Laser Additive Manufacturing of Fe-Based Magnetic Amorphous Alloys

magnetochemistry

Review Laser Additive Manufacturing of Fe-Based Magnetic Amorphous Alloys

Merve G. Ozden * and Nicola A. Morley

Department of Material Science and Engineering, University of Sheffield, Sheffield S1 3JD, UK; n.a.morley@sheffield.ac.uk * Correspondence: mgozden1@sheffield.ac.uk

Abstract: Fe-based amorphous materials offer new opportunities for magnetic sensors, actuators, and magnetostrictive transducers due to their high saturation magnetostriction (λs = 20–40 ppm) and low coercive field compared with polycrystalline Fe-based alloys, which have high magnetostriction but large coercive fields and Co-based amorphous alloys with small magnetostriction (λs = −3 to −5 ppm). Additive layer manufacturing (ALM) offers a new fabrication technique for more complex net-shaping designs. This paper reviews the two different ALM techniques that have been used to fabricate Fe-based amorphous magnetic materials, including the structural and magnetic properties. Selective laser melting (SLM)—a powder-bed fusion technique—and laser- engineered net shaping (LENS)—a directed energy deposition method—have both been utilised to fabricate amorphous alloys, owing to their high availability and low cost within the literature. Two different scanning strategies have been introduced by using the SLM technique. The first strategy is a double-scanning strategy, which gives rise to maximum relative density of 96% and corresponding magnetic saturation of 1.22 T. It also improved the glassy phase content by an order of magnitude of 47%, as well as improving magnetic properties (decreasing coercivity to 1591.5 A/m and increasing magnetic permeability to around 100 at 100 Hz). The second is a novel scanning strategy, which Citation: Ozden, M.G.; Morley, N.A. involves two-step melting: preliminary laser melting and short pulse amorphisation. This increased Laser Additive Manufacturing of the amorphous phase fraction to a value of up to 89.6%, and relative density up to 94.1%, and lowered Fe-Based Magnetic Amorphous coercivity to 238 A/m. On the other hand, the LENS technique has not been utilised as much as SLM Alloys. Magnetochemistry 2021, 7, 20. in the production of amorphous alloys owing to its lower geometric accuracy (0.25 mm) and lower https://doi.org/10.3390/ surface quality, despite its benefits such as providing superior mechanical properties, controlled magnetochemistry7020020 composition and microstructure. As a result, it has been commonly used for large parts with low

Academic Editors: Sabina Lesz complexity and for repairing them, limiting the production of amorphous alloys because of the size and Cristina Favieres limitation. This paper provides a comprehensive review of these techniques for Fe-based amorphous Received: 16 December 2020 magnetic materials. Accepted: 25 January 2021 Published: 29 January 2021 Keywords: selective laser melting; laser engineered net shaping; 3D printing; magnetic glassy alloys; bulk metallic glasses Publisher’s Note: MDPI stays neutral with regard to jurisdictional claims in published maps and institutional afﬁl- iations. 1. Introduction Functional magnetic materials (FMMs) have gained considerable interest for advanced engineering devices owing to their technical beneﬁts for energy conversion, harvesting, transmission, sensing/actuation [1], and more recently for magnetic refrigeration appli- Copyright: © 2021 by the authors. cations, based on their magnetocaloric effects [2]. Fe-based soft magnetic materials are of Licensee MDPI, Basel, Switzerland. great importance for sensors, transformers, and inductive devices because of their superior This article is an open access article magnetic properties such as outstanding magnetic permeability, low coercivity, and good distributed under the terms and corrosion resistance [3–6]. conditions of the Creative Commons In general, Fe-based soft magnetic alloys are used in two distinct categories Attribution (CC BY) license (https:// of applications: creativecommons.org/licenses/by/ 4.0/).

Magnetochemistry 2021, 7, 20. https://doi.org/10.3390/magnetochemistry7020020 https://www.mdpi.com/journal/magnetochemistry Magnetochemistry 2021, 7, 20 2 of 24

i. For producing and utilising electromagnetic energy: due to their low cost and ecological reasons, the usage of soft magnetic materials comprises an important part of these applications because they have high magnetic permeability, low energy losses and high magnetic saturation. Fe-Si-based alloys are considered as the most representative materials for this area. ii. For signal processing: Fe-Ni-based alloys are usually used in informatics, electron- ics, transducers, magnetic recording heads, microwave installations and so on [7]. In the last two decades, significant progress has been achieved in the understanding of alloy design in an attempt to enhance the glass forming ability (GFA) of soft magnetic amorphous materials [8]. Therefore, many new bulk amorphous alloys with large GFA and good magnetic properties have been reported [9–14]. Casting methods including injection moulding have been used for the production of magnetic bulk metallic glasses (BMG). Zhang and his team utilised casting technique to fabricate glassy toroidal cores having good magnetic properties [15]. They consisted of Fe66Co10Mo3.5P10C4B4Si2.5 BMG with an outer diameter of 10 mm and inner diameter of 6 mm, and showed low coercivity (1.0 A/m), high maximum magnetic permeability (450,000), low core loss (0.4 W/kg at 50 Hz) and maximum magnetic flux density (1 T) [15]. Nevertheless, dimensional limitations and poor mechanical properties restrict the use of casting techniques in the production of bulk metallic alloys. Alternatively, to manufacture three-dimensional (3D) amorphous magnetic parts, powder metallurgy (PM), especially hot pressing and spark plasma sintering, has been extensively exploited [16–18]. To reduce the possibility of the deterioration of magnetic softness because of partial crystallisation in the amorphous matrix, it is necessary that consolidation behaviour and thermal stability of the glassy structure at elevated temperatures is controlled [19]. It is promising that the limitations of conventional techniques can be overcome by using additive manufacturing (AM) in the production of glassy magnetic alloys, which is discussed in detail in Section3. This section also mentions the use of powder-bed fusion and directed energy deposition techniques in the fabrication of amorphous magnetic Fe-based alloys, and the effects of process parameters on the final magnetic and mechanical properties. To emphasise the importance of the amorphous magnetic materials, their properties are analysed and compared with those of polycrystalline materials in Section2.

2. Amorphous Fe-Based Magnetic Alloys Amorphous alloys for soft magnetic applications are often fabricated by rapid solidi- ﬁcation of the melt [20,21]. They are generally prepared with the nearly 20% addition of metalloids (Si, B, Al, C and P) for Fe-based and Co-based alloys [7]. Si and B are important metalloids for glass formation and the amorphous structure stabilisation [21,22]. Typical chemical compositions are such that the combined compositions of Fe, Co, Ni elements are 70–85 atomic (at.)% and those of Si and B are 15–30 at.% in total. However, magnetic glassy alloys have a wide variety of compositions. This allows for a large range of soft magnetic properties to be achieved, which depend upon the demands of the application [23]. In magnetic amorphous alloys, the microstructure lacks atomic long-range order and only exhibits short-range order, which is essentially random atomic arrangement of the liquid melt solidiﬁed at a cooling rate of 105–106 K/s. As a result, there are no crystallite- related defects including grain boundaries and dislocations, leading to a decrease in coercivity (see Figure1)[24]. Magnetochemistry 2021, 7, 20 3 of 25 Magnetochemistry 2021, 7, 20 3 of 24

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Figure 1. ComparisonComparison of of the the magnetic magnetic properties properties of ofsoft soft magnetic magnetic materials materials (reprinted (reprinted with with the the permission from [[23],23], Elsevier, 2013).2013).

FigureFigureAmorphous 1. Comparison 2 illustrates Fe-based of the the magnetic alloys, relationship based properties onbetween inexpensive of so theft magnetic saturation raw materials,materials magnetostriction (reprinted have relatively with constant the high andsaturationpermission Si content magnetisationfrom for [23], the Elsevier, FeCuNbSiB (see 2013). Figure alloy1) and system. high magnetostriction The saturation magnetostriction [ 23], which makes (λ thems) has promisingthe highestcandidates value and is for virtually sensors independent and actuators of [25 Si– composition29]. for the amorphous structure, FigureFiguredifferent2 2 illustrates illustrates from the the thenanocrystalline relationship relationship between betweenstate where the the saturation magnetostrictionsaturation magnetostriction magnetostriction is significantly constant constant de- pendentandand Si content on the forSi for concentration the the FeCuNbSiB FeCuNbSiB and alloy alloyits maximum system. system. The value The saturation saturationis as nearly magnetostriction magnetostrictionhalf that of the ( λamor-s) ( hasλs) phoushasthe highest the structure. highest value value and andis virtually is virtually independent independent of Si ofcomposition Si composition for the for amorphous the amorphous struc- structure,ture, different different from from the nanocrystalline the nanocrystalline state state where where magnetostriction magnetostriction is significantly is signiﬁcantly de- dependentpendent on onthe the Si concentration Si concentration and and its maximum its maximum value value is as isnearly as nearly half that half of that the ofamor- the amorphousphous structure. structure.

Figure 2. The saturation magnetostriction constant (λs) of Fe96−zCu1Nb3SixBz−x alloys as a function of Si content for the amorphous and nanocrystalline states (reprinted with the permission from [23], Elsevier, 2013). FigureFigure 2.2. TheThe saturationsaturation magnetostrictionmagnetostriction constant constant ( λ(λs)s) of of Fe Fe9696−−zCuCu11NbNb33SiSixBxBz−zx− alloysx alloys as asa function a function of ofSi Sicontent contentFigure for for3 the indicates the amorphous amorphous the andchange and nanocrystalline nanocrystalline in the saturation stat stateses (reprinted (reprinted polarisation with with the( theJs) permission and the saturation from [[23],23], magnetostrictionElsevier,Elsevier, 2013).2013). constant (λs) of amorphous magnetic alloys as a function of Fe, Ni and

Figure 3 indicates the change in the saturation polarisation (Js) and the saturation magnetostriction constant (λs) of amorphous magnetic alloys as a function of Fe, Ni and

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Co content. Fe-rich alloys possess the highest saturation polarisation and saturation mag- Magnetochemistry 2021, 7, 20 4 of 24 netisation constant, decreasing as Co and Ni concentration increases. Js for amorphous materials is usually lower than polycrystalline ones (Figure 1) because of the addition of

nonmagneticFigure3 indicatesmetalloids the (Si change and B) inrequired the saturation for glass polarisationformation [23]. ( Js ) and the saturation magnetostrictionThe saturation constant magnetostriction (λs) of amorphous of Fe-based magnetic amorphous alloys alloys as a function is typically of Fe, positive, Ni and Co≈ 20–40 content. ppm, Fe-richon the other alloys hand possess for Co-based the highest amorphous saturation alloys, polarisation it is typically and saturation negative, magnetisation≈−5 to −3 ppm. constant,The increase decreasing in λs with as Colowering and Ni Ni concentration concentration increases. is linkedJs tofor a amorphousconcurrent materialsincrease in is J usuallys (|𝜆| ∝ lower𝐽 ). Therefore, than polycrystalline near zero λ oness at high (Figure Ni 1concentrations) because of the only addition occur ofas nonmagneticthe system becomes metalloids paramagnetic (Si and B) [23]. required for glass formation [23].

Figure 3.3. SaturationSaturation magnetisationmagnetisation JJss (dashed lines) and saturationsaturation magnetostrictionmagnetostriction λs (full lines) of amorphous Fe-Ni- based and Fe-Co-basedFe-Co-based alloys with changing Ni and Co content, respectively (reprinted with the permission from [[23],23], Elsevier, 2013).2013).

TheFe-B saturationalloys with magnetostriction a saturation flux of densit Fe-basedy higher amorphous than 1.5 alloys T were is typically the first positive,metallic glass≈ 20–40 developed ppm, on for the the other fabrication hand for of Co-based distributi amorphouson transformers. alloys, The it isaddition typically of negative,Si gave a higher≈−5 to thermal−3 ppm. stability The increase without in λ anys with reduction lowering in Ni the concentration saturation flux is linked density, to a concurrentproducing 2 increasethe ternary in Jalloys (|λs Fe| ∝82BJ12s ).Si6 Therefore,. However, near this zerois proneλs at to high material Ni concentrations oxidation due onlyto air occur pockets as theforming system during becomes the paramagneticproduction process, [23]. lowering magnetic flux density and increasing total Fe-Blosses. alloys Fe81.5 withB13.5Si a3C saturation2 was developed flux density to overcome higher these than limitations 1.5 T were [7]. the first metallic glassFe-Si-B developed glassy for alloys the fabrication possess six of distributiontimes fewer transformers.energy losses Thethan addition traditional ofSi Fe-Si gave al- a higherloys at thermalindustrial stability frequencies. without In any fact, reduction Fe79B13Si in8 has the saturationa higher Curie flux density, temperature producing without the ternarychanging alloy core Fe losses82B12 Siand6. However,flux density, this compared is prone towith material Fe-3%Si oxidation alloys in due the tofabrication air pockets of formingpower transformers during the production[7]. Therefore, process, they are lowering competitive magnetic for applications flux density andwhere increasing Fe-Si al- totalloys are losses. traditionally Fe81.5B13.5 used.Si3C2 was developed to overcome these limitations [7]. Fe-Si-BAfter rapidly glassy quenching alloys possess of amorphous six times feweralloys, energyinternal losses residual than stresses traditional are devel- Fe-Si alloysoped in at the industrial structure, frequencies. altering the In magnetic fact, Fe79 Bbehaviour13Si8 has a(increasing higher Curie coercivity temperature and reducing without changingpermeability) core due losses to the and emergence flux density, of magnetocrystalline compared with Fe-3%Si anisotropy. alloys inStrain the fabricationrelief anneal- of powering has transformers been employed [7]. Therefore,below the theycrystallisati are competitiveon temperature for applications not only whereto achieve Fe-Si higher alloys arepermeability, traditionally lower used. energy losses and smaller coercivity (see Table 1), but also to improve mechanicalAfter rapidly properties quenching by allowing of amorphous atomic alloys,arrangement internal over residual a short stresses distance. are developed Table 1 inclearly the structure,shows that altering although the as-cast magnetic alloys behaviour exhibit very (increasing soft magnetic coercivity behaviour, and reducingstrain re- permeability)lief annealing dueconsiderably to the emergence allows ofthe magnetocrystalline enhancement of magnetic anisotropy. properties Strain relief (reducing annealing co- has been employed below the crystallisation temperature not only to achieve higher permeability, lower energy losses and smaller coercivity (see Table1), but also to improve mechanical properties by allowing atomic arrangement over a short distance. Table1 clearly shows that although as-cast alloys exhibit very soft magnetic behaviour, strain

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relief annealing considerably allows the enhancement of magnetic properties (reducing coercivity, increasing permeability) by reliving the internal residual strains. Amorphous alloys with a wide variety of magnetic properties can be fabricated by annealing in the existence of an applied magnetic ﬁeld [30].

Table 1. Magnetic properties of amorphous alloys under direct current (DC) applications (reprinted with the permission from [30], CRC Press, 2016).

As Cast Annealed 3 3 Alloy Shape Hc (A/m) Mr/Ms µmax (10 ) Hc (A/m) Mr/Ms µmax (10 )

Fe80B20 Toroid 6.4 0.51 100 3.2 0.77 300 Fe40Ni40P14B6 Toroid 4.8 0.45 58 1.6 0.71 275 Fe29Ni44P14B6Si2 Toroid 4.6 0.54 46 0.88 0.70 310 Fe4.7Co70.3Si15B10 Strip 1.04 0.36 190 0.48 0.63 700 (Fe0.8Ni0.2)78Si8B14 Strip 1.44 0.41 300 0.48 0.95 2000 Fe80P16C3B Toroid 4.96 0.4 96 4.0 0.42 130

There are a few limitations of using amorphous magnetic materials in certain applications. Firstly, their low saturation magnetisation restricts their usage in heavy current density engineering. Secondly, their core losses start to rise rapidly at high ﬂux densities. For this reason, they have more use in low-power, low-current applications and specialised small-device applications where transformers are required with moderate ﬂux densities. In these applications, amorphous magnetic materials can be used successfully instead of nickel-iron alloys including permalloy. Amorphous magnetic materials, being manufactured in large quantities, have been utilised in pulsed-power transformers, magnetic sensors, magnetostrictive transducers and communication equipment [30].

3. Additive Manufacturing of Amorphous Fe-Based Magnetic Alloys Additive manufacturing (AM) includes a number of production techniques where components with a variety of structures and complex geometries are manufactured in a layer-by-layer manner directly from 3Reminder: make sure you write from ... to ... D model data [31–35]. The introduction of AM has inﬂuenced the whole fabrication ﬁeld by making the design easy and simplifying the production process, enabling rapid production without needing major change in the fabrication step [36,37]. The economic advantage of AM is revealed especially in low-volume production [37]. Several Fe-based magnetic materials produced by additive manufacturing have been studied owing to their wide range of potential applicability in the energy area [38–44]. Additive manufacturing of amorphous Fe-based magnetic materials is focused on within this review paper. From the range of AM techniques, powder-bed fusion and direct energy deposition have been used for this purpose. This section reviews the AM of amorphous magnetic materials for different laser-based AM techniques, including the research carried out and the advantages and disadvantages of each method [1,45–49]. A summary table is given at the end of the section.

3.1. Powder-Bed Fusion In the powder-bed fusion process, a thin layer (typically 20–100 µm) of very fine powder (with particle size in the range of 20–50 µm) [50] is spread closely packed on a platform. This is where the powder is then fused together with a laser beam or an electron beam. Once the fusion of one layer is completed, another layer of powder is rolled on top of the previous layer and melted together till the targeted 3D part is obtained (Figure4). Powder size distribution and packing both influence significantly the density of the printed component, thus are the most critical factors, along with laser process parameters (laser scan speed, laser power, pulse duration and spot diameter), to the efficiency of this technique [51]. Selective laser sintering (SLS), used generally for various polymers and for some metals with the help of sacrificial binder materials, and selective laser melting (SLM), utilised only for certain metals are the two powder-bed fusion processes, which use a Magnetochemistry 2021, 7, 20 6 of 25

(SLM), utilised only for certain metals are the two powder-bed fusion processes, which Magnetochemistry 2021, 7, 20 use a laser. In SLS process, only sintering occurs between powder particles once temper-6 of 24 ature is increased with a high-power laser above the softening point of polymers. On the other hand, in SLM technique, a relatively higher-powered laser is exploited to fully melt metalliclaser. In powder SLS process, instead only of sintering sintering it. occurs Electron-beam between powdermelting (EBM) particles exploits once temperature an electron beamis increased to melt withthe metal a high-power powder. Unlike laser above SLS, in the the softening SLM and point EBM of processes, polymers. the On laser/elec- the other tronhand, beams in SLM can technique, fully melt athe relatively metal powder higher-powered and fuse them laser together, is exploited leading to fully to meltexceptional metallic mechanicalpowder instead properties of sintering [52]. it. Electron-beam melting (EBM) exploits an electron beam to meltIn the the literature, metal powder. SLM has Unlike been SLS, widely in the utilised SLM andto produce EBM processes, amorphous the magnetic laser/electron ma- terialsbeams due can to fully theirmelt low cost the metaland high powder availability. and fuse Different them together, materials leading and applications to exceptional that exploitmechanical SLM have properties been reviewed [52]. in detail in [53].

FigureFigure 4. 4. SchematicSchematic illustration illustration of of powder-b powder-beded fusion fusion (reprinted (reprinted with with the the pe permissionrmission from from [34], [34], Elsevier, Elsevier, 2018). 2018).

ItIn is theindicated literature, that SLMduring has SLM been processes, widely utilised solidification to produce takes amorphous place relatively magnetic fast withmaterials typical due cooling to their rates low of cost103–10 and4 K/s, high fast availability. enough to Differentprevent crystallisation, materials and [54,55] applications even thoughthat exploit the cooling SLM have rate beenis dependent reviewed on in the detail process in [53 parameters]. [56]. For this reason, SLM is a promisingIt is indicated technique that during to produce SLM processes,bulk amorphous solidification parts. takes place relatively fast with 3 4 typicalFabrication cooling of rates Fe-based of 10 –10bulk K/s,metallic fast glass enough with to SLM prevent has been crystallisation, studied firstly [54, 55by] Jung even andthough his team the cooling [8]. Fe68.3 rateC6.9 is Si dependent2.5B6.7P8.7Cr2.3 onMo the2.5Al process2.1 (at.%) parameters amorphous [56 powder]. For this with reason, different SLM sizesis a promisingin the range technique of 75 µm to and produce 150 µm bulk were amorphous prepared parts. using gas atomisation. The cylindrical Fabricationspecimens ofwith Fe-based a diameter bulk of metallic 2 mm glassand height with SLM of 6 has mm been were studied built with firstly the by SLM Jung technique.and his team The [microstructure8]. Fe68.3C6.9 Si of2.5 theB6.7 finalP8.7Cr pa2.3rtMo was2.5 composedAl2.1 (at.%) of amorphous a mixture of powder amorphous with µ µ anddifferent α-Fe, sizesγ-Fe and in the Fe range23B6 crystalline of 75 m phases, and 150 whichm were are a preparedssociated usingwith impurities gas atomisation. acci- dentallyThe cylindrical included specimens in the master with alloy. a diameter Also, the of effect 2 mm of and the heightlaser power of 6 mm (P) and were scan built speed with (vthe) on SLM the density technique. of the The final microstructure part and magnetic of the properties final part we wasre examined composed in of this a mixture study. α γ Itof was amorphous concluded and that at-Fe, high-Fe scan and speeds Fe23B 6(vcrystalline > 2500 mm/s), phases, bulk which parts could are associated not be gener- with atedimpurities because accidentally the powder includedbed did not in thehave master sufficient alloy. energy Also, input the effect during of the the laser SLM power process, (P ) whichand scan resulted speed in (v incomplete) on the density melting of the and final poor part inter-particle and magnetic bonding properties [57–59]. were At examined lower v in this study. It was concluded that at high scan speeds (v > 2500 mm/s), bulk parts could and higher P, enhanced melting and consolidation of powder brings about the formation not be generated because the powder bed did not have sufficient energy input during the of SLM specimens having high relative densities (Figure 5). To achieve relative densities SLM process, which resulted in incomplete melting and poor inter-particle bonding [57–59]. higher than 99%, it is vital that v lowers to 1500 mm/s and P is higher than 300 W. More- At lower v and higher P, enhanced melting and consolidation of powder brings about the over, hysteresis loops of the atomised powder and SLM specimens produced with differ- formation of SLM specimens having high relative densities (Figure5). To achieve relative ent scan speeds obtained are illustrated in Figure 6. Figure 6 shows that the saturation densities higher than 99%, it is vital that v lowers to 1500 mm/s and P is higher than 300 W. Moreover, hysteresis loops of the atomised powder and SLM specimens produced with different scan speeds obtained are illustrated in Figure6. Figure6 shows that the saturation magnetisation of all three samples are nearly the same within experimental error. The intrinsic magnetic properties including saturation magnetisation (Ms) are dependent on the MagnetochemistryMagnetochemistry 2021 2021, 7, ,7 20, 20 7 7of of 25 25

Magnetochemistry 2021, 7, 20 7 of 24 magnetisationmagnetisation of of all all three three samples samples are are nearly nearly the the same same within within experimental experimental error. error. The The inin- trinsic magnetic properties including saturation magnetisation (Ms) are dependent on the trinsic magnetic properties including saturation magnetisation (Ms) are dependent on the atomicatomic configuration configuration conﬁguration and and the the composition composition of of the the soft soft amorphous amorphous magnetic magnetic materials materials [60], [60], [60 ], indicatingindicatingindicating that that that amorphous amorphous structures structures structures of of of the the the SLM SLM specimens specimens are are identical identical to to those those of of the the startingstarting atomised atomised powder powder in in this this study. study.

Figure 5. Relative density map of the selective laser melting (SLM)-processed Fe68.3C6.9 Figure 5. Relative density map of the selective laser melting (SLM)-processed Fe68.3C6.9 FigureSi2.5B6.7 P5.8.7 CrRelative2.3Mo 2.5densityAl2.1 specimens map of asthe a functionselective of laserlaser powermelting and (SLM)-processed scan speed, the valuesFe68.3C of6.9 Si2.5B6.7P8.7Cr2.3Mo2.5Al2.1 specimens as a function of laser power and scan speed, the values of which Siwhich2.5B6.7P increase8.7Cr2.3Mo in2.5 theAl2.1 directions specimens of theas a arrows. function Figure of laser taken power from and and scan reprinted speed, withthe values the permission of which increase in the directions of the arrows. Figure taken from and reprinted with the permission from increasefrom [8], in Elsevier, the directions 2015 (Common of the arrows. Grading Figure Scale taken (CGS) from to and The reprinted International with Systemthe permission of Units from (SI) [8], Elsevier, 2015 (Common Grading Scale (CGS) to The International System of Units (SI) conver- [8], Elsevier, 2015 (Common 2Grading Scale (CGS) to The International System of Units (SI) conver- conversion 1 emu/g2 = 1 Am /kg and 1 Oe = 79.6 A/m). sionsion 1 1emu/g emu/g = =1 1Am Am/kg2/kg and and 1 1Oe Oe = =79.6 79.6 A/m). A/m).

Figure 6. Hysteresis M-H loops of (a) parent atomised powder and SLM-processed Fe68.3C6.9 Figure 6. HysteresisHysteresis M-H loops of ((aa)) parentparent atomisedatomised powderpowder andand SLM-processed SLM-processed Fe Fe68.368.3CC6.96.9 Si2.5B6.7P8.7Cr2.3Mo2.5Al2.1 specimens generated with (b) v = 2500 mm/s and P = 340 W and (c) v = 1500 Si2.52.5BB6.76.7P8.7P8.7CrCr2.3Mo2.3Mo2.5Al2.52.1Al specimens2.1 specimens generated generated with ( withb) v = ( b2500) v =mm/s 2500 and mm/s P = 340 and WP and= 340(c) v W = 1500 and mm/smm/s(c) v =and and 1500 P P =mm/s =340 340 W, W,and reprinted reprintedP = 340 with with W, the reprintedthe permission permission with from from the [8], permission [8], Elsevier, Elsevier, from2015 2015 (CGS [ 8(CGS], Elsevier, to to SI SI conversion conversion 2015 (CGS 1 1 emu/g = 1 Am2/kg and 1 Oe = 79.6 A/m). emu/gto SI conversion = 1 Am2/kg 1 emu/gand 1 Oe = 1= Am79.62 /kgA/m). and 1 Oe = 79.6 A/m).

Mahbooba et al. tried to fabricate Fe-based BMG samples larger than critical casting thickness with SLM by using gas-atomised FeCrMoB powder with a nominal particle-size

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Magnetochemistry 2021, 7, 20 8 of 24 Mahbooba et al. tried to fabricate Fe-based BMG samples larger than critical casting thickness with SLM by using gas-atomised FeCrMoB powder with a nominal particle-size distributiondistribution ofof 20–8020–80µ µmm [61[61].]. ItIt waswas reportedreported thatthat thethe thicknesses thicknesses ofof thethe produced produced fully fully amorphousamorphous specimens specimens (15 (15 mm) mm) were were much much more more than than the the critical critical casting casting thickness thickness (1 (1 mm)mm) inin all all dimensions. dimensions. In In addition, addition, it wasit was found found that that the the change change in the in SLMthe SLM parameters parameters was notwas enoughnot enough to prevent to prevent stress-induced stress-induced micro-cracking micro-cracking because because of rapid of rapid solidiﬁcation solidification and theand brittlethe brittle nature nature of BMGs. of BMGs. A low A concentration low concentration of localised of localised and isotropic and isotropic nanograins, nanograins, as shown as inshown Figure in7 Figure, was present7, was present in the in microstructure the microstructure of the of producedthe produced part part due due to to mechanical mechanical stress-inducedstress-induced crystallisation. Furthermore, Furthermore, a me a mechanicalchanical test test revealed revealed that thatthe Young’s the Young’s mod- modulusulus of the of SLM the SLM bulk bulkFeCrMoCB FeCrMoCB alloy alloy(220 GPa) (220 was GPa) fairly was close fairly to close the modulus to the modulus of the cast of theamorphous cast amorphous alloys with alloys similar with composition similar composition (190–220 GPa) (190–220 [62]. GPa)This means [62]. This that means amorphous that amorphousmaterials having materials larger having than their larger casting than critic theiral casting thickness critical can be thickness fabricated can using be fabricated SLM with usingmechanical SLM with properties mechanical competing properties with those competing of cast withamorphous those ofalloys. cast amorphous alloys.

FigureFigure 7. 7.Nanograin Nanograin clustersclusters in the cylinder: Electron Electron backscatter backscatter diffraction diffraction (EBS (EBSD)D) micrograph micrograph showing showing a nanograin a nanograin clustercluster in in the the FeCrMoCB FeCrMoCB bulk bulk cylinder cylinder (reprinted (reprinted with with the the permission permission from from [61 [61],], Elsevier, Elsevier, 2018). 2018).

AsAs mention mention in in SectionSection2 ,2, Fe-Si-B Fe-Si-B BMGs BMGs have have attracted attracted much much attention attention as as technological technological materialsmaterials resulting resulting from from their their outstanding outstanding soft softferromagnetic ferromagnetic properties, properties, great great amorphous amorphous phasephase forming forming ability ability (APFA) (APFA) and and exceptional exceptional mechanical mechanical properties properties [63,64]. [63,64]. Consequently, Conse- itquently, is possible it is topossible use them to use as core them materials as core materials sensors [ 65sensors], biosensors [65], biosensors [66] and distribution[66] and dis- transformers.tribution transformers. Mostly, the Mostly, high APFA the high of Fe-B APFA alloys of enablesFe-B alloys the formationenables the of formation the amorphous of the phaseamorphous and Si phase makes and additional Si makes contribution additional contribution to it. A fully to amorphous it. A fully amorphous phase is achieved phase is withinachieved the within range the of 5–26 range at.% of 5–26 B and at.% 0–29 B and at.% 0–29 Si [67 at.%]. Accordingly, Si [67]. Accordingly, Fe92.4Si3.1 FeB92.44.5Sialloy3.1B4.5 wasalloy produced was produced with SLMwith withSLM awith laser a scanlaser speedscan speed of 100–150 of 100–150 mm/s mm/s and and laser laser power power of 90of W90 byW by using using gas-atomised gas-atomised powder powder having having particle particle sizes sizes less less than than 30 30µ µm,m, purchased purchased fromfrom NANOVALNANOVAL companycompany asas FeFe92.492.4Si3.1BB4.54.5 amorphousamorphous powder powder [68]. [68]. Nanocrystalline Nanocrystalline α- α ε Fe-Fe0.950.95Si0.05Si0.05 andand Fe2B Fe phases2B phases in an in amorphous an amorphous ε-FeSi-FeSi type matrix type matrix was observed was observed in the inmicro- the microstructure.structure. The immoderate The immoderate number number of impurity of impurity atoms atoms in the in the interstitial interstitial sites sites is islinked linked to tothe the formation formation of ofthe the amorphous amorphous ε-FeSiε-FeSi type type structure structure by by distor distortingting the the crystal crystal lattice lattice lo- locally.cally. The The strong strong attraction attraction between between Fe Fe and and Si Si atoms atoms indicates indicates that APFA maymay bebe compar-compar- ativelyatively highhigh inin thethe Si Si composition composition aboveabove 2020 at.% at.% Si Si [ 67[67].]. Furthermore, Furthermore, the the large large negative negative − heatheat ofof mixingmixing betweenbetween constituentsconstituents (−2626 kJ/mol kJ/mol for for Fe-Si Fe-Si and and –38 –38 kJ/mol kJ/mol for for Fe-B Fe-B alloys alloys[69]) [ 69allows]) allows short-range short-range order order in the in theliquid liquid upon upon laser laser melting. melting. The The crystallite crystallite sizes sizes of ofα- α-Fe0.95Si0.05 and Fe2B, retained from the starting powder, decreases with increasing laser Fe0.95Si0.05 and Fe2B, retained from the starting powder, decreases with increasing laser scan scan speed due to the high solidiﬁcation rate (106–108 K/s). Such a fast solidiﬁcation of speed due to the high solidification rate (106–108 K/s). Such a fast solidification of droplets droplets hinders the nucleation and growth of the crystallites, bringing about the amor- hinders the nucleation and growth of the crystallites, bringing about the amorphous struc- phous structure. On the other hand, it is found that lower laser scan speeds of 100 mm/s ture. On the other hand, it is found that lower laser scan speeds of 100 mm/s and 400 mm/s and 400 mm/s enable denser microstructure, having lower porosity. enable denser microstructure, having lower porosity. The change in the relative proportions of the phases as a function of laser scan speed The change in the relative proportions of the phases as a function of laser scan speed (Figure8a) depicts the variation in composition during the SLM process and therefore (Figure 8a) depicts the variation in composition during the SLM process and therefore the the atomic motion between distinct phases for Fe92.4Si3.1B4.5 alloys. The relative ratios of atomic motion between distinct phases for Fe92.4Si3.1B4.5 alloys. The relative ratios of ε-FeSi ε-FeSi type amorphous and α-Fe0.95Si0.05 phases show an antagonist behaviour in which the maximum relative ratio of ε-FeSi amorphous phase represents a minimum fraction of α-Fe0.95Si0.05 phase in the sample produced by using laser scan speed of 400 mm/s in

Magnetochemistry 2021, 7, 20 9 of 24

this study. Consequently, it can be said that 400 mm/s is the optimum laser scan speed considering the process parameters used in this study to achieve both densification and high amorphous phase fractions for Fe-Si-B alloys. Furthermore, hardness was measured by a Vickers microhardness tester with a load of 100 gr. The high microhardness values (1654–2273 HV) were observed, linking to the crystallite size refinement as well as the dissolution of boron into the amorphous matrix. Those microhardness values are much Magnetochemistry 2021, 7, 20 higher than those of the Fe Si B produced by the modified melt spinning technique10 of 25 70 10 20 (830–110 HV) [64].

(a) (b)

(c)

Figure 8. (a) Evolution of the relative proportions of the phases, (b) saturation magnetisation alterations and (c) coercivity changes of the selectively laser-melted Fe92.4Si3.1B4.5 specimens as a function of laser scan speed, reprinted with the permis- changes of the selectively laser-melted Fe92.4Si3.1B4.5 specimens as a function of laser scan speed, reprinted with the sion from [68], Springer, 2018 (CGS to SI conversion 1 emu/g = 1 Am2/kg and 1 Oe = 79.6 A/m). permission from [68], Springer, 2018 (CGS to SI conversion 1 emu/g = 1 Am2/kg and 1 Oe = 79.6 A/m).

TheIn the study case conducted of the magnetic by Ouyang performance et al. [71] co ofnfirms the SLM that samples, the properties it is observed of selectively that lasersaturation melted magnetisation parts strongly has depends maximum on the value mechanism of 199 Am of2/kg the formicrostructural the sample fabricated develop- withment thein amorphous laser scan speed alloys of during 700 mm/s the SLM and thenprocess. decreases In this linearly study, toin approximatelyorder to show 188the Ameffect2/kg forof thatSLM with 1500parameters mm/s as canon be seenth ine Figuremicrostructural8b). The alteration evolution in the nearest- of Feneighbour43.7Co7.3Cr conﬁguration14.7Mo12.6C15.5B4.3 ofY1.9 Fe alloy, atoms a map with is their made substitution (Figure 9a). by It nonmagneticindicates that Silow atoms, laser powercauses ( aP decrease < 200 W) in and magnetic high laser moment scan speed per atom, (v > thus700 mm/s) reducing resultMs .in Still, larger the amorphous specimens phasefabricated fractions. by SLM The have relationship much higher (EquationMs values (1)) thanbetween FeSiB amorphous coatings produced content and by using laser theen- ergysame input starting is attained powder (30–40(Figure Am 9b)2/kg) based [70 on]. the laser energy density equation, which is presentedHow by: the coercivity (Hc) changes with respect to laser scanning speed is depicted in Figure8c. It is obvious that Hc increases with increasing laser scanning speed up to 𝑃 400 mm/s and then levels off. Since Hc is𝐸 generally= associated with the size, shape and(1) 𝑣ℎ𝑑

where Ed is the energy input of energy density (J/mm3), P is the laser power (W), v is the laser scan speed (mm/s), h is the hatch distance (mm) and d is the thickness of the powder layer (mm) [72–74]. It is obvious that there is a linear transition from a nearly fully crystalline structure to a fully amorphous one as the Ed decreases. This may be because as energy density decreases, the cooling rate increases, leading to amorphous structure [75]. A novel scanning strategy has been introduced for amorphisation of Fe-based alloys with low glass forming ability with SLM as well as for ensuring enhanced magnetic properties [76]. For this purpose, Fe71Si10C6Cr2 (at.%) was used. The strategy comprises of (1)

Magnetochemistry 2021, 7, 20 10 of 24

the dispersion degree of the crystallites, including lattice distortion and internal stresses, high values of coercivity are attributed to the structural defects, such as vacancies and interstitials that originated from the laser melting process [68]. The study conducted by Ouyang et al. [71] conﬁrms that the properties of selectively laser melted parts strongly depends on the mechanism of the microstructural development in amorphous alloys during the SLM process. In this study, in order to show the effect of SLM parameters on the microstructural evolution of Fe43.7Co7.3Cr14.7Mo12.6C15.5B4.3Y1.9 alloy, a map is made (Figure9a). It indicates that low laser power ( P < 200 W) and high laser scan speed (v > 700 mm/s) result in larger amorphous phase fractions. The relationship (Equation (1)) between amorphous content and laser energy input is attained (Figure9b) based on the laser energy density equation, which is presented by:

Magnetochemistry 2021, 7, 20 11 of 25 P E = (1) d vhd

the preliminary laser melting the loose powder by using a checkboard3 strategy, (2) second wheremeltingEd is the(remelting) energy by inputrandom ofpulses energy (point-random density (P-R) (J/mm strategy)), P andis (3) the finally laser short- power (W), v is the laserpulse scan amorphisation speed (mm/s), as shown hinis Figure the hatch10. In the distance checkboard (mm) strategy, and edged is length the thicknessof 1 of the powdermm layer and laser (mm) power [72 –of74 20]. W It with is obvious 90o rotations that (a therestandard is alternating a linear transition scanning strategy) from a nearly fully were used with an exposure time of 500 µs. In the P-R strategy, minimum distance of 1 crystalline structure to a fully amorphous one as the Ed decreases. This may be because as mm between successive points and maximum laser power of 120 W was exploited. Both energymelting density processes decreases, were carried the cooling out by using rate the increases, same focal leadingdiameter of to 40 amorphous µm. structure [75].

Figure 9. (a) AmorphousFigure 9. (a) Amorphous content content of SLM of SLM Fe43.7 FeCo43.7Co7.37.3CrCr14.7MoMo12.6C12.615.5BC4.315.5Y1.9 Bglassy4.3Y 1.9alloys,glassy 3D-printed alloys, by3D-printed using different by laser using different laser powers andpowers laser and scanlaser scan speeds. speeds. The The data data were were interpolated interpolated between between the measured the measured data points. data (b) The points. graph compares (b) The graphthe compares laser energy density and amorphous content of SLM specimens, reprinted with the permission from [71], Elsevier, 2018 the laser energy(S1, densityS2 and S3 andhave amorphousthe laser energy content densities of of SLM36.8 J/mm specimens,3, 245.3 J/mm reprinted3 and 343.3 withJ/mm3 the). permission from [71], Elsevier, 2018 (S1, S2 and S3 have the laser energy densities of 36.8 J/mm3, 245.3 J/mm3 and 343.3 J/mm3).

Magnetochemistry 2021, 7, 20 11 of 24

A novel scanning strategy has been introduced for amorphisation of Fe-based alloys with low glass forming ability with SLM as well as for ensuring enhanced magnetic properties [76]. For this purpose, Fe71Si10C6Cr2 (at.%) was used. The strategy comprises of (1) the preliminary laser melting the loose powder by using a checkboard strategy, (2) second melting (remelting) by random pulses (point-random (P-R) strategy) and (3) ﬁnally short-pulse amorphisation as shown in Figure 10. In the checkboard strategy, edge length of 1 mm and laser power of 20 W with 90o rotations (a standard alternating scanning strategy) were used with an exposure time of 500 µs. In the P-R strategy, minimum distance Magnetochemistry 2021, 7, 20 of 1 mm between successive points and maximum laser power of 12012 W of was25 exploited. Both melting processes were carried out by using the same focal diameter of 40 µm.

Figure 10. ProductionFigure 10. scheme Production of one scheme layer of of Fe one71Si layer10C6 Crof 2Fealloys71Si10C with6Cr2 alloys the novel with scanning the novel strategy: scanning (a )strategy: loose powder, (b) preliminary laser(a) loose melting powder, of loose (b)powder preliminary by using laser checkboard melting of loose strategy powder (Sample by using A), (c )checkboard remelting with strategy the point-random (P-R) strategy(Sample (Sample A), B)—the (c) remelting numbers with stand the for point-random scanning order—and (P-R) strategy (d) the (Sample completely B)—the re-melted numbers layer stand (reprinted for with the permissionscanning from [76 order—and], Elsevier, 2019).(d) the completely re-melted layer (reprinted with the permission from [76], Elsevier, 2019). It is reported that the novel scanning strategy has enabled the restoration of 89.6% of It is reportedthe amorphous that the novel structure scanning of the stra parenttegy has alloy enabled in spite the of restoration its low glass-forming of 89.6% of ability and the amorphousconsiderable structure of crystallisation the parent alloy during in spite the preliminary of its low glass-forming melting. After ability P-R scanning and (second considerable melting),crystallisation the amorphous during the phasepreliminary appears melting. as a continuous After P-R scanning matrix, rather(second than forming melting), theseparate amorphous regions, phase as appears is observed as a continuous in the specimens matrix, rather produced than withforming a single sep- melting. In arate regions,addition, as is observed second in melting the specimens increases produced the density with of a bulksingle sample melting. from In addition, 78.2% to 94.1%. second melting increasesIt is assumed the density that if of the bulk amorphous sample from phase 78.2% is heated to 94.1%. by a laser pulse to a maximum It is assumedtemperature that if the with amorphous a heating ratephase lower is heated than theby criticala laser pulse value, to it a devitriﬁes. maximum In this study, temperature with a heating rate lower than the critical value, it devitrifies. In this study, the critical heating rate was calculated as 8.77 × 106 K/s, at which the amorphous phase does not devitrify till the melting point of the material is reached. The heating rate for the preliminary melting (pulse duration (exposure time) 500 µs) (Sample A) was estimated as 2.5 × 106 K/s, which is below the critical heating rate value. The heating rate of the second melting (P-R remelting with the pulse duration of 20 µs) (Sample B) was approximated as 6.25 × 107 K/s, which is higher than the critical heating rate [76]. It is indicated that this strategy has a positive effect on soft magnetic properties. Co- ercivity was reduced substantially after the second melting (using P-R scanning) as seen in Table 2. Stress relief annealing at 820 K provided further reduction in coercivity. On the other hand, the M-H hysteresis loops (Figure 11) show that the saturation magnetisation seems to be the same for all samples except sample A within the margin of error.

Magnetochemistry 2021, 7, 20 12 of 24

the critical heating rate was calculated as 8.77 × 106 K/s, at which the amorphous phase does not devitrify till the melting point of the material is reached. The heating rate for the preliminary melting (pulse duration (exposure time) 500 µs) (Sample A) was estimated as 2.5 × 106 K/s, which is below the critical heating rate value. The heating rate of the second melting (P-R remelting with the pulse duration of 20 µs) (Sample B) was approximated as 6.25 × 107 K/s, which is higher than the critical heating rate [76]. It is indicated that this strategy has a positive effect on soft magnetic properties. Coercivity was reduced substantially after the second melting (using P-R scanning) as seen in Table2. Stress relief annealing at 820 K provided further reduction in coercivity. On the Magnetochemistry 2021, 7, 20 other hand, the M-H hysteresis loops (Figure 11) show that the saturation magnetisation13 of 25 seems to be the same for all samples except sample A within the margin of error.

Table 2. Coercivity of feedstock powder and melted samples of Fe71Si10C6Cr2 (at.%) alloy (reprinted Table 2. Coercivity of feedstock powder and melted samples of Fe71Si10C6Cr2 (at.%) alloy (reprinted with the permission from [76], Elsevier, 2019). with the permission from [76], Elsevier, 2019).

SampleSample CoercivityCoercivity (A/m) (A/m) PowderPowder 99 99 Melted onceMelted once 1032 1032 Melted twice 397 Melted twice andMelted annealed twice 238 397 Melted twice and annealed 238

Figure 11. M-H hysteresis loops for Fe 71SiSi1010CC6Cr6Cr2 2parentparent powder, powder, samples samples A A (produced (produced by by melting melting every powder layer only only once) once) and and B B (fabricated (fabricated by by melting melting each each powder powder layer layer twice), twice), and and sample sample B after after annealing, annealing, reprinted reprinted with with the the permission permission from from [76], [76 Elsevier,], Elsevier, 2019 2019 (CGS (CGS to SI to conversion SI conversion 1 emu/g1 emu/g = 1 = Am 1 Am2/kg2/kg and and 1 Oe 1 = Oe 79.6 = 79.6 A/m). A/m).

Most recently, recently, Nam Nam et et al. al. have have utilised utilised a similar a similar scanning scanning technique technique as in as Reference in Refer- [76],ence referred [76], referred to as the to asdouble-scan the double-scan strategy, strategy, to achieve to achievefull verification full verification and densification and den- ofsification Fe-based of BMGs Fe-based having BMGs high having magnetic high satu magneticration [75]. saturation In the double-scan [75]. In the double-scanstrategy, be- forestrategy, the coating before of the a subsequent coating of powder a subsequent layer over powder the build layer area, over every the Fe build73.7Si area,11B11C every2Cr2.28 amorphousFe73.7Si11B11 powderC2Cr2.28 layeramorphous was rescanned powder (remelte layer wasd) with rescanned the linear (remelted) scan method with theusing linear the samescan methodlaser power using and the laser same scan laser speed power as the and fi laserrst laser scan scanning. speed as In the this first study, laser they scanning. com- paredIn this the study, double-scanning they compared strategy the double-scanning with a single scan strategy in terms with of a relative single scan density in terms (Figure of 12b),relative transverse density (Figurerupture 12 strengthb), transverse (TRS) (Figure rupture 12c) strength and magnetic (TRS) (Figure properties 12c) and as a magnetic function of laser energy density, calculated using equation (1). It was found that double scanning enables effective shrinkage of voids formed because of partial melting of powder (Figure 12d), resulting in a relative density of maximum 96% and corresponding to the mechanical strength (TRS) of 75 MPa. This was achieved at the highest Ed of 37.5 J/mm3 (high power of 90 W and low scanning speed of 1200 mm/s) as it is shown in Figure 12d). From the graphs in Figure 13, it is obvious that double scanning improves magnetic properties significantly. Magnetic saturation has a linear dependence on the Ed and a maximum Ms value of 1.22 T was attained with double scanning. Coercivity tends to increase with increasing Ed because higher Ed provides enough time for crystals to nucleate and grow, by lowering the cooling rate. This leads to a reduction in the amorphous fraction (fam), which was found to be 47% for Ed of 29.2 J/mm3. X-ray powder diffraction (XRD)

Magnetochemistry 2021, 7, 20 13 of 24

properties as a function of laser energy density, calculated using equation (1). It was found Magnetochemistry 2021, 7, 20 that double scanning enables effective shrinkage of voids formed because of partial melting14 of 25 of powder (Figure 12d), resulting in a relative density of maximum 96% and corresponding to the mechanical strength (TRS) of 75 MPa. This was achieved at the highest Ed of 37.5 J/mm3 (high power of 90 W and low scanning speed of 1200 mm/s) as it is shown in results (Figure 14) show that the intensity of the α-Fe crystallite peak lowers as Ed de- Figure 12d). creases, implying that the fraction of the crystallite phase reduces.

Figure 12. The graph of (a) macroscopic photograph of 3D-printed Fe73.7Si11B11C2Cr2.28 alloys, (b) relative density and (c) Figure 12. The graph of (a) macroscopic photograph of 3D-printed Fe73.7Si11B11C2Cr2.28 alloys, (b) relative density and (transversec) transverse rupture rupture strength strength (TRS) (TRS) versus versus laserlaser energy energy density density and and (d ()d optic) optic microscope microscope images images of of double-scanned double-scanned alloys (reprinted with the permission from [75], Elsevier, 2020). (reprinted with the permission from [75], Elsevier, 2020).

From the graphs in Figure 13, it is obvious that double scanning improves magnetic properties signiﬁcantly. Magnetic saturation has a linear dependence on the Ed and a maximum Ms value of 1.22 T was attained with double scanning. Coercivity tends to increase with increasing Ed because higher Ed provides enough time for crystals to nucleate and grow, by lowering the cooling rate. This leads to a reduction in the amorphous fraction 3 (fam), which was found to be 47% for Ed of 29.2 J/mm . X-ray powder diffraction (XRD) results (Figure 14) show that the intensity of the α-Fe crystallite peak lowers as Ed decreases, implying that the fraction of the crystallite phase reduces.

Figure 13. The graph of (a) magnetic saturation (Ms), (b) coercivity (in Oe), (c) core loss relativity and (d) relative permeability, µr, taken as at 100 Hz frequency (@ 100 Hz) versus laser power density for 3D-printed Fe73.7Si11B11C2Cr2.28 alloys (reprinted with the permission from [75], Elsevier, 2020).

Magnetochemistry 2021, 7, 20 14 of 25

results (Figure 14) show that the intensity of the α-Fe crystallite peak lowers as Ed decreases, implying that the fraction of the crystallite phase reduces.

Figure 12. The graph of (a) macroscopic photograph of 3D-printed Fe73.7Si11B11C2Cr2.28 alloys, (b) relative density and (c) Magnetochemistry 2021, 7, 20 14 of 24 transverse rupture strength (TRS) versus laser energy density and (d) optic microscope images of double-scanned alloys (reprinted with the permission from [75], Elsevier, 2020).

Figure 13. The graph of (a) magnetic saturation (Ms), (b) coercivity (in Oe), (c) core loss relativity and (d) relative permea- MagnetochemistryFigure 2021, 13. 7, 20The graph of (a) magnetic saturation (Ms), (b) coercivity (in Oe), (c) core loss relativity and (d)15 relative of 25 bility, µr, taken as at 100 Hz frequency (@ 100 Hz) versus laser power density for 3D-printed Fe73.7Si11B11C2Cr2.28 alloys permeability, µ , taken as at 100 Hz frequency (@ 100 Hz) versus laser power density for 3D-printed Fe Si B C Cr (reprinted withr the permission from [75], Elsevier, 2020). 73.7 11 11 2 2.28 alloys (reprinted with the permission from [75], Elsevier, 2020).

Figure 14. X-ray powder diffraction (XRD) patterns of 3D-printed Fe73.7Si11B11C2Cr2.28 alloys with Figure 14. X-ray powder diffraction (XRD) patterns of 3D-printed Fe73.7Si11B11C2Cr2.28 alloys with different laser energy density (reprinted with the permission from [75], Elsevier, 2020). different laser energy density (reprinted with the permission from [75], Elsevier, 2020). In conclusion, SLM has been successfully used to produce Fe-based BMGs because Init conclusion, provides high SLM cooling has been rates successfully 103–104 K/s used [54 to] produce but depends Fe-based strongly BMGs on because the process it pa- 3 4 providesrameters high cooling used such rates as 10 laser–10 power K/s [54] and but laser depends scan speedstrongly [56 on]. Lowthe process laser power parame- and high ters usedlaser such scan as speedlaser power leads toand low laser energy scan inputspeed onto [56]. theLow powder, laser power causing and high high cooling laser rate. scan speedAlthough leads decreasing to low energy the energy input onto input, the enables powder, an increasecausing inhigh the cooling amorphous rate. fractionAlt- of houghthe decreasing magnetic the alloy, energy it affects input, adverselyenables an full increase densification, in the amorphous which requires fraction a highof the energy magneticdensity. alloy, Hence,it affects low adversely energy inputfull densification, results in poor which mechanical requires propertiesa high energy (hardness den- and sity. Hence,mechanical low energy strength). input It isresults crucial in thatpoor optimum mechanical process properties parameters (hardness are defined and me- for both chanicalfull strength). densification It is crucial and amorphous that optimu structure.m process The parameters studies indicateare defined that for double both full scanning densification(remelting) and amorphous helps to increase structure. the glassyThe stud phaseies indicate content that and double relative scanning density, (remelt- and improve ing) helpssoft to magnetic increase and the mechanicalglassy phase properties content and [75 relative]. density, and improve soft magnetic and mechanical properties [75].

3.2. Direct Energy Deposition Direct energy deposition (DED) has been utilised for fabricating high-performance super-alloys. It contains several techniques, which are laser-engineered net shaping (LENS), laser solid forming (LSF), directed light fabrication (DLF), direct metal deposition (DMD), electron-beam AM (EBAM), and wire and arc AM (WAAW). In DED, laser or electron beams are used as a source of energy, which is directed onto a specific location of the substrate and exploited to melt feedstock materials (wires or powder) at the same time. Then, the molten materials are deposited and fused onto the substrate. This process continues layer by layer until the 3D part is built (Figure 15) [77]. Differently from SLM, DED has no powder bed and the feedstock is melted, then deposited into substrate in a layer-by-layer manner, enabling both multiple materials and multiple axis deposition simultaneously.

Magnetochemistry 2021, 7, 20 15 of 24

3.2. Direct Energy Deposition Direct energy deposition (DED) has been utilised for fabricating high-performance super-alloys. It contains several techniques, which are laser-engineered net shaping (LENS), laser solid forming (LSF), directed light fabrication (DLF), direct metal deposition (DMD), electron-beam AM (EBAM), and wire and arc AM (WAAW). In DED, laser or electron beams are used as a source of energy, which is directed onto a speciﬁc location of the substrate and exploited to melt feedstock materials (wires or powder) at the same time. Then, the molten materials are deposited and fused onto the substrate. This process continues layer by layer until the 3D part is built (Figure 15)[77]. Differently from SLM, DED has no Magnetochemistry 2021, 7, 20 powder bed and the feedstock is melted, then deposited into substrate in a layer-by-layer16 of 25

manner, enabling both multiple materials and multiple axis deposition simultaneously.

Figure 15. (a) The experimental setup and (b) the schematics of the laser-engineered net shaping (LENS) process (reprinted Figure 15. (a) The experimental setup and (b) the schematics of the laser-engineered net shaping (LENS) process (reprinted with the permission from [78], Elsevier, 2020). with the permission from [78], Elsevier, 2020).

TheThe LENS LENS technique technique generally has beenbeen utilisedutilised toto produceproduce Fe-based Fe-based BMGs BMGs because because it 3 5 itprovides provides (i) (i) a a high high cooling cooling rate rate (10 (103–10–105 K/s)K/s) [79], [79], (ii) less dilution at at the the substrate-layer substrate-layer interface,interface, restricting restricting crystallisation crystallisation and and (iii) (iii) tailorable tailorable process process parameters parameters with with net net shape shape productionproduction ability ability [80]. [80]. ToTo prevent prevent crystallisation crystallisation in in BMGs BMGs during during the the LENS LENS process, process, laser laser power, power, laser laser scan scan speedspeed and and substrate substrate temperature temperature are are the the main main pa parametersrameters that that can can be be controlled. controlled. The The total total heatheat input input per per unit unit area, area, I,I, isis utilised utilised to to combine combine first ﬁrst two two parameters parameters along along with with the the beam beam diameterdiameter as as follows: follows: P I = (2) vD𝑃 𝐼= (2) where P is the laser power (W), v is the laser scanning𝑣𝐷 speed (mm/s), and D is the laser beam diameter on the substrate (mm) [81]. It is obvious from Equation (2) that reducing laser where P is the laser power (W), v is the laser scanning speed (mm/s), and D is the laser power and increasing laser scanning speed brings about a lower heat input [82], similarly beam diameter on the substrate (mm) [81]. It is obvious from Equation (2) that reducing to the SLM process as mentioned in Section 3.1. The lower the heat input, the higher the laser power and increasing laser scanning speed brings about a lower heat input [82], sim- cooling rate due to high thermal gradients, which restricts crystallisation. Moreover, in ilarly to the SLM process as mentioned in Section 3.1. The lower the heat input, the higher order to maintain high cooling rates and obtain BMGs with maximum amorphous content, the cooling rate due to high thermal gradients, which restricts crystallisation. Moreover, it is crucial that substrate temperature is kept as low as possible. in orderThere to maintain are two wayshigh cooling to estimate rates the and cooling obtain rateBMGs during with LENSmaximum processing. amorphous First con- one tent,includes it is crucial the use that of the substrate following temperature equation: is kept as low as possible. There are two ways to estimate the cooling rate during LENS processing. First one includes the use of the following equation:dT = RG (3) dt 𝑑𝑇 =𝑅𝐺 (3) where dT/dt (K/s) is the cooling rate, R (mm/s)𝑑𝑡 is the solidiﬁcation velocity and G (K/mm) is the local thermal gradient [83]. Through experiments, it has been presented that R is wherein the dT/dt order (K/s) of the is laserthe cooling scanning rate, speed R (mm/s) (v)[ 84is ],the which solidification is 20 mm/s velocity in the and study G (K/mm) by Balla is the local thermal gradient [83]. Through experiments, it has been presented that R is in the order of the laser scanning speed (v) [84], which is 20 mm/s in the study by Balla and his team [74], and G is approximately 100 K/mm [79]. Hence, by using this study, the cooling rate in the melt zone calculated using Equation (3) is approximated to be 2000 K/s, which is more than enough to produce an amorphous NSSHS7574 alloy with a critical cooling rate of 610 K/s [85]. Still, in order to make better estimations of cooling rates, it is necessary to take into consideration the substrate or prior deposit temperature, temperature change in the melt zone, laser input and the thermal conductivity of substrate and deposited alloys. Despite the complicated physics of laser surface melting [86,87], improved approximation of cooling rates during laser surface melting can be achieved by using the Rosenthal solution for a mobile heat source (Equation (4), as presented by Steen [87].

Magnetochemistry 2021, 7, 20 16 of 24

and his team [74], and G is approximately 100 K/mm [79]. Hence, by using this study, the cooling rate in the melt zone calculated using Equation (3) is approximated to be 2000 K/s, which is more than enough to produce an amorphous NSSHS7574 alloy with a critical cooling rate of 610 K/s [85]. Still, in order to make better estimations of cooling rates, it is necessary to take into consideration the substrate or prior deposit temperature, temperature change in the melt zone, laser input and the thermal conductivity of substrate and deposited alloys. Despite the complicated physics of laser surface melting [86,87], improved approximation of cooling rates during laser surface melting can be achieved by using the Rosenthal solution for a mobile heat source (Equation (4), as presented by Steen [87]. dT v = −2πk ∆T2 (4) dt Q where k is the thermal conductivity of the substrate (W/mm.K), v is the laser scan speed (mm/s), Q is the laser power (W) and ∆T is the temperature range during cooling (K). It confirms that reducing the laser power and increasing the laser scan speed provides a higher cooling rate during laser-based additive manufacturing processes (Equation (4)). In addition, by using Equation (4), the cooling rates were quantified for different substrate/prior deposit temperature, which indicates that the cooling rate increases when the prior deposit temperature decreases (Equation (4)) [80]. In the production of Fe-based amorphous alloys with LENS, the initial attempt was made by Balla and Bandyopadhyay [80], who used a gas-atomised NanoSteel NSSHS7574 glass-forming alloy powder of Cr<25Mo<15W<10C<3Mn<5Si<2B<5Febalance (at.%) with a particle size in the range of 53 and 180 µm. Laser power of 250 W and laser scan speed of 20 mm/s were chosen to reduce heat input and to melt amorphous powder. In addition, even though no direct technique exists to control substrate and prior deposit temperature during the LENS process, in order to keep substrate temperature/prior deposit temperature low, a five-seconds delay was inserted between each successive layer with argon gas flowing through the powder delivery nozzles. In this study, the cooling rate for the NSSHS7574 alloy was approximated to be in the range of 3.38 × 103 and 1.72 × 103 K/s using Equation (4), which is higher than the critical cooling rate (610 K/s) required for complete amorphisation of NSSHS7574 alloy. Even though theoretically the process parameters used in this study are ideal to obtain full amorphous alloys, XRD (Figure 16), differential scanning calorimetry (DSC) (Figure 17) and microstructural analysis (Figure 18) of the 3D-printed alloy exhibit partial crystallisation/incomplete amorphisation. It is thought that the reason for that is the coarse feedstock powder (53–108 µm), only partially melted during deposition because of the low heat input (14 J/mm2). Those partially melted particles are found in the deposit with melted and re-solidified surface having amorphous structure, surrounded with crystalline features as shown in Figure 18. It was concluded that to obtain fully amorphous alloys with the help of the LENS technique, the use of a finer starting powder would enable complete melting with the same process parameters, or increasing the heat input would allow for melting particles completely without influencing negatively the critical cooling rate needed to achieve full amorphisation [80]. Magnetochemistry 2021, 7, 20 18 of 25 Magnetochemistry 2021,, 77,, 2020 1718 of 2425

Figure 16. XRDXRD patterns patterns of of the top surface (15 mm for substrate) an andd near substrate (0.8 mm from substrate) of LENS Figure 16. XRD patterns of the top surface (15 mm for substrate) and near substrate (0.8 mm from substrate) of LENS processed Fe-based BMGs compared to NSSHS7574 starting powd powder,er, illustrating that the amor amorphousphous alloys have a mixture mixture processed Fe-based BMGs compared to NSSHS7574 starting powder, illustrating that the amorphous alloys have a mixture of amorphous and crystal crystallineline phases (reprinted with the permissionpermission from [[80],80], Elsevier,Elsevier, 2010).2010). of amorphous and crystalline phases (reprinted with the permission from [80], Elsevier, 2010).

Figure 17. Differential scanning calorimetry (DSC) curves of the feedstock powder and the 3D-printed NSSHS7574 alloy, Figure 17. Differential scanning calorimetry (DSC (DSC)) curves of the feedstock powder an andd the 3D-printed NSSHS7574 NSSHS7574 alloy, alloy, depicting three crystallisation peaks with three-stage crystallisation behaviour, reprinted with the permission from [80], depicting three crystallisation peaks with three-stage crystallisation behaviour, re reprintedprinted with the permission from [80], [80], Elsevier, 2010 (Tg is glass transition temperature and Tx1, Tx2 and Tx3 are crystallisation temperatures, showing that the Elsevier, 2010 (Tg is glass transition temperature and Tx1, Tx2 and Tx3 are crystallisation temperatures, showing that the materialElsevier, undergoes 2010 (Tg is a glass three-stage transition crys temperaturetallisation process and T x1upon,Tx2 heating).and Tx3 are crystallisation temperatures, showing that the material undergoes aa three-stagethree-stage crystallisationcrystallisation processprocess uponupon heating).heating).

MagnetochemistryMagnetochemistry2021 2021, 7, ,7 20, 20 1819 of of 24 25

FigureFigure 18. 18. Cross-sectional SEM micrographs micrographs of of 3D-printed 3D-printed NSSHS7574 NSSHS7574 alloy alloy (a) ( ademonstrating) demonstrating exceptionalexceptional interfacial interfacial bonding bonding between between substrate substrate and and the the produced produced alloy alloy ( b(b)) the the microstructure microstructure at at the top region of the bulk specimen (reprinted with the permission from [80], Elsevier, 2010). the top region of the bulk specimen (reprinted with the permission from [80], Elsevier, 2010).

BorkarThe effect et al. of produced laser process a compositionally parameters on gradedthe hardness Fe-Si-B-Nb-Cu of Fe-based alloy BMGs with produced varying Si/Bby LENS ratios was by usingstudied the by LENS Xie and technique his team [5 [91],]. It waswho observedused Fe41Co that7Cr dendrite15Mo14C15 formationB6Y2 (at.%) occurredamorphous in the spherical amorphous powder matrix, and whichdifferent is associatedlaser energy with densities. constitutional The variation super-cooling in hard- becauseness as ofa thefunction compositional of laser energy differences density between (LED the) is crystal illustrated andamorphous in Figure 19. matrix It is [obvious88–90]. Thethat partitioning hardness increases of Si at the as interfaceLED lowers. of α -Fe/amorphousThe use of the lowest matrix energy causes density the constitutional of 30 J/mm super-cooling,led to the highest promoting hardness the of dendrite1263.7 HV. formation This value owing is significantly to interfacial close instability. to the hardness Fur- thermore, it was found that Fe-Si-B-Nb-Cu alloy with a Si/B ratio of around 1.7 showed

Magnetochemistry 2021, 7, 20 19 of 24

considerably low coercivity (1512 A/m) and higher magnetic saturation (143 Am2/kg) compared with the other FeSiBNbCu samples with different Si/B ratios. The effect of laser process parameters on the hardness of Fe-based BMGs produced by LENS was studied by Xie and his team [91], who used Fe41Co7Cr15Mo14C15B6Y2 (at.%) amorphous spherical powder and different laser energy densities. The variation in hardness Magnetochemistry 2021, 7, 20 20 of 25 as a function of laser energy density (LED) is illustrated in Figure 19. It is obvious that hardness increases as LED lowers. The use of the lowest energy density of 30 J/mm led to the highest hardness of 1263.7 HV. This value is signiﬁcantly close to the hardness of castof cast Fe 41FeCo41Co7Cr7Cr15Mo15Mo1414CC1515BB66Y2 amorphous alloy (1253 HV) [[92],92], whichwhich suggestssuggests thatthat thethe printed specimen retained aa certaincertain amountamount ofof itsits amorphousamorphous phase.phase.

Figure 19. Hardness of the LENS-processed Fe41Co7Cr15Mo14C15B6Y2 amorphous alloys as a func- Figure 19. Hardness of the LENS-processed Fe41Co7Cr15Mo14C15B6Y2 amorphous alloys as a function of laser energy density (LED) (A, B, C, D, E and F denote Fe41Co7Cr15Mo14C15B6Y2 samples tion of laser energy density (LED) (A, B, C, D, E and F denote Fe41Co7Cr15Mo14C15B6Y2 samples that that were produced using different laser parameters, i.e., laser energy density (LED). A and E have were produced using different laser parameters, i.e., laser energy density (LED). A and E have the thesame same LED; LED; however, however, for A, for p A,= 300p = W 300 and W v and = 5 vmm/s= 5 mm/s and for and E, p for = 600 E, p W= 600andW v = and 10 mm/s,v = 10 for mm/s, B: p for= 600 B: Wp = and 600 v W = and5 mm/s,v =5 for mm/s, C: p = for 900 C: Wp = and 900 v W = 5 and mm/s,v =5 for mm/s, D: p = for 300 D: Wp = and 300 v W = 10 and mm/sv = 10 and mm/s for andF: p = for 900 F: Wp = and 900 v W = 10 and mm/s)v = 10 (reprinted mm/s) (reprinted with the withpermission the permission from [91], from Springer, [91], Springer, 2019). 2019).

In general,general, thethe LENSLENS technique technique has has not not been been much much used used in thein the production production of magnetic of mag- amorphousnetic amorphous alloys alloys because because it has it lowerhas lower geometric geometric accuracy accuracy (0.25 (0.25 mm) mm) and and lower lower surface sur- quality,face quality, requiring requiring post-processing post-processing [93] and [93] can and only can fabricate only fabricate less complicated less complicated components com- comparedponents compared with SLM with [77]. Consequently,SLM [77]. Consequent LENS is commonlyly, LENS is utilised commonly for large utilised parts withfor large low complexityparts with low [34 ]complexity and for repairing [34] and them for repairing [94], which them restricts [94], which the fabrication restricts the of amorphous fabrication alloysof amorphous by LENS alloys due to by their LENS size due limitation. to their Allsize the limitation. drawbacks All aside,the drawbacks the beneﬁt aside, of using the LENSbenefit isof theusing ability LENS to is produce the ability parts to produc with superiore parts with mechanical superior mechanical properties, properties, controlled compositioncontrolled composition and microstructure and microstructure [34]. As explained [34]. As in explained this section, in this generating section, amorphous generating alloysamorphous with thealloys LENS with technique the LENS is possible technique by loweringis possible laser by energy lowering input laser (low energy laser power input and high scan speed similar to SLM) and reducing substrate/prior deposit temperature. (low laser power and high scan speed similar to SLM) and reducing substrate/prior de- Table3 gives a summary of the different materials and AM techniques described in posit temperature. this paper. It is observed that Fe-based magnetic alloys with various amorphous content Table 3 gives a summary of the different materials and AM techniques described in up to 100% can be produced by the SLM technique by changing process parameters. It this paper. It is observed that Fe-based magnetic alloys with various amorphous content can be said that even though both of the elemental compositions and powder sizes of the up to 100% can be produced by the SLM technique by changing process parameters. It can feedstock powder in [75,76] are quite similar to each other, the SLM processes alloys have be said that even though both of the elemental compositions and powder sizes of the feedstock powder in [75,76] are quite similar to each other, the SLM processes alloys have different amorphous phase fractions and magnetic properties due to the different process parameters and scanning strategy, as mentioned in Section 3.1. This means that laser process parameters and scanning strategy play critical roles in the production of amorphous alloys with laser additive manufacturing. From the perspective of magnetic properties, while the produced alloy in [8] possesses relatively low magnetic saturation and coercivity, the glassy alloy in Ref. [68] has relatively high saturation magnetisation and coercivity.

Magnetochemistry 2021, 7, 20 20 of 24

different amorphous phase fractions and magnetic properties due to the different process parameters and scanning strategy, as mentioned in Section 3.1. This means that laser process parameters and scanning strategy play critical roles in the production of amorphous alloys with laser additive manufacturing. From the perspective of magnetic properties, while the produced alloy in [8] possesses relatively low magnetic saturation and coercivity, the glassy alloy in Ref. [68] has relatively high saturation magnetisation and coercivity. This may due to the different alloy design and composition, process parameter or problems that occurred during the production, causing structural defects. It is indicated in Table3 that SLM-processed glassy alloy (Fe71Si10C6Cr2 (at.%)) has the best soft magnetic properties (that is, high saturation magnetisation and low coercivity) resulting from the double- scanning strategy (remelting) carried out during the SLM process. The last three rows in Table3 represent the Fe-based amorphous materials produced by the LENS technique. As stated before, the LENS technique has not been studied extensively in fabricating amorphous alloys due to the process-related problems. Considering the hardness values of the FeCrMoCB alloy produced by SLM [61] and the Fe41Co7Cr15Mo14C15B6Y2 (at.%) alloy fabricated by LENS [91], it can be said that LENS provides more mechanical strength than the SLM technique.

Table 3. Amorphous content, magnetic and mechanical properties of Fe-based alloys produced by laser additive manufacturing, also showing process parameters and powder size of starting powder.

Process Parameters Magnetic Properties

Alloy Powder Size (µm) Technique P (W) v (m/s) Amorphous Content (%) Ms Hc (A/m) Hardness (HV) Reference

Fe68.3C6.9 27.9 <150 SLM 340 1.5, 2.5 Unknown 102.8, 103.6 Am2/kg -[8] Si2.5B6.7P8.7Cr2.3Mo2.5Al2.1 31.6 FeCrMoCB <80 SLM 80–200 0.8–5 100 a - - 902 [61] Fe92.4Si3.1B4.5 < 30 SLM 90 0.1–1.5 70–80 188.6–199 Am2/kg 3485.5–5809.2 1654–2273 [68] Fe43.7Co7.3Cr14.7Mo12.6C15.5B4.3Y1.9 (at.%) < 33 SLM 150–350 0.2–1 4.96–100 - - - [71]

Fe71Si10C6Cr2 (at.%) ~24.5 SLM Max 120 - 89.6 ~150 Am2/kg 238 - [76] Fe73.7Si11B11C2Cr2.28 ~25 SLM 50, 70, 90 1.2, 1.6 47 1.22 T (kg/s2A) 1591.5–2387.3 - [75] Fe-Si-B-Nb-Cu <150 LENS - - - 118–150 Am2/kg 1273.2–3819.7 - [5] Cr<25Mo<15W<10C<3Mn<5Si<2B<5Febalance (at.%) <180 LENS 250 0.02 - - - 1421 ± 101 [80] 300, 600, 0.005, Fe Co Cr Mo C B Y (at.%) - LENS - - - 1263.7 [91] 41 7 15 14 15 6 2 900 0.01 a According to the XRD analysis, the part is fully amorphous.

4. Conclusions This paper provides a review of laser additive manufacturing found in the literature for Fe-based amorphous alloys. The literature demonstrates that laser additive manufacturing, which contains directed energy deposition and powder-bed fusion, is a promising process to produce Fe-based glassy alloys with good soft magnetic properties. The SLM technique from the powder-bed fusion category and the LENS technique from the directed energy deposition category have been used for that purpose within the literature, resulting from their high availability, low cost and being less time-consuming compared with the conventional techniques that have been used to produce amorphous alloys, such as casting. In both techniques, amorphous alloys were manufactured by altering the process parameters. Generally, low energy input, achieved via low laser power and high laser scanning speed, brings about increased amorphous phase content due to a higher cooling rate, but decreases relative density and deteriorates mechanical properties. To overcome this limitation, the double scanning (remelting) strategy has been employed, which results in enhancing both magnetic and mechanical properties. Additionally, in the LENS process, it is also necessary to consider the substrate/prior deposit temperature, to ensure that it is kept as low as possible to obtain a high cooling rate. SLM has been utilised more than LENS in the fabrication of Fe-based amorphous alloys, since LENS is not preferred for the production of small parts because of low geometric accuracy and poor surface quality. This contradicts with the size limitation of glassy alloys owing to their high critical cooling rates. Magnetochemistry 2021, 7, 20 21 of 24

Author Contributions: Writing—original draft preparation, M.G.O.; writing—review and editing, M.G.O. and N.A.M.; supervision, N.A.M.; project administration, N.A.M.; funding acquisition, N.A.M. All authors have read and agreed to the published version of the manuscript. Funding: This work was funded by the Ministry of National Education of Turkish Government, the Royal Society Mid-Career Leverhulme Trust Fellowship scheme (SRF\R1\180020) and the Lever- hulme Trust (RPG-2018-324). Data Availability Statement: No new data were created or analyzed in this study. Data sharing is not applicable to this article. Conﬂicts of Interest: The authors declare no conﬂict of interest.

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