstructural features have been identified as relevant concerning mechanical properties. The three separate *α*-phase fractions, Widmanstatten (*Xα<sup>w</sup>* ), grain boundary (*Xαgb* ), acicualr and massive Martensite (*Xα<sup>m</sup>* ) and *β*-phase fraction (*Xβ*) were included in the current model. Though

in the current flow-stress model individual *α*-phase fractions were not included separately, it is possible to incorporate them when more details about their respective strengthening mechanisms are known.

#### *3.1. Phase Transformations*

Depending on the temperature and heating/cooling rates, Ti-6Al-4V undergoes allotropic transformation. The mathematical model for the transformation is described schematically in Figure 1. Transformations denoted by F1, F2, and F3 represent the formation of *αgb*, *αw*, and *α<sup>m</sup>* phases, respectively, and D3, D2 and D1 show the dissolution of the same phases. If the current volume fraction of *β* phase is more than *βeq*, the excess *β* phase transforms into an *α* phase. Here, *αgb* formation that occurs in high temperatures is the most preferred, followed by the *αw*. Remaining excess *β* fraction is transformed to *α<sup>m</sup>* if the temperature is lower than *Tm*, (martensite start temperature) and cooling rate is above 20◦C/s. Conversely, if the current volume fraction of *β* is lower than *βeq*, excess *α* phase is converted to *β*. Primarily, the *α<sup>m</sup>* phase dissolves to *β* and *α<sup>w</sup>* phases in the same proportion as the *αeq* and *βeq*. Remaining excess *α<sup>w</sup>* and *αgb* transform to *β* in that order. The equilibrium fraction of the *β* phase (see Figure 2) is computed by Equation (17), where *T* is the temperature in degrees Celsius.

$$X\_{\tilde{\beta}}^{eq} = 1 - 0.89 \,\mathrm{e}^{-\left(\frac{T^\* + 1.82}{1.75}\right)^2} + 0.28 \,\mathrm{e}^{-\left(\frac{T^\* + 0.59}{0.67}\right)^2}$$

$$T^\* = (T - 927) / 24. \tag{17}$$

**Figure 1.** Phase-change mechanism.

**Figure 2.** Thermal conductivity (K), specific heat capacity (C) and equilibrium-phase fraction (*Xβ*).

*3.2. Adaptation of Johnson–Mehl–Avrami–Kolmogrov (JMAK) Model for Diffusional Transformation*

The JMAK model [34–36], originally formulated for nucleation and growth during isothermal situations, can be adapted to model any diffusional transformation. Employing the additivity principle and using sufficiently small time steps ensures that any arbitrary temperature change can be computed. The JMAK model assumes that a single phase *X*<sup>1</sup> that is 100% in volume from the start transforms to 100% of second phase *X*<sup>2</sup> in infinite time. However, in the case of Ti-6Al-4V, this is not the case, as it is a *α* − *β* dual-phase alloy below *β*-transus temperature. Hence, in order to accommodate for an incomplete transformation, the product fraction is normalized with the equilibrium volume. Conversely, the starting volume of a phase can also be less than 100%, which is circumvented by assuming that the available phase volume is the total phase fraction. Another complication is the existence of the simultaneous transformation of various *α* phases (*αw*, *αgb*, *αm*) to the *β* phase and back. This can be modelled by sequentially calculating each transformation within the time increment [27].

#### *3.3. Formation of α Phase*

During cooling from the *β*-phase, the *αgb* and *α<sup>w</sup>* phases are formed by diffusional transformation. According to the incremental formulation of the JMAK model described by Reference [27], the formation of *αgb* and *α<sup>w</sup>* can be modelled by the set of equations in rows F1 and F2, respectively, of Table 1. The Martensite phase is formed at cooling rates above 410 °C/s by diffusion-less transformation. While cooling at rates above 20 ◦C/s and up to 410 ◦C/s, massive *α* transformation was observed to co-occur with Martensite formation [37,38]. Owing to the similitude in crystal structures between massive-*α* and Martensite-*α*, they are not differentiated here except that, above 410 °C/s, 100% *α<sup>m</sup>* was allowed to form. An incremental formulation of the Koistinen–Marburger equation described by Charles Murgau et al. [27] was used here (see equation set in row F3 of Table 1).

**Table 1.** Models for *α*-phase formation.


#### *3.4. Dissolution of α Phase*

The *α<sup>m</sup>* phase formed by instantaneous transformation is unstable and therefore undergoes diffusional transformation to the *α<sup>w</sup>* and *β* phases on the basis of its current equilibrium composition. The incremental formulation of the classical JMAK model by Reference [27] and its parameters are given in row D1 of Table 2. During heating or reaching nonequilibrium phase composition, *α<sup>w</sup>* and *αgb* can transform into a *β*-phase controlled by the diffusion of vanadium at the *α* − *β* interface. A parabolic equation developed by Kelly et al. [39,40] derived in its incremental form by Charles Murgau et al. [27] was used here (see rows D2 and D3 of Table 2).


**Table 2.** Models for *α*-phase dissolution.

#### **4. Coupling of Phase and Flow-Stress Models**

Young's modulus and Poisson's ratio were assumed to be identical for both phases. The Wachtman model [41] for Young's modulus (*E*), calibrated using measurements from Babu and Lindgren [11], is written as

$$E = 107 - 0.2 \left( T + 273 \right) e^{-\left( 1300/T + 273 \right)},\tag{18}$$

where *T* is the temperature in degrees Celsius applied with a cut-off at *T* = 1050 °C (see Figure 3). A linear model for Poisson's ratio (*μ*) after fitting to measurements by Fukuhara and Sanpei [42] as

$$
\mu \quad = \quad 0.34 + 6.34 \ e^{-5} \ T,\tag{19}
$$

where *T* is the temperature in degrees Celsius (see Figure 3).

**Figure 3.** Poisson's ratio (*μ*), thermal strain ( *th*) and Youngs modulus (E).

Using X-ray diffraction, Swarnakar et al. [43] measured the volumetric expansion of unit cells of *α* and *β* phases during heating. On this basis, the average Coefficient of Thermal Expansion (CTE) of the phase mixture can be calculated using the rule of mixtures (ROM) as in Equation (20), where *αα* and *αβ* give the CTE of *α* and *β* phases, respectively. The linear thermal strain can be computed using Equation (21), plotted in Figure 3. Here, *εadj* makes the ROM (Equation (20)) nonlinear.

$$
\mathfrak{a}\_{\text{avg}} \quad = \; X\_{\mathfrak{a}} \mathfrak{a}\_{\mathfrak{a}} + X\_{\mathfrak{f}} \mathfrak{a}\_{\mathfrak{f}} \tag{20}
$$

$$
\varepsilon^{th} = \pi\_{avg} \Delta T - \varepsilon^{adj} \tag{21}
$$

$$
\epsilon^{\text{adj}} \quad = \quad 1.0 \epsilon^{-8} T^2 - 8.4 \epsilon^{-6} T + 3.0 \epsilon^{-4}. \tag{22}
$$

The thermal conductivity and specific heat capacity of the alloy taken from References [44] and [45], respectively, are given in Figure 2. The latent heat of phase transformation (*α* → *β*) and the latent heat of fusion were measured to be 64 and (290 ± 5) kJ/Kg, respectively [44].

The yield strength of the phase mixture can be written according to the linear rule of mixtures as

$$
\sigma\_{\mathcal{Y}} = X\_{\mathfrak{a}} \sigma\_{\mathcal{Y}}^{\mathfrak{a}} + X\_{\mathfrak{f}} \sigma\_{\mathcal{Y}}^{\mathfrak{f}}.\tag{23}
$$

The plastic strain distribution can be obtained assuming an iso-work principle. According to Reference [46], this can be written as

$$
\sigma\_y^a \dot{\varepsilon}\_\alpha \quad = \quad \sigma\_y^\beta \dot{\varepsilon}\_\beta \tag{24}
$$

$$
\dot{\varepsilon}^p \quad = \quad X\_a \dot{\varepsilon}\_a + X\_{\beta} \dot{\varepsilon}\_{\beta}. \tag{25}
$$

The above formulation ensures that the *β* phase with lower yield strength has a more significant share of plastic strain as compared to the stronger *α* phase. For temperatures above 1100 °C, (*σ*><sup>1100</sup> ◦<sup>C</sup> *<sup>y</sup>* = *σ*<sup>1100</sup> ◦<sup>C</sup> *<sup>y</sup>* ). The stress–strain relationship predicted by the model for varying strain rates and temperatures are given in Figure 4. The rate dependence and flow-softening demonstrated by the model is visible here. A detailed comparison of model predictions and measurements along with model parameters are given in Babu et al. [10,11].

**Figure 4.** Stress–strain–temperature relationship.

#### **5. Additive Manufacturing**

In this article, a DED process described in Reference [47] was simulated using general-purpose Finite Element (FE) software MSC.Marc. A set of subroutines for modelling the AM process were implemented in MSC.Marc, which are explained in Lundbäck and Lindgren [48]. A coupled thermomechanical–metallurgical model described in the previous sections was also implemented as subroutines within MSC.Marc.

The dimensions of the substrate (152.4 × 38.1 × 12.7 mm) and AM component are given in Figure 5. One end of the fixture was held in position by a clamping fixture (see notations in Figure 5).

**Figure 5.** Dimensions of additive-manufacturing (AM) component.

Three beads were added per layer with a total width of 6.7 mm and a height of 38.1 mm; 42 discrete layers and their respective beads are also shown in Figure 5. Figure 6 shows the order of deposition starting with the middle bead, followed by one on each side. Odd layers are deposited in the direction away from the clamping, followed by even layers in the opposite direction. Figure 6 also shows the temperature contours and direction of deposition of bead three of the first layer. After each layer, a dwell time (DT) of 0, 20 and 40 s was applied for studying the effect of varying cooling rates.

**Figure 6.** Temperature contours in weld pool.

#### **6. Modelling of Additive Manufacturing**

In the current work, the scope of the model was to predict microstructure evolution, the overall distortion of the component and residual stresses. This requires a solution where thermomechanical–metallurgical models are coupled by using a staggered approach. Figure 7 shows the coupling of different domains using a staggered approach. The thermal field is solved using the FE implicit iterative scheme by computing heat input and heat losses by conduction, convection and radiation. On the basis of the computed temperature in the increment, the metallurgical model computes the phase evolution for each Gauss point. The computed temperature and phase fractions are input to solve the mechanical-field equations. A large-deformation FE implicit iterative scheme was used here, where mechanical and physical properties are strongly dependent on the temperature and phase composition. Latent heat and volume changes due to phase evolution and deformation energy converted to heat are included here.

**Figure 7.** Coupling of thermomechanical–microstructural fields.

#### *6.1. Heat Source*

Modelling of weld-pool phenomena requires high-resolution discretization and at least one other physics domain, namely, fluid flow. This requires an impractical amount of resources for solving the problem and can be avoided considering the scope of the current work. The heat input can be modelled by using volume heat flux in a geometric region representing the weld pool, and is calibrated using measured temperatures. Goldak's [49] double ellipsoidal heat-input model was implemented in the current work, with efficiency calibrated to be 0.29. The parameters of the heat source are given in Figure 8. See Lundbäck and Lindgren [48] for details on the implementation of this model.

**Figure 8.** Gaussian distribution of density and double ellipsoid shape in xy plane.

#### *6.2. Modelling of Material Addition*

The inactive-element approach was used here, where all elements representing the added metal were deactivated before the start of the simulation, and only activated after meeting certain criteria. In each increment, the set of elements that overlapped and the geometric region represented by the current weld-pool position were thermally activated, whereas mechanical activation occurred only when the thermally active elements cooled below the solidus temperature. Before thermal activation, the elements may have had to be moved to accommodate for the distortion of the substrate and the already activated elements during the process. The moving elements maintained connectivity with the activated elements, and their volume matched the added material during that time step [48].

#### *6.3. Boundary Conditions*

In DED, much of the heat input in the initially deposited layers is absorbed by the substrate. To balance the heat input, losses by free and forced convection, and conduction to fixtures as well as radiation were included in the model. A lumped convective coefficient of 18 W/m2/°C was applied to model the natural and forced convection from shielding gas. Both convective and radiative boundary conditions were applied on the outer surface of all thermally active elements. A surface emissivity of 0.25 was used here. In the current model, heat losses due to cooling by the fixture were achieved by using convective heat transfer with a high coefficient of 250 W/m2/°C.

#### **7. Comparison of Measurements and Simulations**

In situ measurement of temperature and distortion was done during the AM process. Three thermocouples were attached to the bottom of the substrate at the positions shown in the left part of Figure 5. DTs allowed the component to considerably cool down during the process. In Figure 9a,c,e, dots denote the thermocouple measurements and lines, predictions from the model. The thermocouple attached to the middle of the component (TC2) registered the highest temperature, as the other two were closer to the ends that are subjected to higher convective cooling. The thermocouple attached close to the clamping (TC3) had the lowest temperature since the fixture acts as a heat sink. Raising dwell times by 20 s increased cooling, thereby reducing the peaks. As the height of the wall increased, this effect was less detectable, as this thermocouple was beneath the substrate.

**Figure 9.** Comparison of Measurements and Simulations.

The component distortion is measured at the free end by using a laser displacement sensor. In Figure 9b,d,f, red lines denote the measured values, and the blue lines denote predictions from the model. The addition of each layer made the component bend downwards due to the thermal gradient between the top and bottom of the substrate, which is diminished during cooling, producing oscillations. In order to compare measurement and simulation results, these oscillations were smoothed out by using the Savitzky–Golay filter [50], and are plotted in Figure 10. Here, dotted lines denote measurements, and continuous lines denote model predictions. The start and end of the linear region, its slope and the detection of the first peak can be deduced from Figure 10 and are shown in Table 3. The peak-to-peak amplitude of the oscillations at the first peak, and also at the start and end of the linear region are given in Table 3. The measured final bending of the build plate at the outer edge increased by 0.4 mm for each 20 s increase in dwell time. Simulations also gave a similar result. Simulation results are given in Table 3 within brackets.

**Figure 10.** Comparison of distortions.

**Table 3.** Comparison of distortion (computed values given in brackets).


Figure 11 shows the residual stress in the welding direction along with its spread measured by Reference [47] by using hole drilling. The testing location was in the middle of the specimen at the bottom of the substrate. Results showed that, for the case with dwell times of 0 and 20 s, simulation results were close to the measurements or within the margin of error; for 40 s, it was slightly below the margin. Residual-stress distribution after cooling to room temperature in the welding direction and the von Mises effective stress for the case with dwell time of 0 s are plotted in Figure 12. The predicted temperature for the case with dwell time of 0 s at the top surface of the substrate above the location of TC2 is given in Figure 13. The computed *α*-phase fraction is also provided here. The addition of each bead is denoted as B1-3, and grey areas in between are the cooling time. In total, five layers are shown in Figure 13.

**Figure 11.** Residual stress.

**Figure 12.** Residual stress for 0 s dwell time (model clipped longitudinally in midplane).

**Figure 13.** Prediction of phase fraction and temperature.

#### **8. Discussion**

The yield strength of the material is very much dependent on state variables of dislocation density and excess vacancy concentration. The density of these state variables changes by many orders of magnitude during heating and deformation. Deformation increases dislocation density and results in hardening, whereas an increase in vacancy concentration (due to heating and deformation) assists in the remobilization of dislocations, thereby material recovery. The advanced material model described here combining the metallurgical and flow-stress models has proven to be suitable for AM simulation. Diffusional and instantaneous transformations are included in the metallurgical model. This model was formulated in a way that it could be implemented in any standard kind of finite-element software. Temperature measurements and results from the simulations demonstrated a good overall fit. This model predicted the final distortion of the component with good accuracy except for the case with 40 s dwell time. This trend was also evident in the residual-stress measurements. The comparison of distortion before the onset of cooling showed a larger difference between model and measurements. This might be because the stress-relaxation behavior was less accurately predicted by the model. The computed phase fraction in Figure 13 showed that, after the addition of the fourth layer, the substrate underwent no significant phase evolution. Temperature peaks after 120 s had slightly less magnitude, and were therefore below the cutoff levels to introduce any phase change. Denlinger and Michaleris [51] performed the simulation of all the three cases described here. They used an approach where the plastic strain was reset to zero at a temperature of 690 °C which is a parameter calibrated for that particular AM case. This transformation-strain parameter made it possible for Denlinger and Michaleris [51] to include the effects of dwell time, whereas in the current work, mechanisms of dislocation climb and globularization resulted in the restoration of the lattice.

#### **9. Conclusions**

One of the challenges involved in the AM process is residual deformation and stresses due to the thermal dilatation of the substrate and added structure. The final properties of the AM structure are strongly influenced by the microstructure, which is dependent on the thermomechanical-processing history of the component. For the industry to fully adopt additive manufacturing and be able to qualify titanium parts for critical applications, such as in aerospace, a complete understanding of the microstructure properties and mechanical behavior is necessary. This paper showed the implementation and application of a coupled microstructural–thermal–mechanical model to an AM process. A physically based constitutive model was explicitly coupled to the microstructural model. The phase composition predicted by the microstructural model therefore affected the mechanical properties, namely, flow strength, in a direct way. Validation of the thermal and mechanical model was performed by comparing the simulation results with the available measurements in the literature. The comparison had good agreement between the results from the model and the measurements.

**Author Contributions:** Conceptualization, B.B. and A.L.; methodology, B.B. and A.L.; software, B.B. and A.L.; validation, B.B. and A.L.; formal analysis, B.B. and A.L.; original-draft preparation, B.B.; visualization, B.B.; supervision, L.-E.L.

**Funding:** This research was partly funded by the Swedish Foundation for Strategic Research (SSF), Development of Processes and Material in Additive Manufacturing, Reference number GMT14-0048.

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

#### **Abbreviations**

The following abbreviations are used in this manuscript:


#### **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* **Fatigue Crack Growth of Electron Beam Melted Ti-6Al-4V in High-Pressure Hydrogen**

**M. Neikter 1,4, M. Colliander 2, C. de Andrade Schwerz 3, T. Hansson 3,4, P. Åkerfeldt 1, R. Pederson 4,\* and M.-L. Antti <sup>1</sup>**


Received: 26 February 2020; Accepted: 10 March 2020; Published: 12 March 2020

**Abstract:** Titanium-based alloys are susceptible to hydrogen embrittlement (HE), a phenomenon that deteriorates fatigue properties. Ti-6Al-4V is the most widely used titanium alloy and the effect of hydrogen embrittlement on fatigue crack growth (FCG) was investigated by carrying out crack propagation tests in air and high-pressure H2 environment. The FCG test in hydrogen environment resulted in a drastic increase in crack growth rate at a certain ΔK, with crack propagation rates up to 13 times higher than those observed in air. Possible reasons for such behavior were discussed in this paper. The relationship between FCG results in high-pressure H2 environment and microstructure was investigated by comparison with already published results of cast and forged Ti-6Al-4V. Coarser microstructure was found to be more sensitive to HE. Moreover, the electron beam melting (EBM) materials experienced a crack growth acceleration in-between that of cast and wrought Ti-6Al-4V.

**Keywords:** fatigue crack growth (FCG); electron beam melting (EBM); Ti-6Al-4V; hydrogen embrittlement (HE)

#### **1. Introduction**

Oxygen and hydrogen are used as fuel for space rockets [1]. Hydrogen is combusted rapidly in an oxygen-rich environment and a high propelling force can be achieved, as the energy density for hydrogen is high (142 MJkg<sup>−</sup>1) [2]. Although hydrogen is excellent for combustion it can cause hydrogen embrittlement (HE), which deteriorates the mechanical properties. HE occurs at or ahead of crack tips because at these locations tri-axial stresses are present [3]. These stresses render slightly expanded lattices, making it more energetically favorable for the hydrogen to diffuse to this location. Once at the crack tip the hydrogen causes degraded properties due to one or several HE mechanisms [4,5].

Additively manufactured Ti-6Al-4V has achieved large interest within the space industry since it can reduce weight and lead time. Ti-6Al-4V is a dual-phase alloy, consisting of both α and β phase [6]. In the α phase, which has a hexagonal close-packed crystal structure, the diffusion rate of hydrogen is lower compared to the body-centered cubic β phase [7]. The amount of β phase at room temperature is not as high compared to the α phase. The microstructure is different for Ti-6Al-4V material built with the additive manufacturing (AM) process electron beam melting (EBM) compared to conventionally wrought or cast material [8,9]. In Ti-6Al-4V manufactured with EBM, the prior β grains grow epitaxially towards the heat source, which renders a columnar structure perpendicular to the built layers [9–12], a unique morphology that is observed in neither cast nor wrought titanium. At the grain

boundary of the β phase, there is nucleation of α phase when the temperature is reduced below the β transus temperature (995 ◦C for Ti-6Al-4V [8]). Within these columnar β grains, EBM built Ti-6Al-4V typically has a basketweave microstructure. Wrought Ti-6Al-4V microstructure consists of primary α combined with Widmanstätten microstructure [8,13], called bimodal or duplex microstructure. Cast microstructures normally consist of coarse prior β grains with large α colonies, where the α laths within the colonies are oriented in the same crystal orientation.

Diffusion is an important part of the HE mechanisms [4] where microstructure plays an important role. Tal-Gutelmacher et al. [14] investigated the effect of hydrogen solubility for Widmanstätten and bimodal microstructures. The conclusion was that fully lamellar Widmanstätten microstructure had several orders of magnitude higher hydrogen solubility than the bimodal microstructure, which was explained by the continuous β phase in the fully lamellar structure.

Texture can affect the ingress of hydrogen and the hydride formation [15,16] but the texture of EBM built Ti-6Al-4V has been shown to be weak [17]. Residual stresses, which are known to be present in various AM processes [18–21], can furthermore affect the fatigue crack growth (FCG) rate. However, Maimaitiyili et al. [22] did show that there are no or small residual stresses in as-built EBM Ti-6Al-4V.

Rozumek et al. [23] showed that the FCG rate in Ti-6Al-4V is highly depending on post processing. By performing a hardening and ageing heat treatment five times higher fatigue life was obtained.

Relevant for space applications are the cryogenic properties and temperature has been shown to have a strong effect on FCG properties. Increased temperature renders an increased diffusion rate of hydrogen [14], and two temperatures were investigated by Pittinato [24]; −129 ◦C and −73 ◦C. At −129 ◦C there was no difference in the FCG rate in helium and hydrogen atmospheres, whereas at −73 ◦C there was an increase in FCG rate in hydrogen environment.

FCG experiments have previously been performed on a wide range of Ti-6Al-4V material [24–27], but few of these studies concerned hydrogen environment and additive manufacturing. In previous studies, the FCG properties of conventionally manufactured Ti-6Al-4V in high-pressure hydrogen was performed by Gaddam et al. [26,28] showing that microstructure has an effect on the FCG properties in hydrogen. Cast Ti-6Al-4V with coarser microstructure was shown to have inferior FCG properties compared with forged material with fine microstructure.

In this work, the effect of hydrogen embrittlement on FCG properties of EBM built Ti-6Al-4V has been investigated. Samples have been exposed to either air or high-pressure (150 bar) hydrogen atmosphere at room temperature. The EBM built material has been compared to previous results [26,28] for cast and forged materials, and the differences in FCG properties have been linked to the different types of microstructures. To further investigate the hydrogen embrittlement, fractography was performed along with crack profile characterization.

#### **2. Experimental Method**

#### *2.1. Material*

EBM Ti-6Al-4V samples were manufactured with an Arcam Q20+ machine, using a layer thickness of 90 μm. Cylinders were manufactured having a length of 135 mm and a diameter of 25 mm. The layers were oriented perpendicular to the major axis of the cylinders i.e., the applied load during the FCG tests was perpendicular to the AM built layers. The powder used was Ti-6Al-4V B110 (Virgin Hoeganaes) in accordance with the aerospace material specification AMS 4992. Prior to testing, the samples received a hot isostatic pressure (HIP) treatment at 920 ± 10 ◦C for two hours with a pressure of 1020 bar, followed by a heat treatment at 704 ± 10 ◦C for two hours. Both HIP and heat treatment were conducted in argon. Out of the manufactured cylinders, Kb bars were machined (see Figure 1) using low stress grinding and polishing of the gauge section. A tensile test was performed on the post treated material and the result showed a yield strength of 890 MPa and tensile strength 990 MPa.

**Figure 1.** Sketch of Kb bar specimen. Dimensions in mm (printed with permission from Metcut Research).

#### *2.2. FCG Experiments*

The FCG test in air was conducted at Metcut Research Inc., Cincinnati, OH, USA, while the hydrogen testing was conducted at The Welding Institute (TWI) in Cambridge, UK. One Kb bar was tested in a hydrogen-rich atmosphere with a pressure of 150 bars and two Kb bars were tested in air at ambient temperature. The samples were designated H-A, Air-A, and Air-B, respectively. All tests were stress-controlled and performed at room temperature with the max loads; 645 MPa for H-A, 534 MPa for Air-A, and 528 MPa for Air-B. The FCG testing fulfills the requirements for plane strain conditions. The pre-cracking was performed using a frequency of 10 Hz. Tint temperatures between 450 ◦C to 350 ◦C were used to mark the crack propagation. The fatigue tests were performed with a uniaxial load perpendicular to the AM built layers with R = σmin/σmax = 0 and a test frequency of 0.5 Hz, using a triangular waveshape. The crack propagation was measured using potential drop. The pre-crack and final crack sizes were measured using heat tinting and these sizes were correlated to the potential drop signal. The translation from the potential drop signals to crack sizes were made using a calibration curve [29]. Corrections were made so that the measured pre-crack and final crack sizes and corresponding potential drop values were consistent with the calibration curve. The FCG rate was then computed per data point using the secant method (ASTM E647-15e1). Once the FCG crack reached a certain length the remaining material was heat tinted to reveal the final crack length and fractured in tension using a monotonically increasing load. The equation used for the experiment was:

$$K = S \sqrt{\pi \frac{a}{Q} F\_s \left(\frac{a}{c}, \frac{a}{t}, \frac{a}{b}, \Theta\right)}\tag{1}$$

where *S* is the tensile strength, *a* the crack depth, *Q* elliptical crack shape factor, *Fs* boundary correction factor, *t* thickness of the sample, *b* half-width of the sample, and *c* half-width of the crack. See [30] for full solution.

#### *2.3. Fractography and Microstructural Characterization*

For overview images of the fracture surfaces a stereomicroscope (Nikon SMZ1270) was utilized, whilst for fractography a scanning electron microscope (SEM, Jeol IT300LV) was used. Crack profiles were made on the hydrogen and air-tested samples, first by cutting cross-sections parallel to the *x*-*x'* plane according to Figure 1. Then, by grinding and polishing carefully, the desired positions were reached in the plane perpendicular to the *x*-*x'* plane (edge of notch). The grinding was monitored using a stereomicroscope. The crack profile was characterized with light optical microscope (LOM, Nikon eclipse MA200) and SEM. For microstructural characterization, a representative cross-section was ground and polished according to the conventional sample preparation techniques for titanium, then etched using Kroll's etchant (see ASTM Standard E 407). To investigate the microstructure at low magnification a LOM was used. The software Image J version 1.52a [31] was used to measure the width of the α laths and prior β grains; 100 measurements were performed for each microstructural feature to obtain the average size.

#### **3. Results**

#### *3.1. Microstructure*

The EBM Ti-6Al-4V material consisted of columnar prior β grains. The grain boundaries are illustrated as black dotted lines in Figure 2a and they were elongated parallel to the build direction, with lengths up to approximately 2 mm. In the plane perpendicular to the build direction, the prior β grains were equiaxed, with widths of ~100 μm. The prior β grains were partially separated by discontinuous grain boundary α (GB-α) and the microstructure within the prior β grains was basketweave, with an average α lath width of 2.2 ± 0.6 μm, see Figure 2b.

**Figure 2.** In (**a**) columnar prior β grains and in (**b**) the basketweave microstructure are shown. The black arrow in the bottom left corner in (**a**) points towards the build direction (BD).

#### *3.2. Fatigue Crack Growth*

Different plots of FCG test data are shown in Figures 3–5; crack length versus number of cycles, crack growth rate versus ΔK, and ratio between crack growth rate in hydrogen (specimen H-A) and air (specimen Air-B). The two air-tested samples followed the Paris law with similar inclination (see Figure 4). The hydrogen-tested sample, on the other hand, has a fluctuating crack growth rate in the first stage of the test. After the fluctuating stage, at <sup>Δ</sup>K ~23 *MPa* <sup>√</sup> *m*, a sudden increase in crack growth rate was observed. The resulting increase in crack length is shown in Figure 3: After ~5600 cycles there is a sudden offset in crack growth rate for H-A, whereas for Air-A/B no such abrupt offset is present. Figure <sup>5</sup> shows that below ~23 *MPa* <sup>√</sup> *m* the relative crack propagation rate is around one, indicating no difference in air and hydrogen tests. Around this value of <sup>Δ</sup>K (23 *MPa* <sup>√</sup> *m*), the crack propagation rate ratio started to increase, and continues to do so at a constant rate up to ~32 *MPa* <sup>√</sup> *m*, where it reaches a plateau. By then, the crack propagation rate of the material tested in hydrogen is over 12 times higher than in air.

**Figure 3.** Crack length versus the number of cycles for all three tested electron beam melting (EBM) samples. At ~5600 cycles the crack length of the hydrogen-tested sample accelerates.

**Figure 4.** Crack growth rate versus <sup>Δ</sup>K. The hydrogen-tested material fluctuates below 23 *MPa* <sup>√</sup> *m*, while accelerating above. The air-tested material follows Paris law.

**Figure 5.** Relative crack propagation rate between hydrogen and air-tested material versus ΔK. At 23 *MPa* <sup>√</sup> *<sup>m</sup>* the relative crack propagation increases steadily while reaching a plateau at 31 *MPa* <sup>√</sup> *m*.

#### *3.3. Fractography*

Figure 6 shows the fracture surface of the hydrogen-tested sample (H-A). A change in macroscopic appearance was observed at a crack length ~1 mm (this crack length is shown in Figure 3 to coincide with the accelerated crack growth), which corresponded to <sup>Δ</sup>K 23 *MPa* <sup>√</sup> *m* and the location of this ΔK is shown by a white dashed semi-ellipse in Figure 6 (in-between X and Y line). The lines X, Y, and Z shown in Figure 6, following semi-elliptic paths, correspond to fractographic locations where the fracture surface has been characterized extra carefully. The X line surrounds the pre-crack, Y is in the middle of the crack, and Z is at the end of the crack. Table 1 presents lengths and widths of notch, pre-crack, and the fatigue crack of all samples. The lengths a and 2c of the final fatigue crack are illustrated in Figure 6.

**Figure 6.** Overview of H-A fracture surface. The two tint temperatures, 450 and 350 ◦C, resulted in the two different golden color contrasts shown in the figure. "2c" represents the width and "a" the length of the crack. The three semi-elliptical dashed lines X, Y, and Z show profiles where fractography was performed in higher detail. The white dotted line in-between the X and Y lines corresponds to the ΔK <sup>23</sup> *MPa* <sup>√</sup> *m*. The vertical white line indicates the cross-section where crack profiles were investigated.


**Table 1.** Notch, pre-crack, and final fatigue crack lengths (a), widths (2c), and their ratio (a/c), for the three investigated samples Air-A/B and H-A.

In Figure 7a,b representative areas along Y and Z in Figure 6 of the fracture surface of the hydrogen-tested sample are shown, in Figure 7c the position of the crack profile in the ground specimen (white vertical line in Figure 6) is shown. A transition in fracture surface is shown in Figure 7a, below the white dashed semi-ellipse corresponding to crack length where the <sup>Δ</sup>K was 23 *MPa* <sup>√</sup> *m*. At this ΔK the fracture surface starts to transition from flat to rough. In Figure 7b (region Z) large cracks, exceeding lengths of 100 μm, are observed on the fracture surface.

**Figure 7.** In (**a**) the transition zone from flat to rough fracture surface (the white dotted semi-ellipse shows <sup>Δ</sup>K 23 *MPa* <sup>√</sup> *m*), whereas (**b**) shows a rough fracture surface with large cracks. In (**c**) a section of H-A fracture surface is shown, cut as indicated by the vertical white line in Figure 6.

Figure 8 shows the characteristic fracture surface of the air-tested material, in Figure 8a higher magnification, whereas in Figure 8b lower magnification. The transition area observed in Figure 7, from flat to rough, did not exist in the air-tested material, neither do the large cracks.

**Figure 8.** (**a**) 100× magnification fracture surface image of air-tested material (sample Air-B) in the region ~Y. (**b**) lower magnification image that shows an overview of the whole fracture surface. The black arrow points towards the crack propagation direction.

In Figure 9 representative images of the fracture surface are shown along the profiles X to Z (see Figure 6 for illustration of the locations on the fracture surface). In the air-tested samples, striations were observed along the whole fatigue crack, becoming increasingly larger the greater the ΔK became. In Air-B, section X, striations were only observed at higher magnification, whereas in section Z they were clearly visible for the present magnification (2000×).

**Figure 9.** Images of the three sections X-Z (see Figure 6 for locations of the areas on the fracture surface) for one air (sample Air-B) and one hydrogen (sample H-A) tested sample. The images are in the same magnification and the white arrow in the bottom right indicates the crack propagation direction. In section Z of Air-B, examples of striations are indicated with the white arrows. In sections Y and Z of H-A, white arrows point at cracks; their dimensions increase from the former section to the latter.

In the hydrogen-tested sample, the fracture surface on the first stage of the fatigue crack (section X) appeared flat. With increased <sup>Δ</sup>K (above 23 *MPa* <sup>√</sup> *m* i.e., section Y) an increase in fracture surface tortuosity aroused, along with the appearance of small cracks. Then, towards the end of the fatigue crack (section Z), larger cracks were observed with dimensions of ~100 μm. At the flat area of the H-A's fracture surface, i.e., below 23 *MPa* <sup>√</sup> *m*, features that resembled crack arrest marks (CAM) were observed. Crack propagation stops at the CAM interface, causing the formation of these marks. The origin could for example be the cleavge of a hydride. In Figure 10 features that resemble CAMs are marked by white arrows, where each plateau is the end of a CAM, which according to literature could indicate the interface between the hydride and titanium base metal [4,32]. Note that the CAM locations were not numerous.

**Figure 10.** An area with features that resembled crack arrest marks (CAM). Each white arrow indicates a possible CAM, where each plateau could be a hydride titanium interface between cleaved hydrides. The black arrows show the crack growth direction.

Figure 11 is a high magnification image of a fracture surface cross-section in the hydrogen tested material. Secondary cracks were present across α/β interfaces. The crack path can be seen on the upper right corner, showing that the crack propagated along an α lath, seemingly in the α/β interface.

**Figure 11.** Image of a crack profile for sample H-A. It shows secondary cracks that grow across α/β interfaces. In the upper right corner, the crack propagated along an α lath, seemingly in the α/β interface.

In Figure 12 two crack profiles are shown (one for Air-B and one for H-A). The crack profiles correspond to the cross-section illustrated as a white vertical line in Figure 6. Crack profiles covering the whole fracture surface from notch to final tensile fracture were obtained in LOM, while the remaining images, Figure 12a–e, were obtained in an SEM using backscattered electrons.

The crack profile showed a similar pattern as was observed on the fracture surfaces. For H-A, in the region from the start of crack to <sup>Δ</sup>K 23 *MPa* <sup>√</sup> *m*, few cracks are observed i.e., image (c), from 23 *MPa* <sup>√</sup> *m* and onwards large vertical (perpendicular to the crack direction) cracks were observed. With higher magnification also numerous small secondary cracks were observed after the <sup>Δ</sup>K 23 *MPa* <sup>√</sup> *m*, see Figure 12d. Note that these smaller secondary cracks had no connection to the main crack. For the H-A sample, three areas are shown in higher magnification; in-between pre-crack and 23 *Pa* <sup>√</sup> *m* i.e., (c), after 23 *MPa* <sup>√</sup> *m* i.e., (d), and before final crack i.e., (e). In (c) the crack profile was less tortuous and no vertical cracks or secondary cracks were observed. However, directly after 23 *MPa* <sup>√</sup> *m*, secondary cracks were observed as shown in Figure 12d. After area (d), deep vertical cracks appeared as shown in (e). The maximum major axis for the secondary cracks was in the dimensions of 10 to 20 μm, whereas the vertical cracks reached about 50 to 70 μm.

The hydrogen-tested sample had a tortuous crack profile, whereas the air-tested sample had a comparably straight crack path, with a homogeneous appearance throughout the various ΔK. The cracks propagated both parallel to the direction of α laths in the α/β interface and perpendicular to the laths, for both the air and hydrogen-tested material. Due to the build-orientation of the samples in regard to the load, the propagation of the fatigue crack was perpendicular to the major axis of the prior β grains i.e., perpendicular to the heat gradient, thus no propagation along grain boundary α could be observed.

For the air-tested material, no deep cracks were observed along the crack profile and features that appeared to be cracks on the fracture surface were shown to be superficial. In Figure 12b, the largest identified crack in Air-B sample is shown. The vertical cracks shown in (e) both propagated parallel to an α lath in the α/β interface, then switched to propagation across several α laths. The location of the notch, pre-crack and final crack were approximately the same for both H-A and Air-B, as was shown in Table 1.

**Figure 12.** Crack-profiles of Air-B and H-A samples. The crack profiles correspond to a cross-section illustrated in Figure 6 as a white vertical line. In H-A, many larger cracks that grew perpendicular to the fracture surface were observed, along with numerous smaller secondary cracks that grew parallel. In H-A the crack path was tortuous, whereas in Air-B it was comparably smooth. Compared to H-A no cracks were present in Air-B. Images (**a**–**e**) show magnified areas.

#### **4. Discussion**

#### *4.1. Comparison between Hydrogen and Air Atmospheres*

As discussed and observed by Lynch [4,33] and others [26,34–39], the mechanical properties of titanium are expected to deteriorate when exposed to hydrogen, due to HE [3]. Such deterioration of the mechanical properties was observed in this work. When exposing the EBM built Ti-6Al-4V to hydrogen-rich atmosphere, hydrogen can be absorbed and diffused throughout the material to the crack tip. At the crack tip the hydrogen can interact with the titanium either by the nucleation of hydrides, resulting in HE through brittle fracture of titanium hydrides, or through one or several other HE mechanisms [4,34–36,40]. On the first stages of the test the FCG rate fluctuated (as shown in Figure 4 and this is believed to be due to temporarily arrests of the crack i.e., CAMs. The nucleation and cleavage of hydrides are suggested in literature to be repeated, forming a relatively flat region on the fracture surface with the presence of features such as CAMs (see Figure 10) [26,32].

At <sup>Δ</sup>K 23 *MPa* <sup>√</sup> *m*, the FCG rate abruptly increased, which was likely due to one or more of the non-hydride forming mechanisms. These mechanisms are called hydrogen enhanced local plasticity (HELP) [34], adsorption induced dislocation emission (AIDE) [4], and hydrogen enhanced de-cohesion (HEDE) [35,36], thoroughly explained in a review paper by Lynch [4]. From the fractography and the crack profiles, it could be shown that cracks appeared once the <sup>Δ</sup>K increased above 23 *MPa* <sup>√</sup> *m*. The crack profile showed that two types of cracks were present, vertical and secondary cracks. The vertical cracks were both deeper and wider than the secondary cracks. These two types of cracks were not present in the air-tested samples, thus it is evident that the hydrogen influenced the formation and subsequent propagation of these different types of cracks. The presence of secondary cracks across the interface between α laths and residual β phase is most probably related to the faster diffusion rate of hydrogen in the β phase and its larger hydrogen storage capacity. This would be coupled with the formation of hydrides since the α phase is a strong hydride forming phase [39]. Underneath the crack path, stress fields were present as well, causing further cracking of these hydrides. Thus, the presence of hydrogen influenced the crack tortuosity.

The relative crack growth rate between the hydrogen and air-tested samples, as shown in Figure 5, clearly shows the critical effect of the presence of hydrogen on titanium: At high ΔK, the FCG rates were roughly 10 times larger in a hydrogen atmosphere. As previously discussed, there was an acceleration in the FCG rate at 23 *MPa* <sup>√</sup> *m*. However, Figure 5 also shows that this acceleration did not continue until final failure. At <sup>Δ</sup><sup>K</sup> <sup>&</sup>gt; <sup>31</sup> *MPa* <sup>√</sup> *m*, the relationship between the hydrogen and air-tested samples seemed to have reached a plateau, indicating that increased ΔK in this range and increased hydrogen content ahead of the crack tip did not promote additional acceleration of the FCG rate for EBM built Ti-6Al-4V. Apart from the presence of secondary cracks, the fracture surface cross-section of the hydrogen and air-tested material differed regarding crack path. The crack path in hydrogen-tested material was tortuous in nature, whereas it was relatively flat in air-tested material. Crack paths of both specimen types crossed through α laths, as well as propagated parallel to α laths in the α/β interface.

#### *4.2. Comparison with Cast and Wrought Ti-6Al-4V*

In Figures 13 and 14, the FCG properties of the investigated Ti-6Al-4V EBM samples are compared to those of forged and cast material of the same alloy (data source: Gaddam et al. [26,28]), both in air and hydrogen atmospheres. The data in these references have been obtained through testing with similar ΔK.

**Figure 13.** Fatigue crack growth (FCG) data for EBM, forged [26], and cast [28] Ti-6Al-4V. The FCG properties of material exposed to air and hydrogen atmosphere are shown, where the crack growth rate da/dN is plotted on the y-axis and the ΔK on the x-axis.

**Figure 14.** FCG data for EBM, forged [26], and cast [28] Ti-6Al-4V. The y-axis shows the crack propagation rate in hydrogen divided by the crack propagation rate for air, the x-axis shows ΔK. The dashed box is a magnified image to more clearly show the critical points in crack growth rate for the EBM and forged material.

In Figure <sup>13</sup> the crack growth rate (da/dN) is plotted against <sup>Δ</sup>K (*MPa* <sup>√</sup> *m*). In Figure 14 the crack propagation rate in hydrogen environment relative to that in air is plotted as a function of ΔK. A difference in FCG rate behavior was observed between the material exposed to the high-pressure hydrogen and air: For the air-tested material, the FCG rate followed Paris law. For the hydrogen-tested material the FCG rate fluctuated at first, then at certain ΔK, depending on the microstructure, the FCG rate increased fast. The dotted square in Figure 14 is a magnified area to make it easier to distinguish

the difference between the EBM and forged materials. In cast Ti-6Al-4V, this occurrence appears at a ΔK of ~17 *MPa* <sup>√</sup> *<sup>m</sup>*; in EBM-built Ti-6Al-4V, at ~23 *MPa* <sup>√</sup> *<sup>m</sup>*; in forgings of the same alloy, at ~26 *MPa* <sup>√</sup> *m*. In Figure 14 it is furthermore shown that the maximum FCG rate difference between the hydrogen and air atmosphere was 160 times for the cast material, whereas for the EBM material it was 10 times and the forged material 5 times, all values at <sup>Δ</sup><sup>K</sup> <sup>=</sup> <sup>29</sup> *MPa* <sup>√</sup> *m*.

The micrographs in the two cited papers by Gaddam et al. [26,28] were utilized as a basis for microstructural comparison. It has been observed that the microstructure of forged Ti-6Al-4V consisted of islands of primary α grains with basketweave microstructure surrounding them, i.e., bimodal microstructure. The microstructure of cast Ti-6Al-4V [28] consisted of coarse prior β grains with large α colonies with α laths with the same crystal orientation, i.e., α laths that grow parallel to each other and not perpendicular such as in the basketweave microstructure. The cast material also had prior β grains that appeared equiaxed, being surrounded with a continuous layer of grain boundary α. By comparing these results with the microstructural features of the EBM material, it can be concluded that the cast material had the coarsest microstructure of the investigated materials.

The diffusion rate has a key role in the HE mechanisms [4]. The main parameters that determine the diffusion of hydrogen in a given material are hydrogen concentration gradients, temperature, presence of hydrostatic stresses, and microstructure [14,41]. The phase distribution and grain size of the microstructure are well linked to the diffusion properties in the material [41]. As discussed by Yazdipour et al. [41] and Ichimura et al. [42] a two-fold effect exists where the number of grain boundaries affects the diffusion rate. Finer grain structure implies an increased amount of grain boundaries, which are the fastest diffusion paths and thereby enhance hydrogen diffusion rate. On the other hand, increased amounts of grains and grain boundaries render an increased grain boundary triple junction density, sites that act as hydrogen traps and decrease hydrogen diffusion rate. These effects compete, increasing or decreasing the hydrogen diffusion.

Differences in hydrogen diffusion in the materials analyzed in this work might also be due to proportions of β and α phases. Hydrogen is much more soluble in body centered cubic (BCC) β than in hexagonal close packed (HCP) α [14], due to its preferential absorption in tetrahedral sites [43], which are more abundant in BCC crystal structures than in HCP. As a result, relatively low hydrogen content generates hydrides in α titanium.

The FCG behavior of cast Ti-6Al-4V illustrated in Figure 13 can be explained as follows. Cast Ti-6Al-4V has been shown to have a higher β phase fraction than the same alloy produced by EBM [44]. In addition, the casting's coarse microstructure means a lower number of hydrogen traps. Both these reasons favor faster hydrogen diffusion. Then, due to faster diffusion rate sufficient hydrogen is diffused ahead of the crack tip, for the material to experience at least one of the HE mechanisms at 17 *MPa* <sup>√</sup> *m*, which accelerates the FCG rate even further, reaching a relative FCG rate that is ~160 times higher in hydrogen than in air. Then the same phenomena happen for the EBM and forged material at 23 *MPa* <sup>√</sup> *<sup>m</sup>* and 26 *MPa* <sup>√</sup> *m*, respectively. The results are in well accordance to Tal-Gutelmacher et al. [14], that also found the bimodal microstructure to be less sensible to HE than the Widmanstätten microstructure with its more continuous network of residual β phase.

#### **5. Conclusions**

By performing FCG experiments of EBM built Ti-6Al-4V in hydrogen and air atmosphere and then comparing the results with already published data of cast and forged Ti-6Al-4V the following conclusions can be made:


predominantly parallel to the main crack direction and large cracks that grew perpendicular to the main crack direction, being connected to the main crack.


**Author Contributions:** M.N. wrote the paper and performed all the characterization work. P.Å., M.-L.A., R.P. and M.C. gave feedback on the work, helped with organization and proofread the manuscript. T.H. and C.d.A.S. were the contact persons at GKN Aerospace and helped with the fatigue crack growth experiment, they also helped with proofreading the manuscript. All authors have read and agreed to the published version of the manuscript.

**Funding:** The National Aviation Research Program (NRFP), the EU funded Space for innovation and growth (RIT) and the Graduate School of Space Technology at Luleå University of Technology have contributed with financial support in this project.

**Acknowledgments:** The author would like to acknowledge the dedicated work of Géraldine Puyoo, Staffan Brodin and Clas Andersson at GKN-Aerospace, for both the sample manufacturing that was conducted at GKN Filton but also for organizing the testing.

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

#### **References**


© 2020 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* **Residual Lattice Strain and Phase Distribution in Ti-6Al-4V Produced by Electron Beam Melting**

**Tuerdi Maimaitiyili 1,2,\*, Robin Woracek 3,4, Magnus Neikter 5, Mirko Boin 6, Robert C. Wimpory 6, Robert Pederson 7, Markus Strobl 3,4,8, Michael Drakopoulos 9, Norbert Schäfer <sup>10</sup> and Christina Bjerkén <sup>2</sup>**


Received: 29 January 2019; Accepted: 20 February 2019; Published: 23 February 2019

**Abstract:** Residual stress/strain and microstructure used in additively manufactured material are strongly dependent on process parameter combination. With the aim to better understand and correlate process parameters used in electron beam melting (EBM) of Ti-6Al-4V with resulting phase distributions and residual stress/strains, extensive experimental work has been performed. A large number of polycrystalline Ti-6Al-4V specimens were produced with different optimized EBM process parameter combinations. These specimens were post-sequentially studied by using high-energy X-ray and neutron diffraction. In addition, visible light microscopy, scanning electron microscopy (SEM) and electron backscattered diffraction (EBSD) studies were performed and linked to the other findings. Results show that the influence of scan speed and offset focus on resulting residual strain in a fully dense sample was not significant. In contrast to some previous literature, a uniform αand β-Ti phase distribution was found in all investigated specimens. Furthermore, no strong strain variations along the build direction with respect to the deposition were found. The magnitude of strain in α and β phase show some variations both in the build plane and along the build direction, which seemed to correlate with the size of the primary β grains. However, no relation was found between measured residual strains in α and β phase. Large primary β grains and texture appear to have a strong effect on X-ray based stress results with relatively small beam size, therefore it is suggested to use a large beam for representative bulk measurements and also to consider the prior β grain size in experimental planning, as well as for mathematical modelling.

**Keywords:** residual stress/strain; electron beam melting; diffraction; Ti-6Al-4V; electron backscattered diffraction; X-ray diffraction

#### **1. Introduction**

Titanium-based alloys have been widely used as engineering materials in many industries because of their excellent combination of a high strength/weight ratio and good corrosion resistance [1]. Application of the titanium and its alloys can be found in many different industries ranging from the aerospace to consumer goods, e.g., heat exchangers, automotive, offshore petroleum, gas exploration and medical implants [1]. Among various titanium alloys, Ti-6Al-4V is by far the most commonly used and accounting for some 50% of the total titanium output in the world [1]. However, extracting high purity titanium and the production of usable raw Ti-6Al-4V is a difficult and expensive process.

Many additive manufacturing (AM) techniques have been developed to fabricate geometrically complex and fully dense metal parts for a variety of applications. The most common AM methods use a laser or electron beams as the power source. In comparison with conventional methods, the electron beam melting (EBM) method has the advantage of increased component complexity with limited manufacturing expertise, shorter lead times, reduced material waste and minimum or almost zero tooling cost [2–4]. In the EBM, electromagnetic lenses are used for focusing and guiding of high-energy electron beams to selectively melting the ingredient powder to form a three-dimensional solid object point by point and layer-by-layer bases. The process is a fully computer controlled automatic system. Electron beam melted Ti-6Al-4V and its microstructure have previously been described by [3–6]. Despite the young age of the EBM process, it has already been demonstrated that it can create defect-free complex products with good mechanical properties, such as functionally graded cellular structures [7–10]. However, the EBM process is complex, and the results depend upon the variables of the system, such as beam power (beam current), beam size, scan speed, and scan direction/scan strategy. These are collectively referred to as process parameters. Thus, each set of process parameter settings produces a different built environment and cooling conditions, and as a consequence, different microstructures are observed [2,11] as well as possible different states of residual stress/strain (RS) and texture [12–15].

RS is internal, self-balanced stress that exists in alloy systems without any external applied forces and may appear during mechanical, thermal or thermochemical processing [16–18]. Depending on the compressive or tensile nature and magnitude of the RS, it significantly affects the mechanical properties of the materials. Lack of understanding RS distribution and the effect of different processing on the stress may create serious consequences [19]. Therefore, it is crucial to know the nature and magnitude of RS in the material for safe and economical operation. AM processes, such as EBM have already proven to be a potential manufacturing process in various fields with many possibilities. However, during AM processes, because of the repetitive selective point by point melting and solidification process, strong heat gradients are generated between different parts of the built material, which potentially can lead to unfavourable residual stresses. Despite several decades of development of AM processes, metal AM is still in its infancy and the relation between different AM process combinations with RS, microstructure, texture and material properties still lack significant understanding. Quality and performance of additively manufactured material can only be reliably controlled and optimized by understanding the effects of different process combinations on the RS and microstructure of the built material. There is limited research concerning RS on titanium alloys produced with conventional methods [17], wire/arc based additive manufacturing [20,21] and selective laser melting (SLM) [22–24]. However, there are only a few concerning RS in EBM Ti-6Al-4V [25–27]. Therefore, this study aims to explore the effect of a subset of EBM process parameters on residual strain/stress in Ti-6Al-4V. RS can be investigated by e.g., hole drilling techniques, ultrasonic and magnetic methods and diffraction-based methods [16,18,28]. Each RS measurement technique has its own advantages and disadvantages. Among them, diffraction (e.g., X-ray and neutron diffraction) is one of the most accurate and well-developed methods of quantifying local and global residual stresses in the material [16,18]. Compared with conventional RS measurement techniques, the diffraction method offers many advantages. The residual stresses in the material are calculated from the strain measured in the crystal lattice. Generally, RS diffraction measurements are not significantly influenced by material properties such as hardness, the degree of

cold work, or preferred orientation. In fact, all can be quantified by diffraction techniques. In addition, they are non-destructive, have a high spatial resolution (ranging in the order if μm to mm depending on the radiation source)), almost no specimen geometry restrictions, and can be applied to all types of crystalline material. Furthermore, with the high-energy synchrotron X-ray diffraction setup used in this work, it is possible to characterize and conduct quantitative studies on multi-component systems containing phase quantities as low as 0.1–0.7% [29–31] with a high temporal resolution (0.2 s) [31]. This means that phase specific RS in the target system can be determined dynamically with high accuracy. Last but not the least, diffraction-based RS techniques are already well established—there is good level of expertise, and standards have been developed [16,18,32].

To investigate the effect of EBM process parameters on residual strain and stress development, various specimens produced with different beam sizes, scan speeds and build thicknesses were investigated. In addition, original powder specimens that were used to produce these solid specimens have also been tested. Diffraction measurements were carried out using the high energy (50–150 KeV) beamline I12-JEEP (joint engineering, environment and processing) at the Diamond Light Source (UK) with an energy dispersive X-ray diffraction (EDXRD) setup [33], as well as the dedicated RS neutron diffractometer (ND) E3 [34] at Helmholtz Zentrum Berlin (Germany). After data collection, the Pawley pattern fitting method was performed for the EDXRD data and single peak fitting for the ND data using the structure analysis software package TOPAS-Academic [35].

#### **2. Experimental Setup and Data Analysis**

#### *2.1. EBM Process and Material*

The EBM process is described in greater detail e.g., in [36]. However, the EBM process is quite complicated owing to numerous parameters (e.g., feed material, beam current, feed rate, build layer thickness and scan speed), which potentially change the build environment and cause a different micro-structure in the final product that may not be wanted. The solidification rate, surface smoothness, and microstructural homogeneity of EBM processed parts are strongly influenced by the process parameters.

The electron beam parameters, such as beam speed, beam current and scan length can be varied in a controlled sequence throughout the build according to algorithms developed by the manufacturer. According to the literature, the resolution of the build in EBM is influenced by layer thickness [37,38] powder size [39] and spot size [2]. It has been reported that the size of the powder and layer thickness has a direct influence on build part quality [38,39]. Commonly, a smaller powder gives finer surface finish and it has been demonstrated that the EBM can process powder with size 25–45 μm [39]. The benefit with smaller particles is finer surface finish and more compact layer i.e., no/little shrinkage. Nonetheless, due to the charging from the electron beam, smaller powder particles are not possible for EBM i.e., a phenomenon called smoking can otherwise occur. A benefit with larger particle size and layer thickness, however, is an increased build rate. To achieve the highest quality, the current standard EBM layer thickness has been reduced from 100 to 50–70 μm and powder with size 45–100 μm was used [40]. The powder used in this study was gas atomized and size within the range of 45–100 μm. The chemical composition is Ti-6Al-4V-0.1Fe-0.15O-0.01N-0.003H (in wt %).

The mechanical properties of the EBM built Ti-6Al-4V are affected by scan speed [41]. It has been reported that the rapid heating and cooling, and repetitive thermal cycles of AM processes in laser-based additively manufactured material create RS in as-built material [42,43]. According to Klingbeil et al. [43], preheating of the build platform and powder layer can reduce RS and negative effects, such as part warping. In the EBM process, for Ti-6Al-4V the whole build chamber is heated to approximately 700 ◦C prior to building, and each layer of powder is preheated before melting. The minimum temperature of the build chamber is maintained ≥ 600 ◦C throughout the build process [40]. Therefore, the EBM built component is naturally annealed immediately after building,

and thus remaining residual stresses may be limited in magnitude. However, the relation between scan speed, offset focus and RS has not yet been verified experimentally to the authors' knowledge.

Ti-6Al-4V blocks of 50 mm × 50 mm × 5 mm were built and measured at the centre of each specimen both along the build direction (vertical, "V") and in the build plane (horizontal, "H") by neutron diffraction. To investigate the finger prints of the phase specific residual stresses, specimens were cut from the as-build blocks in varied thicknesses (see Table 1) and investigated with synchrotron XRD in two directions (Figure 1). For the analysis of the RS state and microstructure, three different standard controlling factors were chosen: build thickness, scan speed and offset focus. The details of specimen production parameters are tabulated in Table 1. Three specimens were made for each set of processing parameters. Because of limited beam time, only one specimen was measured during neutron diffraction experiment.


**Table 1.** Specimen process parameters.

\* is the as-built specimen thickness that was used in the neutron diffraction studies.

**Figure 1.** Beam line setup and illustration of measurements. "V" is the vertical direction and "H" is horizontal direction. Each blue and green rectangle represents the measurement locations. Embedded image is a photo of one of the as-built specimens.

#### *2.2. Synchrotron X-ray Measurements (EDXRD)*

Synchrotron X-ray diffraction is a well-known and powerful tool for structural analysis, quality characterization [18,44] and conducting quantitative studies on multi-component systems [30,31]. Herein reported experiments were performed at the high energy (50–150 KeV) beamline I12-JEEP at Diamond Light Source in the UK with EDXRD setup [33]. All specimens were irradiated with a white/continuous X-ray beam with a photon flux range from 1.8 × <sup>10</sup><sup>11</sup> to 9.4 × 1010 photons s−<sup>1</sup> depending on the energy of the X-ray beam. The EDXRD data were recorded with a "horseshoe" detector, consisting of 23 liquid-nitrogen-cooled germanium (Ge) energy-sensitive detector elements, positioned 2 m behind the specimen position with a take-off angle of 5◦ from the incident beam. The energy resolution of Ge-detector ranges from 7 × <sup>10</sup>−<sup>3</sup> at 50 keV to 4 × <sup>10</sup>−<sup>3</sup> at 150 keV. As the 23 detector elements are equally spaced in steps of 8.2◦, they have full azimuthal coverage from 0◦ to 180◦. During the measurements, each detector element independently records diffraction patterns under a different Bragg angle. As the sequence of energy values of each detector element is approximately the same for all, the 23 diffraction patterns can be summed to form one pattern after normalization of each data set for detailed phase analysis. The actual EDXRD setup with schematics of measurements and specimen can be seen in Figure 1.

In the experiment, a 2-mm thick copper filter and slit sizes of 0.1 × 0.1 mm<sup>2</sup> and acquisition time of 100 s were used. To map strain variations in the material along build direction and in the build plane as complete as possible in the limited synchrotron beam time, measurements were made at an average of 55 points in the centre of the specimen along the vertical, and 10 in the horizontal direction. The volume of the material contributing to the diffraction pattern corresponds to the intersection of the incident and diffracted beams, typically defined by slits and collimators, respectively. To ensure the gauge volume is fully inside the specimen, the centre of the measurements points close to the edges was positioned about 300 μm from the edge during horizontal and vertical line scans in the centre of the specimen. The clearance of corner scan, however, is about 0.5 mm from the corner edge. Each detector element in the setup measures the lattice spacing in a specific direction, which means that strains in 23 different directions in 0–180◦ range can be measured in a single measurement without specimen rotation. The measured strain direction can be defined by the X-ray scattering vector, which bisects the incident and diffracted beams. With this EDXRD setup, the three-dimensional strain/stress field can be calculated by including additional information or by adopting assumptions, such as the plane strain criterion. For that reason, only strains along the longitudinal and transversal directions are presented in this work. Detector elements 1 and 23 were used to measure transverse strain, and elements #11 or #12 were used to measure longitudinal strain. Compared with the angle dispersive diffraction setup used in the lab X-ray source, the JEEP (I12) at Diamond has many orders of magnitude higher X-ray flux with high energy. At JEEP it is possible to probe metal specimens in millimeter to centimeter thickness in relatively short times. This gives the possibility to determine the RS state of the various phases in one experiment. This was important for the present study as thicker specimens allowed to obtain better statistical average (larger diffracting volume, i.e., more grains) for proper RS state probing.

#### *2.3. Neutron Diffraction Studies*

To avoid setup or instrument dependent error and justify the observations, the same specimens were investigated with the neutron diffractometer E3 [34] and the ESS test beam line V20 [45,46] at the research reactor BER-II in the Helmholtz-Zentrum Berlin (HZB). E3 is a constant wavelength and angular dispersive instrument (detector is moved around 2θ), whereas V20 is a time of flight beamline, where a wavelength dispersive setup with a fixed detector position was used.

#### *2.4. Microstructure Studies*

All specimens used for microscopic studies were prepared by using standard metallographic preparation routines. To investigate the microstructure of the specimens, visible light microscope (Nikon ECLIPSE L150, Amsterdam, The Netherlands) and scanning electron microscope (SEM, ZEISS EVO LS10, Zeiss, Oberkochen, Germany) were used. The SEM was operated using an acceleration voltage of 25 keV with magnifications ranging between 1000 and 5000×. Electron backscatter diffraction (EBSD) scans were performed in an SEM equipped with a field emission source and an automated EBSD acquisition system. Here, the SEM was operated at an accelerating voltage of 20 kV and a beam current of 5 nA. The EBSD scans were acquired with the specimen tilted to 70◦ at a working distance of 10 mm. Two EBSD maps have been obtained from two different locations at the centre of the specimen along the build direction.

#### *2.5. Diffraction Data Analysis*

After measurement, the software MATLAB was used to extract the data in xy binary format, which is friendly to many structure refinement programs. Then, the structure analysis software packages TOPAS-Academic was used to carry out least-square refinement (Pawley method) and standard single peak fitting routines [35] for stress/strain analysis. Phase quantities of the sample were determined by the Rietveld analysis [35]. The basic crystal structure information of the various phases needed for the Pawley and the Rietveld refinement were obtained from the Inorganic Crystal Structure Database (ICSD) [47]. All structural data used, such as reference and refined parameters, are tabulated in Table 2.



As overall diffraction patterns (or peak shapes) are a convolution of background, specimen and instrument, a standard cerium oxide specimen with lattice parameters 5.41165(1) Å was measured under the same configurations for calibration and instrument profile function extraction purpose. Then, the Pawley [35] refinements were performed to determine the profile function. After a good fit was obtained from cerium oxide data, the instrument profile function was set fixed together with zero error correction, and then a Pawley batch fitting was carried out to refine the structure of various phases to define best possible unit cell parameters of each measurement. The Pawley fitting was performed on the full diffractogram, which spans from 0 to 145 degrees in 2θ. The peak profiles were modelled with a modified Thompson–Cox–Hastings pseudo-Voigt (pV-TCHZ) profile function [35]. The background was fitted with a Chebyshev function with six coefficients and the zero-shift error together with axial divergence calibrated with a standard cerium oxide reference specimen. During batch refinement all instrument related parameters were kept fixed and only unit cell parameters together with background were refined. To avoid any complication introduced by oxides, elemental composition variation, and phase distribution to the measured strain, a single peak fitting was also performed together with the Pawley fitting in limited diffraction ranges, i.e., 1.19639–1.79480 Å.

#### *2.6. Strain Calculations*

Diffraction patterns are routinely used as a fingerprint of a material's crystalline structure. Any external stimulus, such as applied load or heat gradient can alter the interplanar lattice (*d*) spacing of the material, causing a change of the diffraction pattern. Therefore, the atomic lattice spacing in crystalline materials can be used as a natural strain gauge [16,48]. For instance, tensile stress will cause an increase of the lattice spacing in a given direction and compressive stress leads to the opposite. The average elastic lattice macrostrain (*εhkl*) in the sampled volume can be calculated by comparing the measured lattice spacing *dhkl* with that of the unstrained (stress-free) lattice spacing *d*0 *hkl* as [16,48]:

$$
\varepsilon\_{hkl} = \frac{d\_{hkl} - d\_{hkl}^0}{d\_{hkl}^0} \tag{1}
$$

If the target material is isotropic and all grains/domains in the material have the same responses to external stimulus, then the Equation (1) is sufficient to determine the stress/strain state in the material for a given direction of the stress/strain tensor. However, most crystalline materials have anisotropic properties, including the stiffness of the unit cell, and hence accurate stress/strain analysis requires consideration of multiple reflections. While individual peak fitting can be used to investigate the elastic and plastic anisotropy of individual lattice plane families, a fit of a complete diffraction profile—as used for results reported herein—can be used to determine the average lattice parameter from which the strain can be calculated that is closely representative of the bulk macroscopic strain (for the same direction).

The strain in a cubic material/phase (e.g., β-Ti) can be calculated by [48,49]:

$$
\varepsilon\_{cubic} = \frac{a - a\_0}{a\_0} \tag{2}
$$

Similarly, the strain in a hexagonal material can be calculated by [48,49]:

$$\varepsilon\_{\text{hexagonal}} = \frac{2\left(\frac{a}{d\_0} - 1\right) + \left(\frac{c}{c\_0} - 1\right)}{3} \tag{3}$$

where *a* and *c* are the lattice parameter of the given phase; *a*<sup>0</sup> and *c*<sup>0</sup> are the strain-free lattice parameters. Unstrained lattice parameter can be extracted by several methods, e.g., summarized by Withers et al. [32]. In this work, the *d*<sup>0</sup> *hkl*, *a*<sup>0</sup> and *c*<sup>0</sup> are obtained by (1) measuring the raw powder used for building these specimens used in the study; (2) measuring from the corner of each specimen; (3) taking the mean values of all scans of each sample.

#### **3. Results and Discussion**

#### *3.1. Microstructure*

Depending on processing and alloying elements, the Ti-6Al-4V microstructure may include α-Ti (hcp), β-Ti (bcc), α (hcp martensite), as well as different amount of dislocations, substructures and twinning [1,4]. The EBM built Ti-6Al-4V, however, commonly consists of two phases, namely α-Ti (hcp) and β-Ti (bcc) twinning [2,4,5,50]. Typical visible light micrographs of one of the specimens from build-direction and build plane are shown in Figure 2a,b, respectively. All materials studied show a microstructure which consists of intertwined α lath colonies where individual α-laths are separated by a thin layer of retained β-phase. This type of microstructure is formed when fast cooling rates from the β-phase field are achieved [51], the presence of grain boundary α does, however, indicate that the cooling rate was not fast enough to suppress the grain boundary α nucleation, which can occur

for AM processes with even faster cooling rates. High-resolution SEM micrographs corresponding to Figure 2a,b are shown in Figure 2c,d, respectively. In these SEM micrographs, the dark colour corresponds to the alpha phase and white colour corresponds to the β phase. The microstructure shown in Figure 2 is similar to the microstructure of EBM built Ti-6Al-4V reported by others [2–4,11,15]. The columnar nature of prior β grains that is shown in Figure 2 is a direct consequence of the thermal gradient that exists in the build direction [15].

**Figure 2.** Light optical microscopy image of S7V (**a**), and S7H (**b**). In (**c**) and (**d**), the images are taken with a SEM where the specimens used are S8V and S8H respectively. Z indicates the build direction and etching procedure employed to obtain the alloy microstructures was by immersion in Kroll's etchant for 10 s.

Synchrotron X-ray and neutron diffraction work confirmed the α+β microstructure observed by microscopy. Figure 3 shows the raw XRD diffraction patterns of as-received Ti-6Al-4V powders used for building all the specimens investigated in this study together with a representative diffraction pattern of EBM built material collected at the centre and one of the top corners (close to the final layer) of the same specimen. The colour coded, small vertical tick marks in Figure 3 (and all other figures throughout) represent the hkl peak positions of labelled phases in the same colour. As expected, the diffraction patterns from the corner and centre of the built material showed close similarity, but they both differ significantly from the powder pattern. According to Figure 3, it is evident that the as-received Ti-6Al-4V powders and the EBM specimens (independent of build location) contain predominantly α-phase with the HCP structure, consistent with previous studies [52]. In addition, the system also contained detectable amount of BCC β-phase (1.5–10 wt %) together with limited quantity of Ti-oxide (<1 wt %). The peak position and the intensity of the oxide diffraction peaks measured at different parts of the specimen did not show observable variation. Therefore, the effect of the oxide is excluded from the discussion. The presence of the β-phase can be clearly seen by the corresponding high angle diffraction peaks at d-spacing = 2.26 Å and d-spacing = 1.60 Å for the (011) and (002) reflections as reported in [53,54].

**Figure 3.** Diffraction pattern obtained by EDXRD of Ti-6Al-4V powder and EBM built material. The colour coded, small vertical tick marks represent the hkl peak positions of labelled phases in the same colour (α = purple, β = magenta).

It is known that the mechanical properties of the Ti-alloys strongly depend on the distribution of alloying elements and phases present in the system [1]. Generally, Ti-alloys show higher strength and higher density but lower creep strength with an increase in β-Ti phase content [1]. To ensure and control the properties of the material for a specific application, it is important to understand how different material processing affect the material phase composition of the alloy system. The microstructure and mechanical properties of Ti-6Al-4V for specific applications can be tailored by post heat treatment after manufacturing. Commonly, the post thermal process operations will take place in the α+β phase and/or single β-phase field. The volume fraction of β-phase increases with temperature and therefore governs the mechanical properties [55]. Therefore, it is important to understand how different AM processes will influence the formation/distribution of the β-phase in the built material. However, because of the limited phase quantity at room temperature of the β-phase in Ti-6Al-4V and its low diffraction strength, the β-phase has not been observed [56] or has been excluded from the analysis. While studying Ti-6Al-4V fabricated by the selective laser melting process (SLM), Chen et al. [56] did not observe a clear β-phase from either the as-received powder or the built solid specimens from their XRD work. According to Chen et al. [56], the reason behind weak or absent β-phase in their system is the high cooling rate used during the build. They argued that [56], SLM processed specimens maintain the original composition and crystal structure of the powders; all possible existing β-phase at high temperatures might all transform into the α or martensitic α' phase via rapid solidification during the SLM process. However, it is interesting that even the as received Ti-6Al-4V powder used in their work did not show detectable β-phase. We think that such phase absence might be explained by the sensitivity of the conventional XRD method and the textures present in the system in addition to the rapid cooling of SLM processing and powder manufacturing.

Ti-6Al-4V has been shown to have texture [1,12], and its development strongly depends on processing type and temperature [55]. Because of localized melting and solidification together with repetitive heating processes, the microstructure and texture formed in additively manufactured material are different from that in cast or wrought materials. In order to emphasize the complexity and variation in textures in EBM built Ti-6Al-4V, two grain orientation maps together with corresponding inverse pole figures obtained from the same specimen measured along the build direction at two different areas in the specimen are shown in Figure 4. Figure 4a is close to the last melted layer and Figure 4b is a few mm below. As can be seen in the figure, the material shows preferred grain orientation (texture), and it varies from one location to another despite that these two measurements were done in the same vertical plane along the same line and only a few mm apart. Microstructure and textures presented in Figures 2 and 4 indicate the complexity of the Ti-6Al-4V microstructure produced by EBM. Because of this highly varying microstructure and its strong correlation with the AM process parameters and build geometry, the mechanical properties of the material reported in the literature shows a large scatter [5,50,52].

**Figure 4.** Inverse pole figure maps of Ti-6Al-4V (S7V). The build direction is indicated by the white arrows. (**a**) is close to the last layer; (**b**) few mm below from (**a**); Images under (a,b) is the corresponding inverse pole figures.

#### *3.2. Peak Intensities and Phase Compositions*

All observed diffraction patterns from one diffraction angle (from detector element 15) collected during a vertical scan of one of the specimens, and the corresponding specimen patterns obtained from all 23 directions from the same specimen are presented in Figure 5a,b, respectively. As shown in Figure 5a, some of the α- and β-phase diffraction peaks are absent in certain parts of the build material from the diffraction patterns collected from only one scattering direction. From Figure 5a, one can see that there is no correlation between the absences of α- and β-phase diffraction peaks. In addition, such an absence does not coincide with the additive build layer thickness. Similar observations have also been observed for all other specimens studied in this work. The summed diffraction patterns of all measurements in all 23 scattering directions, however, showed all corresponding α- and β-phase diffraction peaks with similar peak intensity and width with a sign of α- and β-phases homogeneity as shown in Figure 5b. This difference is thought to be a result from the preferred orientation of the different grains (i.e., the α-phase is also present in probed volumes that do not show any α-peak for a single detector, but the peaks are observed in other scattering angles from the same volume). In Figure 5, HKL indices of α- and β-phases with significant changes are given at corresponding peak Bragg peak positions and interesting regions where β-phase peaks are absent in detector element 15 also marked with (1), (2) and (3) with green curly brackets.

Similarly, diffraction patterns from one of the EBM Ti-6Al-4V specimens collected along the build direction and in the build plane are shown in Figures 6a and 7a as a two-dimensional plot, viewed down the intensity axis, with the diffraction pattern number along the ordinate, and peak positions along the abscissa. For clarity and emphasizing the existence of β-Ti in the system, the peak intensity of well separated (002) β-Ti peak from the measurement presented in Figures 6a and 7a is plotted in Figures 6b and 7b, respectively. Peak intensity is colour coded to maintain consistency with previous figures. Note that the first and last data acquisition location (gauge volume) in Figure 7 does not exactly correspond to the edges of the specimen. There is about 300-μm clearance from each side. The peak intensity drop is shown in Figure 7b might be caused by the scan strategies used in this work. The outermost boundaries between the build and powder bed were melted first to form a contour. Thereafter powders inside the contour were melted in the manner of serpents' path (from

one side to another, then one step downwards followed by return, then one step downwards etc.), which is also known as hatching. As shown in Figures 5b and 6, none of the diffraction patterns showed any sign of α' martensite in the top layer of the build material unlike that reported in [15]. According to the Galarraga et al. [13], the observation of a martensitic structure is indicative of a high cooling rate imposed during solidification and subsequent cooling in the solid state of the last layer, and a cooling rate of >410 K/s may result in such a structure in last layer. Al-Bermani's [15] and Galarraga et al. [13] have also reported that faster cooling rates after solution heat treatment produce a greater amount of α' martensitic phase, with water-cooling at a rate of 650 K/s resulting in a fully α' martensitic microstructure. However, apart from the last layer, the first layer can also vary in microstructure compared to the bulk, due to the faster cooling rate close to the thermal conductive substrate. And these pronounced effects of cooling rates on microstructure also render microstructural differences. Nonetheless, the volume of these microstructural extremes is relatively small. Interestingly, Galarraga et al. [13] did not observe α' martensite in their unprocessed EBM built Ti-6Al-4V ELI (Extra Low Interstitial) specimens. However, based on the weakening of the (200) β phase diffraction peak at d-spacing = 1.60 Å and change in microstructure, tensile test results and microhardness values after post build heat treatment, they concluded that even air-cooled specimens may contain α' martensite. The difference between the observation reported by Al-Bermani's [15] and this study might be because of differences in specimen build geometry, process parameters and post-build heat treatments used in the two different studies.

**Figure 5.** Assembled diffraction patterns of a vertical scan (along the build length) of one of the EBM Ti-6Al-4V samples (S7V). (**a**) is from the detector element 15 alone and (**b**) is the sum of all 23 elements of the same scan. The black arrow indicates the build direction from bottom to top.

During EBSD and SEM studies of EBM built Ti-6Al-4V, Al-Bermani et al. [15] observed 100% β-phase within the build height 0–300 μm and mostly α-phase at other parts. According to Al-Bermani et al. [15], the β phase occurs due to the co-melting and diffusion of the alloying elements in the austenitic stainless-steel base plate and the initial Ti-6Al-4V layers. The β phase stabilizing elements, such as Cr, Fe, and Ni were provided by melting of stainless-steel plate. Since all measurements presented in this work have been performed from 13 mm upwards, we are unable to confirm that observation. Nonetheless, based on observations presented in Figures 4–7 it is safe to conclude that globally there is no significant phase difference in different parts of the EBM built Ti-6Al-4V parts at some distance away from the base plate despite the varied microstructure. This conclusion can be further supported by neutron diffraction studies. The deep penetration power into Ti and the larger

sampling volume of neutrons provide the possibility to evaluate the bulk average phase compositions and residual stress/strain of all present phases in the system.

**Figure 6.** (**a**) Accumulated diffraction patterns of vertical line scan along the build direction of specimen S8, and (**b**) corresponding peak intensity of (002) β-Ti (magenta) and (012) α-Ti. (purple).

**Figure 7.** (**a**) Accumulated diffraction patterns of horizontal line scan along the build plane of specimen S8, and (**b**) corresponding peak intensity of (002) β-Ti (magenta) and (012) α-Ti (purple).

To cross-check for potential variations among larger sampling volumes, data sets collected at the ESS test beamline at HZB on the same specimen (S8) are presented in Figure 8a. The beamline operates in time of flight mode, where neutrons start to travel from a source at time t = 0 and travel tens of meters to the specimen at which point they have separated by their different velocities and hence wavelengths. In this measurement, an incident beam size of 10 mm × 10 mm was used, with the detector positioned at a scattering angle of 2θ = 90◦. The results in Figure 8a present a relative direct comparison (the setup remained unchanged besides translating the specimen) between the upper half (red) and lower half (blue) of the specimen. Both diffraction patterns are almost identical to each other with the indication that there is no significant phase variation between the upper and lower part of the specimen. These results agree well with those in Figure 5b and supports the conclusion that phase distributions in the material are uniform.

**Figure 8.** Time-of-flight neutron data for specimen S8 (**a**). Raw diffraction patterns depicting a relative comparison between the upper and lower half of specimen S7, supporting the observations in Figure 6. (**b**) Collection of diffraction patterns for seven different EBM Ti-6Al-4V specimens.

To investigate whether different EBM process parameter combinations may lead to phase differences, an EDXRD pattern from the middle of different specimens is presented in Figure 8b. As shown in the figure, there is no phase difference between specimens other than intensity. As expected all specimens contain both α- and β-phases. Figure 9 represents a typical Pawley refinement result of data presented in Figures 6 and 7. In that figure, the blue dots represent observed intensities, the red line represents the calculated, and the green line is the difference curve on the same scale. The color-coded tick marks indicate the calculated positions of Bragg peaks. As seen from the difference curve, the fit is highly satisfactory. All diffraction patterns contained some of the major α-Ti and β-Ti peaks, which are commonly reported in the literature [57] together with many other minor and high angle peaks which have not been included in any other reported studies. The matches of the interplanar spacings (d-spacings) obtained from these peaks and unit cell parameters of α-Ti and β-Ti were not in perfect agreement with the data of the ICSD (see, Table 2). This difference might be caused by that different EBM process parameters were used to manufacture the materials investigated in these studies, which may influence the lattice strains and hence the overall lattice parameter.

**Figure 9.** One typical Pawley refinement result: Rwp 1.73, Goodness of fit (GOF) 1.52.

#### *3.3. Residual Strain*

One of the biggest challenges in diffraction-based residual strain measurements of additively manufactured material is obtaining the strain-free reference lattice parameters. Commonly, the reference can be determined from a stress-free specimen of identical material. Because of the

localized heating effect, microstructure variation and possible texture, specimens built by AM are expected to show some differences and even at different locations in the same specimen. Considering the microstructure difference between the ingredient powder, the corner of the build material and the center of the build material, only mean value-based strain calculations have been selected as the suitable strain reference extraction method. The lattice strain of α- and β-Ti calculated by using unit cell parameters (a and c) obtained from Pawley fitting of diffraction data from vertical and horizontal scans are shown in Figures 10 and 11, respectively. Each coloured line in these figures corresponds to one specific scan belonging to a different specimen. The colour code is the same for both figures. The error bars are equal to or smaller than the marker size.

Recently, Tiferet et al. [27] have reported that none of SLM or EBM specimens contain residual macro- nor micro-strains, instead the ingredient powder contains residual strain. Interestingly, lattice parameters of α-Ti obtained from ingredient powder from our studies show almost identical result as their bulk specimen. In addition, in contrast to Tiferet et al. [27], lattice parameters obtained from specimen corners, ingredient powder and bulk specimens from our measurements are similar. Such small lattice parameter differences indicate that the residual strain in EBM built material is low. The difference between our results and Tiferet et al. [27] may be related to the different processing approaches that were used.

In addition to lattice strains calculated from unit cell parameters of α- and β-phases obtained from Pawley pattern fitting, a lattice specific strain was also calculated for selected α- and β-phase hkl reflections after single peak fitting. Despite the availability of other hkl reflections, only selected α-phase ((010), (012), (110), (004)), and β-phase ((011), (002), (211)) hkl reflections, which do not overlap or are well-separated from other phases, were used for calculation. The general trends of the strains from single peak-based calculations are similar to whole pattern fittings shown in Figures 10 and 11. However, each lattice plane families yielded some disparity from one another. For simplicity and considering the superiority of strain calculation by using lattice parameters obtained by whole pattern fitting over single peak fitting, the results of single peak fitting are not shown.

From Figure 10, it is shown that the measured strain in both longitudinal and transverse directions in α- and β-Ti are fluctuating and they do not show any obvious trend with respect to deposition layers. However, closer inspection of the strain result of the α-Ti phase shows a slight compressive to tensile transition trend in both strain measurement directions (Figure 10) along the build. Even though the α-Ti strain result fluctuation does not match with the deposition layer, the frequency of the fluctuation appears to correlate with the length of the columnar grains. Therefore, to obtain reliable RS results from the experiment and modelling, we recommend considering the size of the primary β-grains in the measurement strategy, data interpretation and in theoretical simulations. Unlike α-Ti, the β-Ti did not show any observable transition and the strain appears to fluctuate more than the α-Ti (Figure 11). The reason for such a strong fluctuation might be that the β-Ti phase is the minority phase with high crystallographic symmetry and strong texture. Hence, the β-Ti phase in the Ti-6Al-4V system could be working as buffer phase, which causes local equilibrium of the stress/strain distribution in the system.

The material is built layer by layer via selectively melting the raw material. Therefore, each deposited layer in an additive manufactured structure undergoes multiple heating and cooling cycles with repetitive stress relaxation and accumulation processes. Hence, large thermal RS differences are to be expected at different heights in the build. However, at same build height or in same deposition layer, the heat treatment is expected to be the same. For that reason, the RS in the same layer should be the same. As predicted, the RS in the centre of the specimen measured along the scan direction showed linearity in both longitudinal and transverse directions (Figure 11). The magnitude and nature of the strain measured near the edge of most of the measurements appear similar to the result obtained from the centre or has a compressive nature.

**Figure 10.** The lattice strain from the central vertical line of specimens perpendicular to the EBM build plane. Transverse direction at left and longitudinal direction at right. Color code of each specimen can be seen from the legend at top of the figure.

**Figure 11.** The lattice strain from the central horizontal line of a specimen that is parallel to the EBM build plane. Transverse direction at left, and longitudinal direction at right.

The transverse strain in α-Ti, however, appears compressive on one side and tensile on the other. The EBM scan strategy used for building is expected to induce a sharp strain gradient from the surface/edge to the centre/inside of the build. During the EBM process, each layer of the target object is built in two steps. Firstly, the outer boundary is melted (contour). Then, in the second step, the actual part is built within the contours. The contour provides an interface between the actual build and the surrounding powder. Since molten metal is deposited on a colder contour wall, thermal contraction of the solidified material occurring during solidification creates a tensile stress in the deposit and compressive stress in the contour wall. The origin of this is not currently known. Whether it is build thickness or beam power requires further investigation. These observed differences within and between specimens may be attributed to the differences in their structure at the temperatures experienced during EBM processes. As shown in Figure 2a, the material microstructure consists of large prior β-grains with α-lamellas. Therefore, the scatter in the strain results presented could be understood to be a result of the grain statistics as the synchrotron X-ray experiment was conducted with a relatively small gauge volume. In order to justify whether this is the case, more measurements have been carried out on selected specimens with the E3 instrument at HZB. Similar to the synchrotron result, neutron diffraction measurements also did not show any significant difference from one another at different locations within the specimens (Figure 12). From Figure 12, one can see that there is neither any directional peak shift of the (112) peak of the α-Ti phase along the build direction nor peak position differences between specimens. This observation agrees with the results presented in Figures 10 and 11. The main difference between the neutron diffraction results in Figure 12 and the X-ray diffraction results in Figures 10 and 11 is that the neutron diffraction results seemed more linear than the X-ray diffraction results. This could indicate that there is indeed some grain size effect on the X-ray results. Therefore, to overcome similar problems, a larger beam size should be used in X-ray based measurements.

**Figure 12.** Variation of Ti(112) peak position with respect to build height by neutron diffraction. (**a**) Longitudinal direction, (**b**) transverse direction. It can be seen that variations are within the error bars.

Residual stresses are important to control with regard to quality and integrity of parts built by AM. Therefore, it is crucial to know the status and magnitude of any residual stresses to be able to reduce any detrimental effects. Since residual stresses in additive manufactured materials arise mainly from temperature gradients, the stress levels formed in the parts could possibly be modified by changing the heat gradients during manufacturing. The scan speed, beam current and offset focus of the EBM process influences the heat gradients and are likely to cause different states of residual stresses. However, according to the present observations, the influence of these process parameters is not that obvious (at least not for the parameters selected in this work). The reason for not observing any strain/stress trends or the effects of different process parameters is most likely because of the high preheating temperature with slow cooling used in the EBM process. In SLM it was reported that heating the build platform ≥ 200 ◦C can reduce the RS levels [58]. Therefore, in the future RS studies, the role of the preheating temperature should be considered.

There are some limitations in the current work that could be addressed by future work, concerning process parameter diversity. In AM there are many processes parameters that affect the cooling rate of the build and consequently the microstructure and RS. Therefore, the authors propose that further investigations be performed with other processes parameters in mind. In Table 1 the process parameters that have been used in this work are shown. An example is that in this work, only a current of 8 mA has been used, this process parameter greatly affects the cooling rate and would undoubtedly render different RS results.

#### **4. Conclusions**

Residual strains in Ti-6Al-4V specimens built by AM from nominal 45–100 μm diameter gas atomised powder using an electron beam melting technique have been evaluated by using state-of-the-art neutron and synchrotron X-ray diffraction techniques. Based on the present studies of the EBM built Ti-6Al-4V materials, the following conclusions can be made:


**Author Contributions:** Conceptualization, T.M.; Data curation, T.M.; Formal analysis, T.M.; Funding acquisition, R.P., M.S. and C.B.; Investigation, T.M., R.W., M.B., R.C.W., M.D. and N.S.; Methodology, T.M. and R.W.; Project administration, T.M.; Resources, T.M. and C.B.; Supervision, R.P., M.S. and C.B.; Visualization, T.M.; Writing—original draft, T.M.; Writing—review & editing, T.M, R.W., M.N., M.B., R.C.W., R.P., M.S. and C.B.

**Funding:** InterReg ESS & MaxIV: Cross Border Science and Society financially supported the project (MAH-003) along with "Nationellt rymdtekniskt forskningsprogram" (NRFP), the EU funded "Space for innovation and growth" (RIT), the "Graduate School of Space Technology" at Luleå University of Technology and the OP RDE, MEYS, under the project "European Spallation Source–participation of the Czech Republic—OP", Reg. No. CZ.02.1.01/0.0/0.0/16 013/0001794. Diamond light source (EE7858) and HZB provided beam time.

**Acknowledgments:** We acknowledge in-depth discussions with Adnan Safdar, J. Blomqvist, M.S. Blackmur, Axel Steuwer and their help during the synchrotron experiment. EBSD measurements were performed at the HZB Corelab Correlative Microscopy and Spectroscopy (CCMS) and authors thank Christiane Förster for EBSD sample preparation.

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

#### **References**


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