Laser additive manufacturing of biodegradable magnesium alloy WE43: A detailed microstructure analysis

Acta Biomaterialia 98 (2019) 36–49 Acta Biomaterialia xxx (xxxx) xxx

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Laser additive manufacturing of biodegradable magnesium alloy WE43: A detailed microstructure analysis q Florian Bär a, Leopold Berger a, Lucas Jauer b, Güven Kurtuldu a, Robin Schäublin a, Johannes H. Schleifenbaum b,c, Jörg F. Löffler a,* a b c

Laboratory of Metal Physics and Technology, Department of Materials, ETH Zurich, 8093 Zurich, Switzerland Fraunhofer Institute for Laser Technology ILT, 52074 Aachen, Germany Digital Additive Production, RWTH Aachen, 52074 Aachen, Germany

a r t i c l e

i n f o

Article history: Received 14 December 2018 Received in revised form 30 April 2019 Accepted 21 May 2019 Available online 25 xxxx May 2019 Keywords: Laser powder bed fusion Magnesium Bone scaffolds Biodegradable implants WE43 Microstructure Electron microscopy Rapid solidification

a b s t r a c t WE43, a magnesium alloy containing yttrium and neodymium as main alloying elements, has become a well-established bioresorbable implant material. Implants made of WE43 are often fabricated by powder extrusion and subsequent machining, but for more complex geometries laser powder bed fusion (LPBF) appears to be a promising alternative. However, the extremely high cooling rates and subsequent heat treatment after solidification of the melt pool involved in this process induce a drastic change in microstructure, which governs mechanical properties and degradation behaviour in a way that is still unclear. In this study we investigated the changes in the microstructure of WE43 induced by LPBF in comparison to that of cast WE43. We did this mainly by electron microscopy imaging, and chemical mapping based on energy-dispersive X-ray spectroscopy in conjunction with electron diffraction for the identification of the various phases. We identified different types of microstructure: an equiaxed grain zone in the center of the laser-induced melt pool, and a lamellar zone and a partially melted zone at its border. The lamellar zone presents dendritic lamellae lying on the Mg basal plane and separated by aligned Ndrich nanometric intermetallic phases. They appear as globular particles made of Mg3Nd and as platelets made of Mg41Nd5 occurring on Mg prismatic planes. Yttrium is found in solid solution and in oxide particles stemming from the powder particles’ shell. Due to the heat influence on the lamellar zone during subsequent laser passes, a strong texture developed in the bulk material after substantial grain growth. Statement of Significance Additively manufactured magnesium alloys have the potential of providing a major breakthrough in bone-reconstruction surgery by serving as biodegradable porous scaffold material. This study is the first to report in detail on the microstructure development of the established magnesium alloy WE43 fabricated by the additive manufacturing process of Laser Powder Bed Fusion (LPBF). It presents unique microstructural features which originate from the laser-melting process. An in situ transmission electron microscopy heating experiment further demonstrates the development of two distinct intermetallic phases in additively manufactured WE43 alloys. While one forms already during solidification, the other precipitates due to the ongoing heat treatment during LPBF processing.  2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

1. Introduction Because magnesium alloys are inherently biodegradable, they provide key advantages as temporary implant material compared

q Part of the Special Issue associated with the 10th International Conference on Biodegradable Metals, 10th Biometal 2018, held at the University of Oxford, 26–31 Aug. 2018, organized by Professors Diego Mantovani and Frank Witte. ⇑ Corresponding author. E-mail address: [email protected] (J.F. Löffler).

to established materials such as titanium or stainless steel [1,2]. Good mechanical properties and degradation characteristics also make them superior to common biodegradable polymer implant solutions, such as polylactic acid [3–5]. While magnesium-based cardiovascular stents and implants for bone fracture fixation have been intensively studied and also successfully introduced to the market [6,7], the potential of magnesium in patient-specific implants or porous scaffolds for bone regeneration has not been exploited to a similar extent, mainly due to the lack of adequate production routes.

https://doi.org/10.1016/j.actbio.2019.05.056 1742-7061/ 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Please cite this article as: F. Bär, L. Berger, L. Jauer et al., Laser additive manufacturing of biodegradable magnesium alloy WE43: A detailed microstructure analysis, Acta Biomaterialia, https://doi.org/10.1016/j.actbio.2019.05.056

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The selective laser melting or laser powder bed fusion (LPBF) technique is a promising additive manufacturing technology for the production of geometrically complex and porous structures. Here the base material, supplied as a powder, is spread in layers onto a substrate and subsequently melted by a scanning highenergy laser beam, fusing it to the preceding layer. This enables the direct formation of a dense material into shapes whose complexity cannot be achieved by conventional methods [8,9]. The laser impact induces a short-lived molten volume, a so-called melt pool, extending a few hundreds of microns across the powder bed surface. Its lifetime ranges from 0.5 to 25 ms [10]. This rapid localized heating, followed by immediate quenching, generates an extreme cooling rate in the range of 106–108 K s1 [11]. It drives the material away from thermodynamic equilibrium, sometimes with unknown consequences for the microstructure and ultimately for the resulting physical properties, such as mechanical or corrosion properties. LPBF-processed parts exhibit the imprint of the laser process, and the morphology of the melt pools created by the laser-scan tracks is clearly observable in the optical microscope [12,13]. Their width and depth usually greatly exceed the hatch spacing and layer height, respectively. This implies that each layer gets remolten and solidified multiple times and that the material adjacent to the melt pool experiences a certain heat treatment, likely inducing further microstructural changes. In addition to the influence on the material’s base properties, the microstructure is sensitive to numerous process parameters. These include the laser power and intensity distribution, the powder bed temperature, the scanning speed, the scan pattern, the scan-line distance (hatch spacing), the layer thickness, and other factors which influence the material’s cooling rate and thus grain morphology, size, and orientation. LPBF has already been applied to titanium or stainless steel in the production of conventional implants [9,14], and is now envisaged for biodegradable Mg alloys. However, compared to established materials in LPBF processing, magnesium poses certain difficulties which are related to its close melting (650 C) and boiling points (1091 C), and the resulting particularly high vapour pressure of its melt [15]. Some studies on the LPBF processing of Mg alloys have already been conducted on pure magnesium [16,17], Mg–Zn–Zr [18,19], Mg–Y–RE [20–22], where RE stands for rare earths, and Mg–Al [23,24] alloys. The Mg–Y–RE alloy WE43, with a nominal composition of 3.7–4.3 wt% yttrium, 2.0–2.5 wt% neodymium, 0.4–1.0 wt% zirconium, and a maximum of 1.9 wt% of other rare-earth elements [25], appears especially promising for LPBF because of WE43’s potential as structural and biocompatible material [5,21,26,27]. A very similar, conventionally manufactured Mg–Y–RE–Zr alloy has been reported to feature good biocompatibility and osteoconductivity with no signs of toxicity [7,28], resulting in the successful market introduction of screws for the treatment of hallux valgus and fractures of small bones in the European Union. The main alloying element, yttrium, is known to improve the alloy’s overall degradation behavior by decreasing galvanic coupling due to intermetallic phases, and, depending on the environment, producing a net increase in corrosion resistance by forming a protective surface oxide layer [29–31]. Rare-earth elements are also added to the alloy to enhance its mechanical properties [32,33]. The most advanced Mg-based structures so far made by LPBF were reported for WE43 and involved the successful realization of implant demonstrators with interconnected pore structure [20,34]. The goal of these efforts has been the realization of biodegradable metal scaffolds for bone regeneration, which is certainly one of the most promising potential applications for LPBF of magnesium alloys, with WE43 being the logical first candidate. Conventionally synthesized WE43 exhibits various stable and metastable intermetallic phases, which to a large extent are gov-

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erned by aging conditions [35]. Nie et al. [36,37] reported the coexistence of b’ (orthorhombic Mg12NdY), b1 (face-centered cubic Mg3RE) and b (face-centered cubic Mg14Nd2Y) phases. At higher Nd content, the Mg–Nd system reportedly forms Mg2Nd, MgNd [38] and Mg12Nd [39]. In LPBF-manufactured WE43, Zumdick et al. and Li et al. found yttrium oxide particles and proposed the potential presence of the intermetallic phases Mg3Gd, Mg3Nd, Mg24+xY5, Mg41Nd5, and Mg45.9Gd9.08 based on XRD measurements [21,22], but were not able to identify them with certainty. In the ZK60 magnesium alloy, with a nominal composition of Mg5.2Zn-0.5Zr in wt% [18], grains were found to be equiaxed in the center of the melt pool while exhibiting a columnar morphology at the melt-pool borders; this was attributed to the higher cooling rate in the latter [19]. In laser-processed AZ91, with a nominal composition of Mg-8.95Al-0.44Zn-0.19Mn in wt%, Wei et al. [23] observed a different grain morphology, and also a b-Mg17Al12 eutectic phase, compared to the cast alloy. Instead of large grains, and with the eutectic phase located as lamellae at the grain boundaries, the laser-processed microstructure exhibited unconnected b-phase particles which surround smaller and equiaxed a-Mg grains. A clear description of the microstructure evolution upon LPBF and, in particular, of texture development and phase formation is thus still lacking for WE43 and other Mg alloys. This study aims to provide detailed insights into the microstructure of LPBF-fabricated WE43 in terms of grain morphology, and precipitation composition and structure. This is achieved by employing scanning electron microscopy (SEM) and transmission electron microscopy (TEM) for imaging, coupled with energy dispersive X-ray spectroscopy for chemical analysis, and both electron diffraction and X-ray diffraction for structural analyses. These techniques allow us to acquire spatial, chemical and crystallographic information with nanometer resolution. We aim to clearly identify the disputed intermetallic phases in LPBF-processed WE43 alloys, and try to understand and explain the processes that occur during solidification and thus generate the observed microstructure. Cast WE43 is used as a reference. TEM in situ heating was also performed for the first time to gain insights into the dynamics of the phase transitions that occur during LPBF of WE43. The results are discussed in terms of microstructural evolution under out-ofequilibrium conditions.

2. Materials and methods 2.1. Production of WE43 specimens The LPBF samples were made from Mg-alloy powder MAP + 43 (composition in Table 1) provided by Magnesium Elektron. The nominal powder-particle size ranged from 20 to 63 mm. LPBF was performed at the Fraunhofer Institute for Laser Technology ILT in Aachen, Germany. A laboratory LPBF device consisting of a single-mode ytterbium fiber laser and a process chamber, based on AconityMINI system technology, was employed to fabricate 10  10  5 mm3 cuboids. The substrate was made of commercially pure magnesium, with a diameter of 53 mm and a thickness of 11 mm. No preheating was performed. All samples were made at a laser power of 200 W, with a beam diameter of approximately 125 mm, a scan speed of 700 mm s1, a hatch spacing of 40 mm, and a layer thickness of 30 mm. The scanning was performed according to a bidirectional scan pattern rotated by 90 at each new layer (Fig. 1). Each added scan track causes elevated temperatures in the adjacent, already built-up material [40]. Consequently, every location of a LPBF-built material experiences multiple heat treatments of various degrees. By interrupting the laser-scanning process at the center of the specimen’s top layer at the end of sample manufacturing, we obtained a small volume

Please cite this article as: F. Bär, L. Berger, L. Jauer et al., Laser additive manufacturing of biodegradable magnesium alloy WE43: A detailed microstructure analysis, Acta Biomaterialia, https://doi.org/10.1016/j.actbio.2019.05.056

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Table 1 Chemical composition in wt% of the investigated WE43 material as determined by energy-dispersive X-ray spectroscopy (EDS) in SEM. Gd, Dy, La, and other rare earths are displayed combined as ‘Other RE’. The slight increase in the relative content of alloying elements for LPBF compared to the powder is caused by partial Mg evaporation during the melting process.

Cast Powder LPBF

Mg

Y

Nd

Other RE

Zr

92.1 92.1 91.3

4.2 4.0 4.2

2.5 2.3 3.2

0.5 1.0 0.6

0.7 0.6 0.7

Fig. 1. Schematics of the LPBF-process building strategy, also illustrating the interruption of the last laser-scan track (left). Resulting specimen (middle) and SEM micrograph of the last melt pool (right).

(termed ‘‘last melt pool”) in which the material has not been exposed to any secondary heat influence (see blue dot in Fig. 1). By comparing this last melt pool with the bulk of the material (termed ‘‘heat-affected zone”), we were able to study the evolution of microstructure that results from the repeated heat treatments during material build-up. Material of the same nominal composition, but cast into a 25mm-diameter mould, was used as a reference. Table 1 presents the materials’ compositions as measured by energy-dispersive X-ray spectroscopy (EDS). 2.2. Microstructure analysis Sections for metallography were ground manually to SiC grinding paper grade P4000, followed by polishing with a 1/4 lm diamond paste on a Struers RotoPol-21 device using propyl alcohol as lubricant to avoid oxidation. Nital was used to reveal microstructural features. The specimens’ surfaces were investigated using light microscopy (Reichert–Jung Polyvar MET) and SEM (Hitachi SU-70 SEM with a field-emission gun and Oxford Instruments EDS detector). Phase fractions were determined by ImageJ threshold analysis on optical micrographs. Electron backscattered diffraction (EBSD) analyses were performed using a JEOL JSM 7000F instrument at 15 kV and 30 nA probe current with a tilt angle of 70. A step size of 0.2 mm was used for the preparation of the EBSD maps. Grain size was determined by the linear intersection method on optical micrographs and EBSD maps for cast and LPBF-processed materials, respectively. For detailed microstructure investigations of the last melt pool, TEM samples were prepared by focused ion beam (FIB) with a Zeiss NVision 40 FIB-SEM. The lamellae were first thinned using 30 keV Ga+. Subsequently, they were tilted 10 into the beam and showered with a Ga+ beam of 5 keV to reduce the thickness of the surface layer damaged by the ion beam. TEM specimens from the LPBF-heat-affected zone and from the cast material were prepared from approximately 400 mm thin plates that were cut with a diamond wire saw. 3 mm discs were punched from these plates and ground manually down to about 50 mm with SiC grinding paper. The final thinning was achieved

using a GATAN PIPS II precise ion polishing system. This consists of two Ar+ beams that are directed to the disc, one from the top and one from the bottom under a shallow incidence angle, until a hole is created in the center of the disk, whose edge is transparent to electrons. The beam’s angle was set to 3.5 to the surface and the ion energy was set to 4 keV until the hole appeared, then to 2.5 keV for 20 min to improve the surface quality and 1 keV for 10 min as a final cleaning step. Milling was performed under liquid-nitrogen cooling in order to minimize thermally induced changes in the microstructure. TEM investigations were conducted on a FEI Talos F200X operated at 200 kV and equipped with a field emission gun. Bright field (BF) TEM images and high-angle annular dark field (HAADF) images in scanning mode (STEM) were obtained. HAADF images present contrasts related to the atomic mass, thus revealing secondary-phase precipitates. EDS analysis was performed using FEI’s Super-X detector. Chemical maps were acquired in 10–60 min and data were interpreted with Bruker’s Esprit software using the Cliff-Lorimer scheme. Crystallographic analyzes were carried out using selected-area electron diffraction (SAED) and nanodiffraction (ND). Diffraction patterns were interpreted using the electron diffraction analysis software CrysTBox [41] and additionally confirmed by jEMS V4 software [42]. A number of possible phase candidates were tested, i.e. Mg12Nd and Mg17Nd2 [43], Mg41Nd5 and MgNd [44], Mg5Nd [45], Mg14Nd [46], Mg3Nd [47], Mg2Nd [48] and Mg24Y5 [49]. The chemical maps of the precipitates and their diffraction patterns were generally made in areas thin enough so that they were not covered by the matrix. In situ heating in TEM was performed using the windowless heating chip and double-tilt sample holder ‘Wildfire’ from DENSsolutions on a sample prepared via FIB. The FIB lamella was made with a thickness of about 1 mm in order to minimize the effects of the presence of free surfaces and of the potential evaporation of Mg at high temperatures on the studied phase transitions occurring in the bulk. Thanks to the low density of Mg (1.7 g cm3) the lamella was transparent to 200 kV electrons. In the in situ TEM experiment the specimen was rapidly heated from room temperature to 100 C, and then ramped up from 100 C to 400 C at a constant rate of 20 C min1. During the heating 1024  1024 pixel

Please cite this article as: F. Bär, L. Berger, L. Jauer et al., Laser additive manufacturing of biodegradable magnesium alloy WE43: A detailed microstructure analysis, Acta Biomaterialia, https://doi.org/10.1016/j.actbio.2019.05.056

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micrographs were acquired in HAADF STEM mode (to reveal precipitation), with a dwell time of 1 ms for an acquisition time of 1.3 s per image.

3. Results 3.1. Cast material Fig. 2 presents light microscopy, STEM imaging and chemical mapping, and XRD data of cast WE43. The cast material has an average grain size of 10.2 ± 2.1 mm, as determined from light microscopy micrographs by the linear intersection method

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(Fig. 2a). The grains show no evident shape anisotropy, but a large fraction of the observed grains reveals an irregular hexagonal shape. Intermetallic particles in two sizes were observed, namely large ones of a few micrometers in size and small ones of about an order of magnitude smaller (Fig. 2a, b). The large particles represent a total volume fraction of 5.1% and were found to be uniformly distributed throughout the samples. STEM-EDS mapping analysis (Fig. 2c–f) coupled with XRD measurements (Fig. 2g) indicates that the large intermetallic particles are made of Mg41Nd5. The small particles were found to be present in two different morphologies. The first set of particles has a globular shape of about 300 nm in diameter (Fig. 2h) and they are located at grain

Fig. 2. (a) Light micrograph of cast WE43 displaying the intermetallic phase Mg41Nd5 (dark particles). (b) HAADF-STEM image showing globular particles at grain boundaries and dispersed platelet-shaped particles. Corresponding STEM-EDS maps (not quantitative) revealing (c) Nd, (d) Y, (e) other rare earths (mainly Gd, Dy and La), and (f) Zr. (g) XRD spectrum of cast WE43 identifying the intermetallic phase particles seen in (a) as Mg41Nd5. (h) Globular and (i) platelet particles imaged in BF-TEM magnified view.

Please cite this article as: F. Bär, L. Berger, L. Jauer et al., Laser additive manufacturing of biodegradable magnesium alloy WE43: A detailed microstructure analysis, Acta Biomaterialia, https://doi.org/10.1016/j.actbio.2019.05.056

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boundaries. The second set appears in the shape of polygonal platelets of approximately 500 nm in length and around 50 nm in thickness (Fig. 2i). These particles are homogeneously distributed throughout the grains. Both particle types are rich in Nd and other RE though to a lesser extent for the latter (Fig. 2c and e, respectively). Y was distributed rather uniformly throughout the material. However, both particle types present a slightly higher Y content relative to that of the matrix (Fig. 2d). 3.2. LPBF-processed material The LPBF-processed specimens could be produced with a density greater than 99.9% (see also Supplementary Fig. 1). Fig. 3 presents the material in a longitudinal cross section (zy plane with respect to the schematics shown in Fig. 1). The micrographs were taken perpendicular to the laser-scan direction of the last scan track at exactly the last melt pool. The deposited powder-layer thickness of nominal 30 mm could be clearly confirmed. No difference between the x- and y-directions was observed, an expected result due to the 90 scan-pattern rotation at each new layer. The micrographs reveal typical signs of the additive manufacturing process (Fig. 3a). From these we defined three zone types which differ in their microstructural evolution; these are summarized in Fig. 3d. Some features visibly originate from a partially melted zone (Fig. 3a and c), which encloses the completely molten volume (i.e. the melt pool) during the short melting time. The depth and width of the last melt pool (Fig. 3a, b) were found to be about 260 mm and 250–300 mm, respectively. At the melt pools’ bottom,

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thin, near-parallel lines spaced by approximately 500 nm could be observed, forming a lamellar zone (Fig. 3a, c). Towards the melt-pool center equiaxed grains were found, forming the equiaxed zone (Fig. 3b). With the exception of the last melt pool, which is the only volume not affected by a heat treatment through subsequent laser-scan tracks, the material is described as heat-affected zone (Fig. 3a). Some of the microstructural features that are observed in the heat-affected zone are similar to those seen in the last melt pool. In effect, the heataffected zone also includes partially melted zones, and lamellar and equiaxed zones from all previous laser tracks. The microstructure’s gradual change from lamellar to equiaxed morphology is related to the decreasing cooling rate and thermal gradient from the meltpool border to its center during solidification [50]. Fig. 4 shows an EBSD map of the last melt pool and its surroundings. The color code of the map displays the misorientation angle between the c-axis of the hexagonal crystal and the z-direction. The melt pool can be roughly identified by the initiation of elongated columnar grains and is indicated by a line in Fig. 4. It should be noted that most of the columnar grains clearly grow from the existing solid with nearly the same orientation (similar color on EBSD map). During subsequent solidification, equiaxed grains nucleate in the undercooled liquid ahead of the columnar zone and a columnar-to-equiaxed transition occurs in the melt pool [50]. The equiaxed grains have an average grain size of 4.7 ± 0.4 mm. As mentioned above, the columnar zone is termed ‘lamellar zone’ in this study because of its lamellar microstructure. There is a clear microstructure difference between the right and

Fig. 3. (a) Top of the LPBF-processed material, showing a cross section of the very last laser-scan track. (b) and (c) are magnified images of the regions indicated by the white rectangles in (a). (d) Three different types of microstructure are defined: (i) a partially melted zone, which is the transition zone from the melt to the solid material; (ii) a lamellar zone, representing the region at the melt pool’s bottom where regular line patterns are observed; and (iii) an equiaxed zone, where equiaxed grains are visible from the center to the top of the melt pool.

Please cite this article as: F. Bär, L. Berger, L. Jauer et al., Laser additive manufacturing of biodegradable magnesium alloy WE43: A detailed microstructure analysis, Acta Biomaterialia, https://doi.org/10.1016/j.actbio.2019.05.056

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Fig. 4. EBSD map of the last melt pool and its surroundings. The color scheme given below on the left represents the orientation of three directions of hexagonal Mg with respect to the z-direction (build direction). The last melt pool is indicated by a solid line together with three representative regions of the map. Three [0001] pole figures for the three different regions are also shown with their intensity scales.

left-hand sides of the melt pool, which nicely illustrates the effect of the laser-scanning direction on the microstructure. It should be remembered that the laser movement of the last melt pool was along the x-direction (see also Fig. 1). Thus, the area to the right of the melt pool was generated via the last laser scan in xdirection, and columnar grains elongated in y-direction can be identified. On the other hand, the area to the left of the melt pool, where the last layer has not yet been deposited, represents the microstructure after deposition of the second-last layer. Here the laser scanning has been along the y-direction, and equiaxed-like grains may be identified on the existing cutting plane. In reality, however, these grains are not equiaxed, but elongated in xdirection, i.e. columnar grains grew perpendicular to the observation plane during the second-last laser scan (see, for example, region II in Fig. 4 and the corresponding pole figure discussed below). The final grain size and geometry of the additively manufactured bulk material can be seen in the lower part of the EBSD

map in Fig. 4. There are two different types of grain present in the bulk: (i) small grains with an average size of 1.4 ± 0.6 mm, and (ii) large grains in irregular shapes with an average size of 20.4 ± 6.3 mm. The individual laser-scan tracks underneath the last melt pool cannot be recognized in the EBSD map, because of extensive grain growth that occurred in the bulk due to the successive prior laser tracks. The individual layers that form after each scan with a thickness of 30 mm (see Fig. 3) are roughly indicated by wavy blue lines in the EBSD map, to track the grain evolution after each scan. For further clarification, three different regions (as indicated in Fig. 4) were selected to analyse the development of texture via the recording of pole figures. Region I contains the lamellar zone in the melt pool; region II includes a lamellar zone for which the laserscanning movement was along the y-direction; and region III encapsulates the large grains in the bulk material. The pole figures along Mg’s c-axis, presented in Fig. 4, clearly demonstrate the texture development during LPBF-processing of the WE43 alloy for

Please cite this article as: F. Bär, L. Berger, L. Jauer et al., Laser additive manufacturing of biodegradable magnesium alloy WE43: A detailed microstructure analysis, Acta Biomaterialia, https://doi.org/10.1016/j.actbio.2019.05.056

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each of these regions. The pole figure in region II shows most intensity in its center, which clearly indicates the occurrence of columnar grains with extension along the x-direction. In region I, where the columnar grains are elongated along the z-direction, the pole figure shows most of its intensity near the circumference of the pole figure. This is in particular also true for region III, where the c-axis of the large Mg grains is clearly oriented along the zdirection, i.e. the building direction of the additively manufactured material. Fig. 5a presents the TEM microstructure in the central region of the last melt pool where equiaxed grains form. Those grains are of dendritic nature, with RE-rich intermetallic phases being observed

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in the interdendritic regions (Fig. 5b, c). Furthermore, the presence of randomly distributed yttrium oxide particles is revealed (red and yellow in Fig. 5e, f), which originate from the oxide shell of the original WE43 powder particles. Zr (but also Gd and Nd) were found to accompany Y and O, suggesting the presence of additional oxides based on Zr, Gd, and Nd. Fig. 6 shows a STEM image of the melt-pool border (Fig. 6a), with inserts depicting a chemical map of Nd and Y (Fig. 6b), and a diffraction analysis of the Mg matrix in the lamellar zone (Fig. 6c). A melt-pool border noticeable by a partially melted zone is observed stretching from the bottom left in Fig. 6a. The lamellae, of about 500 nm in width, are separated by lines of Nd-rich

Fig. 5. LPBF-processed WE43 microstructure in the last melt pool, as revealed in (a) HAADF STEM imaging and (b-f) EDS chemical mapping for (b) neodymium, (c) gadolinium, (d) zirconium, (e) yttrium, and (f) oxygen. The microstructure is dendritic with intermetallic phases being observed in the interdendritic regions.

Please cite this article as: F. Bär, L. Berger, L. Jauer et al., Laser additive manufacturing of biodegradable magnesium alloy WE43: A detailed microstructure analysis, Acta Biomaterialia, https://doi.org/10.1016/j.actbio.2019.05.056

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Fig. 6. (a) Bright-field image of LPBF-processed WE43. The cross section shows a partially melted zone and a lamellar zone (labeled in the figure). The building direction of the specimen and the directions of the lamellae are indicated by white arrows. (b) A chemical mapping of the lamellar zone and partially melted zone reveals Nd-rich particles on the lamellae boundaries. The corresponding line scan shows varying Nd concentration across adjacent lamellae. Concentration peaks in the lamellae’s centers are indicated by red arrows. (c) BF-TEM image showing similar crystallographic orientation of adjacent lamellae in the lamellar zone. The corresponding SAED pattern reveals that the lamellar planes are oriented perpendicular to Mg’s c-axis. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

particles (Fig. 6b). The compositional profile across the lamellae reveals that an Nd segregation in the Mg matrix exists from one side of a lamella to the other (inset to Fig. 6b). While some Nd dissolves in the Mg matrix, a concentration peak is observed in the lamellae, which forms a neodymium-rich band (concentration profile in Fig. 6b). The Mg lamellae coarsen during partial melting with a visible depletion in Nd at the melt-pool border. In this case, Nd in the Mg matrix diffuses into the Nd-rich particles present between the lamellae (Fig. 6b). BF-TEM reveals that neighboring lamellae form groups which possess a similar crystallographic orientation, as they exhibit the same diffraction contrast (Fig. 6c). An SAED pattern that was acquired over an area which includes several of such lamellae indicates that they are aligned perpendicular to the c-axis of the hexagonal Mg (Fig. 6c), or in other words form on the basal plane of the hexagonal crystal. Fig. 7a shows the microstructure in the heat-affected zone, revealing two types of Nd-rich particle morphologies: (i) platelets and (ii) globular particles. The platelets are about 100 nm in width and 5–10 nm in thickness (Fig. 7a, d, f). They are randomly distributed within the Mg matrix (Fig. 7a) and contain a high content of Nd and Gd (Fig. 7b, c). These elements, however, behave similarly in Mg and form similar binary intermetallics with Mg [51].

Our investigations also found no indication otherwise. In the following, we therefore present the results by displaying Nd only, with the statements given applying also to Gd, the main other rare-earth element present. Platelets orthogonal to the image plane in thin enough regions of the TEM specimen allowed us to perform an EDS analysis without any interference from the matrix. The composition of these platelets is Mg85.6Y2.8RE11.6, in atomic percent. Fig. 7d presents a typical diffraction pattern of a platelet. Note that the bright spots in the pattern relate to hexagonal magnesium, the matrix surrounding the platelet. The weak additional reflections (indexed in Fig. 7d) come from the platelet and could only be matched by a simulated diffraction pattern along the zone axis [1 0 1] with the phase Mg41Nd5. This binary phase has a tetragonal crystal structure with lattice parameters a = b = 14.74 Å and c = 10.40 Å. The SAED pattern of the Mg matrix in Fig. 7f shows that the platelets are preferably aligned along the c-axis of the hexagonal Mg. Fig. 7g also demonstrates clearly the morphology of the platelets along Mg’s c-axis, with an 120 angle between them. Nanodiffraction confirms that they are located on the prismatic planes of the hexagonal Mg. The globular particles exhibit sizes ranging from 40 to 300 nm (Fig. 7a, e). They are also rich in Nd (Fig. 7b). However, unlike the

Please cite this article as: F. Bär, L. Berger, L. Jauer et al., Laser additive manufacturing of biodegradable magnesium alloy WE43: A detailed microstructure analysis, Acta Biomaterialia, https://doi.org/10.1016/j.actbio.2019.05.056

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Fig. 7. (a) Heat-affected zone in LPBF-processed WE43. STEM imaging and chemical mapping in (b) and (c) reveal the presence of two intermetallic particle types, i.e. platelets and globular-shaped particles. (d) and (e) provide close-up morphologies and corresponding diffraction patterns for the platelets (Mg41Nd5) and globular particles (Mg3Nd), respectively. The relevant spots for the crystallographic identification of the platelets and the globular particles are indexed. (f) The microstructure and diffraction pattern of the Mg-matrix demonstrate that the platelets are aligned along the c-axis of hexagonal Mg. (g) The platelets have an angle of 120 between them and are parallel to the prismatic planes of the hexagonal Mg crystal.

platelets they are located mainly at the grain boundaries. Fig. 7e presents a typical diffraction pattern for the globular particles. It could be related only to Mg3Nd by the simulated diffraction pattern for a zone axis [0 0 1], which is a face-centered cubic crystal with a lattice parameter a0 = 7.41 Å. 3.3. In situ heating of the last melt pool A sample was extracted from the last melt pool and annealed in situ in the TEM. This enabled the real-time observation of the heat-induced structural changes that could take place during the transition from the as-solidified state to the heat-affected zone. Fig. 8 presents a sequence of micrographs at different temperatures. Fig. 8a displays the initial condition, observed at 100 C, with yttrium oxide debris at the bottom left, and an Nd-rich band across the image. This microstructure is stable up to about 320 C, where precipitation was observed in the Mg matrix, as indicated by the arrow in Fig. 8b. Simultaneously, the already-present Nd-rich regions form globular particles (indicated by a white arrow in Fig. 8c). Both particle types coarsen with a temperature increase from 320 C to 400 C. At 400 C (Fig. 8c), the size of the particles is with approximately 200 nm greater than that observed in the heat-affected zone of the previous microstructure investigations, where particles of 100 nm in size were identified (see Fig. 7d–f). While temperatures up to the solidus (590 C for WE43 [52]) certainly occur close to the partial melting zone during the LPBF process, the duration of this state is much shorter (seconds) than what

the material experiences in this in situ heating experiment (minutes). Thus, while the transformation path of the secondary phases could be clearly identified, a detailed comparison of the precipitate sizes is not possible. The spatial distribution of Nd and Y before and after the in situ TEM experiment was scrutinized by EDS mapping, presented here for two different regions (Fig. 8d, e). In the melt pool (before the in situ experiment) Nd was found in the intermetallic particles but also as solid solution in the Mg matrix. After heating, the Mg matrix depletes in Nd, which diffuses into the intermetallic particles (Fig. 8d, e). Conversely, the Y-rich features did not evolve (Fig. 8d), which relates to the fact that they are made of yttrium oxide, which is stable at these temperatures. This is also supported by the concentration of solutes in the Mg matrix in the last melt pool and the heat-affected zone of the LPBF-processed material (Table 2), which indicates a lower Nd content, but unchanged Y content after the heat influence resulting from the subsequent laser-scan tracks. 4. Discussion The microstructures of cast WE43 and LPBF-processed WE43 differ from each other due to different initial alloy qualities and thermal histories. While the WE43 powder used for LPBF contains a significant amount of yttrium oxide such that yttrium is somewhat depleted from the matrix by these oxides, we note that cast WE43 lacks any noticeable yttrium oxide particles. The spatial

Please cite this article as: F. Bär, L. Berger, L. Jauer et al., Laser additive manufacturing of biodegradable magnesium alloy WE43: A detailed microstructure analysis, Acta Biomaterialia, https://doi.org/10.1016/j.actbio.2019.05.056

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Fig. 8. In situ heating of the last melt pool of LPBF-processed WE43. (a)–(c) display the same region at different temperatures. At about 320 C precipitation starts visibly, and at 400 C it proceeds significantly. Simultaneously, the Nd-rich regions present in the last melt pool grow and coarsen, and become globular. (d, e) STEM-EDS maps of two other regions reveal that Nd diffuses out of the matrix towards Nd-rich phase particles, while yttrium oxide particles remain unaffected by the heating.

Table 2 Concentration in wt% of elements dissolved in the magnesium matrix of the investigated WE43 material, as determined by EDS in TEM for randomly selected areas free of secondary phases. Elements with concentrations of less than 0.1 wt% are not presented.

Cast LPBF last melt pool LPBF heat-affected zone

Y

Nd

Gd

Dy

Zr

3.6–3.8 1.6–2.1 1.7–2.3

0.2–0.9 1.1–1.4 0.8–1.1

0.1 0.3–0.4 0.3–0.5

0.1–0.25 0.2–0.4 <0.1

0.6–0.7 0.5–0.8 0.6–0.8

distribution of Zr, Gd and Nd in LPBF-processed WE43 suggests the presence of additional oxides formed by these elements. XRD measurements (Supplementary Fig. 2) were found to be in accordance with previously reported data on LPBF WE43 [21,22] and allowed to identify Y2O3 clearly, but not the other oxides. The WE43 powder surface is oxidized deliberately to protect the particles from oxidation for the sake of safety during transportation and handling. Throughout the LPBF process, the particles’ oxide shells are then violently cracked and scattered by the rapid thermal expansion and partial evaporation of the magnesium. Melting of the yttrium oxide does not take place because reaching or even exceeding its melting point of 2438 C [53] can be excluded due to the relatively low boiling point of magnesium (1091 C). Excessive additional energy input by the laser source would simply lead to increased vaporization of magnesium, what is certainly not desired. For future applications of LPBF-processed WE43 as biodegradable implant material in particular, the reduced Y content in the matrix should be taken into consideration. Whereas Y increases the alloy’s corrosion resistance in conventionally manufactured WE43 parts by forming yttrium oxide at the surface, this beneficial effect may be less pronounced in LPBF-produced parts because Y is already partly bound in the form of oxide particles. However, this assumption is contradicted by recent literature. Li et al. published

polarization curves, and measurements of open circuit potential and corrosion current density of a comparable LPBF-processed WE43 material immersed in SBF [21]. Comparing this data with reported values of cast WE43 immersed in SBF [54] reveals a lower corrosion current density of the LPBF-synthesized WE43 (3.1  105 A/cm2) compared to the cast alloy (3.3  104 A/cm2), indicating an actually improved corrosion resistance for the LPBF-processes sample. Improved corrosion resistance, and therefore reduced degradation rate, is in particular favorable for the application of scaffolds in bone regeneration, which feature large exposed surface areas. Future investigations are, however, needed to understand in more detail the corrosion mechanisms occurring in LPBF-processed Mg alloys. Although rapid solidification of the melt pool produced by the laser beam generates fine equiaxed and columnar grains, as can be seen from the EBSD data in Fig. 4, the grains grow with a strong [0 0 0 1] texture along the building direction during heat treatment that results from the subsequent laser scans throughout the LBPF process. The wavy lines in Fig. 4 indicating the layer boundaries of the LPBF process reveal that grain growth occurs extensively after a single-layer deposition and that the grains reach their ultimate size after the generation of two layers. The detailed mechanism which produces such extensive grain growth and texture development is unknown, but interestingly, the aforementioned

Please cite this article as: F. Bär, L. Berger, L. Jauer et al., Laser additive manufacturing of biodegradable magnesium alloy WE43: A detailed microstructure analysis, Acta Biomaterialia, https://doi.org/10.1016/j.actbio.2019.05.056

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recent publication on LPBF-processed WE43 by Zumdick et al. [22] did not report on this grain growth. In fact it reported a significantly lower average grain size of 1.0–1.1 mm, which is similar to the smallest grains found in this study. While the composition, size and distribution of the powder particles and the processing were basically identical, a difference was in the supplier of the powder (Materials Science and Engineering Werkstoffzentrum Clausthal UG versus Magnesium Elektron). The authors also reported yttrium oxide particles incorporated in the material and speculated about a Zener-pinning mechanism prohibiting grain growth. This seems plausible, as the oxide particles can even influence the microstructure formation during solidification. Because of their high melting point they are the possible nucleation sites in the equiaxed zone [55], although we did not investigate the magnitude of interfacial energy between the oxide particles and liquid. As existing solids in the liquid, they may also affect the grain growth in the lamellar zone. The lamellar structure in the melt pool observed in this study is a typical microstructure observed in eutectic alloys. However, the formation of a lamellar eutectic microstructure during solidification is not expected for WE43 because its composition is far from the eutectic points of binary Mg systems. Here, the Mg phase grows from the liquid primarily in the equiaxed zone, with the formation of hexagonal dendrites revealing the underlying symmetry of the solid Mg (Figs. 3b and 5). Nonetheless, one can understand the formation of the lamellar microstructure when the growth directions of the Mg dendrites are considered, especially the growth competition between the directions within the hexagonal basal plane and those away from it. The anisotropy of the solid–liquid interfacial energy generally determines the growth directions and growth rates of dendrites [56]. Mg dendrites were found to  0 > in the basal plane in a comgrow predominantly along < 1 1 2 mercial Mg alloy (Elektron 21 with a composition of Mg-2.8Nd1.5Gd-0.4Zn-0.4Zr in wt% [57]), which is very similar to the com-

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position of WE43. The latter contains Y, but this is reduced in our LPBF-processed material due to the formation of yttrium oxide particles. Dendrite growth in the basal plane was also observed in other Mg alloys [58–60]. Growth rates that are lower off-basal plane than along this plane generate lined-up dendrite trunks on the basal plane [60]. In the current case the lamellae form on the hexagonal basal plane as demonstrated by the SAED pattern and TEM image of Fig. 6c, and Nd-rich particles create boundaries between the lamellae. Therefore, while the dendrites grow from the partially melted zone on the hexagonal basal plane and form lamellae, the limited growth between these lamellae generates an interdendritic liquid enriched in solute elements, mainly in Nd. When the liquid temperature reaches the eutectic temperature, the eutectic phase mixture, i.e. Mg and Nd-rich particles, solidifies and builds the lamellae boundaries. An eutectic microstructure does not form in the lamellae, which may be due to the growth-rate difference within the on and off-basal plane directions, the low solute content of the alloy, and the rapid solidification of the LPBF process. The last two reasons imply that a supersaturated Mg solid-solution forms instead, which is indeed confirmed by the compositional analyses shown in Figs. 6b and 8. However, not knowing the direction of the thermal gradient in the lamellar zone makes it impossible for us to deduce the crystallographic growth directions. Nevertheless, a dendrite growth direction in the basal plane may explain the different lamellae directions observed in Fig. 6, such as the one near-perpendicular and the other near-parallel to the melt-pool border (partially melted zone) as indicated by the two arrows in Fig. 6. Lamellae that form in the basal plane may have any direction rotated around the growth direction. Immediately after solidification, Nd-rich bands can be found separated by wide regions of lower Nd content (Fig. 5). Based on our observations in the heat-affected zone, we assume that during heat treatment in the subsequent heating cycles those irregularly

Fig. 9. Mg–Nd binary phase diagram. Mg41Nd5 (red) and Mg3Nd (blue) are highlighted [44]. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

Please cite this article as: F. Bär, L. Berger, L. Jauer et al., Laser additive manufacturing of biodegradable magnesium alloy WE43: A detailed microstructure analysis, Acta Biomaterialia, https://doi.org/10.1016/j.actbio.2019.05.056

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shaped intermetallics convert into globular particles of composition Mg3Nd, driven by an energetically favorable surface minimization. In the in situ-heating experiment performed on the last melt pool, we observed that the Nd-rich narrow bands evolve into a string of globular particles, supporting our assumption. In the in situ-experiment we also observed at around 320 C the precipitation of particles from solid solution. In the heat-affected zone these were confirmed to be Mg41Nd5. Both confirmed structures (Mg3Nd and Mg41Nd5, respectively) are equilibrium phases, i.e. they are low-temperature stable binary phases of the Mg–Nd system (see Fig. 9). By studying the age-hardening response of Mg–Nd alloys, Nie and Muddle [36] found a face-centered cubic phase, denominated b1, whose composition is that of Mg3Nd and was also observed in the globular particles of the present study. Interestingly, this composition is close neither to the nominal composition of the melt, nor to the closest thermodynamically stable composition (Mg41Nd5), as depicted in the Mg–Nd phase diagram (Fig. 9). We suppose that due to the rapid solidification this phase cannot form, and that the Mg melt is thus enriched in Nd to even higher extent, allowing the formation of Mg3Nd while still retaining Nd in solid solution within the solidified alloy. Due to the subsequent laser-scan tracks, the Nd that is retained in solid solution precipitates during reheating as Mg41Nd5, forming the platelet-shaped particles on prismatic planes, as seen in Fig. 7d, f, g. Based on high-resolution TEM and nanodiffraction, we found indications of coherency. In fact, it  3 2] for hexagonal  2 3] and [2 appears that the zone axes [5 7 Mg and tetragonal Mg41Nd5, respectively, can be aligned. Definite clarification, however, requires equipment of even higher spatial resolution, and therefore remains a subject of future research. It is notable that a clear gradient of Nd concentration can be determined in the lamellar zone between each agglomerate plane (displayed in Fig. 6). It is thus expected that during the period of elevated temperatures, Nd from the supersaturated solid solution diffuses towards the Nd-rich Mg3Nd particles, causing a depletion in between. The aforementioned Mg41Nd5 platelets were also found to precipitate in the very center of the lamellae where the Nd content is highest (Fig. 6). Yttrium has a strong tendency to form stable yttrium oxide (Y2O3). This is evident in the powder-based material. In the LPBFprocessed samples, large fractions of oxide fragments can be found in which approximately half of the whole yttrium content was detected. This, and the extraordinary thermal history of LPBFfabricated materials, may explain the fact that we were not able to detect binary Mg–Y or ternary Mg–Nd–Y phases as occasionally reported in literature for solution-heat-treated and aged Mg–Y–RE alloys produced by casting [37,46]. In cast WE43, yttrium is found in both the Mg matrix and in the nanometer-sized precipitates or secondary-phase particles. Its distribution there matches that of Nd. No crystallographic difference was found between Y-depleted phases in the LPBF-processed material and the Y-containing phases in cast material. Because the morphology of the observed globular and platelet-shaped particles in cast WE43 matches that found in LPBF WE43, it is assumed that Nd and Y can partially substitute each other in binary compounds due to their similar atomic size. The LPBF process certainly generates a thermally complex system in which solidification and annealing conditions are very dynamic, highly anisotropic and difficult to define. However, despite the complexity of this system we were able to classify the microstructure of the LPBF-produced material into three distinct regions, with a clear precipitation scheme. We were also able to identify Mg3Nd and Mg41Nd5 as secondary-phase precipitates. Additional insights may be gained through simulations of the LPBF process to unveil (1) the spatial distribution and time evolution of the thermal condition in the melt during laser-beam impact, and

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(2) the heat load on the already built regions generated by the subsequent layers. 5. Conclusions  Two Nd-rich particle types were identified as the principal building blocks for the microstructure of LPBF-produced WE43: platelet-shaped Mg41Nd5 and globular Mg3Nd. They were also found in cast WE43 with similar morphology, but they differ in size, density and spatial distribution in the two materials.  In contrast to cast WE43, LPBF-produced material incorporates a large number of yttrium oxide particles originating from the powder state of the starting material. Therefore, a limited residual amount of Y remains as an active alloying element dissolved in the matrix.  Three fundamentally different microstructural features can be differentiated for LPBF-processed WE43. These are a partially melted zone, a lamellar zone and an equiaxed zone, which appear in both the last melt pool (as-solidified state) and in the heat-affected zone.  The microstructural features present in the last melt pool differ from those in the heat-affected zone. The heat exposure of solidified material by subsequent laser-scan tracks causes a heat treatment that promotes both a morphological change in the as-solidified secondary phases and the precipitation of Nd-rich binary intermetallics. It also generates substantial grain growth and a strong texture in the bulk material — [0 0 0 1] of hexagonal Mg is parallel to the building direction during the LPBF process.  LPBF-processed WE43 has great potential for application as porous scaffold material for bone regeneration after further understanding of the microstructural features generated by additive manufacturing. In this way, microstructural tuning can be applied to optimize WE43’s properties for biomedical applications.

Acknowledgements The authors thank the Scientific Center for Optical and Electron Microscopy (ScopeM), ETH Zurich, for access to the instruments and Dr Joakim Reuteler from ScopeM for preparation of the FIB lamellae. We also thank Dr Thomas Weber of ETH Zurich for his support in measuring and analyzing the XRD data, and Dr Alexander Schwedt from the Central Facility for Electron Microscopy at RWTH Aachen for his help with EBSD imaging. The authors further acknowledge financial support from the Swiss National Science Foundation via an SNF Sinergia Grant (CRSII5-180367). Declaration of Competing Interest The authors declare no conflict of interest. Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi.org/10.1016/j.actbio.2019.05.056. References [1] M.P. Staiger, A.M. Pietak, J. Huadmai, G. Dias, Magnesium and its alloys as orthopedic biomaterials: a review, Biomaterials 27 (2006) 1728–1734, https:// doi.org/10.1016/j.biomaterials.2005.10.003. [2] S. Agarwal, J. Curtin, B. Duffy, S. Jaiswal, Biodegradable magnesium alloys for orthopaedic applications: a review on corrosion, biocompatibility and surface

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Please cite this article as: F. Bär, L. Berger, L. Jauer et al., Laser additive manufacturing of biodegradable magnesium alloy WE43: A detailed microstructure analysis, Acta Biomaterialia, https://doi.org/10.1016/j.actbio.2019.05.056