**1. Introduction**

In the course of evolution, nature has developed a broad range of nanostructured surfaces with many different purposes: to optimize friction behavior (e.g., [1–3]) or the behavior of liquids (e.g., [4,5]), to achieve antimicrobial (e.g., [6,7]) or antireflective effects (e.g., [8–10]) and to achieve decorative effects (e.g., [11–14]), just to name a few. Many of these nanostructures are also of high interest for technical applications and have been investigated accordingly (e.g., [1,15–17]).

Nanoimprinting [18–20] is a replication technology, capable of replicating nanoscale (e.g., [21,22]) as well as microscale features [23] on large areas [24–26].

Due to the fact that nanoimprinting can—in a single process step—replicate complex multilevel structures [1,24,27,28], it is ideally suited for the replication of such biomimetic or bio-inspired nanostructures. A big challenge for nanoimprinting, however, is the replication of undercut features, since it is usually not possible to remove the stamp from such features, especially if they are of nanoscale. It has, however, been shown in the literature that this can be achieved for larger features in the several µm range [29,30].

The nanostructures we are interested in are responsible for the metallic blue color of the Morpho butterfly. These structures have attracted considerable interest and their functioning has been researched in detail [31,32]. Based on the understanding of the optical properties of these nanostructures, it has been possible to mimic the optical effect using nanostructured surfaces and multilayer stacks of deposited materials [31]. We wanted to

**Citation:** Muehlberger, M.; Ruttloff, S.; Nees, D.; Moharana, A.; Belegratis, M.R.; Taus, P.; Kopp, S.; Wanzenboeck, H.D.; Prinz, A.; Fechtig, D. Nanoimprint Replication of Biomimetic, Multilevel Undercut Nanostructures. *Nanomaterials* **2021**, *11*, 1051. https://doi.org/10.3390/ nano11041051

Academic Editor: Konstantins Jefimovs

Received: 24 February 2021 Accepted: 9 April 2021 Published: 20 April 2021

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2021 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 (https:// creativecommons.org/licenses/by/ 4.0/).

investigate if it might be possible to avoid such complicated and expensive partly vacuumbased processes and replace them with a single nanoimprint process. The aim of the work presented in this paper was, however, not to replicate the full optical effect, but to demonstrate the basic possibility. The mastering process would have been significantly more complicated if we would have included the necessary randomness to its full extent as it would have been necessary to replicate the full optical effect. The capabilities of the nanoimprint process can, however, also be shown with a reduced master design.

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

#### *2.1. Master Design and Fabrication*

Figure 1 schematically shows cross-sections of the structures that we dealt with in this work. The very left structure represents the structure as it is found in nature and is drawn after SEM images as they are found in the literature (e.g., [12,33]). For our purpose, we repeated the individual structures periodically along the x-direction. The structures extend along the y direction, as can also be seen later in Figure 6 (left image).

As far as we understand from the literature [34], to obtain the full optical effect of the Morpho butterfly in our model structures, several aspects are missing, most notably randomness in the lateral placement of the individual "trees" as well as randomness in the height of the "trees". Furthermore, the branches of the trees are not positioned at the same height on the left and on the right side of the trunk in the real Morpho butterfly structure. Taking these facts into account for our master design would have significantly increased the challenges in master fabrication. At the same time, the nanoimprint challenge would have been basically the same. Therefore, since the goal was to investigate if nanoimprinting can be used to replicate such a structure, we focused on the model structures as sketched in Figure 1 in the center and on the right and did not include the randomness.

Details on the masters and the mastering process are described in [35]; here, only the basic facts should be mentioned. All masters were fabricated on silicon wafers, first using a multilayer deposition followed by lithography and anisotropic etching to etch through the multilayer stack. Additionally, one of the layer materials was selectively etched to create the undercut. The layer heights of the multilayer stack were 85 nm and 50 nm. The lateral width of the single T-structure was around 400 nm after processing, while tree structures with different lateral dimensions were fabricated. Undercuts in the range of 100 nm were typically achieved.

#### *2.2. Nanoimprint Material*

The imprint resin primarily used in this work is an in house development of JOAN-NEUM RESEARCH Forschungsgesellschaft mbH. In the development process, 13 different acrylate resins have been formulated and tested, with the main focus on low viscosity, high elasticity and low surface energy in order to enable the fast and complete filling of the undercut features and easy demolding of the cured structures. Some of these resins had

problems during demolding and stuck to the stamp, e.g., leaving residual material in the undercut areas. Others demolded easily but did not replicate the undercut features. Typical defects were collapsed T-structures. Out of the tested ones, the resin named NILcure® JR5 (JOANNEUM RESEARCH, Weiz, Austria) proved to be optimal for high fidelity replication of undercut structures.

This elastic UV curable acrylate resin can be efficiently cured within seconds by UV-A light between 350 nm and 400 nm at a moderate irradiance of 50–100 mW/cm2 . Additional properties facilitating the UV imprinting of undercut features are low viscosity (<20 m Pas), low glass transition temperature (below room temperature), high crosslinking and low surface energy after curing down to 20 mN/m—depending on the surface energy of the stamp [25]. Furthermore, this imprint material shows high scratch and abrasion resistance as well as chemical and weathering stability.

A key feature of NILcure® JR5 that was used in the course of this project is also that it can be used as a stamp material for NILcure® JR5 itself, meaning that it is self-replicable. The nanoimprint process using this material is described in the next section.

The challenge was to find the right balance between surface energy and mechanical properties, i.e., Young's modulus, tensile strength and elongation at break of the cured imprint resin. The surface energy of the resins was tuned by doping the resins with highly surface active anti-adhesive additives—such as 1H,1H,2H,2H-perfluorooctyl-acrylate—as well as a JR-proprietary related molecule.

For tuning the viscosity of the uncured and the mechanical properties of the cured imprint resins, various formulations containing multifunctional acrylate oligomers with aliphatic poly-urethane and poly-siloxane backbones as well as mono- and bifunctional acrylate monomers were investigated. The acrylate oligomers were chosen according to their mechanical properties, ranging from Young's Modulus = 20 MPa–100 MPa, tensile strength = 5–25 MPa and elongation at break = 10–100%. For formulating the UV-imprint resins, reactive diluents had to be added to the highly viscos acrylate-functionalized oligomers. Mono- as well as bi-functional acrylate monomers were chosen. Whilst the addition of mono-functional acrylate monomers largely preserves low Young's modulus and high elongation at break, the tensile strength is generally lowered. The addition of bifunctional acrylate monomers, on the other hand, decreases the elongation at break and increases the Young's modulus and tensile strength of the cured resins [36,37].

NILcure® JR5 featuring optimal mechanical properties and the JR proprietary fluorinated anti-adhesive additive turned out to enable the highest imprint fidelity for the undercut structures of all investigated resin formulations.

A universal testing machine (Instron 3342 [38]) was used to determine the mechanical parameters of the most successful material NILcure® JR5 (see Figure 2). This machine can be used to perform classical tensile tests (Young's modulus, tensile strength, elongation at break) as well as other important measurements such as friction and adhesion measurements (peel test). For the determination of the Young's modulus and the elongation at break (e.g., [39]), a classical tensile test was performed. The rectangular tensile specimens had dimensions of 40 mm × 10 mm × 150 µm (clamping length, width, thickness). The test speed was 5 mm/min. The following figure shows an exemplary measurement of JR5. The following parameters were obtained: Young's modulus: 246 MPa, elongation at break: 8%. The surface energy of the liquid material is 25.5 mN/m as determined by the pendant drop method [40,41], using a DSA100 from Krüss [42].

Initially, materials such as h-PDMS (hard Polydimethylsiloxane) [43] and MD700 [44] as stamp materials and OrmoComp [45] as an imprint material were used. However, the results were not encouraging in terms of defectivity (see later in Figure 4), which is why our own materials were formulated and tested. The results are discussed below in the results section.

**Figure 2.** Tensile stress plotted versus tensile strain for the NILcure® JR5 material.

#### *2.3. Nanoimprint Process*

The goal of our nanoimprint process was to replicate a complex undercut nanostructure in a single imprinting step. To do so, we fabricated a stamp from a nanoimprint master and used this stamp in a nanoimprint process. For all process steps, we used UV-based nanoimprinting, i.e., UV-curable materials. The sample fabrication sequence is sketched in Figure 3.

**Figure 3.** Sample sequence and nomenclature for this work: From a master, a nanoimprint stamp was fabricated (stamp material on backplane), which was then used to create the imprint (imprint material on substrate).

All masters were treated with an anti-sticking layer (BGL-GZ-83 from PROFACTOR GmbH [46,47]) to allow for easy removal of the stamp from the master structure. To fabricate the stamp, NILcure® JR5 resin was manually deposited on the master. A UV-transparent flexible backplane was used and placed gently onto the droplet in a rolling motion alongside the direction of the T-structures (y-direction in Figure 1). The whole stack was UV-exposed with a wavelength of 365 nm and an intensity of about 100 mW/cm2 for a minimum of 30 s. Both UV-LED (light emitting diode) light sources (self-made at PROFACTOR, 365 nm) and UV fluorescent tubes (supplier: biostep GmbH [48]) can be used. After curing, the stamp was very gently removed from the master, again in the y-direction, to ensure easier demolding and also to prevent the tree branches from being destroyed.

As backplanes, different polymer films have been used. Two hundred micrometerthick commercial PVC (polyvinyl chloride) films worked well, as well as PET (polyethylene terephthalate) films with a thickness of 125 µm. The PVC was treated with an oxygen plasma before applying an adhesion promoter (HMNP12 from PROFACTOR GmbH [49]). The PET films were already supplied with an adhesion promoter, such as Kemafoil HSPL80 from Coveme [50,51] and Melinex ST506 from DuPont [52]. No backplane was used for the imprints using h-PDMS stamps.

After having obtained the stamps in the way described above, the imprints were made using those stamps in the same way. The imprint material was manually placed on the substrate, the stamp was carefully brought into contact with the coated substrate and, after curing, the stamp was separated from the finished imprint. The substrates that were used were the same polymer foils as for the backplanes for the stamps, but also glass substrates were used (e.g., Superfrost Ultra Plus [53]).

For both processes—the stamp fabrication as well as the actual imprinting—no additional pressure was applied. The imprinting was performed manually and was left to the capillary forces to obtain good contact between the stamp and substrate. The imprinting area was always several mm2 , typically 5 × 5mm2 .

There are two reasons for using polymer foils in this process chain. Firstly, it was necessary for sample characterization for both stamps and imprints, as described below. Secondly, and this is only valid for the stamp, the flexibility of the stamp allowed for easy separation of the stamp from the imprint in a peeling-like motion. The flexibility and UV-transparency will also allow the use of such stamps as imprinting plates in roller-based nanoimprint processes [25,26], which would pave the way for industrial implementation.

#### *2.4. Sample Characterisation*

Sample characterization was performed using atomic force microscopy (AFM) (Bruker Dimension Edge, Bruker Corporation, Billerica, MA, USA) (PROFACTOR) and scanning electron microscopy (TU Wien (Zeiss Neon 40 EsB (Zeiss Microscopy, Jena, Germany) and JOANNEUM RESEARCH (SM-IT 100 (JEOL, Tokyo, Japan)). For the characterization of the master, focused ion beam (FIB) cutting (TU Wien using a Zeiss Neon 40 (Zeiss Microscopy, Jena, Germany) was also used. It proved to be challenging to obtain reliable information on the undercut on the polymer samples from these methods. AFM imaging naturally cannot provide direct evidence of any undercut features, but also indirect evidence gathered by comparing the lateral dimensions of master, stamp and imprint proved to be not conclusive. One reason for this is tip convolution effects. Additionally, standard SEM imaging caused problems since the polymeric nanostructures tend to deform under electron beam bombardment or during FIB cutting.

High-quality cross sections of the nanostructures (in x-direction) on polymer substrates were made by means of cryo-ultramicrotomy (PowerTome XL RMC Products, Boeckeler®, Tucson, AZ, USA [54]). The samples were physically hardened by cooling with liquid nitrogen below their glass transition temperature at −60 ◦C. Nanomaterial imaging and analysis were performed by utilizing the scanning electron microscope RAITH e-line at 3 kV. Cryo-ultramicrotomy at JOANNEUM RESEARCH for sample preparation proved to be the only variant that led to reliable results.

#### **3. Results**

#### *3.1. Single Layer Undercut Structures*

Using the imprinting and characterization procedures described above, we were able to show the full nanoimprint chain from master to final imprint, for both the single "Ts" as well as the "trees".

In the course of the project, and before the NILcure® JR5 material was used, several other material combinations were tested. Results are shown in Figures 4 and 5. Initially, our approach was to use soft stamps and more rigid imprint materials, such as h-PDMS or MD700 as soft stamp materials and OrmoComp as an imprint material. However, due to

the limited tuning possibilities with these commercial materials and the high number of defects (see Figure 4) observed already in the stamps for our special application, we started to test our own materials. In Figure 4, two different material combinations are shown. On the left, the result of an imprint using an h-PDMS stamp and OrmoComp as the imprinting material is shown. As can be seen, the T-shapes are not well defined and the edges of the Ts are rough and irregular, indicating problems during the imprinting step—either non-filling of the undercut or mechanical tear-off of the overhanging part of the nanostructures, or possibly both. A lot of defects can also be seen in the right image of Figure 4, which shows an SEM image of an OrmoComp imprint using an MD700 stamp. Additionally, here the overhanging structures are not well replicated. Here, defects on the bottom can also be observed, which are, however, related to the defective master which was used for these experiments and are not related to the undercut structures.

**Figure 4.** Left: cross-sectional SEM image of an imprint using an h-PDMS stamp and OrmoComp as imprinting material. Right: SEM image imprint using MD700 as stamp material and OrmoComp as imprint material.

**Figure 5.** Cross-sectional SEM images of stamps made from different materials. For NILcure® JR2 and JR4 it can be seen that the undercut is not replicated at all, while for JR3 the undercut is only partly replicated.

Switching to the self-made materials described above initially resulted in stamps which did not replicate the undercut, either completely or only partly, as shown in Figure 5, where cross-sectional SEM images of three stamps made from three different experimental materials are shown. It can be seen that the undercut is not replicated (left image and right image) at all, or only partly (center image).

The images in Figures 4 and 5 give a representative overview of the typical defects we observed: either non-replication of the undercut or only partial replication. We attribute the non-replication of the undercut either to non-filling of the undercut part of the structures, which is related to the wetting properties of the material and its viscosity, or to a material which was too soft, so that the nanostructures collapsed. The second type of defect we observe is when the undercut was only partly replicated. This we attribute to an imprint material that was mechanically too brittle to allow for a defect-free separation. The defects described above appear on the whole area with the same density, meaning that the full imprint area cannot be used. This is in contrast to other typical nanoimprint defects induced by dust particles or air bubbles, which are very local. In summary, the mechanical stability of the stamp and the imprint material have to be right as well as the surface energy and wetting properties.

The material that fulfilled all requirements was the NILcure® JR5 material decribed above. Figure 6 shows SEM images of the full sequence from master over stamp to imprint using the NILcure® JR5 nanoimprint material. The left images show the master structure, the center images the stamp, which was replicated from the master, and to the very right the imprint is shown, which was made using the stamp. It can be seen that the undercut features are nicely replicated. Looking at the master structure in more detail, a slight tapering of the T crossbar can be noticed, which is indicated by the dashed line in the second row of images in Figure 6. The same tapering can be observed in the SEM crosssection of the stamp (top-down) and the imprint as well. This shows that in the whole process chain the replication of the undercut structures works very well.

We successfully used the same stamp for several imprints. Figure 7 shows SEM crosssectional images of the first imprint (left) and of the fifth imprint (right), all made with the same stamp. This shows that the stamp remains intact during the imprinting process, which is of course a necessary prerequisite for a nanoimprint process. The difference in the shape of the edges of the T-structures compared to the ones shown in Figure 6 results from the fact that a different master was used, which exhibited a slightly different geometry.

**Figure 6.** SEM cross-sectional images of the sample sequence from the master structure [35] (left) over the stamp in NILCure® JR5 (center) and the imprinted structures in NILcure® JR5 (right). The second row shows close-ups from the images in the upper row.

**Figure 7.** SEM cross-sectional images of the first imprint from (left) the master structure and the fifth imprint (right) using the same stamp.

#### *3.2. Mulitlayer Undercut Structues*

In addition, the multilayer structures could also be replicated very well with the described procedure. Figure 8 shows the sequence from the master over stamp to the imprint for the multilayer structures. Especially in the images of the stamp and the imprint, it can be seen that the tree branches partially stick together, and this can be attributed to the flexibility of the imprint in connection with the small size of the structures. It cannot be excluded that this is also an artifact of SEM imaging or sample preparation. Nevertheless, it can be seen that the nanoimprint replication of such complex features is possible.

**Figure 8.** SEM cross-sectional images of the multilayer structure sequence from the master structure in Si (left) over the stamp in NILCure® JR5 (center) and the imprinted structures in NILcure® JR5 (right).

#### **4. Discussion**

Taking the Morpho butterfly structure as an example, we could show that nanoimprinting is capable of replicating complex undercut nanostructures. For such a process to work, it is crucial to find a material which exhibits the right wetting properties as well as suitable mechanical properties. This broadens the potential of nanoimprinting significantly and opens up the possibility for a cost effective single-step replication of, e.g., bioinspired Morpho-blue nanostructures. We tested several material combinations to arrive at a process–material combination that reliably allowed for the replication of such structures. Although all the results shown in this paper were obtained using manual nanoimprinting, we designed our processes in such a way that the transfer to roller-based processes [25,26] is possible. The roller-based nanoimprint processes will also mimic the peeling-like separation of the stamp from the imprint, which was used successfully in our experiments.

Further research will be necessary to see how the material formulation and material properties and the dimensions of the undercut nanostructures are related, i.e., if the same material can be used for significantly larger or smaller structures. The research presented here is, in our opinion, a first important step since it shows for the first time that the replication of undercut nanostructures is feasible using UV-based nanoimprint lithography.

Regarding the Morpho butterfly structure, it will be interesting to introduce the necessary randomness in the master design to faithfully reproduce the optical effects. Using nanoimprinting, it then should be possible to replicate these structures in a single imprinting step in a cost efficient way. This will be facilitated in more automated processes which will also reduce defectivity.

Depending on the application, it might be interesting or necessary to protect the imprinted nanostructures, which would require a material with a proper refractive index contrast to the imprinting resin—a topic which we did not address so far.

Furthermore, we hope to come across many more interesting nanostructures, for all types of applications, which could be enabled using this or similar processes.

**Author Contributions:** Conceptualization, M.M.; investigation, S.R., D.N., A.M., M.R.B., S.K., P.T., H.D.W., A.P., M.M.; writing—original draft preparation, M.M.; writing—review and editing, S.R., D.N., S.K., A.M.; project administration, M.M.; funding acquisition, M.M., D.F. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded by the rollerNIL project (FFG, grant 843639).

**Data Availability Statement:** The data presented in this study are available on request from the corresponding author.

**Acknowledgments:** The authors acknowledge administrative support by Hannes Fachberger and Eva Mühlberghuber (PROFACTOR), as well as highly interesting and very helpful discussions with Jan Stensborg and Leif Yde (Stensborg A/S), Christof Neuhauser (D. Swarovski KG), Stephan Trassl (Hueck Folien GmbH), Jan Klein (micro resist technology GmbH), Uwe Richter (SENTECH Instruments GmbH) as well as Martin Panholzer, Jean-Pierre Perin and Kurt Hingerl (Johannes Kepler University Linz).

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

#### **References**


*Edited by Dennis Wat and Dilip Nazareth*

*Cystic Fibrosis - Heterogeneity and Personalized Treatment* provides the latest research and clinical evidence for clinicians, scientists and researchers involved in the care of patients with cystic fibrosis (CF). This book outlines the burden of the CF microbiome, utilisation of CF registries to impact future care, the sequelae of hepatobiliary complication, the use of upcoming technologies to provide patient-centred care, and provides an overview of cystic fibrosis transmembrane regulator (CFTR) modulators. Looking after patients with CF is highly rewarding, allowing those of us to combine our dedication and problemsolving skills to create a personalized approach. This book is invaluable for those involved in the care of CF patients.

ISBN 978-1-78984-145-9

Cystic Fibrosis - Heterogeneity and Personalized Treatment

Published in London, UK © 2020 IntechOpen

© undefined undefined / iStock


**Giovanni Ferrara**

Giovanni Ferrara is full professor at the Deparment of Industrial Engineering at Universita degli ` Studi di Firenze, and he is the coordinator of the REASE research group. He is an active member of the EU wind energy community with many years of experience, and he has authored more than 80 international publications in wind energy related to his many scientific contributions in the area of machinery and energy engineering.


**Acknowledgements**

Funding was provided by a graduate scholarship grant from CONACYT and a postdoctoral research grant from DGAPA at Universidad Nacional Autónoma de México (UNAM). My ideas about biogeography and diversification were greatly benefited from discussions with Enrique Martinez Meyer, Oscar Flores Villela, Luna Sanchez Reyes and Mark. E. Olson. Thanks to Mark E. Olson to read several draft versions of this manuscript and for his continuous thoughtful comments. All remaining errors are my own.


**3. Ladder Operators and the Degeneracy of the Spectrum of Operators**

We identify the degeneracy of the spectrum of the operator L1 as a consequence of the existence of a particular kind of operators, called ladder operators. We provide a general definition of ladder operator after proving the following result, which can be indicated as the *shift theorem*.

**Theorem 5** (**Shift theorem**)**.** *Let* O *be an operator acting on a Hilbert space, and let* **v** *be an eigenfunction of* O *having an eigenvalue λ. If another operator* T *satisfies the condition* [O, T]**v** = *μ*T**v***, where the coefficient μ is a real number, then: either* T**v** *is the null function or* T**v** *is another eigenfunction of the operator* O *with eigenvalue λ* + *μ.*

**Proof.** If such an operator T exists, we have that, by linearity and and since *λ* is an eigenvalue of O, the above relation becomes:

$$[\mathbb{O}, \mathbb{T}]\mathbf{v} = \mu \mathbb{T}\mathbf{v} \qquad \Longleftrightarrow \qquad \mathbb{O}\mathbb{T}\mathbf{v} - \mathbb{T}\mathbb{O}\mathbf{v} = \mathbb{O}\mathbb{T}\mathbf{v} - \lambda\mathbb{T}\mathbf{v} = \mu \mathbb{T}\mathbf{v}$$

implying the new eigenvalue equation:

$$\mathbf{OTv} = (\lambda + \mu)\mathbf{Tv}\_{\prime}$$

meaning that either T**v** = **0** or *λ* + *μ* is an eigenvalue of O associated with the eigenfunction T**v**, so the proof is complete.

**Definition 4** (**Ladder operators**)**.** *An operator* T *satisfying the hypothesis of the shift theorem is called the ladder operator for the operator* O*. In particular,* T *is a:*

A very interesting case in which Theorem 5 is applied occurs when there exists a complete set of *n* self-adjoint operators O1, O2, ..., O*n* acting on a Hilbert space such that, by virtue of Theorem 3, there exists a basis of the space formed by all their simultaneous eigenfunctions {*y*1, *y*2,..., *yn*}.

If there exists an operator T commuting with the *k* operators O*i*1 , O*i*2 , ... , O*ik* and satisfying the *n* − *k* relations of the shift theorem with the remaining *n* − *k* operators O*j*1 , O*j*2 ,..., O*jn*−*k* for some certain eigenfunction *y*¯, the following relations hold:

$$\begin{cases} \mathbb{O}\_{i\_1}(\mathbb{T}\mathcal{Y}) = \mathbb{T}\mathbb{O}\_{i\_1}\mathcal{Y} = \mathbb{T}\lambda\_{i\_1}\mathcal{Y} = \lambda\_{i\_1}(\mathbb{T}\mathcal{Y})\\ \mathbb{O}\_{i\_2}(\mathbb{T}\mathcal{Y}) = \mathbb{T}\mathbb{O}\_{i\_2}\mathcal{Y} = \mathbb{T}\lambda\_{i\_2}\mathcal{Y} = \lambda\_{i\_2}(\mathbb{T}\mathcal{Y})\\ \vdots\\ \mathbb{O}\_{i\_k}(\mathbb{T}\mathcal{Y}) = \mathbb{T}\mathbb{O}\_{i\_k}\mathcal{Y} = \mathbb{T}\lambda\_{i\_k}\mathcal{Y} = \lambda\_{i\_k}(\mathbb{T}\mathcal{Y}) \end{cases} \tag{12a}$$

and:

$$[\mathbb{G}\_{\text{j}\prime}, \mathbb{T}]\bar{y} = \mu\_{\text{j}} \mathbb{T}\bar{y}, \quad [\mathbb{G}\_{\text{j}\prime}, \mathbb{T}]\bar{y} = \mu\_{\text{j}} \mathbb{T}\bar{y}, \quad \cdots, \quad [\mathbb{G}\_{\text{j}\prime}, \mathbb{T}]\bar{y} = \mu\_{\text{j}} \mathbb{T}\bar{y}, \tag{12b}$$

from which we obtain that the function T*y*¯ is either the null function or a simultaneous eigenfunction of O*i*1 , O*i*2 , ... , O*ik* with respect to the same eigenvalues *λi*1 , *λi*2 , ... , *λik* , respectively. Therefore, by Theorem 5, that function is a simultaneous eigenfunction of O*j*1 ,..., O*jn*−*k* with respect to the shifted eigenvalues *λj*1 + *μj*1 ,..., *λjn*−*k* + *μjn*−*k* .

**Remark 1.** *The degeneracy of the spectrum of a given operator* O *can be clarified (we precisely use the term 'clarification' if it is viewed in terms of ladder operators). Basically, we can consider the complete set of operators as a necessary tool to eliminate the degeneracy of the spectrum of* O *and to identify all the operators* T1, ... ,T*p that satisfy the relations* (12a) *together with the operator* O *and the relations* (12b) *of the shift theorem with the remaining operators of the complete set.*

The operator L1 has a degenerate spectrum because its eigenvalues *λn* given in (11a) are independent of the parameters and *m*. Since the operator L1 belongs to the complete set of operators {L1, *A*(*θ*, *ϕ*), M1 3}, in order to clarify the whole degeneracy in terms of ladder operators, it is sufficient to find the ladder operators T1 and T2 commuting with L1. Besides commuting with L1, such operators also satisfy the relations (12b) with the operators *A*(*θ*, *ϕ*) and M1 3, in such a way that the functions T1*ψnm*(*r*, *θ*, *ϕ*) and T2*ψnm*(*r*, *θ*, *ϕ*) are eigenfunctions of L1 associated with the same eigenvalue *λn* and eigenfunctions of *A*(*θ*, *ϕ*) and of M1 3 associated with a shifted eigenvalue with respect to and *m*, respectively.


**Separation of the Formation Mechanisms of Residual Stresses in LPBF 316L**

#### **Alexander Ulbricht 1,\*, Simon J. Altenburg 1, Maximilian Sprengel 1, Konstantin Sommer 1, Gunther Mohr 1,2, Tobias Fritsch 1, Tatiana Mishurova 1, Itziar Serrano-Munoz 1, Alexander Evans 1, Michael Hofmann 3 and Giovanni Bruno 1,4**

Received: 29 July 2020; Accepted: 3 September 2020; Published: 14 September 2020

**Abstract:** Rapid cooling rates and steep temperature gradients are characteristic of additively manufactured parts and important factors for the residual stress formation. This study examined the influence of heat accumulation on the distribution of residual stress in two prisms produced by Laser Powder Bed Fusion (LPBF) of austenitic stainless steel 316L. The layers of the prisms were exposed using two different border fill scan strategies: one scanned from the centre to the perimeter and the other from the perimeter to the centre. The goal was to reveal the effect of different heat inputs on samples featuring the same solidification shrinkage. Residual stress was characterised in one plane perpendicular to the building direction at the mid height using Neutron and Lab X-ray diffraction. Thermography data obtained during the build process were analysed in order to correlate the cooling rates and apparent surface temperatures with the residual stress results. Optical microscopy and micro computed tomography were used to correlate defect populations with the residual stress distribution. The two scanning strategies led to residual stress distributions that were typical for additively manufactured components: compressive stresses in the bulk and tensile stresses at the surface. However, due to the different heat accumulation, the maximum residual stress levels differed. We concluded that solidification shrinkage plays a major role in determining the shape of the residual stress distribution, while the temperature gradient mechanism appears to determine the magnitude of peak residual stresses.

**Keywords:** additive manufacturing; Laser Powder Bed Fusion; LPBF; AISI 316L; online process monitoring; thermography; residual stress; neutron diffraction; X-ray diffraction; computed tomography

#### **1. Introduction**

In recent years, Additive Manufacturing (AM) has evolved from a method for rapid prototyping to a mature production process for certain parts in industries, such as the aerospace industry [1]. Among the different AM manufacturing techniques, Laser Powder Bed Fusion (LPBF) is an important technique for the production of net shaped metallic parts [2]. Early research conducted by Mercelis and

Kruth [3] showed that metallic parts made by LPBF inherently contain residual stresses (RS). They had described two driving mechanisms for the formation of RS: the Temperature Gradient Mechanism (TGM) and the Solidification Shrinkage Mechanism (SSM). The two mechanisms are interlinked and their combined effect on RS in AM 316L is a topic of current research [4–7]. Wang et al. [6] showed within long bars of LPBF 316L that scan strategies using shorter scan tracks reduced RS and attributed this to lower solidification shrinkage. Roehling et al. [7] observed a decrease of RS in samples with a bridge geometry manufactured by LPBF of 316L due to post-solidification heating of each layer during the build job using selective large-area diode surface heating. This method aimed to decrease the cooling rate. Each of the two publications had mainly utilized one of the two mechanisms to reduce RS: Wang et al. [6] mainly utilized the SSM, whereas Roehling et al. [7] mainly utilized the TGM. In both cases, a reduction of RS was observed. However, there is still a level of uncertainty on the magnitude of the influence of each mechanism onto the shape of the resulting RS field.

Diffraction is a well-known non-destructive method to evaluate RS [8–10]. Determining elastic strains by measuring the variation of lattice spacing provides a powerful method to identify RS. This is achieved at the surface by Lab X-ray Diffraction (XRD), up to a depth of about 5 μm in metals, as well as in the bulk by Neutron Diffraction (ND) up to a depth of about a few mm to a few cm [11–13]. In this work, the bulk triaxial RS state was determined using ND and was combined with the Lab XRD biaxial RS state at the surface. This methodology allows for the mapping of the RS distribution across the complete cross-sectional plane. Such residual stress tends to be compressive in the bulk and tensile near the surface [14]. ND enables the non-destructive determination of the triaxial RS state over a complete two-dimensional (2D) plane or three-dimensional (3D) volume. Destructive methods, such as incremental or deep hole drilling, slitting, or contour method, would also yield stress depth profiles, but it would be extremely difficult to determine triaxial stress states over a complete cross-section.

To exploit the benefits of lightweight, load driven structural designs for LPBF metallic parts, it is necessary to understand RS in those parts, since their effect on fatigue life can be significant [15]. In order to understand RS, it is necessary to decouple the contributing mechanisms, especially if we aim at modelling the manufacturing process.

Therefore, this study aims at unravelling the contributions of the two mechanisms, to provide a basis for discussion on the length scale of RS introduced into parts by the TGM and the SSM. Therefore, the specimen design was chosen to provide similar solidification shrinkage, but at the same time different cooling rates without changing the volumetric energy density (VED) of specimens. Based on this design, similar RS results should be assigned to equal solidification shrinkage, whereas differences should be caused by the different cooling rates. Online monitoring by thermography during the build process was used in order to assess these cooling rates and their effect on RS formation.

The TGM is mainly related to process parameters (e.g., VED) and the SSM is mainly related to the length of shrinking scan tracks. Therefore, the results of this study might help to decide which approach is more suitable to reduce RS for a specific part design and its expected load profile.

Additionally, the results from Micro Computed Tomography (μCT) and Optical Microscopy (OM) were evaluated to link RS fields and defect distributions, with the aim to produce a holistic approach towards the analysis of the interconnection of TGM and SSM.

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

#### *2.1. Material and LPBF Processing Conditions*

Austenitic stainless steel 316L powder was processed by the commercial LPBF system SLM280 HL (SLM Solutions Group AG, Lübeck, Germany). The powder was characterised by its supplier: it has an apparent density of 4.58 g cm−3 and a mean diameter of 34.69 μm. The cumulative mass values of the particle size distribution are: *D*10 = 18.22 μm, *D*50 = 30.50 μm, *D*90 = 55.87 μm. The LPBF system uses a single 400 W continuous wave ytterbium fibre laser with a spot size of approx. 80 μm in a focal position. The processes were conducted in an argon gas atmosphere with an oxygen content of less than 0.1%. The parts were manufactured on a stainless steel substrate plate, which was heated up to 100 ◦C as a preheating temperature before the start of the build process. Two prismatic specimens of the dimension 24 mm × 36 mm × 24.5 mm were manufactured in two separate built processes. In order to remove specimens from the base plate a band saw was used. This reduced the height to a final value of 21 mm. A specimen design of low aspect ratio was chosen for this experiment in order to prevent significant RS relaxation due to distortion after the removal from the substrate plate. Such a distortion had been reported in literature for LPBF 316L [16,17]. Although the removal from the substrate plate may have caused a degree of RS relaxation, the overall relative trend between the specimens was considered to be mainly unaffected. The specimens were placed close to the border of the base plate to fit within the field of view of the thermography camera setup. The specimens were manufactured using the following process parameters: layer thickness *t* = 50 μm, scanning velocity *v* = 700 mm s−1, laser power *P* = 275 W, and hatch distance *h* = 0.12 mm. Two different so-called border fill scanning strategies were applied, which scan along the edges of the rectangular cross sections of the parts: for one the scanning sequence starts in the centre of the part with growing rectangles towards the perimeter and the other has a converse scanning sequence (see Figure 1). The interlayer time (according to Mohr et al. [18]) was approximately 27.6 ± 1.0 s due to time variations between re-coating forwards and backwards. Therefore, the total time for each build process of 490 layers added up to 3.76 h.

**Figure 1.** Schematics of both border fill scan strategies. (**a**) Describes the Centre to Parameter (CtP) strategy indicated by the green arrow, while (**b**) shows the Parameter to Centre (PtC) scan strategy indicated by the blue arrow. The black arrows show the direction of scan of the laser around each border fill scan.

#### *2.2. Thermography*

An ImageIR 8300 hp camera (Infratec GmbH, Dresden, Germany) working in the spectral range of 2–5.7 μm was used for thermographic measurements. It was mounted on top of the SLM280 HL machine's build chamber, observing the build plate through a sapphire window. The chosen subframe image had a size of 160 px × 114 px featuring a geometric resolution on the build plate of 360 μm px−1. The acquisition frame rate was set to 1000 Hz. The camera was calibrated for black body radiation. Due to the fact that the emissivity of the used material is well below unity [19] and the process was observed through optical elements, the calibration is not valid for quantitative evaluation of the obtained thermography data. Nonetheless, assuming that the emissivity remains (approximately) constant during the build process, the obtained apparent temperatures enable comparisons within a single build process and between the two different build processes. The thermography data for the two specimens were obtained during the build process while using two different calibration ranges: 673 K to 1073 K for the CtP specimen (- ) and 623 K to 973 K for PtC specimen (- ). Several overlapping ranges were being tested during this experiment to find the optimum for these specimens. Values within the overlapping apparent temperature range of 673 K to 973 K can be compared between the two build processes. For quantitative evaluation of the process, it would also be necessary to address the

additional error in temperature estimation introduced by the limited spatial resolution of 360 μm px−1. This error diminishes once the spatial temperature gradients have decreased due to lateral heat flow. These limitations are discussed in more detail by Mohr et al. [20]. Despite these limitations, qualitative analysis of the thermography data revealed results that contribute to the understanding of the presented RS results.

#### *2.3. Lab X-ray Diffraction*

A StressTech Xstress G3 X-ray diffraction instrument (Stresstech GmbH, Rennerod, Germany) was used in order to determine the RS distribution at the surface of the specimens according to the sin2 *ψ*-method. Based on the assumption that that principal stresses are aligned with the geometrical axes of the specimens and the normal stress component can be neglected at the surface, the RS could be calculated from the slope of the linear fit of the lattice spacings over sin2 *ψ*-plot [9,21]. The *ψ*-tilt was carried out in the angular range of *ψ* = −45° to *ψ* = 45° in 19 steps. The specimens were tilted around two perpendicular axes, to yield two perpendicular stress components. On the 36 mm surface, this corresponds to seven measurement positions of the prisms' normal and longitudinal stress component (see blue circles in Figure 2b). On the 24 mm surface, five measurement positions correspond to the normal and transversal component of the RS distribution of the prisms. The exposure time for each acquisition was 5 s. The 311 diffraction line at a corresponding 2*ϑ* angle of 152.26° was acquired using a Mn Kα radiation source and a 2 mm diameter collimator. Further details of the setup were described by Thiede et al. [13]. The software Xtronic (Stresstech GmbH, Rennerod, Germany) was used for data processing. The peak fitting process was performed using the Pearson VII function and the background was fitted with a parabolic function. The diffraction elastic constants (DEC) were calculated for austenitic steel 316L based on the Eshelby–Kroener model [22]. The calculated Young's modulus of *E*311 = 184 GPa and Poisson's ratio *ν*311 = 0.294 agree with values reported by Rangaswamy et al. [23], as well as with results from measurements and simulations of the DEC values of LPBF 316L reported by Chen et al. [24].

#### *2.4. Neutron Diffraction*

Stress determination by neutron diffraction (ND) was carried out at the STRESS-SPEC diffraction instrument [25] at the neutron facility FRM II in Munich, Germany (Figure 2a).

**Figure 2.** Beamline setup and measurement positions.

A bent Si400 single crystal monochromator was used to select the wavelength of 1.550 Å. The Fe311-peak was selected in order to characterise the RS distribution due to the low accumulation

*Metals* **2020**, *10*,

of inter-granular stresses reported for this reflection in conventional face-cubic-centred (fcc) iron materials [10].

To map the RS in the cross sectional plane at the mid build height of the specimen, a gauge volume of 2 mm × 2 mm × 2 mm was used in a grid of 7 × 5 measurement points. (see grey cubes in Figure 2b). The coordinate system is depicted in Figure 2b. *σL* represents the RS along the *y*-direction, *σT* along the *x*-direction and *σN* along the *z*-direction, which was also the build direction.

A stress-free reference was needed to calculate strains from the measured *d*311 lattice-spacing. A small cube with the size of 3 mm × 3 mm × 3 mm was sectioned from the bottom corner of a separate test build job of the PtC specimen (- ). This cube was regarded as free of Type I macro-stresses [9], due to the mechanical relaxation during sectioning. The strain can subsequently be derived from the measured *ϑ* angles using Bragg's law [9].

$$\varepsilon = \frac{d^{311} - d\_0}{d\_0} = \frac{\sin \theta^{311}}{\sin \theta\_0} - 1 \tag{1}$$

Assuming that the principal geometric directions correspond with the principal stress directions Hooke's law reads as the following, as described by Holden et al. [26]:

$$
\sigma\_{L,T,N} = \frac{E^{311}}{\left(1 + \nu^{311}\right)\left(1 - 2\nu^{311}\right)} \left[ \left(1 - \nu^{311}\right)\varepsilon\_{L,T,N} + \nu^{311} \left(\varepsilon\_{T,N,L} + \varepsilon\_{N,L,T}\right) \right] \tag{2}
$$

The *d*0 value was derived from the average of the *ϑ*311 measurements of the cube in longitudinal (*L*), transversal (*T*) and normal (*N*) direction, where the normal direction corresponds to the build direction (as depicted in Figure 2b). The same DECs that were derived from the Eshelby–Kroener model were applied to both Lab XRD and ND results (see Section 2.3).

#### *2.5. Micro Computed Tomography*

The small reference cube for ND was studied using Micro Computed Tomography (μCT) to obtain a detailed dataset of the internal defect structure. The μCT measurements were performed at a GE v|tome| × 180/300 CT scanner (GE Sensing & Inspection Technologies GmbH, Wunstorf, Germany) using the 180 kV source at a voltage of 150 kV and a current of 40 μA without any metal pre-filter. A voxel size of (3 μm)3 was achieved. The analysis of the data was performed using the commercial software VG Studio MAX version 3.2.1 (Volume Graphics GmbH, Heidelberg, Germany). A lower threshold limit of 8 voxels was used for pore detection.

#### *2.6. Optical Microscopy*

For Optical Microscopy (OM) investigations of the microstructure, the bottom faces of the samples were ground, polished, and etched. Emery papers with 180, 320, 600 and 1200 grits followed by clothes with 3 μm and 1 μm were used. For etching the Bloech & Wedl II method [27] (a solution of 50 mL H2O, 50 mL HCl, 0.6 g K2S2O5) was applied. The microstructure was captured using a Olympus BX53M microscope with a DP74 camera module (Olympus Corporation, Tokyo, Japan). The analysis was performed using the software Olympus Stream Essentials (Olympus Corporation, Tokyo, Japan).

#### **3. Results**

#### *3.1. In-Situ Thermography*

The build jobs of both samples were supervised in-situ by thermography in order to receive more detailed information on the local variation of the temperature gradient and cooling rates. Figure 3a,b display the maximum apparent temperatures at the mid-height layer.

The thermography data were averaged over 40 layers at the mid build-height to reduce noise and the influence of smoulder and spatter. In relation to the height of the specimens of 21 mm, these 40 layers represent an average of the height range from 9.35 mm to 11.35 mm. The number of 40 layers (=2 mm) was chosen, since this corresponds to the size of the used gauge volume for ND (2 mm × 2 mm × 2 mm).

The scan strategy of the PtC specimen (- ) resulted in an increased heat accumulation in the center, compared to the CtP specimen (- ). This could be observed as an increase in the maximum temperature at the centre, when comparing Figure 3d,c.

Figure 3d shows the different apparent temperature values for the four sections of the plane. These differences were caused by a combination of the different surface roughness of each section and of the shadowing effects from the smoulder.

**Figure 3.** Apparent (uncalibrated) maximum temperature of the two specimens obtained from thermography data acquired during the build process at the sample's mid build height. Each border fill scan started and ended at the bottom right hand laser turn position of the images. (**a**,**b**) display the mid-height layer, (**c**,**d**) display an average of 40 layers at the mid-height.

Figure 4 shows the cooling rate of both specimens. The cooling rate—dT/dt was obtained by comparing two images at *t* = 1 ms and *t* = 2 ms after an apparent temperature of 700 K was reached for the last time at the surface. Typical times for cooling from the maximum temperature to 700 K were between 2 ms and up to 15 ms. The cooling time of 15 ms was observed at the centre of the PtC specimen (-).

As displayed in Figure 4d, the PtC specimen (- ) also featured a lower cooling rate in the centre of the plane in addition to the higher maximum temperature depicted in Figure 3d. The CtP specimen (- ) shows a low cooling rate at the edges of the specimen indicating that the surrounding metal powder served as an heat insulator for conduction, as assumed in modelling [28].

**Figure 4.** Cooling rates measured by the different between images taken at *t* = 1 ms and *t* = 2 ms after an apparent temperature of 700 K was observed for the last time. Thermography data for both samples was obtained during the build process at the middle of the total build height. (**a**,**b**) display the mid-height layer, (**c**,**d**) display an average of 40 layers around the middle of the total build height.

#### *3.2. Combined Neutron and Lab X-ray Diffraction*

The ND results (bulk RS) were combined with Lab XRD results (surface RS) to show the complete stress distribution across the full cross section of the specimens' middle plane. The contour plot function of the commercial software Origin 2018 (OriginLab Corporation, Northampton, USA), which is based on the Delaunay triangulation, was used to visualise the combined results. Figure 5 shows the combination of ND and XRD results. The RS distribution is visualised in longitudinal, transversal, and normal direction, where the normal direction corresponds to the build direction (Figure 2b). The XRD technique used only allows for the determination of stress components which are parallel to the surface (i.e., in-plane). Therefore, to visualise the third orthogonal stress component (perpendicular to the surface) in Figure 5 the following boundary condition was used: at the surfaces of the specimens corresponding to positions at *y* = 0 mm and *y* = 36 mm, the value of the longitudinal stress component was assumed to *σL* = 0 MPa. For the value of the transversal stress component at *x* = 0 mm and *x* = 24 mm it was assumed *σT* = 0 MPa, since these are free surfaces in these corresponding stress directions. In general, for the two scan strategies, a similar RS distribution was observed in the longitudinal, transversal and normal direction. This distribution is characterised by compressive RS within the bulk, balanced by tensile RS at the surface. Instead, the RS distribution in the longitudinal and transversal direction of each specimen is similar in terms of shape and magnitude, the RS distribution in the normal (i.e., building) direction differed from the other directions in terms of shape and magnitude.

**Figure 5.** Comparison of RS maps of the tw scan strategies including results from lab X-ray diffraction at the surface (The big semi-translucent squares in the bulk represent the almost cubic ND gauge volume (orientated differently for different stress components), whereas the small semi-translucent squares at the edges represent the lab X-ray measurement positions).

The highest magnitude of RS of each specimen (maximum tensile or maximum compressive) were observed in the normal direction. However, the PtC specimen (- ) displayed higher bulk compressive stresses in all three orthogonal directions. The different stress values in the normal direction, as compared to the longitudinal and transversal direction, are also reflected in the lattice spacing of the the *d*0 cube. Whereas, the *d*-spacing in the longitudinal and transversal direction are relatively similar to each other, in the normal direction a larger lattice spacing was measured (see

Table 1). The *d*0 value was obtained from averaging the three measured directions. This averaged value was used as a stress-free lattice parameter to calculate the RS.

**Table 1.** Distribution of orthogonal *d*-spacing values of the reference cube.

#### *3.3. Micro Computed Tomography*

The reference cube was analysed by μCT to provide an example of local defect distribution in the specimen. Because it was cut from the corner where the laser path started and ended, it represented the area where the highest amount of defects was expected. The μCT results presented in Figure 6 reveal a network of defects at the location where the laser started and ended, as well as between the hatches. Because the same scan vector was used on each hatch and layer, the projection of defects onto one plane (Figure 6c) reveals the lack of fusion between neighbouring hatches. As reported in literature [29,30] alternating the orientation of scan vectors between layers prevents the formation of lack of fusion defects. Because the effect of shrinkage on RS was the subject of this study, scan vectors were not altered between layers to magnify such effect.

The largest defects were situated close to the edge of the sample (Figure 6b,c). A total porosity of 0.28% was observed.

**Figure 6.** μCT reconstructions of the ND reference cube sectioned from a twin PtC specimen (-).

#### *3.4. Optical Microscopy*

Optical microscopy images of the polished and etched specimen's bottom surface are shown in Figure 7. Defects at the turn location of the laser are visible (Figure 7d,e). The bottom surface of both specimens are expected to be less effected by heat accumulation due to the smaller build height at the time the microstructure *froze*, and a better heat conduction into the build plate as compared to the top layers of the specimens. Nonetheless, defects (pores and voids) were observed in both specimens at the positions, where the laser turns by 90°.

(**a**) -full sized image of CtP specimen (**b**) -

(**c**) - magnified upper right section of CtP specimen

(**e**) - magnified section of the centre of CtP specimen

full sized image of PtC specimen

(**d**) - magnified upper right section section of PtC specimen

(**f**) - magnified section of the centre of PtC specimen

**Figure 7.** Optical microscopy images of the samples' bottom surface after polishing and etching.

#### **4. Discussion**

As mentioned in the introduction, Mercelis and Kruth [3] described two major driving mechanisms for the formation of RS in AM metallic parts made by LPBF: the temperature gradient mechanism (TGM) and the solidification shrinkage mechanism (SSM). Both of the mechanisms induced compressive stresses into the bulk material. The combined data of ND and Lab XRD showed the typical stress distribution pattern for metallic AM parts: compressive stresses in the bulk and tensile stresses close to the surface.

Because of the same geometric dimensions, it can be assumed that the solidification shrinkage effects in the two samples were equal. This implies that any differences in the RS distributions of the two samples should be attributed to the TGM.

According to the SSM, the induced RS distribution by solidification shrinkage depends on the length of the scanned laser tracks. The shrinkage of longer laser tracks introduces higher RS than shorter laser tracks [31]. Therefore, the longer laser tracks parallel to the *y*-axes of the specimens presented in Figure 5a–d) appear to be the main reason for the higher compressive stresses in the longitudinal stress distributions of the two specimens, as compared to the shorter laser tracks parallel to the *x*-axes, which seemed to result in lower compressive stresses for the transversal RS distributions of the two specimens.

The distribution of the normal component of the RS was assumed to be mostly independent of the SSM, since there were almost no restrictions to solidification shrinkage due to the free top surface during the layer-wise production. Therefore, the TGM was assumed to be the main mechanism to shape the RS distribution of the normal component. The RS distribution has the shape of a *butterfly* (see Figure 5e,f). The spikes of this *butterfly* pattern match the location where the laser turned by 90° (see Figure 8).

**Figure 8.** Normal (i.e., build direction) stress component (contour lines) in MPa (Figure 5e,f) overlayed with thermography results (Figure 3c,d) to highlight the compressive stresses at the laser's turn locations.

It is also observed that the normal component exhibits the highest tensile stresses (see Figure 5 and Table 2). The determined RS values are below the tensile yield strength ranges for LPBF 316L (450 MPa to 590 MPa) as reported by Wang et al. [32].

**Table 2.** Max. and min. values of the orthogonal stress components in the plane.

Figure 8b shows that higher compressive stresses in the PtC specimen (- ) are localised close to the zone of the highest heat accumulation. This is in contrast to the CtP specimen (- ), which shows lower and more evenly distributed compressive stresses in the plane (see Figure 5). The higher maximum temperatures at the centre of the PtC specimen (- ) in combination with the slower cooling rate seem to result in larger RS in the centre of the plane as compared to the CtP specimen (- ). Lower contributions from the TGM to the compressive stress profile of the CtP specimen (- ) might be a result of the lack of heat accumulation in the centre of its plane (see Figures 3c and 8a) and a faster cooling rate (see Figure 4c).

Line stress profiles (Figure 9) that were derived from the in-plane data presented in Figure 5 at a middle line of the plane at *x* = 12 mm emphasise the conclusion of SSM being the major mechanism to define the *shape* of the RS distribution. However, the TGM appears to influence the *magnitude* of the peak compressive stresses (see Figure 9).

It should be noted that, in Figure 9, only measured values are displayed. Therefore, the boundary condition values of 0 MPa were excluded from the longitudinal stress profile at *y* = 0 and *y* = 36 (see Figure 9a). The measured bulk values close to these surfaces support the assumption of zero stress at these surfaces.

(**a**) Longitudinal RS component at *x* = 12 mm

(**b**) Transversal RS component at *x* = 12 mm

(**c**) Normal RS component (=build direction) at *x* = 12 mm

**Figure 9.** Line scans in both samples for all three orthogonal directions using the combined data from ND and lab XRD.

Since the TGM does not require melting [33], it can be assumed that the compressive stress components induced by the TGM were formed in subjacent layers within the heat affected zone of the melt pool but below the layers that were remolten. It should be noted that a linear interpolation was applied to combine the surface measurement positions and the neutron data in Figure 9. Therefore, any possible sub-surface tensile peaks that had been reported by Mishurova et al. [11,34] in LPBF Ti-6Al-4V were not be taken into account in this study.

The μCT results (Figure 6) reveal a network of defects at the laser's start and stop position. OM results (Figure 7d,e) show defects at the turn locations of the scan track. These regions have been measured to be under higher tensile RS (in particular, *σT*), which is assumed to arise from the

#### *Metals* **2020**, *10*,

longitudinal and transversal solidification shrinkage, as depicted in Figure 5. Therefore, any defects in these regions might have served as micro-crack initiators. The combination of pores and tensile stress at corners could also explain why the defects observed by μCT (Figure 6b) seem to be larger towards the outer edge: the tensile stress due to solidification shrinkage increases towards the edges of the specimens due to an increase length of the laser tracks. The analysis of the optical microscopy and μCT data indicates that an unknown amount of RS might had been relaxed by micro-cracks at the laser's start and stop location.

#### **5. Conclusions**

Two prismatic AISI316L specimens using a border fill scan strategy were produced in order to differentiate the effect of the temperature gradient mechanism from the solidification shrinkage mechanism in AM metallic parts produced by LPBF. The following conclusions could be made:

#### **6. Outlook**

In-situ Thermography recorded a detailed data set of apparent surface temperatures and enables an analysis of cooling rates. These data can be used in future simulations in order to model the RS field at the specimens' mid height. In future studies, the RS distribution in the subsurface region might be resolved using additional techniques, such as hole drilling, slitting, X-ray with layer removal, deep hole drilling, or contour method.

Additionally, twin specimens of the two prisms will allow further experiments to study the effect on maximum values of the RS distribution, if the reference cube is cut from different positions within these twin specimens. In addition, these twins will allow a systematic μCT analysis of different positions within the two specimens, in particular to relate RS and defect distributions.

**Author Contributions:** Conceptualization, A.U., T.F., G.M. and S.J.A.; methodology, A.U., G.M., M.S., K.S., S.J.A., T.M., I.S.-M., T.F., M.H.; formal analysis, A.U., S.J.A., M.S., K.S.; investigation, A.U., M.S., K.S., S.J.A., T.M., I.S.-M., T.F., M.H.; writing—original draft preparation, A.U., G.M., M.S., K.S., S.J.A.; writing—review and editing, A.U., G.M., M.S., S.J.A., T.M., I.S.-M., M.H., A.E., G.B., T.F.; visualization, A.U., S.J.A.; supervision, A.E., G.B.; project administration, A.E., G.B.; funding acquisition, G.B. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research received no external funding.

**Acknowledgments:** This work has been funded by the BAM Focus Area Materials project AGIL "Microstructure Development in Additively Manufactured Metallic Components: from Powder to Mechanical Failure" and ProMoAM "Process monitoring of Additive Manufacturing". We are thankful for the financial support and the fruitful cooperation with all partners. This work is based upon experiments performed at the STRESS-SPEC instrument operated by FRM II at the Heinz Maier-Leibnitz Zentrum (MLZ), Garching, Germany.

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

*Metals* **2020**, *10*,

#### **Abbreviations**

The following abbreviations are used in this manuscript:

#### **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 Stress and Microstructure of a Ti-6Al-4V Wire Arc Additive Manufacturing Hybrid Demonstrator**

**Tatiana Mishurova 1,\*, Benjamin Sydow 2, Tobias Thiede 1, Irina Sizova 2, Alexander Ulbricht 1, Markus Bambach 2,3 and Giovanni Bruno 1,4**

Received: 9 April 2020; Accepted: 23 May 2020; Published: 26 May 2020

**Abstract:** Wire Arc Additive Manufacturing (WAAM) features high deposition rates and, thus, allows production of large components that are relevant for aerospace applications. However, a lot of aerospace parts are currently produced by forging or machining alone to ensure fast production and to obtain good mechanical properties; the use of these conventional process routes causes high tooling and material costs. A hybrid approach (a combination of forging and WAAM) allows making production more efficient. In this fashion, further structural or functional features can be built in any direction without using additional tools for every part. By using a combination of forging basic geometries with one tool set and adding the functional features by means of WAAM, the tool costs and material waste can be reduced compared to either completely forged or machined parts. One of the factors influencing the structural integrity of additively manufactured parts are (high) residual stresses, generated during the build process. In this study, the triaxial residual stress profiles in a hybrid WAAM part are reported, as determined by neutron diffraction. The analysis is complemented by microstructural investigations, showing a gradient of microstructure (shape and size of grains) along the part height. The highest residual stresses were found in the transition zone (between WAAM and forged part). The total stress range showed to be lower than expected for WAAM components. This could be explained by the thermal history of the component.

**Keywords:** residual stress; WAAM; Ti-6Al-4V; additive manufacturing; neutron diffraction; hybrid manufacturing

#### **1. Introduction**

The titanium alloy Ti-6Al-4V is intensively used in lightweight applications for the aviation and space industry because of its high specific strength [1]. Hot forging is the usual manufacturing route; this allows better material formation and control of the microstructure. Subsequent machining is used to attain the desired tolerances regarding dimensions, shape, and surface condition [2]. One positive effect of hot forging is that the microstructure and the fiber flow can be optimized, which results in excellent mechanical properties. However, forging is limited as far as the achievable geometries is concerned: as an example, due to friction and heat transfer, the material cannot flow into cavities of the

forging die with a large depth-to-diameter ratio. Moreover, multistage hot forging operations require expensive heated die sets and lead to parts with large tolerances; this requires further machining to create the final shape. In fact, the "buy-to-fly" ratio (the mass ratio between the originally purchased stock material used to produce a part and the mass of the final finished part) ranges from 12:1 to 25:1 for aircraft titanium components made by traditional manufacturing techniques [3]. This means that 12–25 kg of raw material are required to produce just 1 kg of parts. In this way, more than 90% of the material is machined away. Such a high buy-to-fly ratio is unacceptable from many points of view. For a typical titanium component produced by additive manufacturing (AM) techniques the buy-to-fly ratio drops to 3–12:1 [4] and can even be close to 1:1 [5]. By using AM for production of aerospace components not only costs for expensive titanium alloys can be reduced, but also the machining time can be drastically cut.

AM is an innovative manufacturing technology that allows manufacturing almost arbitrarily complex shapes. In powder-bed metal AM processes, the powder is distributed layer-wise on a substrate or already deposited material and fused by melting the material using an electron (electron beam melting–EB-PBF) or laser beam (laser powder bed fusion–L-PBF). Because of the small (focused) beam, the deposition rates are quite low. PBF AM processes have in general a small layer thickness of a few tens of μm up to 2 mm [6]. This is necessary due to the local bonding of adjacent powder particles and the limited beam penetration depth. Although beam scanning systems can achieve high laser or electron beam scanning speeds, large parts require a high number of layers, which causes high production times. The part size is limited due to the small process chamber (typically 200–300 mm edge size). Also, as the powder-bed has to be established before the AM process begins, no existing parts can be extended by means of PBF. This makes this technology only attractive for small parts. An increasingly popular AM technology that may allow overcoming this issue is Wire Arc Additive Manufacturing (WAAM). WAAM utilizes common and well-known welding technologies. A wire is fed to the desired position and molten by an electric arc. The melt pool solidifies and forms a weld bead. One advantage of WAAM compared to other AM technologies is the high deposition rate. This is especially important for large scale products such as landing gear parts or wing ribs, which are difficult to forge due to their size or have a large material waste when machined. WAAM achieves deposition rates as high as several kg/h [7] Also, a wide range of materials is available as wires, which offers multiple possible applications. The microstructure and mechanical performance of WAAM manufactured Ti-6Al-4V parts will not reach the level of forged material without additional (often expensive) post-processing techniques such as hot isostatic pressing (HIP) or thermomechanical processing [8]. One strategy is to combine the benefits of the two technologies (forging and WAAM) to produce so-called hybrid parts with the desired mechanical properties (even at local scale). This approach is of particular interest for the production of the components, where the forged sections are heavily loaded during service and the added AM features should be capable to withstand only moderate loads. Moreover, in aerospace industry many parts possess a symmetry axis. To produce such components by means of conventional processing a number of different forging die sets should be used. Instead, the hybrid approach can be more efficient, since the required structural or functional features can be added to the forged part without using additional tools for every part design. In so doing, both the cost of forging tools and material waste can be reduced. Figure 1 represents a possible hybrid manufacturing route, where the basic structure can be produced by hot forming and the missed structural features can be added using AM.

Basically, two process routes are possible: (i) AM and subsequent forming or (ii) forming and subsequent AM. The first hybrid processing route was extensively studied by authors of the present work in regard to Ti-6Al-4V [9–11]. It was observed that Ti-6Al-4V pre-forms for forging made either by L-PBF or by WAAM show a good hot workability. Moreover, voids could be closed applying compressive stresses at elevated temperature, so that the final microstructure and mechanical properties could be improved. Extensive work on the second hybrid processing route can also be found in the literature.

**Figure 1.** Illustration of the possible hybrid manufacturing route combining metal hot forming and WAAM.

Bambach et al. [11] examined the same demonstrator shape and process route (forging with subsequent WAAM) used in this study and proved that its mechanical properties exceeded the minimum requirements for forged parts regarding yield strength (YS = 837 MPa > 830 MPa) and ultimate tensile strength (UTS = 934 MPa > 900 MPa). Also, only minor anisotropic behavior was observed. In another work, Bambach et al. [12] showed that laser cladding technology could be used to produce flexibly applicable local patches to locally increase the stiffness or the thickness of sheet metal components. Papke et al. [13] analyzed tensile bonding strength of hybrid parts made of Ti-6Al-4V produced by combination of laser beam melting and warm bending. It was reported, that sheet material and AM material possess different hardness values, and the sheet thickness strongly influences the bonding strength. Moreover, it was confirmed, that the contact zone between sheet and AM is the most significant for the strength of hybrid components [14]. Hirtler et al. [15] investigated the production of modifying stiffening ribs made of AlSi12 on conventionally produced pre-forms (EN-AW 6082). The authors proved the feasibility of the process combination between WAAM and hot forming. In general, most of the existing work on hybrid manufactured of Ti-6Al-4V concentrates on the increase of the strength of the sheet metal components using AM.

To produce a sound Ti-6Al-4V hybrid part the microstructure should be thoroughly controlled. During WAAM of Ti-6Al-4V β-grains (BCC) grow epitaxially, similar to PBF AM techniques [16]. Such microstructures are hard to avoid because at the low concentration Al and V have a high solubility in the alloy and, thus, do not partition ahead of the solidification front [17]. The length of prior β grains along the solidification direction can reach a few cm or even cover the whole height of the sample, possessing a strong <001> fiber texture [18]. After cooling below the β-transus temperature (approximately 995 ± 25 ◦C) β grains typically transform to fine α laths (HCP) retaining β lamellar structure (Widmanstätten structure). Here the thickness of α-platelets and the sizes of β-grains are the fundamental parameters that affect the mechanical performance of Ti-6Al-4V alloy. Very high cooling rates (>410 ◦C/s) promote martensitic transformation. Slower cooling rates (<20 ◦C/s) promote formation of Widmanstätten microstructures [19]. Furthermore, the cooling rate determines the thickness and the presence of α-phase on prior β grain boundaries, which strongly affect the mechanical performance, because they induce anisotropy [20]; such anisotropy is unwanted in many applications. In general, the Widmanstätten-type microstructure is characterized by relatively low tensile ductility, good creep resistance, moderate fatigue properties and crack growth resistance [21].

Together with the microstructure, residual stress (RS) is one of the major obstacles in the development of the AM techniques [22]. It may lead to cracking and geometrical distortion of the parts, thereby limiting the design freedom of the component [23,24]. In the case of WAAM, RS is also an important factor influencing the structural integrity [20]. Colegrove et al. have reported that in as-deposited state the RS values in steel WAAM part can reach the yield stress of the material [25]. Also, the presence of RS highly affects the mechanical behavior of the part. Zhang et al. have reported [26] that RS facilitates crack propagation from a WAAM Ti-6Al-4V part to its substrate. One of the strategies for the reduction of RS and the refinement of the microstructure in WAAM is high pressure rolling after every deposited layer (or a few of them) [25,27]. This approach, however, imposes some limitations on the part design and degrades the main advantage of the WAAM technique: the high deposition

rates. In general, according to ASTM F2924 [28], stress-relief heat treatment is mandatory for Ti-6Al-4V AM parts. The connection between microstructure and RS for welded and AM materials has been reported in several studies [29–32]. Due to high cooling rate needed for martensitic transformation (β→α') in Ti-6Al-4V, the martensitic microstructure is often linked with high tensile RS [30]. It has been shown that the presence of β phase in PBF Ti-6Al-4V, introduced by intrinsic heat treatment, leads to the RS relaxation [31]. Thus, for the general application of WAAM, it is important to understand the microstructure and the RS in as-manufactured WAAM components. Such properties depend also on the location in the hybrid part. An investigation of the microstructure and the RS at the interface between forged and AM features is here of particular importance. In AM parts, high tensile RS is usually reached at the interface between the substrate and the deposited material [33]. This could be critical for the utilization of hybrid parts.

The main aim of the current study is to investigate the microstructure and the RS distribution in a hybrid part produced by the combination of conventional hot forming and WAAM. A T-section geometry manufactured by means of WAAM on a hot forging pre-form is investigated. The paper is structured as follows: Section 2 presents the manufacturing of hybrid Ti-6Al-4V parts using WAAM and hot forging. Also, the procedures of microstructural analysis and of determination of RS by means of neutron diffraction are presented. Section 3 gives an overview of results of metallographic examinations and RS. Finally, the results are discussed, and conclusions are presented.

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

#### *2.1. Sample Manufacturing*

The considered hybrid process route encompassed forging and subsequent WAAM of Ti-6Al-4V. A T-shaped pre-form was forged, as shown in Figure 2. It contained a 10 mm wide and ~90 mm long rib, which was milled flat at a total height of 42 mm. The WAAM process was used to increase the height of the rib by 66 mm.

Hot forging was performed in the α + β-temperature range using a 2500 t crank press at OTTO FUCHS KG (Meinerzhagen, Germany). The forged part was machined on the top side to obtain a flat surface to ensure a stable WAAM process.

**Figure 2.** Drawing of the final pre-form with dimensions. Front view (**a**), side view (**b**), final appearance of the hybrid near-net shape demonstrator (**c**).

WAAM was performed using a Fanuc six-axis robot (FANUC Europe Corporation S. A., Echternach, Luxembourg) and a TPS 500i welding power source (Fronius®, Wels, Austria) utilizing the Cold Metal Transfer (CMT) variant of the Gas Metal Arc Welding (GMAW) process. The welding was carried out in an argon filled, sealed chamber to avoid oxidation. The chamber was equipped with an O2 sensor. Pre-heating of the sample was not applied. The used WAAM set-up is displayed in Figure 3a. A Ti-6Al-4V wire with a diameter of 1.0 mm was used. One hybrid demonstrator was produced for

investigations. The shielding gas flow rate at the torch was 15 L/min, also the build chamber was flooded with 50 L/min before the deposition process. During deposition the chamber gas flow rate was reduced to 20 L/min to compensate the small gas leakage. The argon gas purity was ≥99.99%. The chemical composition for as-built Ti-6Al-4V WAAM is shown in Table 1.

**Figure 3.** For the WAAM process (**a**) used WAAM machine, (**b**) applied sinus wave movement and (**c**) swapping of the starting point each layer.

**Table 1.** Chemical composition of Ti-6Al-4V as-built by WAAM (max. weight %).

The deposition pattern is shown in Figure 3b,c. The starting point (A/B) changed for every layer using a bidirectional tool path on a single bead. A sinus wave was superposed onto the translational motion of the torch to obtain the desired wall thickness. The layer thickness was approximately 4.4 mm. The number of deposited layers was 15.

The welding parameters are shown in Table 2. The first three layers were welded with a higher electrical current to achieve a good fusion with the forged substrate and to increase the average temperature (this ensures a stable WAAM process). The current was than reduced to 100 A for the following layers.

**Table 2.** WAAM manufacturing parameters.

#### *2.2. Microstructural Characterization*

In order to investigate the microstructure, the manufactured hybrid part (including the forged region) was cut along the build direction. The sample was mounted with the cut cross-sections on top and ground flat with successively finer grades (from 320 to 1200 μm) of silicon carbide (SiC) papers. The specimen was then polished with 0.05 μm silica solution (OP-S Suspension, Struers GmbH, Willich, Germany) with the addition of H2O2, HNO3, and HF. Sample was etched with Kroll's agent solution to reveal the microstructure. The sections were analyzed using an optical microscope Carl Zeiss Axiotech by Carl Zeiss Microscopy (Jena, Germany).

#### *2.3. Residual Stress Analysis*

The neutron diffraction experiment was conducted on the instrument E3 at BER II reactor (Helmholtz Zentrum Berlin, Germany) [34]. A monochromatic neutron beam of wavelength λ = 1.476 Å was used. The three orthogonal strain components, assumed to be principal directions based on the sample geometry, were measured (longitudinal, transversal, normal, see Figure 4). The longitudinal and normal directions coincided with the deposition and the build direction of WAAM, respectively. The acquisition time was set to 75 min for each stress component. Additionally, to improve the diffraction signal, a constant ω oscillation of ±10◦ around the scattering vector (this is the bisectrix between the incident and the diffracted beam directions and corresponds to the direction in which the strain is measured) was performed. For the longitudinal stress component, a primary slit with opening of 4 mm × 4 mm was used. For the normal and transversal component, a primary slit with vertical opening of 18 mm and horizontal opening of 2 mm was used. This allowed at the same time increasing the signal and keeping high spatial resolution along the build direction of the sample. A secondary collimator with a focus of 2 mm was used for all measurements. The strain was measured along two lines: along the deposition direction at the middle height of the WAAM part (L-line) and along the build direction (N-line). The measured points and coordinate system are schematically presented in Figure 4a.

**Figure 4.** Photo of sample during neutron diffraction experiment on E3 with the schematics of (**a**) the measured points and the coordinate system, (**b**) the set-up. (Note that the sample is aligned for the measurement on the normal strain component).

The measurements were performed at the diffraction angle 2θ = 76◦, which allowed the simultaneous detections of three diffraction peaks for α-Ti: 1122, 2021 and 0004. However, only the 1122-α reflection appeared at every measured point and sample direction, therefore the RS was calculated for this crystallographic family. The lattice spacing *d* was calculated according to Bragg's law:

$$d\_{11\overline{2}} = \frac{\lambda}{2 \cdot \sin \theta\_{11\overline{2}}} \tag{1}$$

The lattice spacing (averaged within the gauge volume) was determined by fitting the diffraction peak with a Gaussian function using the software StressTexCalculator 1.53 (TU-Clausthal, Germany). In order to calculate lattice strains ε a strain-free reference *d*0 has to be used, since:

$$
\varepsilon = \frac{d\_{11\overline{2}2} - d\_0}{d\_0} = \frac{\sin \theta\_0}{\sin \theta\_{11\overline{2}2}} - 1 \tag{2}
$$

One of the main problems in diffraction-based RS analysis is the estimation of such a reference [35,36]. There are different approaches for the determination of *d*0: measurements using a reference powder or stress-relieved coupons; or calculations using a global average or stress balance/boundary conditions. In the case of AM materials the variation of the microstructure inside the part and the different thermal history of the part from the raw material (power, wire) makes the determination of *d*0 even more challenging [35]. In our case, a global *d*0 (or θ0) was taken as the average of all measured points, as proposed in [37]. This value was 2θ0 = 72.856◦ with an average error of ±0.007◦ (coming from the fit of all diffraction peaks). As an alternative, the stress balance condition was applied to the L-line scan for the normal component (Figure 4): σ*NormdL* = 0. For this estimation the stress profile was extrapolated to the surface and interpolated between the points with a B-spline function. Symmetry of the stress profile was assumed, i.e., σ*N*(+*L*) = σ*N*(−*L*). Using this approach, we obtained 2θ0 = 72.859◦. This value lies within the error range of the 2θ0 obtained using the global average approach and shifts the resulting stress values only by +10 MPa. Therefore, only stress profiles calculated satisfying the above-mentioned stress balance conditions will be reported.

RS were calculated according to the tensorial Hooke's Law: σ = *C*ε, with *C* as the stiffness tensor. With the assumption that the principal geometric directions are also principal stress directions, Hooke's Law in the case of a quasi-isotropic solid reads:

$$
\sigma\_{L,T,N} = \frac{E^{11\tilde{2}2}}{\left(1 + \nu^{11\tilde{2}2}\right)\left(1 - 2\nu^{11\tilde{2}2}\right)} \left[ \left(1 - \nu^{11\tilde{2}2}\right)\varepsilon\_{L,T,N} + \nu^{11\tilde{2}2} \left(\varepsilon\_{T,N,L} + \varepsilon\_{N,L,T}\right) \right] \tag{3}
$$

where σ*L*,*T*,*N* and ε*L*,*T*,*N* are stresses and strains along the longitudinal, transversal and normal direction, and *E*1122 = 112.7 GPa and ν1122= 0.321 are the diffraction elastic constants for α-Ti 1122 reflection calculated by the Kröner's model [38].

Hydrostatic stress σ*H* and von Mises stresses σ*vM* were calculated according to:

$$
\sigma\_H = \frac{\sigma\_T + \sigma\_L + \sigma\_N}{3} \tag{4}
$$

$$
\sigma\_{\rm \varepsilon M} = \sqrt{\frac{1}{2} \left[ (\sigma\_L - \sigma\_T)^2 + (\sigma\_T - \sigma\_N)^2 + (\sigma\_N - \sigma\_L)^2 \right]} \tag{5}
$$

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

Figures 5 and 6 show the macro- and microstructure of the hybrid part, respectively. To interpret the microstructural features, we must recall that the first layers were produced with higher current (see Table 2) to compensate the heat sink effect of the substrate. Then, the heat input (i.e., current) was gradually decreased, and after the fourth layer it was set to a constant value. The macrostructure of the WAAM Ti-6Al-4V is characterized by the epitaxial growth of large columnar prior β-grains, which stretch through several deposited layers (Figure 5). The average dimensions of such grains are about 1.3 mm perpendicular and 8 mm parallel to the build direction.

**Figure 5.** (**a**) Micrograph showing prior β-Ti grains in the middle section of the sample with enlarged view for (**b**) top of WAAM part, (**c**,**d**) middle height of WAAM part, (**e**) transition region between WAAM and forged parts.

**Figure 6.** Typical micrographs obtained (**a**) in the substrate, (**b**) in the transition zone close to the substrate, (**c**) in the center of the WAAM part.

The forged Ti-6Al-4V T-section shows, instead, a bi-modal α+β microstructure (Figure 6a), typically obtained in Ti-6Al-4V after conventional thermomechanical processing [21]. The first layers of the WAAM material are characterized by a higher cooling rate compared to the following layers, due to heat conduction to the substrate. Therefore, in the transition zone (e.g., in the first layers, up to approximately 7 mm distance from the seam line), a fine martensitic α' microstructure is found within small globular prior β-grains with average dimensions of approximately 50 μm (Figure 6b).

At sample mid-height, the micrographs show an αw-Widmanstätten microstructure with a thickness of the α-platelets of approximately 2 μm within the columnar prior β-grains (Figure 6c). Such microstructures are often reported in the literature [3]. Moreover, the α-phase also appears as a thin layer on the prior β-grain boundaries with a thickness comparable to the thickness of the α-platelets. This microstructure was generated by the repeated rapid heating and cooling that occur during the WAAM process. It is important to stress that the coarse prior β grains may critically affect the mechanical properties (since β is more ductile than α).

A flat panel detector was used in the neutron diffraction experiment. The 2D distributions of the diffraction signal give valuable information about the microstructure (Figure 7): especially in the longitudinal direction, diffraction spots (instead of the typical Debye rings observed for polycrystalline materials) were observed. This is called the "coarse grain effect" (Figure 7a). The grain size of α-Ti laths is quite small, with an average thickness of about 2 μm and an average length of about 25 μm (see Figure 6). However, the origin of α-Ti laths create a distinct crystallographic texture. All the grains of α-Ti inside the prior β-grain are only variants of the same crystallographic orientation, according to Burgers orientation relation [39]. Thus, from the crystallographic point of view, such laths form coarse grains, which scatter similarly to single crystals in the longitudinal direction. In spite of this effect, the projections of the few spots inside a 2D detector image could be analyzed as a Bragg peak, and the lattice parameter could be reliably determined.

**Figure 7.** Image of 2D detector obtained during neutron diffraction experiment on the WAAM part (at 46 mm) for (**a**) longitudinal direction, (**b**) normal direction, (**c**) transversal direction. Note: the images are normalized to the maximum intensity, the pixel size is 1.17 mm × 1.17 mm.

The neutron diffraction patterns (i.e., integrated along the Debye ring) collected along the L-line for three orthogonal directions are presented in Figure 8. The patterns look similar at every position, thereby showing no microstructural variations along the deposition direction. The peak intensities (of both 1122 and 2021 reflections) do not vary for different orientations, highlighting the absence of a strong crystallographic texture.

**Figure 8.** Diffraction patterns along the L-line for (**a**) transversal component, (**b**) normal component, (**c**) longitudinal component.

The intensity ratios follow the theoretical predictions for the neutron diffraction pattern of HCP-Ti powder (i.e., assuming random crystallographic texture, calculated by PowderCell [40]): the intensity of the 1122 reflection is the highest out of the three reflections observed, followed by 2021 and 0004 reflection.

In contrast to the diffraction patterns acquired along the deposition direction (L-line, Figure 8), a decrease of the intensity towards the substrate for all reflections of the transversal strain component can be observed along the build direction (Figure 9a). As shown in Figure 2, the thickness of the sample is the same as the thickness of the substrate (in the analyzed region up to 15 mm of substrate), so this cannot be caused by higher neutron absorption. In other words, the path of the neutron beam stays the same for the WAAM and the forged part. Therefore, this intensity variation is an evidence of the microstructural changes reported in Figure 6 and, in particular, of the crystallographic texture. The ratio between the 1122 and 2021 peak intensities changes with the height for the longitudinal component (Figure 9c). The 0004 peak increases its intensity for the points in the transition zone and inside the substrate for the normal component (Figure 9b). Thus, an increase in intensity from the prismatic crystallographic plane (1000) for the transversal component could be expected. Since we mostly observed mixed reflections (1122 and 2021), it is hard (and anyway out of the scope of this study) to quantify the effect of crystallographic texture from the current experiments.

**Figure 9.** Diffractograms along the N-line for (**a**) transversal component, (**b**) normal component, (**c**) longitudinal component.

The normalized integrated intensity and the full width at half maximum (FWHM) of diffraction peaks along the N-line are presented in Figure 10. Although the coarse grain effect was observed for the longitudinal component (Figure 7a), the FWHM and the integrated intensity of the peak did not vary within the sample (Figure 10b). As for the transition zone and for the substrate, the intensity drops for each measured strain component. In the substrate (Figure 6a), the intensity increases for the normal component. This can be caused by microstructural and crystallographic texture changes (see Figures 5 and 6). It should be noticed that most changes happen in the transition region (from +10 mm to −10 mm) and stabilize inside the WAAM section and the substrate. An increase of the FWHM is often connected to plastic deformation of the material, however, in the present case the FWHM remains constant for the whole height of the sample. Only the transversal component of some points in the substrate shows an increase of the FWHM; this is an artifact due to the low intensity of the peaks (i.e., higher fitting error), see Figure 9a.

**Figure 10.** (**a**) FWHM and (**b**) normalized integrated intensity of the 1122 -Ti diffraction peak along the build direction.

The hydrostatic σ*H* and von Mises stress σ*vM* profiles along L- and N-lines are presented in Figure 11. The stress range for both is around 100 MPa. The hydrostatic stresses cannot be released by post-processing heat treatment and, hence, could be critical for structural integrity of the part under load. In the present case, σ*H* reaches its maximum value of around 75 MPa, this is only 10% of the yield stress of WAAM Ti-6Al-4V (around 750 MPa according to [41]). While this value is not critical for the part integrity, it should be taken into account when assessing the life of the component (e.g., under fatigue).

**Figure 11.** Hydrostatic and von Mises stress along (**a**) L-line, (**b**) N-line. Error bars are contained in the markers.

The longitudinal RS profile along the sample L-line shows a maximum (about 100 MPa) in the middle of the sample (Figure 12a). The profile is symmetric and is similar to that simulated by Ding et al. [42] for steel WAAM parts. The stress profiles along N-line are more complicated (Figure 12b). In the literature the longitudinal component of the RS (coinciding with the deposition direction) usually shows the highest values and the highest gradients. Several experimental RS determination methods (such as the contour method [43] and neutron diffraction [44]) and modeling [42] have revealed the typical longitudinal RS profile generated along the height of WAAM parts. There is some consensus that high tensile stress appears in the region near the substrate and constantly decreases to compressive towards the top of the sample (see schematic in Figure 12c). This is also typical for stress profiles across heterogeneous junctions (see e.g., [45]). The RS rapidly decreases from the transition zone (WAAM/substrate) to the bottom of the substrate. In fact, the whole part should satisfy the stress balance conditions. In our case, the RS decreases to around −50 MPa at 20 mm and then increases to slightly tensile values and decreases again at the very top of the sample (Figure 12b). This follows typical RS profiles (see Figure 12c), with the exception of the point at 20 mm. As mentioned above, during production the energy input was decreased for the first four layers (around 18 mm, see Table 2); therefore, this region of the sample had different thermal input compared to the rest of the WAAM part. This may also be the reason why in this region low RS is found. The sensitivity of RS to variations of the production parameters has been reported for many AM materials [31,36].

Along the build direction all components of the RS lie in the range between −100 MPa and 100 MPa (Figure 12b). This stress range is lower than that reported for WAAM Ti-6Al-4V (reaching around 600–700 MPa [33,37,43,44], with a maximum tensile stress value of around 500 MPa). High RS could be induced by large cooling rates, which are favored by large contact areas (WAAM fabrication, Figure 1) and small substrate thickness, as they were used in the cited works. In the present study, the part of the pre-form used as substrate (a thin upright wall, as in hybrid fabrication, Figure 1), allowed a larger heat accumulation. In this case the geometry of the hybrid part played an important role in RS formation. Furthermore, the used deposition strategy with a sinusoidal path induced a higher energy per unit length (Figure 3b). Thus, the average temperature level was higher during production. This caused a lower temperature gradient, lower flow stresses, and, therefore, smaller RS. Furthermore, the repeated sinusoidal passes induced a so-called intrinsic heat treatment (IHT), analogous to certain scan strategies for L-PBF (see e.g., [31,36]), thereby offering a mechanism for RS relaxation.

**Figure 12.** RS profile (**a**) along the L-line and (**b**) along the N-line, (**c**) schematic of the typical stress profile along the WAAM sample height.

#### **4. Summary**

We investigated a so-called hybrid Ti-6Al-4V part, made of a WAAM wall and a forged T-shape substrate. Such a component showed the following microstructure along its height: bi-modal (α+β) in the forged part, martensitic in the transition zone, and Widmanstätten laths inside the WAAM part (as observed by optical microscopy). Neutron diffraction patterns revealed the presence of preferential crystallographic orientation, changing along the sample height (build direction) but not within deposited layers. Residual stress analysis in the bulk of such hybrid part (by means of neutron diffraction) showed a strong stress gradient in the transition region for every stress component. The hybrid forging+WAAM production induced lower residual stress (maximum values around 100 MPa) compared to purely WAAM components reported in the literature (with maximum stress around 500 MPa). We explained such a low stress by the higher heat accumulation and lower cooling rate during hybrid production (the thin substrate hinders heat accumulation), as well as by the sinusoidal shape of the torch movement. An intrinsic heat treatment (successive sinusoidal passes) also assists stress relaxation. The residual stress profiles along the build direction are also expected to be affected by the heat input variation within the part during production. Simulation work is ongoing to understand the formation of residual stress and microstructure in such hybrid WAAM parts.

**Author Contributions:** Conceptualization, T.T., B.S.; neutron diffraction experiment, T.T., A.U., T.M.; microstructural investigation, I.S.; data curation, I.S., T.M.; writing—original draft preparation, T.M., G.B.; writing—review and editing, all authors.; supervision, G.B., M.B. All authors have read and agreed to the published version of the manuscript.

**Funding:** The authors gratefully acknowledge the financial support provided by Federal Ministry for Economic Affairs and Energy (BMWi) for the LUFO SAMT64 Project "Forging and additive manufacturing as a process combination for the resource-efficient production of aerospace structural components made of TiAl6V4 on flexible production scales" (20W1719D).

**Acknowledgments:** The authors would like to thank Robert Wimpory for the support during the beamtime. The authors are also grateful to Frank Meiners from OTTO FUCHS KG for providing the material used in this study.

**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*