Microstructural control of an Al–W aluminum matrix composite during direct laser metal deposition

Journal of Alloys and Compounds 813 (2020) 152208

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Microstructural control of an AleW aluminum matrix composite during direct laser metal deposition A. Ramakrishnan, G.P. Dinda* Department of Mechanical Engineering, Wayne State University, Detroit, MI, 48202, USA

a r t i c l e i n f o

a b s t r a c t

Article history: Received 14 May 2019 Received in revised form 27 August 2019 Accepted 8 September 2019 Available online 9 September 2019

This study focuses on the evolution of a design strategy for aluminum matrix composites (AMC) in additive manufacturing, using a mixture of 82 wt % Al and 18 wt % W (82Ale18W). Direct laser metal deposition was used to fabricate and assess the 82Ale18W AMC using three different scanning speeds (12, 6, and 1.5 mm/s). This AMC was examined by optical microscopy, scanning electron microscopy, energy dispersive X-ray spectroscopy, calculation of phase diagrams (CALPHAD), X-ray diffraction (XRD), and microhardness measurements. The hardness of the AMC increased by 50% compared to pure aluminum. The AMC showed remarkable densification during processing with uniform distribution of fine intermetallic phases and undissolved W reinforcement in the Al matrix. A notably fine microstructure was observed within the solidified melt pool, developing phases of supersaturated solid solutions of primary a-Al FCC elongated cellular/doublon microstructure in the near edge of the melt pool, and fine cellular/doublon structure in the top of the solidified melt pool. The formation mechanism of the intermetallic (Al12W, Al4W, Al5W) phases during rapid solidification from the melt, and the fine eutectic structure evolving at the outer solid-liquid boundaries of the intercellular/doublon structure that is due to the presence of Ni, and Fe impurities in W is well documented. This communication reports the influence of cooling rate on the solidification structures, stable, and metastable phases observed in the 82Al e18W AMC system activating future alloy design by LMD processing. © 2019 Elsevier B.V. All rights reserved.

Keywords: Aluminum matrix composite Additive manufacturing Tungsten (W) AleW intermetallics Microstructure control

1. Introduction Aluminum-based metal matrix composites (AMC) have been the subject matter of scientific investigation for decades. AMC have an attractive combination of low density, high specific strength and elastic modulus, improved stiffness, abrasion, oxidation resistance, wear resistance, workability, high temperature properties, and controlled thermal expansion coefficient. It is widely used in areas such as aerospace, automotive, electronic packaging, precision instruments and sporting equipments [1,2]. AleW alloys is used as radiation shields in both aerospace and medical applications, sports equipment (golf club heads) and automobile motor pistons because of its low weight, high hardness, excellent wear resistance and high strength [3e5]. Moreover, binary alloys with the combination of high electrically conductive metals and refractory transition metals seem to be promising materials according to their respective specific properties [6]. Tungsten as a reinforcement has a variety of

* Corresponding author. E-mail address: [email protected] (G.P. Dinda). https://doi.org/10.1016/j.jallcom.2019.152208 0925-8388/© 2019 Elsevier B.V. All rights reserved.

special applications that depend on its unique properties owing to the very large difference between the melting points of W (3410
C) and Al (660.45
C), and the density of W (19.25 g/cm3) and Al (2.70 g/cm3) [7]. The alloying of such incompatible elements with a homogeneous distribution of the reinforcement particles is challenging. The WeAl AMC system was chosen because they can produce several intermetallic phases such as Al12W, Al4W with retained W particles that inhibit localized corrosion such as pitting, and grain boundary attack [8e10]. Moreover, an increase in the amount of solute in solid solution during non-equilibrium cooling and the presence of intermetallic phases improves the corrosion resistance, thermal stability, and specific strength by several orders of magnitude [11,12]. Addition of transition metals as reinforcement into the Al matrix has drawn attention indicating a boost in properties of the AMC [1,13,14]. Furthermore, the inclusion of intermetallic particulate reinforcement in the Al matrix have shown signs of the significant improvement in property owing to the development of strong bonding between the Al matrix and the intermetallic phases [15,16]. Mechanical, wear, and corrosion properties are shown to

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always improve when Al is reinforced with W when compared with pure Al in various fabrication routes [17e20]. In the current study, a processing route is incorporated to manufacture quick complex shaped components with the outcome of both W and AleW intermetallic particulates acting as reinforcement in the Al matrix leading to the superior performance of the AMC. Feng et al. fabricated 97 vol% Al and 3 vol% W (Al97W3) composites by reaction sintering and observed all the W to dissolve and form the Al12W intermetallic produced due to diffusion between Al and W. Compared to pure Al, Al97W3 observed an increase in UTS by 84.5% with an elongation of 7.3% which suggests the composite possess excellent ductility [21]. Porosity, limited process control, and difficulty to achieve uniform size and shapes are the drawbacks during reaction sintering [22]. Rajamure et al. spray coated W on Al 1100 substrate followed by laser melting the surface of W and observed majority of the second phase particles consisted of Al4W along with Al matrix and unmelted W particles as the major alloying phases. The corrosion resistance of Al 1100 substrate revealed inferior corrosion resistance compared to the laser treated AleW composite, which resulted in increased corrosion resistance [19]. Spray coating has the disadvantage of void formation, oxidized particles, unmelted particles and rough surface followed by laser melting is time consuming and is limited to one layer at a time. Dixit et al. processed AleW composites that consisted up to 3.79 at.% W by friction stir processing technique and revealed the presence of homogeneous distribution of W reinforcements and the second phase Al12W particles in the Al matrix. A maximum average matrix hardness value of 65 HV, yield strength of 129 MPa and 21.4% elongation was measured with an increase in electrical conductivity with an IACS % higher than several Al alloys from 2000 to 7000 series [23]. The composite region and design complexity during friction stir processing is limited to the stir zone. Manufacturing processes for AMC at an industrial scale can be classified into two main groups that are solid state processes such as powder blending, diffusion bonding, and physical vapor deposition, and liquid state processes using stir casting, vacuum arc melting, infiltration process, spray deposition, etc. [17,24,25]. These processes limit the production ability to achieve quick complex shapes and components, the synthesis of transition metals is difficult with casting due to very low solid solubility in Al, and vapor deposition limits coating thickness (>50 mm). On the other hand, (AM) introduces several opportunities through its capacity to fabricate complex geometrical components and presents a selective design to control the microstructure by altering the thermal conditions. Nevertheless, AM demonstrates specific challenges, such as solidification cracking, compositional variations due to the presence of low boiling point elements, anisotropy in material property pertaining to textural outcome during processing, and other defects such as porosity, and lack of fusion. However, the drawbacks in AM can be manipulated by tuning the processing parameters, scanning pattern, postdeposition processing by proper heat-treatment cycles and hotisostatic pressing, and in some cases developing alloys suitable for AM. The LMD process is equipped to be a non-equilibrium method of producing AMC and increases the solid solubility of transition metals in Al. In this work, the deposits of AleW AMC (82Ale18W composition in wt. %) has been fabricated by LMD using three scanning speeds (12, 6, and 1.5 mm/s). The solidification mechanism of the 82Ale18W AMC system in integration with the non-equilibrium processing during LMD resulted in the presence of the Al matrix reinforced by W and AleW intermetallics. The AMC revealed a considerable increase in density and a homogeneous distribution of reinforcement particles in the Al matrix. The microstructure and intermetallic phases that evolved in the melt region of the LMD

specimens was assessed by scanning electron microscopy. The present work uses the calculation of phase diagrams (CALPHAD) by Thermocalc to analyze and identify the phases present in the 82Ale18W AMC systems and compares the finding with the experimental results. The differences in the morphology, size, and growth motif of the primary and intermetallic phases due to the variation in scanning speed were investigated.

2. Experimental procedure Laser metal deposition process starts with creating a CAD model. The CAD model is imported as an STL file into any slicing software in which the scanning raster and build height is specified and an output tool path is obtained. Typically, the build height is about 1/3rd to 1/4th of the laser beam diameter. A 6-axis robot (ABB IRB 1410 M2004) is used to navigate the tool path of the nozzle that is controlled by a robot controller (IRC 5 M-2004). The coaxial nozzle is mounted onto the robot, and a 2 KW diode laser (Laserline LDM 2000-40) is used as a source of heat. The laser beam with a 978 nm wavelength and circular spot size of 2 mm diameter was used in the current study. Inert gas (Ar) is used as a medium to transport the metal powder and provides a protective environment to prevent oxidation. The metal powders are delivered such that the powder stream converges at the same point with the focused laser beam as schematically shown in Fig. 1. Fig. 2a-b reveals the micrographs of the commercially available elemental gas atomized Al (Ampal Inc, 45e75 mm, 99.9% purity) and milled W (Technon, 30e120 mm, 99.5% purity, 0.5% of Ni, Fe) powders that are used to fabricate the 82Ale18W AMC. The elemental powders Al and W were fed into the powder hoppers 1 and 2, as shown in Fig. 1. A block of 20 mm  20 mm  4.8 mm was deposited on a rolled 6.35 mm thick Al 6061 substrate. Hatched scanning raster was employed alternating by 90
is schematically shown in Fig. 3b with a total of 8 layers of 0.6 mm/layer. The achieved average total height of the deposition measured is 4.75 mm. The laser power of 900 W, scanning speeds (V) and corresponding energy density (E) of 12 mm/s (68.2 J/mm3), 6 mm/s (136.4 J/mm3), and 1.5 (545.45 J/ mm3) mm/s was employed. Total of 5 g/min of powder was mixed in-situ during deposition from hopper 1 and 2. 4.1 g/min of Al and 0.9 g/min of W constituted the 82Ale18W wt.% AMC. Samples were prepared by following standard metallographic procedures for microstructure characterization. The Kellers reagent was used to

Fig. 1. Schematic of the LMD set up illustrating the deposition process from multiple powder hoppers.

A. Ramakrishnan, G.P. Dinda / Journal of Alloys and Compounds 813 (2020) 152208

Fig. 2. (a) Morphology of as-received gas-atomized Al powder particles, and (b) milled W powder particles.

etch the samples. Microstructural analysis of the as-deposited samples was conducted using scanning electron microscopy (JOEL-7600 FE SEM) with energy dispersive X-ray spectroscopy (EDS). For transmission electron microscopy (TEM) and selected area diffraction (SAD)patterns, 3 mm in diameter samples were mechanically punched out. Thereafter samples were thinned by ion milling with low accelerating voltage of 3.5 (kV) to minimize the beam effects on the microstructure. Crystalline phases were analyzed using a BRUKER D8 X-Ray diffraction (XRD) system. The XRD was performed using Cu Ka radiation at 40 kV and 40 mA.

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Diffraction patterns were collected from 25
to 100
with a sampling interval of 0.01
. Image analysis was performed using ImageJ software to measure the cellular spacing, and amount of W particles dissolved in the Al matrix to form intermetallic phases. Microhardness measurements were performed with a Vickers microhardness tester (Instron 9100) using 100 g load and dwell time of 10 s. Each of the hardness value reported resulted from an average of 5 indentations. Solidification sequence and intermetallics formed during LMD of 82Ale18W AMC experimentally were compared with the thermodynamic calculations employing the recognized CALPHAD method using Thermocalc software [26,27]. 3. Results and discussion 3.1. Optical microscopy In the current study for all the process parameters investigated, no cracking or significant porosity was observed. Fig. 3c shows the micrograph of the melt pool geometry during layer-by-layer deposition at 12 mm/s scan speed. The deposits cross-section of the final five layers processed at 12 mm/s reveals alternating patterned micrograph along the longitudinal (X-Z plane), and transverse (Y-Z plane) cross-sections. The bands observed are

Fig. 3. (a) As-deposited 82Ale18W block processed at 720 mm/min, (b) schematic of the deposition strategy, and (c) cross-sectional micrograph of the as-deposited 82-Al-18W AMC illustrating the layer 4 to the final deposited layer.

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circular straight tracks along the longitudinal section in odd layers. Whereas, it is parallel circular arc pattern boundaries along the transverse section in even layers. This clearly shows the deposition direction alternated by 90
after every layer to maintain some level of isotropy. The average measured height in the topmost layer of the deposit was 1.1 mm and the remaining layers were 0.59 mm. The top layer was not re-melted, and the bottom layers were remelted by about 50% during the addition of layers. The width of the beads in the transverse section in layers 8 and 6 was measured to be 1.2 mm. This is due to the overlapping of the 2 mm beam by re-melting 40% of the previous deposited track. The overall microstructure also reveals the distribution of the bright W particles in the Al matrix is intact and mostly observed in the bottom of each layer due to the enormous difference in the density between Al and W presents an extreme situation. Therefore, the influence of most of the W particles sinking to the bottom of the melt pool is observed. Moreover, the total vol. % of W added during deposition was 2.98% out of which 2.25% remained unmelted and the remaining 0.73% formed intermetallic phases. Optical micrographs in Fig. 4a-c reveal the cross-section of the as-deposited 82Ale18W AMC and compare the differences observed in the solidified melt pools processed using 12, 6, and 1.5 mm/s scanning speeds. As the scanning speed decreased from 12 mm/s to 1.5 mm/s, the width of the melt pool increased by approximately 5.3%. The three samples processed at different scanning speeds exhibit microstructural morphologies varying with specific location in the deposit. The microstructure just above the bead or layer boundaries appears to be elongated cellular structure. These elongated cellular structure transition to fine cellular structure in the middle of the solidified melt pool. In the top of the melt pool, extremely-fine cellular structures are observed in comparison to the middle and bottom of the solidified melt pool. Essentially from the bottom of the solidified melt pool to the top, the microstructural morphology transformed from elongated cellular structures to fine cellular structure in the top half of individual layers. Note that the bright regions are the cellular region, the blue regions are the interdendritic region, and a closer look in the interdendritic region reveals green particles that consist of the AleW intermetallic phases. The unmelted W particles peripheral structure also reveals a thin green film that consists of the AleW intermetallic phase. LMD is a typical high cooling rate process (103e105 K/s) [28], and the fine and coarse microstructure in the deposit can be controlled by altering the processing parameters. Solid-state diffusive transformation during LMD is usually repressed. However, with subsequent layers deposited the bottom layers undergo periodic heating and cooling cycles that triggers some level of solid-state diffusion. Henceforth, the microstructure and properties of the resulting deposit are the outcomes of solidification strategy. Fig. 5a shows the effect of the temperature gradient in liquid (GL) and the interfacial velocity or growth rate (R) on the development of various solidification microstructural morphology [29]. It also shows the microstructural refinement owing to increasing cooling rate (K ¼ GLR). Kurz et al. and coworkers have studied in detail the microstructure of laser processed materials [30,31]. The local solidification condition in the melt pool critically defines the solidification microstructure that is a function of R, and GL at the solid-liquid interface and sequentially rely on the heat and mass transfer in the framework. During LMD, a convoluted 3-dimensional profile of the melt pool is resolved by three distinct isotherms: The liquidus temperature of the substrate or the previously deposited layer ahead of the laser spot as the previous layer or the substrate is melted to establish a good metallurgical bond and the solidus temperature of the AMC. A straightforward relationship prevails between R and the scanning speed (V) as shown in equation (1), where q is the angle between R and V.

Fig. 4. Optical micrographs (a) melt pool morphology of the as-deposited 82Ale18W AMC processed at 720 mm/min, (b) melt pool morphology of the as-deposited 82Ale 18W AMC processed at 360 mm/min illustrating the fundamental principle direction of decreasing thermal gradient (G)and increasing growth rate during LMD, and (c) melt pool morphology of the as-deposited 82Ale18W AMC processed at 90 mm/min displaying the changes in morphology from the bottom layer to the top surface.

R ¼ V cos q /

(1)

R increases swiftly from zero at the bottom of the melt pool as cos q ¼ 0, therefore, R ¼ 0 and the value eventually increases as it approaches the top of the melt pool surface and the value remains nearly constant in the majority of time during solidification. Similarly, GL is highest in the near edge or bottom of the melt pool and reduces swiftly to a constant value as it approaches the top surface. GL/R value to a great degree controls the S/L interface morphology, and K influences the microstructures dimension. Fig. 5b shows the influence of cooling rate on size variation of primary doublons/ cellular spacing (lP) as a function of the distance from the layer boundary or bottom of the solidified melt pool to the top surface in the current study. Note that doublons are observed during optical microscopy and is discussed in detail below in the current study. Quantitative analysis of lP variation and statistics were carried out with image analysis software (ImageJ). The average value of the lP was obtained by measuring the distance between the nearest

A. Ramakrishnan, G.P. Dinda / Journal of Alloys and Compounds 813 (2020) 152208

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Fig. 5. (a) A graph representing the temperature gradient (GL) as a function of solidification rate and the resulting influence of GL/R on the microstructures morphology [26], and (b) change in average primary doublon/celluar spacing versus the distance from the layer boundary as measured in 82Ale18W system under three different processing conditions.

primary doublons/cellular structure tips from Fig. 4a-c by using equation (2) where L and n, respectively, are the length of the experiment line and the number of counted doublon/cellular structure.

lP¼L/n /

during 12 mm/s, 6 mm/s, and 1.5 mm/s were measured to be 65.96
, 74.22
, and 85.79
, respectively. Chen et al. revealed a critical solidification velocity exists where the evolution of cellular spacing change abruptly from a strained free structure to a highly strained structure during laser surface melting of FeeNi alloys [32].

(2)

The lP value approximately linearly reduced with increase in the distance from the near edge boundary. Concurrently, the cooling rate increases from zero at the bottom to a maximum value at the top surface of the layer. It is also observed that with higher energy density and increasing heat input under the influence of varying the processing parameter coarse structures (doublons/cellular structures and eventually crystals) form at the layer boundary and gradually refined structures form approaching the top surface. Fig. 5c shows the evolution of the average primary doublons/ cellular structure lP as a function of solidification velocity R. The solidified melt pool reveals different microstructural morphology depending on the orientation of the cut section as shown in Fig. 3c, which reveals both longitudinal and transverse section. Along the transverse section it is observed the microstructure translates from cellular to fine equiaxed structure and varies locally depending on the heat flow direction. The q values were measured from the longitudinal section as the cellular/doublon structure is viewed along the deposition direction. The average q values which is the angle between the scanning direction and the solidification velocity

3.2. SEM investigation Fig. 6a-b, 6(c-d), and 6(e-f) show a comparative microstructural feature of the 82Ale18W samples processed using 12, 6, and 1.5 mm/s samples at the near edge of the bottom layer boundary. Similarly, Fig. 7a-b, 7(c-d), and 7(e-f) reveal the microstructural morphologies of the samples processed using 12, 6, and 1.5 mm/s, respectively, in the middle, and the top surface of the layer. A clear microstructural anisotropy in the deposit is observed from the bottom layer boundary to the top surface in the deposit. This is mainly due to the local solidification condition and heat flow direction from the superheated melt into the cooler solid. From Fig. 6b, d, and 6f reveal directional solidification as the substrate or the previously deposited layer act as a heat sink. A typical feature of the laser processed material was observed in the solidified microstructure of the 82Ale18W AMC [33,34]. Solidification initiated by nucleation and epitaxial growth on the Al 6061 substrate and moves unilaterally depending on the melt pool geometry opposite to the heat flow direction from the bottom layer to the top surface. At the bottom of the melt pool due to the constitutional

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Fig. 6. 82Ale18W as-deposited morphology during LMD at the bottom layer of the solidified melt pool, (aeb) 720 mm/min, (ced) 360 mm/min, and (eef) 90 mm/min.

Fig. 7. 82Ale18W as-deposited morphology during LMD in the middle and top surface of the solidified melt pool.(aeb) 720 mm/min, (ced) 360 mm/min, and (eef) 90 mm/ min.

undercooling criterion of high GL/R a plane front solidification zone evolves as represented in Fig. 5a. The plane front gradually transforms into a cellular structure as GL/R decreases. At the bottom of the melt pool, columnar grains developed and grew parallel to the heat flux direction as observed in Fig. 6a. During scanning speed of 12 mm/s due to very high thermal gradient and low interfacial velocity at the solidification front planar growth is observed in Fig. 6b, respectively. As the G/R ratio is slightly less than needed for the stable planar front, therefore, cellular morphology with splitting tips represented as doublons can be observed close to the S/L interface as shown in Fig. 6b and f, respectively. Seaweed structures have been known since 1940 and it is identified as a new steady state growth structure known as doublons in two-dimension and triplon in three-dimension. Usually, the cellular structure transforms into dendrites when R increases and with further increase in R columnar dendrites are observed growing parallel to the heat extraction direction [30,35]. However, Fig. 6b and f reveal a stable seaweed and fractal seaweed structures. In the current study, only cellular structure and doublons are observed and they develop at faster R and become less responsive to GL [36]. Cellular structures have a <100> growth direction oriented nearly parallel to the heat flow direction and thus successfully presenting anisotropic grain growth orientation. In case of doublons they do not exhibit a clear growth direction such as in cellular structure and lack specific dendritic morphology and secondary arms along the grain growth direction [37]. However, seaweed or doublon structures have been studied in detail by Henry et al. and the classification of a wide range of growth morphologies that form as a function of undercooling and anisotropy is still under wide investigation [38]. All the scanning conditions reveal a transition from an elongated cellular/ doublon structure in the bottom layer boundary to fine cellular/ doublon structure in middle and the top surface of the layer. As discussed earlier, the solidification morphology depends on GL and R. Usually, high GL and low R produce elongated cellular structure observed in the bottom layer, and in the middle and top surface of the solidified melt pool low GL and high R resulted in fine cellular structure. These results observed in the current study revealed similar developments during the solidification of the melt pool, and the GL and R extensively depends on location in the melt pool. Moreover, the fine cellular microstructure of the Al cellular/doublons observed in the middle and top layer is fundamentally a cross-sectional micrograph of the elongated cellular/doublon microstructure of the bottom layer. This is due to the S/L interface gradually transforming from horizontal at the bottom of the melt pool to vertical at the top surface of the trailing edge of the melt pool. Effectively the heat flux direction progressively changes from a vertical direction at the bottom of the melt pool to a horizontal direction at the top surface of the melt pool. As a result, the primary elongated Al cellular/doublon microstructure changes from the vertical direction at the bottom of the melt pool to a horizontal direction at the top of the melt pool as shown in Figs. 4 and Fig. 6ab, respectively. Dinda et al. investigated similar microstructural transformation while processing Al 4047 alloy by direct metal deposition [39]. An important difference to be noticed is in the size of the fine cellular/doublon structure. The magnitude of refinement in the top surface compared to the bottom and the middle of the solidified melt pool during 12, 6, and 1.5 mm/s scanning speeds are observed from figure (6-7), respectively. Similarly, with increasing energy density in 6, and 1.5 mm/s samples the size of the fine cellular/doublon structure increased in size compared to 12 mm/s specimen. Note that in figure (6-7) the grey regions are the a-Al and the periphery around the bright W particles and the intercellular regions have several bright AleW intermetallic phases that formed during solidification. In order to study the morphology, composition, and size distribution of the intermetallic phases during LMD

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process under 12, 6, and 1.5 mm/s scanning speeds, a detailed investigation using the calculation of phase diagrams (CALPHAD), SEM, and EDS was carried out. 3.3. Analysis of 82Ale18W system by CALPHAD Computational solidification simulations in the present study were carried out using Thermocalc to produce a valid approximation of the phases formed and transformations calculated in the 82Ale18W system. Practically solidification is categorized into equilibrium condition as per the lever rule and non-equilibrium condition as per Scheil-Gulliver calculation. In both equilibrium and non-equilibrium condition, infinite diffusion in liquids is assumed as liquids have extreme convection and high diffusion. During the non-equilibrium condition, diffusion in solid state is assumed to be negligible. Whereas, equilibrium condition demand infinitely slow cooling, substantial diffusion in solid state, and, therefore, the respective proportion of liquid and solid at any given temperature can be resolved by phase diagrams. However, LMD is a rapid solidification process and the resulting microstructure is nonequilibrium in nature. Although the solidification behavior is nonequilibrium, repeated reheating and cooling of the deposit during deposition of successive layers gives rise to back-diffusion (solidstate diffusion) involving equilibrium condition to exist to some extent. Fig. 8a illustrates the phase fraction as a function of temperature during equilibrium simulation in the 82Ale18W system. The first phase to form during cooling in the liquid is the Al4W intermetallic that forms at 1197.33
C and exists until 846
C. At 846
C Al5W intermetallic phase forms and exists until 691.02
C. Moreover, Al4W and Al5W phases are known to melt peritectically,

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and requires several hours of heat-treatment to form these structures [40]. Al12W starts to form at 691.02
C in the liquid and exist even after the primary a-Al starts to solidify at 658.12
C. Al3Ni and Al9FeNi phases form due to the presence of 0.5% of Fe, and Ni impurity in the received W powder. Fig. 8b illustrates the isothermal transformation curves (TTT) of the 82Ale18W system that form Al4W, Al5W, and Al12W intermetallic phases. The TTT curves indicate an exothermic reaction takes place when a-Al melts at 660.32
C and the intermetallic phases Al4W, and Al5W occurs immediately depending on the heating rate. Moreover, with the holding temperature between 400 and 650
C and additional time indicate that the reaction enthalpy reduces and Al and W forms intermetallic phases even before the melting point of Al is reached by solid-state diffusion. Al12W is the first phase to form during heating, and subsequently as the temperature increases to 650
C with longer holding time Al4W also form the equilibrium intermetallic phase in the 82Ale18W system. Zhang et al. investigated on the influence of the rapid heating rate and holding time by differential scanning calorimeter resulted in the formation of equilibrium Al4W, and Al12W intermetallic phases well below the melting point of Al. Fig. 8c reveals the Scheil simulation which represents the non-equilibrium transformation of phases of the 82Ale18W AMC under the condition of infinite diffusion in liquid and no solid-state diffusion. Yang et al. studied the effect of Scheil simulation temperature step size and the limit of the residual liquid amount for AleMgeZn ternary system using CALPHAD (Pandat software) techniques [41]. The simulation results during numerical procedures may not be determined uniquely due to the presence of several numerical models. To understand the influence of temperature step (dT), which is how much the temperature will decrease at

Fig. 8. Thermodynamic simulation using Thermocalc (a) Equilibrium phase fraction as a function of temperature of the 82Ale18W system, and (b) time-temperaturetransformation curves of the Al4W, Al5W, and Al12W intermetallic phases, and (c) non-equilibrium Scheil simulation at a temperature step (dT) of 1 and FCC_L12 represents a-Al.

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Table 1 Scheil simulation results with different temperature step sizes (dT) for 82Ale18W AMC. Temperature Step Size (dT) (
C)

Temperature (
C) End of Liquidus

G(Liquid þ a-Al)  100

G(a-Al)  100

0.05 0.5 1 10

660.32 660.43 660.35 666.35

0.9995 0.158 0.158 Not-Present

Not-present 0.204 0.8 0.157

Table 2 Comparison of overall segregation of primary a-Al, AleW intermetallic phases, and eutectic phase in the as-deposited 82Ale18W system. Phase

Element (at %) Al

W

Ni

Fe

1e720 mm/min - a-Al 1e360 mm/min a-Al 1e90 mm/min a-Al 2 - Al4W 3 - Al5 W 4 - Al12W 5 - Eutectic

99.55 99.7 99.81 79.99 83.78 92.22 94.54

0.45 0.3 0.19 20.01 16.22 7.78 0.15

e e e e e e 2.72

e e e e e e 2.59

each step in the simulation during Scheil simulation employing Thermocalc four different temperature step of 0.05, 0.5, 1, and 10, respectively, were simulated. Table 1 reveals the influence of temperature step on the end liquidus temperature, fraction of liquid þ a-Al, and a-Al, respectively. During most of the minute variations of dT studied, change in fraction solid of formation occurred in the transformation of liquid þ a-Al, and a-Al phases. At a temperature step of 0
C the simulation runs for infinite time and a very high-temperature step eliminates some of the major phases that form. The results from Table 2 reveal at a very low dT of 0.05
C complete solid a-Al does not form. At a very high dT of 10
C liquid þ a-Al does not exist and a-Al forms as the liquid þ Al12W phase ends. Further increase in the temperature step eliminates the presence of liquid þ Al12W phase. Similarly, the default residual amount of liquid is 0.01 and several variations were tested until 0.000001 and no changes were observed for the current binary 82Al18W system. However, for complex alloy systems with more than two elements a larger cut-off limit of the residual liquid may end the simulation early and a smaller cut-off limit will produce several low melting point phases. Therefore, careful attention of dT and residual solid/liquid cut-off limit should be considered. To further understand the formation mechanism of the intermetallic phases during the LMD process due to rapid solidification high-resolution backscatter electron (BSE) microscopy was carried out. 3.4. Intermetallic phases AleW intermetallics amalgamate good strength, wear resistance, and thermal resistance of W and oxidation resistance, and low density of Al. The growth behavior and pattern of the intermetallic phases affect the mechanical properties of the material significantly. Therefore, the influence of three different scanning speeds of 12, 6, and 1.5 mm/s on the solidification morphology of the intermetallic phases, size, and shape changes are examined. At 12 mm/s scanning speed, a refined intermetallic morphology is observed and as the scanning speed reduces coarser intermetallic morphology is observed. Fig. 9a shows the BSE micrograph and one can observe a uniform distribution of the AleW light grey intermetallic phases around the W particle periphery, intercellular regions, and grain boundaries in the a-Al matrix. The bottom part of the solidified melt pool has most of the undissolved W particles. The W particle size in the middle of the melt pool is much smaller

Fig. 9. Morphology of the intermetallic phases developed during LMD in the asdeposited 82Ale18W AMC system when processed at 720 mm/min. Regions (1), (2), (3), (4), and (5) represent the composition of a-Al, Al4W, Al5W, Al12W, and the eutectic phase, respectively.

compared to the W particles in the near edge of the solidified melt pool that indicates a sign of more dissolution. Tonejc et al. and several others found supersaturated Al-rich AleW solid solutions can be extended up to 1.87 at.% W by rapid quenching the liquid at a rate of 107e108 K/s [42,43]. It is well known that rapidly cooling an alloy can result in significant extended (non-equilibrium) solid solubility. Indicated as region 1 in Fig. 9b shows the matrix composition of the supersaturated Al-rich solid solutions with a solid solubility of 0.4 at.% W while processing at 12 mm/s. The amount of solute W concentrations is far in excess of the equilibrium solid solubility due to rapid cooling producing nonequilibrium microstructure. Fig. 9b-c illustrates a clear interface morphology of the intermetallic phase around the periphery of the solid W particle that grew outwards into the liquid Al matrix forming a fine interface. Several different intermetallic shapes are observed such as hexagonal, faceted, oblong, dendritic, and petal morphologies. EDS investigation measurement from Table 2 confirmed the phases marked as 2 is the Al4W phase shown in Fig. 9c, respectively. There is two distinct morphology of the Al4W phase that was faceted, and dendritic shaped observed in Fig. 9c at the middle of the melt pool near to the W particle. The faceted

A. Ramakrishnan, G.P. Dinda / Journal of Alloys and Compounds 813 (2020) 152208

morphology has edges resembling secondary arms and the dendritic morphology with primary and secondary arms were joined to form a circular pattern of branches growing from a central ridge. Fig. 9d shows the morphology of the intermetallic phases in the top of the solidified melt pool. EDS confirmed the petal-shaped morphology with six to eight branches identified as region 3 is the Al5W phase. The size of the Al4W intermetallic was coarse near the middle and bottom of the melt pool in comparison to fine intermetallic phases in the top surface. In both Fig. 9c-d the dendrite and petal branches of the Al4W and Al5W are not equal in size due to the irregular partitioning of solutes and heat flow. In the liquid melt coarse W powder are partially dissolved, and the fine W powder is completely dissolved and mixed with the liquid matrix. In the event of rapid cooling and solidification from the liquid phase, the partially, and fully dissolved W nuclei provide heterogeneous nucleation sites resulting in precipitation of various shapes, faceted and non-faceted dendritic intermetallics depending on the local composition and temperature gradient. Bogels et al. revealed the growth process of faceted crystals is controlled by volume diffusion as the crystal morphology is sensitive to external growth environments and in the current study it depends on the location of the melt pool [44]. In the top of the melt pool, the corners and edges of the crystal grow more easily than the faces of the crystal as observed in the Al5W petal and consequently faster growth from the sides and corners are observed. Moreover, the top surface is remelted and heat-treated during subsequent deposition of layers that resulted in inter-diffusion taking place between a-Al matrix and the W particles forming Al5W, and Al12W intermetallics. From thermodynamic computation from Fig. 8a and EDS analysis from Table 2, it can be inferred that as the temperature decreases in melt pool Al4W is the first phase to form in the liquid phase at 1197.33
C. It provides nucleation sites as the temperature drops to 846
C and 691.02
C and as a result the Al5W and Al12W phase nucleates and grows. The low volume fraction of very fine Al12W phase about 10e100 nm was observed in the intercellular region during 12 mm/s scanning speed. Fig. 9e-f show the BSE micrograph of the eutectic structure in the deposit and the sample is not etched. Average eutectic spacing at 12 mm/s scan speed was found to be 160.50 nm. The eutectic reaction is observed due to the impurity of Fe and Ni of 0.5 wt % present in the W particle. From the ternary phase diagram of AleFeeNi at the eutectic temperature of 638
C the following reaction shown in equation (3) takes place [45,46]. L 4 a-Al þ Al3Ni þ Al9FeNi /

(3)

On cooling down the liquid at the composition range of 96.72 at.% Al, 0.105 at.% Fe and 3.174 at% Ni, the eutectic mixture consisting of a-Al solid solution and Al3Ni precipitates and/or FeNiAl9 and a-Al form the eutectic structures due to undercooling. In Fig. 9f a region denoted as 5 during EDS confirms the composition range of the eutectics observed and is shown in Table 2, respectively. Thermodynamic simulations in Fig. 8a illustrate the aAl, Al3Ni, and Al9FeNi form at the same time and confirm the reaction shown in equation (3), respectively. Fig. 10a reveals the overall micrograph of the bottom, middle, and the top layer of the melt pool regions during 6 mm/s scanning speed, respectively. Most of the undissolved W particles are seen settled in the bottom of the solidified melt pool. Similar to 12 mm/s scan speed there is no evidence of W particles in the top layer. Fig. 10b show higher dissolution of the W particles and bigger sized faceted intermetallic phases in 6 mm/s scanning speed compared to fine intermetallics observed during 12 mm/s scanning speeds surrounding the W particles. Fig. 10d shows the BSE micrograph of a non-etched sample revealing a distinct contrast between the bright (Al4W) and light grey (Al12W) intermetallic phases. EDS analysis

9

Fig. 10. Morphology of the intermetallic phases developed during LMD in the asdeposited 82Ale18W AMC system when processed at 360 mm/min. Regions (1), (2), (3), and (4) represent the composition of a-Al, Al4W, Al5W, and Al12W, respectively.

from Table 2 confirmed the composition of regions 2 and 4 indicated in Fig. 10(cee) as the Al4W and Al12W intermetallic phases, respectively. Some of the bigger Al4W phase boundaries are surrounded by the Al12W intermetallic phase. Moreover, a closer view of the surrounding W particles revealed the formation of Al4W and Al4W surrounded by the Al12W intermetallic phase, respectively, during 6 mm/s scanning speed. The presence of W enveloped by Al4W and Al4W enveloped by Al12W was not observed at higher cooling rates during processing at 12 mm/s scanning speeds. Moreover, higher dissolution of the W particles was measured, and the majority of the intermetallic phases belonged to Al4W and Al12W intermetallics, respectively. Al5W was also measured during EDS analysis. However, a very low volume fraction formed at 6 mm/ s scanning speed. The formation mechanism of the Al4W and Al12W phases proceeds in the beginning at a specific intersection between the solid W and the liquid Al. The contact area grew rapidly as Al melts easily at a much lower temperature compared to the W particle. Initially, the liquid Al atoms are activated and diffuse into the W matrix, and the W particles establish contact with the liquid Al. As the energy density and the input heat increases the W dissolves and establish contact with the Al atoms and an interface is created between Al and W. Furthermore, adhesion of the Al occurs on the solid surface of the W particles and creates an AleW intermetallic film on the surface of the adsorbent W particle driven by capillarity and gravity. Al4W is the first phase to form because of its low formation enthalpy, and Gibbs free energy as Al surpasses its solid-solubility limit in W. During processing at 6 mm/s more heat input is supplied and with a larger melt pool size, cooling is a lot slower compared to 12 mm/s deposits that give enough time for the Al4W interface to react with molten Al and form the Al12W intermetallic phase. Al12W forms at 691.02
C and then finally the a-Al

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phase solidifies along with eutectic reaction due to undercooling in the intercellular boundaries as shown in equation (3). Note that cooling rate plays a key role in the formation and size of the intermetallic phases. Fig. 10e reveal the three-dimensional morphology of the Al4W and Al12W intermetallic phase. In Fig. 10e the hexagonal plates indicated as 2 is the Al4W phase with average sides of 2e3 mm, and thickness of 200e300 nm, respectively. Several areas of the deposit that have oblong and needleshaped morphology was found to be the Al4W phase, however, the set of basal planes {0001}is observed in Fig. 10e, and f0110g is the side that appears to be needles. The particle indicated as 4 in Fig. 10e is the Al12W phase that has an icosahedral shape [10]. Fig. 10f reveals a cluster of Al4W and Al12W intermetallic phase in the top of the melt pool. Fig. 10f shows a high magnification image of the Al5W intermetallic phase with a particle size of 1.85 mm. However, a low volume fraction of the Al5W intermetallic phase was present during 6 mm/s scan speed. This shows most of the Al5W intermetallic phase transformed to Al12W intermetallic phase. Fig. 11a shows the intermetallic particles distributed in the top and near edge of the solidified melt pool during 1.5 mm/s scanning speed. The near edge of the solidified melt pool during processing at 1.5 mm/s has a coarse distribution of the intermetallic phases and most of the W particles are settled in the bottom of the layer. However, in the top surface below the layer boundary fine intermetallic phase is observed. Fig. 11b shows the topmost layer of the 1.5 mm/s deposit that shows a denuded region almost free of the intermetallic phases. This shows that most of the W particles settle at the bottom of the layer and due to large heat input the melt pool remains liquid for a longer period forcing the denser intermetallic

phases to settle down. Note that Al5W and Al12W have a calculated density of 5.71 and 3.88 g/cc, respectively [10,47]. Fig. 11c-d shows the near edge and middle of the solidified melt pool. Similar presence of Al4W and Al12W intermetallic phase indicated as 2 and 4 is observed. However, more dissolution of the W particles compared to 12, and 6 mm/s scanning speed is observed. Fig. 11e shows the intermetallic phases observed in the top surface and the region indicated as 2 and 3 shows the presence of Al4W and Al5W intermetallic phases. However, Al12W intermetallic phase was not observed in the top surface. Most of the Al12W phase formed in the middle of the solidified melt pool. Fig. 11f shows the eutectic phase present during 1.5 mm/s scanning speed with an average eutectic spacing of 327 nm. This shows that at lower R in 12 mm/s specimens finer eutectic spacing was observed, and with increasing the R in 1.5 mm/s specimen, the amplitude of these eutectic waves also increased with larger eutectic spacing. Similar observations were made during the study of AleCu eutectic alloy by laser melting [48]. 3.5. TEM analysis Fig. 12 shows the TEM micrographs of the as-deposited 82Ale18W AMC. Fig. 12a shows the bright-field (BF) TEM image of the Al4W particle diffused in the Al matrix during solidification. Several dislocations were observed in the Al matrix. Fig. 12b reveals the BF and SAD patterns of the fine eutectic structure observed during SEM investigation and confirmed the presence of Ni and Fe forming the eutectic structure with the supersaturated Al matrix. Fig. 12c shows the BF and SAD pattern of the Al12W intermetallic phase. Finally, Fig. 12d reveals the BF image of the Al5W intermetallic phase in the 82Ale18W AMC. Eutectic, Al4W and Al5W phases is not expected under equilibrium conditions. However, due to the high degree of undercooling eutectic phases (Al3Ni precipitates and/or FeNiAl9) are reported to form at 638
C, and thermal excursion during subsequent layers deposited results in the formation of high melting point Al4W and Al5W metastable phases, which are retained and precipitated in the deposit. 3.6. Phase analysis by X-ray diffraction

Fig. 11. Morphology of the intermetallic phases developed during LMD in the asdeposited 82Ale18W AMC system when processed at 90 mm/min. Regions (1), (2), (3), (4), and (5) represent the composition of a-Al, Al4W, Al5W, Al12W, and the eutectic phase, respectively.

Fig. 13 shows the XRD pattern of the laser deposited 82Ale18W samples obtained for the processing parameters of 12, 6, and 1.5 mm/s scanning speeds. The main phases found is the facecentered-cubic Al phase and the average diffraction peaks in the as-deposited specimens in all the samples did not have any significant variation and occurred at 2q ¼ 38.46
, 44.71
, 65.07
, 78.21
, and 82.41
that corresponds to the (111), (200), (220), (311), and (222)planes, respectively. Table 3 shows the XRD peak intensity characteristics of 82Ale18W AMC under different processing conditions. In the bottom layer of the 12 mm/s processed specimen significant peak intensities from the W lattices are observed. However, the top layer clearly shows a tremendous decrease in the intensity of the W peaks. This shows that W particles are mostly settled in the bottom layers than the top layers during the solidification process. Moreover, XRD pattern of the W peaks in both 6 mm/s, and 1.5 mm/s was low, however, an increase in the Al4W, Al5W, and Al12W is observed. The top layer of the 12 mm/s sample displayed smooth a-Al peaks, whereas, the remaining peaks under different processing conditions showed convoluted a-Al peaks indicating an overlap of the Al4W peak. The overlapping of the Al4W peak with the W (110) peak is clearly observed in 6, and 1.5 mm/s samples. Therefore, an indication of more dissolution of the W particles and consequently, higher presence of the Al4W intermetallic phase is validated. Al5W has weak intensity peaks and were detected in all the processing condition. Strong Al12W peaks are observed in all the samples that confirms the presence of this

A. Ramakrishnan, G.P. Dinda / Journal of Alloys and Compounds 813 (2020) 152208

11

Fig. 12. (a)Bright-fieldimageoftheAl4WphaseintheAl-matrixalongwiththeselectedareadiffraction(SAD)patternoftheAlmatrix,(b)brightfield(BF)andSADpatternoftheeutecticstructureinthe82Al-18WMMCalongwiththeEDScompositiondata,(c)BFandSADpatternofAl12WobtainedfromthematrixandWparticlediffusedzone,and(f)brightfield(BF)andSADpatternoftheAl5Wphaseinthe82Al-18WAMCalongwiththeEDScompositiondata.

Fig. 13. XRD profiles of the as-deposited 82Ale18W AMC under 720 mm/min, 360 mm/min, and 90 mm/min processing conditions.

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Table 3 XRD examination of samples under different processing conditions revealing its peak characteristics. Sample Type

720 mm/min (bottom layer) 720 mm/min (top layer) 360 mm/min 90 mm/min

Peak intensity ratio (%) [(I/Imax)(hkl)] (111)

(200)

(220)

(311)

(222)

100 100 100 100

31.59 30.72 16.60 25.15

12.03 15.36 44.86 6.15

16.61 13.58 9.27 10.19

7.54 9.71 11.15 4.34

Fig. 14. Micro-hardnessmeasurementscarriedoutfortheas-deposited82Al-18Wsystem duringLMDinthreedifferentprocessingconditionsfromthebottomlayertothetopmostlay erinthedeposit.

intermetallic phase during the LMD process.

samples that contained an area of the denuded zone. In 6 and 1.5 mm/s specimens the top layers during BSE study from Fig. 11b revealed a large area of the denuded region. The lack of intermetallic phases present in the top surface and the lower supersaturated matrix is responsible for the low hardness value in the topmost layer of the deposit. 4. Conclusion The motivation of this study was to confirm an evaluation method to approximate the solidification microstructure, and intermetallic phases developed in the novel 82Ale18W system during LMD by varying the processing conditions. 82Ale18W AMC with a composition of the low melting point of Al and the dispersion of high melting point W was successfully developed using the LMD process. The laser power was 900 W, and the scanning speeds of 12, 6, and 1.5 mm/s were used. The LMD processed 82Ale18W AMC revealed fine microstructure during 12 mm/s scan speeds in comparison to 6 mm/s, and 1.5 mm/s scan speeds due to the high growth rates attained in higher scanning speeds. Similarly, very fine AleW intermetallic phases were measured in the 12 mm/s scan speeds, and these intermetallic phases gained sufficient time to grow in the melt pool with increasing energy density in 6, and 1.5 mm/s scanning speeds due to slower cooling rates and solidstate diffusion during reheating, and cooling of the previously deposited layers. During laser melting of the AleW powders, high cooling rates are achieved that lead to rapid solidification and a steady interface at an atomic level is achieved between the a-Al matrix, W, and the intermetallic phases. During processing at 12 mm/s majority of the Al4W intermetallic phase formed with variation in size at different location of the solidified melt pool and formed in the interface of the W particles, intercellular regions, and grain boundaries. During slower scanning speeds, the major intermetallic phase of Al4W, and a well-defined Al12W intermetallic phase was observed. The Al matrix in the AMC also exhibited higher strength compared to pure Al.

3.7. Micro-hardness measurements Fig. 14 shows the Vickers micro-hardness plot as a function of distance from the bottom layer to the top layer of the deposit. Fig. 13 reveals an increase in matrix hardness in all the processing parameters compared to pure Al that has a hardness of 30.6 HV [20]. The undissolved W reinforcement embedded in the a-Al matrix measured the hardness of 451.27 HV. A uniform trend is observed in all three processing conditions. Near the heat-affected and dilution zone a high hardness is measured that reduces with subsequent deposition of layers. A high hardness value in the dilution zone of the deposit is due to the alloying of the deposit with Al 6061 substrate that has an average hardness of 148 HV. The as-deposited 12 mm/s scanning speed sample has the highest hardness compared to 6 mm/s and 1.5 mm/s specimens. The high degree of supersaturated solid solution of W atoms in the a-Al matrix, fine grains developed during higher cooling rates, the very fine intermetallic phases of the intercellular regions, and grain boundaries are namely the advantages. The fine intermetallic phases pin the grain boundaries and act as obstacles to retard dislocation motion resulting in difficult flow during plastic deformation. However, the 6 and 1.5 mm/s scanning speeds show a gradual decrease in hardness that vastly indicates the effect of lower cooling rates. During slow cooling, a lower degree of supersaturated solid solution, coarser cellular structures, grains and intermetallic phases have been observed during microstructural evaluation and the resulting hardness in 6 and 1.5 mm/s measured is lower compared to 12 mm/s scanning speed. The topmost layer of the deposit measured a small decrease in hardness in 12 mm/s

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