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Enhancing hardness of Inconel 718 deposits using the aging effects of cold metal transfer-based additive manufacturing
Materials Science & Engineering A 776 (2020) 139005
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Materials Science & Engineering A journal homepage: http://www.elsevier.com/locate/msea
Enhancing hardness of Inconel 718 deposits using the aging effects of cold metal transfer-based additive manufacturing Donghyun Van a, G.P. Dinda b, Jaewoong Park a, Jyoti Mazumder c, Seung Hwan Lee a, * a
School of Aerospace and Mechanical Engineering, Korea Aerospace University, 76 Hanggongdaehang-ro, Deokyang-gu, Goyang-si, Gyeonggi-do, 10540, Republic of Korea b Department of Mechanical Engineering, Wayne State University, Detroit, MI, 48202, USA c Center for Laser and Plasmas for Advanced Manufacturing, University of Michigan, 2041 G. G. Brown, 2350 Hayward Street, Ann Arbor, MI, 48109-2125, USA
A R T I C L E I N F O
A B S T R A C T
Keywords: WAAM CMT Inconel 718 Strengthening phase Hardness
The effects of successive deposition using continuous heat from a plasma arc on a multi-layer Inconel 718 deposit was investigated using a wire-arc additive manufacturing (WAAM) process with cold metal transfer (CMT). Deposits consisting of 10 layers were manufactured, where interpass time was the only process parameter varied. The shape of the deposits was determined by the interpass time, while microstructural characteristics (growth of dendrite, segregation of chemical component, formation of various phases, etc.) developed due to successive deposition. An aging effect was observed, which caused precipitation of the γ00 and γ0 strengthening phases and the δ phase. The highest hardness value was measured in the middle section of each deposit due to precipitation of the strengthening phases. The cooling rate was calculated using the thermal profile of the deposition process and was found to be slower using the CMT-based WAAM process than that of other additive manufacturing (AM) processes that use laser or electron beams. The results presented here also demonstrated that in-situ heat treatment during deposition is possible using this CMT-based WAAM approach.
1. Introduction Additive manufacturing (AM) technology is an easy means of manufacturing complex three-dimensional designs and utilizes less material than traditional manufacturing methods. These advantages led to the use of AM in the production of metal products, where recent applications include expensive superalloy products for use in the fields of aerospace, energy, defense and medicine [1–5]. Many studies have recently focused on the application of AM to the alloy Inconel 718. Dinda et al. [6] investigated microstructural morphology of deposits produced using a CO₂ laser beam and metal powders, and demonstrated that the microstructure and laser beam scanning pattern were correlated. Furthermore, changes in the hardness were related to the aging temperature difference. Zhu et al. [7] studied the changes and the strength of the deposit’s microstructure resulting from manufacture using a fiber laser and metal powder. This study related the grain size, dendrites and Laves phases to the laser power and beam diameter. A tensile test was conducted on the deposits produced under various process conditions. Chen et al. [8] identified the process parameters that affect the cooling rate of deposit using a disc laser and
metal powder. The microstructure and the hardness underwent changes due to the separation of chemical components. Empirical formulae related to process parameters were applied to evaluate shape-related formation qualities, including deposit width and height, cooling rate and hardness. Sames et al. [9] produced Inconel 718 deposits using electron-beam melting. The effects of solution treatments, aging and in-situ heat treatment on the mechanical properties and microstructures of the deposits were analyzed. Kumara et al. [10] studied the homoge nization effect of in-situ heat treatments during electron-beam AM processing. Overall, heat treatments, aging effects and precipitate be haviors of the Inconel 718 alloy deposits produced using metal powder with laser or electron beams have been widely studied [9,11]. However, AM based on laser and electron beams have limitations for the commercial implementation of the production of large deposits, as initial investment cost is high and a relatively low deposition rate is generally expected. The wire-arc additive manufacturing (WAAM) process [12] has been developed to address these issues. The method utilizes plasma arcs and wires, where the plasma arc is typically generated using a conventional welding machine and the consumable electrodes produced for welding are used as the feedstock wires. Only a
* Corresponding author. E-mail address: [email protected] (S.H. Lee). https://doi.org/10.1016/j.msea.2020.139005 Received 7 January 2020; Received in revised form 21 January 2020; Accepted 23 January 2020 Available online 25 January 2020 0921-5093/© 2020 Elsevier B.V. All rights reserved.
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few studies have focused on the WAAM process, but have demonstrated promise for the production of simple designs using a tungsten inert gas (TIG) welding machine and wire. Xiangfang et al. [13] produced Inconel 718 components using the WAAM process and interpass rolling, with TIG as the WAMM heat source. The Inconel 718 components produced by the WAAM process alone were compared with components produced using both WAAM and interpass rolling. The effects of both methods on the mechanical properties and microstructure of the components was reported. Asala et al. [14] manufactured deposits using WAAM with a TIG heat source and ATI 718 Plus feedstock material. The study iden tified a deposited metal heat affected zone (DMHAZ), which is a softened zone created during the deposition process. The microstructures and strengthening phases, including the γ00 and γ0 phases of the Inconel 718 deposits, were also investigated. An issue related to the use of a conventional TIG heat source is that the heat input cannot be as precisely controlled as a laser heat source. This can result in a non-uniform melt pool and unstable deposition conditions, which contribute to defects (e.g. liquation cracking) that deteriorate the mechanical properties of the deposit [15,16]. Control of the arc heat source is essential to addressing this issue and ensuring the desired quality of the products. Cold metal transfer (CMT) [17] provides controllable heat input using a high speed digital control and a high frequency wire feeding system. CMT may be implemented in the WAAM process to provide accurate heat input control. Unlike the AM processes that use a metal powder, applications of the CMT process using the Inconel 718 alloy are very limited. Most of the research involving CMT has focused on the welding process and materials other than the Inconel 718 alloy. In-situ heat treatment using the Inconel 718 alloy and the CMT method has not yet been investigated. Before CMT is applied in the WAAM process, the microstructure characteristics of the deposits after accumulated heat input during solidification must be studied. This study aimed to investigate the changes in microstructure in Inconel 718 multi-layer deposits after successive deposition and continuous heat input by CMT. Ten thin layers were produced with various interpass times as a process parameter. The thermal profile of the deposition process was measured. The effects of the interpass time on the shape of the deposit and the microstructural changes per layer were investigated. Furthermore, the cooling rates of the precipitations, the corresponding microstructures and the behavior of the precipitate were evaluated according to thermal profiles measured at a fixed point on the deposit.
XY stage and Z stage, respectively. The scanning and building directions of the deposit are shown in Fig. 1b. 100% Ar gas with a flow rate of 20 ℓ/min was applied to shield the plasma arc and the deposit and the contact tip to workspace distance (CTWD) was fixed at 10 mm. The experimental conditions are listed in Table 1. The temperature was measured in real-time by using a pyrometer (OPTRIS, CTlaser 3MH3) positioned at the center of the deposit 4 mm from the surface in the z-direction. Different emissivity values were applied for the various materials when measuring temperatures using the pyrometer. A preliminary experiment was conducted to compare the K-type thermocouple temperature with the pyrometer temperature to confirm the surface emissivity of the Inconel 718 deposit. The surface emissivity of the pyrometer was 0.92. Thin walled samples comprising 10 layers were deposited using three interpass times (0, 30 and 60 s), which means that the deposit was aircooled for either 0, 30 and 60 s between layers. Deposition was per formed in one direction along the X axis. When the interpass times are 0, 30, and 60 s, the corresponding deposition rates are about 0.6, 1.7 and 3.5 kg/h, respectively. Deposits were cut along the center of the YZ plane to reveal the sample’s microstructure. The test samples were mounted and polished using standard metallographic methods. The microstructure of the Inconel 718 deposit was observed using a field emission scanning elec tron microscope (FE-SEM, Hitachi, S-4800) and an optical microscope (OM, Olympus, BX51RF) after corroding the surface with Keller’s etchant (190 ml water, 5 ml nitric acid, 3 ml hydrochloric acid and 2 ml hydrofluoric acid). A FE-SEM with an energy dispersive X-ray spec trometer (EDS) was used to analyze the chemical compositions of the samples. Hardness measurements were performed from the bottom to the top of the deposit using a Vickers hardness tester (Matsuzawa, PMTX7) by pressing the test sample for a holding time of 10 s with a force of 1000 gf. The hardness value was calculated as the mean value of three indentations. Lattice parameters of test samples were examined using an X-ray diffraction (XRD) system (Malvern Panalytical, X’Pert PRO MPD). The XRD measurements were conducted on the polished surface of the deposit samples. Cu Kα radiation with a wavelength of 1.5406 Å was Table 1 Experimental parameters for the wire-arc additive manufacturing (WAAM) process with cold metal transfer (CMT). Parameter Wire (diameter) Current/voltage Wire feed speed Welding speed CTWD Shielding gas Welding direction Length (layer) Substrate Interpass time
2. Materials and methods The deposition material was Inconel 718 alloy, and AH36 was chosen the substrate due to its low cost. The deposit was produced using a 1.2 mm diameter Inconel 718 wire on a substrate with the dimensions 250 x 80 x 15 mm. A CMT machine (FRONIUS, TPS 4000) was used during deposition, and a XYZ 3-axis automatic stage (Six Degrees Inc.) with a 1μm resolution was used for transition of the sample and welding torch (Fig. 1a). The deposition sample and the welding torch were fixed to the
Fig. 1. Schematic of (a) the experiment, (b) the deposit showing the scanning and building direction. 2
Inconel 718 (1.2 mm) 120 A/13.4 V 4.8 m/min 0.5 m/min 10 mm Ar 100%, 20 ℓ/min One direction 100 mm (10 layers) AH36 (250 x 80 x 15 mm) 0 s, 30 s, 60 s
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applied and measured in the range 30–100� at intervals of 0.09� with an acquisition time of 2 s per increment. The XRD patterns were analyzed using the HighScope Plus v3.0 software package. Phase analysis of mi crostructures was conducted using an electron backscatter diffraction (EBSD) detector (Bruker, CrystAlign 200).
near top sections. The top of the 10th layer was not subjected to the heat associated with depositing the next layer and exhibited a faster cooling rate and smaller thermal gradient. This resulted in the formation of an equiaxed dendrite structure in the uppermost layer of the deposits. When the first layer was deposited to form the base of the deposit, no heat had yet accumulated. Thus, this first layer had a faster cooling rate than the deposits in the middle and top of the deposit. This was asso ciated to the formation of fine primary dendrite structures in the bottom section. As deposition proceeded, the cooling rate slowed down due to continuous arc heat accumulation. The spacing between the primary dendrites formed during deposition of the middle and top sections became greater than the spacing between the primary dendrites at the bottom. A secondary dendrite arm began to appear, and a welldeveloped secondary dendrite arm was observed in the near top part of the deposit due to the low cooling rate. SEM-EDS analysis was conducted at the dendritic and inter-dendritic regions in the middle section of the deposit produced using an interpass time of 30 s (Fig. 4). The weight percentage (wt%) of each chemical component was investigated at points 1, 2, 3 and 4 indicated in Fig. 4. Point 1 was taken during SEM-EDS analysis of the dendritic region (Fig. 4a), where the wt% of Nb and Mo was substantially lower than at the other points with values of 2.51 and 2.08%, respectively. At point 2, taken in the inter-dendritic region, this increased to 6.30 and 3.41% for Nb and Mo, respectively. This observation was associated with the segregation of chemical components, resulting in a higher concentration of Nb and Mo in the inter-dendritic region than the dendritic region. Elements such as Nb, Mo and Ti tend to form segregation zones during the last stages of solidification in the boundary and inter-dendritic re gions [10,18]. Nb has a particularly small diffusion coefficient and tends to remain in the liquid phase until the final stage of solidification, allowing for easy segregation. From the weight percent of elements listed in Table 2, the phases of the point 3 and 4 were further investigated to analyze the precipitates and segregation due to constitutional liquation. The Nb and Mo contents at points 3 and 4 are higher than at point 1, but the content of Ni, Fe and Cr had decreased. Laves phase is a topologically closed-packed (TCP) phase with a composition of (Ni,Fe,Cr)₂(Nb,Ti,Mo) that contains more Cr and Nb than the γ matrix. Antonson et al. [19] found that the Laves phase is formed when the wt% of Nb in the inter-dendritic region ex ceeds 20. In addition, it is known that SEM microstructures of a nickel-based superalloy such as Inconel 718 displayed lamellae in the Laves phase [20]. As the wt% of Nb was 26.43 and the precipitate had a lamellar form, the phase of point 3 shown in Fig. 4c is regarded as Laves. Fig. 4d indicates that the precipitate at point 4 had a polygonal form
3. Results 3.1. Macrostructural features The Inconel 718 alloy was deposited in ten layers using an interpass time of either 0, 30 or 60 s, and an OM macrostructure photograph was taken of the YZ plane cut along the center in the X direction (Fig. 2). The grain growth and remelting zone in the Z direction was identified in the macrostructure photograph of the YZ plane. The remelting zone was formed when the upper part of the solidified air-cooled layer was remelted by the arc heat applied as the following layer was deposited. As deposition proceeds from the bottom to the top direction (Z direction), the arc heat continues to accumulate in the deposit. Coarse grains were observed below ca. 3 mm from the substrate surface due to this heat accumulation. Interpass time was the only process parameter that was varied during depositing. As the interpass time was increased, the height of the deposit tended to increase while the maximum width decreased (Fig. 2). In addition, the inclination angle increased and approached 90� with longer interpass time. These changes were caused by the interlayer temperature control effect, where a shorter interpass time resulted in less air-cooling and the heat accumulation from previous depositions resulted in a higher preheated temperature for the following layer. Increased preheating temperature resulted in a higher temperature and decreased viscosity of the generated melt pool, which increased the velocity of the fluid. This caused the deposit with the shortest interpass time to be wider and shorter than the others. 3.2. Microstructural features The SEM micrographs of the deposits revealed that a columnar dendrite had grown from the bottom to the top in all layers except the top of the 10th layer (Fig. 3). Solidification of the deposit was initiated with the formation of a primary dendrite. The grain grew in the opposite direction of the heat flow, and crystal growth proceeded in the Z di rection when deposited from the bottom to the near top point. The columnar dendrite at the top of the 10th layer was observed to grow in a different direction than the crystal growth in the bottom, middle and
Fig. 2. OM macrographs of the Inconel 718 deposit samples produced using interpass times of (a) 0 s, (b) 30 s and (c) 60 s. 3
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Fig. 3. SEM micrographs of the Inconel 718 deposit deposits produced using interpass times of 0, 30 and 60 s.
on these weight percentages and geometrical shape, the main pre cipitates observed in the inter-dendritic region in the deposited SEM microstructures given in Figs. 4c and 4d were the Laves phase and MC-type carbide, respectively. Differences in the shape of the Laves phase was observed along the height of the deposit produced using an interpass time of 30 s were observed (Fig. 5). The SEM microstructure of the Inconel 718 exhibited lamellar morphology in the Laves phase. Discrete Laves phases were observed in the bottom sections of the deposit, which becomes more continuous in shape towards the middle section. This was attributed to the change in cooling rate, which was highest at the bottom of the de posit. The cooling rate at the very top of the deposit was also high due to convective air-cooling, thus a similar discrete form was observed. The Laves shape changed from a discrete to a lamella shape from the bottom to the top of the deposit and changed back to discrete at the very top. 3.3. Hardness measurement results The Vickers hardness test found that the hardness of the deposits formed using interpass times of 0, 30 and 60 s varied from 250 to 306 HV (Fig. 6). A previous study found that the Vickers hardness of a deposit produced using laser beams and metal powder was ca. 240 to 250 HV, and thus lower than those in the current study [21]. The maximum hardness was measured at the bottom of the deposits in the previous study, and decreased along the height of the deposit [8,21]. In this study, when the interpass time was 0, 30, and 60 s, the maximum hardness was
Fig. 4. SEM-EDS analysis points of the (a) dendritic region, (b) inter-dendritic region, (c) Laves phase in the inter-dendritic region, and (d) MC carbide in the inter-dendritic region.
which the MC-type carbide is [20]. The wt% of Nb at point 4 was 25.48, the second highest chemical component after Ni. In addition, the wt% of C is 3.33, which was considerably higher than at any other point. Based Table 2 Weight percent of elements at points 1, 2, 3 and 4 (see Fig. 4). Position
Ni
Fe
Cr
Nb
Mo
Ti
C
Ta
Others
Remark
Point 1 Point 2 Point 3 Point 4
53.69 52.81 38.76 36.69
21.02 18.08 12.89 13.27
18.17 17.26 12.03 11.44
2.51 6.30 26.43 25.48
2.80 3.41 7.74 6.98
0.89 1.18 1.01 1.08
– – – 3.33
– – – 1.49
0.92 0.96 1.14 0.24
Wt%
4
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Fig. 5. Shape changes of the Laves phase depending on the height in the z-direction of deposit (interpass time ¼ 30 s).
Fig. 6. Vickers hardness measurement of the Inconel 718 deposit produced using an interpass time of (a) 0 s, (b) 30 s, and (c) 60 s.
291, 306, and 300 HV, respectively, which were within the range of measurement error. An increase of hardness value from the bottom was expected for the deposits in the current study due to precipitation of the strengthening phases such as the γ00 and γ0 phases. γ00 is a strengthening phase of the Inconel 718 alloy, and is a disc-shaped precipitate with a chemical composition of Ni3Nb and has a DO₂₂ ordered body-centered tetragonal (BCT) crystal structure. γ0 is a spherical fine precipitate with a chemical composition of Ni₃(Al,Ti) that also acts as a strength ening phase, and has a L1₂ ordered FCC crystal structure [14,22]. The increased hardness in each sample may be attributed to precipitates such as γ00 and γ0 formed due to aging through successive deposition because Inconel 718 is a typical precipitation hardening alloy and strengthening is usually achieved through post-heat treatment, tempering and aging. The details of the aging effect related with the profile of hardness will be discussed in section 3.5 and 4. After passing the maximum hardness, the decrease of the hardness values was observed from 4 mm to 14 mm in Fig. 6a, from 8 mm and 14 mm in Fig. 6b and from 8 mm and 14 mm in
Fig. 6c. This is caused by the coarse grain induced by the low cooling rate during the solidification. The low cooling rate was caused by the heat accumulation during the multi-layer depositions. The increase in hardness values at the top of the three samples was attributed to convective heat transfer that allowed for rapid cooling and quenching. 3.4. XRD analysis results XRD analysis was conducted at the positions where the maximum hardness was measured in each deposit (Fig. 7). Peaks were mainly attributed to the γ-matrix of the deposits. It is difficult to confirm the precipitation of the strengthening phases γ00 and γ0 using XRD without prolonged annealing or coarsening of the precipitates. However, the precipitation can be indirectly confirmed according to the reduced lat tice parameter [23]. The lattice parameters of the deposits produced in this study were lower than the lattice parameter of the γ-matrix of Inconel 718 (360.05 p.m.) (Table 3) [11]. The lattice parameter of the 5
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Fig. 7. The XRD patterns of the Inconel 718 deposits produced using interpass times of 0, 30 and 60 s.
sample produced using an interpass time 30 s was the lowest, and is indicative of precipitation [23]. These results were consistent with the hardness measurement results in Fig. 6. Strengthening phases such as the γ00 and γ0 dissolved in the γ-matrix that was partially formed and precipitated during deposition.
Fig. 8. (a) EBSD orientation map, (b) EBSD phase map of the Inconel 718 deposit sample (interpass time ¼ 30 s) in middle part, which shows precipita tion morphology (8 mm from the substrate surface).
3.5. EBSD analysis results
Table 4 Phase area fraction in the bottom, middle and top sections of an Inconel 718 deposit (interpass time ¼ 30 s).
EBSD analysis was performed to examine the distribution of the microstructure and precipitation in the deposits [24]. The EBSD orien tation map and the EBSD phase map ca. 8 mm from the substrate surface (z-direction) of the deposit produced using an interpass time of 30 s was evaluated, as this point exhibited the maximum hardness (Fig. 8). The EBSD orientation map was shaded according to the main growth di rection of the columnar grains, with red shading indicating the z-di rection and purple shading indicating precipitates. The growth direction of the precipitation phase was different to that of the γ-phase (Fig. 8a). The EBSD phase map analysis of the distribution of precipitates such as γ00 , γ0 and δ revealed that the area fraction of the strengthening phase γ00 was larger in the middle section (8 mm from the surface) than the bot tom and top sections (Fig. 8b). As shown in Table 4, the area fraction of the strengthening phases (sum of γ00 and γ0 fraction) in the middle section was larger than the other areas. The previous hardness profile of the deposit is related to the content of strengthening phases in the micro structure, and these findings confirmed the hardness findings. As shown in Table 4, the area fraction of the δ phase is listed. In the Inconel 718 alloy, the δ phase is known to precipitate at the interface of the existing Laves phase or NbC in the temperature range from 850 to 1000 � C [25]. In addition, the precipitation of δ phase is also dependent on the residual strain and the amount of formed γ00 precipitates [26,27]. These are the reasons why the small area fraction of the δ phase might exist compared with γ00 precipitates in Table 4.
Phase
γ(Ni–Cr–Fe) γ’(Ni₃(Al, Ti)) γ’’(Ni₃Nb) δ(Ni₃Nb)
Interpass time (s)
Lattice parameter (pm) 359.75 359.22 359.69
Bottom (2 mm from substrate surface)
Middle (8 mm from substrate surface)
Top (14 mm from substrate surface)
90.44 8.49
89.30 8.52
97.22 1.89
0.90 0.17
1.91 0.27
0.61 0.28
4. Discussion The temperature profiles of the deposition process were greatly varied due to the layer by layer nature of the AM method (Fig. 9a). The cooling rates were calculated using the temperature profile of the de posit measured at a fixed point using a pyrometer, and ranged 5.4 to 50.0 � C/s. The minimum value of the cooling rate was observed when an interpass time of 0 s was used, and the maximum value was associated with an interpass time of 60 s. The cooling rate was 16.3 to 20.7 � C/s using an interpass time of 30 s. As shown in the thermal profile of Fig. 9a, the deposited area which stay at the aging temperature band between 600 and 1000 � C produces strengthening phases γ’’ and γ’ [18]. The maximum and minimum cooling rates are indicated in the continuous cooling transformation (CCT) phase diagram of the Inconel 718 sample in Fig. 9b. The CCT phase diagram is referenced from the simulation results of Kumara et al. [18]. The solid lines in Fig. 9b represent 0.5% precipitation in the dendrite core and the dashed lines represent the 0.5% precipitation close to the Laves phase. The changes in the pre cipitation kinetics caused by the segregation of chemical components such as Nb and Mo during deposition and solidification could be simu lated [18,21]. The CCT phase diagram indicated that the precipitation of the strengthening phases γ00 and γ0 were altered by the chemical
Table 3 Lattice parameters of the samples where the Inconel 718 deposit have maximum hardness at interpass times of 0, 30 and 60 s. 0 30 60
Area fraction (%)
6
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Fig. 9. (a) The temperature profile 4 mm from deposit substrate surface (b) continuous cooling transformation phase diagram for Inconel 718 alloy created from simulations [18]. Solid lines represent 0.5% precipitation in the dendrite core and dashed lines represent the 0.5% precipitation near the Laves phase. Cooling curves were based on the thermal profile in (a).
compositions of the Laves phase and dendrite core regions (Fig. 9b). In the simulation, the level of segregation near the Laves phase was Nb 40.85 wt% and Mo 0.88 wt%, and Nb 2.41 wt% and Mo 2.49 wt% in the dendrite core. In other words, the CCT diagram can shift depending on the level of segregation. Therefore, precipitates such as γ00 , γ0 were formed in the Laves phase or MC-type carbides of the inter-dendritic region at lower cooling rates due to the elemental segregation of ele ments including Nb and Mo. This mechanism may account for the different fractions of strengthening phases throughout the deposit (Table 4). SEM micrographs of the Laves phase taken using high-resolution FESEM of the bottom and top sections of the deposit revealed that the Laves phase had a typical lamellar form (Figs. 10a and 10b). It was difficult to identify the strengthening phases or the δ phase around the Laves phase in Figs. 10a and 10b, but the strengthening phases γ ’’ and γ ’ were observed close to the Laves phase in Fig. 10c [10,28]. The δ phase had precipitated and grown around the Laves phase in the inter-dendritic region, and was observed to be needle-shaped in a 2-D cross-section [13]. A high-resolution FE-SEM micrograph of inter-dendritic region in the middle section (8 mm from the substrate surface) exhibited blocky precipitates with a needle-shaped δ phase (Fig. 11a). Further magnifi cation of the high-resolution FE-SEM micrograph revealed disc-shaped
γ00 and spherical γ0 near to the blocky precipitate (Fig. 11b). The γ00 and γ0 phases in these micrographs were similar to those from near the Laves phase (Fig. 10c). The use of a plasma arc as a heat source in the WAAM process instead of a laser or electron beam allowed for the application of high heat to the deposit [28]. Sustained thermal effects by the successive deposition steps resulted in an aging effect on the deposit similar to heat treatment. The SEM micrographs given in Figs. 10c and 11b showed the formation of γ00 , γ0 and δ precipitates due to the aging effect. This aging effect cannot be produced in Inconel 718 products manufactured using con ventional casting or forging without the use of an additional process (e. g. heat treatment), and the phenomenon has only been observed in Inconel 718 deposits using the CMT method. The cooling rates using the CMT process were slower than that of AM processes that used laser or electron beams, and by slowing this cooling rate, the aging effect was maximized. The strengthening phases such as the γ00 , γ0 and δ phases were generated more quickly using the CMT-based WAMM method than AM methods that use laser or electron beams as heat sources. The method is expected to perform well in in-situ heat treatment applied simultaneously to deposition.
Fig. 10. SEM micrographs of the Laves phase in the Inconel 718 deposit sample (interpass time ¼ 30 s) in the (a) bottom section (height ¼ 2 mm), (b) top section (height ¼ 14 mm), and (c) middle section (height ¼ 8 mm). 7
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Fig. 11. SEM micrographs of the middle section of the Inconel 718 deposit (interpass time ¼ 30 s) showing (a) blocky precipitates with needle-shaped δ phase, and (b) a magnified view of the inter-dendritic region containing the precipitate.
5. Conclusions
References
Inconel 718 deposits were produced using a CMT-based WAAM method. The microstructures of the deposit samples were observed using SEM. Analysis was performed using SEM-EDS, EBSD and XRD. Columnar dendrite microstructures were observed, which were formed along the entire length of the sample (z-direction), with the exception of the top layer. Chemical components such as Nb and Mo were segregated in the inter-dendritic region, and the Laves phase, MC-type carbide, and γ00 , γ0 and δ phase precipitates were identified by SEM, SEM-EDS and EBSD. Successive deposition resulted in an aging effect that was observed for the first time using deposition with CMT. Strengthening phases such as γ00 and γ’ were formed, which contributed to a microhardness of 250 to 306 HV. The highest hardness value of each deposit was observed in the middle sections due to the high precipitation of these strengthening phases. The XRD analysis showed a reduced value of the lattice parameter, which indicated the formation of strengthening phases such as the γ00 , γ0 and δ phases. The thermal profile at a fixed point during deposition was used to estimate the cooling rate that resulted in the formation of strengthening phases. The cooling rate was 5.4 to 50.0 � C/s, which is slower than the cooling rate of deposits produced using an AM process with laser or electron beams. The aging effect was maximized in the WAAM process using CMT when the cooling rate was controlled. The study demon strated that in-situ heat treatment during deposition is possible. WAAM offers the industry with a means to produce large deposits without the need for expensive equipment, specifically the use of a conventional welding machine instead of a laser or electron beam. The addition of CMT will further enhance the strength of the products. This combined approach has great potential for future practical implementation.
[1] T.M. Pollock, S. Tin, J. Propul. Power 22 (2006) 361–374, https://doi.org/ 10.2514/1.18239. [2] T. DebRoy, H. Wei, J. Zuback, T. Mukherjee, J. Elmer, J. Milewski, A.M. Beese, A. Wilson-Heid, A. De, W. Zhang, Prog. Mater. Sci. 92 (2018) 112–224, https://doi. org/10.1016/j.pmatsci.2017.10.001. [3] V. Popovich, E. Borisov, A. Popovich, V.S. Sufiiarov, D. Masaylo, L. Alzina, Mater. Des. 114 (2017) 441–449, https://doi.org/10.1016/j.matdes.2016.10.075. [4] A. Strondl, R. Fischer, G. Frommeyer, A. Schneider, Mater. Sci. Eng., A 480 (2008) 138–147, https://doi.org/10.1016/j.msea.2007.07.012. [5] A. Odabas¸ı, N. Ünlü, G. G€ oller, M.N. Eruslu, Metall. Mater. Trans. A 41 (2010) 2357–2365. [6] G. Dinda, A. Dasgupta, J. Mazumder, Scripta Mater. 67 (2012) 503–506, https:// doi.org/10.1016/j.scriptamat.2012.06.014. [7] L. Zhu, Z. Xu, P. Liu, Y. Gu, Optic Laser. Technol. 98 (2018) 409–415, https://doi. org/10.1016/j.optlastec.2017.08.027. [8] B. Chen, J. Mazumder, Rapid Prototyp 23 (2017) 919–929, https://doi.org/ 10.1108/RPJ-05-2016-0071. [9] W.J. Sames, K.A. Unocic, G.W. Helmreich, M.M. Kirka, F. Medina, R.R. Dehoff, S. S. Babu, Addit. Manuf. 13 (2017) 156–165, https://doi.org/10.1016/j. addma.2016.09.001. [10] C. Kumara, D. Deng, F. Hanning, M. Raanes, J. Moverare, P. Nyl� en, Metall. Trans. A 50 (2019) 2527–2537. [11] A. Mostafa, I. Picazo Rubio, V. Brailovski, M. Jahazi, M. Medraj, Metals 7 (2017) 196, https://doi.org/10.3390/met7060196. [12] D. Ding, Z. Pan, D. Cuiuri, H. Li, Robot. Comput. Integrated Manuf. 31 (2015) 101–110, https://doi.org/10.1016/j.rcim.2014.08.008. [13] X. Xu, S. Ganguly, J. Ding, C.E. Seow, S. Williams, Mater. Des. 160 (2018) 1042–1051, https://doi.org/10.1016/j.matdes.2018.10.038. [14] G. Asala, A. Khan, J. Andersson, O. Ojo, Metall. Mater. Trans. A 48 (2017) 4211–4228. [15] S. Singh, J. Andersson, Sci. Technol. Weld. Join. 23 (2018) 568–574, https://doi. org/10.1080/13621718.2018.1429238. [16] R. Thavamani, V. Balusamy, J. Nampoothiri, R. Subramanian, K.R. Ravi, J. Alloys Compd. 740 (2018) 870–878, https://doi.org/10.1016/j.jallcom.2017.12.295. [17] O. Ola, F. Doern, Mater. Des. 57 (2014) 51–59, https://doi.org/10.1016/j. matdes.2013.12.060. [18] C. Kumara, A. Segerstark, F. Hanning, N. Dixit, S. Joshi, J. Moverare, P. Nyl�en, Addit. Manuf. 25 (2019) 357–364, https://doi.org/10.1016/j.addma.2018.11.024. [19] T. Antonsson, H. Fredriksson, Metall. Mater. Trans. B 36 (2005) 85–96. [20] G. Ko, K.M. Seo, H.J. Kim, H. Hong, Weld. World 61 (2017) 945–953, https://doi. org/10.1007/s40194-017-0477-2. [21] A. Segerstark, J. Andersson, L.-E. Svensson, O. Ojo, Mater. Char. 142 (2018) 550–559, https://doi.org/10.1016/j.matchar.2018.06.020. [22] G. Knorovsky, M. Cieslak, T. Headley, A. Romig, W. Hammetter, Metall. Trans. A 20 (1989) 2149–2158. [23] G. Dinda, A. Dasgupta, J. Mazumder, Mater. Sci. Eng., A 509 (2009) 98–104, https://doi.org/10.1016/j.msea.2009.01.009. [24] M. Ni, C. Chen, X. Wang, P. Wang, R. Li, X. Zhang, K. Zhou, Mater. Sci. Eng., A 701 (2017) 344–351, https://doi.org/10.1016/j.msea.2017.06.098. [25] Y.-L. Kuo, S. Horikawa, K. Kakehi, Mater. Des. 116 (2017) 411–418, https://doi. org/10.1016/j.matdes.2016.12.026. [26] Y. Mei, Y. Liu, C. Liu, C. Li, L. Yu, Q. Guo, H. Li, J. Alloys Compd. 649 (2015) 949–960, https://doi.org/10.1016/j.jallcom.2015.07.149. [27] H. Zhang, C. Li, Q. Guo, Z. Ma, Y. Huang, H. Li, Y. Liu, Mater. Char. 133 (2017) 138–145, https://doi.org/10.1016/j.matchar.2017.09.032. [28] F. Kong, J. Ma, R. Kovacevic, J. Mater. 211 (2011) 1102–1111, https://doi.org/ 10.1016/j.jmatprotec.2011.01.012.
CRediT authorship contribution statement Donghyun Van: Data curation, Investigation, Writing - original draft. G.P. Dinda: Data curation, Conceptualization. Jaewoong Park: Resources, Investigation. Jyoti Mazumder: Supervision, Investigation. Seung Hwan Lee: Data curation, Supervision, Conceptualization, Project administration, Writing - review & editing. Acknowledgement This work was supported by Korea Aerospace University (No. 201901-006) and Korea Evaluation Institute of Industrial Technology (KEIT) grant, which was funded by the Korean government (MOTIE) (No. 20000258).
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Materials Science & Engineering A journal homepage: http://www.elsevier.com/locate/msea
Enhancing hardness of Inconel 718 deposits using the aging effects of cold metal transfer-based additive manufacturing Donghyun Van a, G.P. Dinda b, Jaewoong Park a, Jyoti Mazumder c, Seung Hwan Lee a, * a
School of Aerospace and Mechanical Engineering, Korea Aerospace University, 76 Hanggongdaehang-ro, Deokyang-gu, Goyang-si, Gyeonggi-do, 10540, Republic of Korea b Department of Mechanical Engineering, Wayne State University, Detroit, MI, 48202, USA c Center for Laser and Plasmas for Advanced Manufacturing, University of Michigan, 2041 G. G. Brown, 2350 Hayward Street, Ann Arbor, MI, 48109-2125, USA
A R T I C L E I N F O
A B S T R A C T
Keywords: WAAM CMT Inconel 718 Strengthening phase Hardness
The effects of successive deposition using continuous heat from a plasma arc on a multi-layer Inconel 718 deposit was investigated using a wire-arc additive manufacturing (WAAM) process with cold metal transfer (CMT). Deposits consisting of 10 layers were manufactured, where interpass time was the only process parameter varied. The shape of the deposits was determined by the interpass time, while microstructural characteristics (growth of dendrite, segregation of chemical component, formation of various phases, etc.) developed due to successive deposition. An aging effect was observed, which caused precipitation of the γ00 and γ0 strengthening phases and the δ phase. The highest hardness value was measured in the middle section of each deposit due to precipitation of the strengthening phases. The cooling rate was calculated using the thermal profile of the deposition process and was found to be slower using the CMT-based WAAM process than that of other additive manufacturing (AM) processes that use laser or electron beams. The results presented here also demonstrated that in-situ heat treatment during deposition is possible using this CMT-based WAAM approach.
1. Introduction Additive manufacturing (AM) technology is an easy means of manufacturing complex three-dimensional designs and utilizes less material than traditional manufacturing methods. These advantages led to the use of AM in the production of metal products, where recent applications include expensive superalloy products for use in the fields of aerospace, energy, defense and medicine [1–5]. Many studies have recently focused on the application of AM to the alloy Inconel 718. Dinda et al. [6] investigated microstructural morphology of deposits produced using a CO₂ laser beam and metal powders, and demonstrated that the microstructure and laser beam scanning pattern were correlated. Furthermore, changes in the hardness were related to the aging temperature difference. Zhu et al. [7] studied the changes and the strength of the deposit’s microstructure resulting from manufacture using a fiber laser and metal powder. This study related the grain size, dendrites and Laves phases to the laser power and beam diameter. A tensile test was conducted on the deposits produced under various process conditions. Chen et al. [8] identified the process parameters that affect the cooling rate of deposit using a disc laser and
metal powder. The microstructure and the hardness underwent changes due to the separation of chemical components. Empirical formulae related to process parameters were applied to evaluate shape-related formation qualities, including deposit width and height, cooling rate and hardness. Sames et al. [9] produced Inconel 718 deposits using electron-beam melting. The effects of solution treatments, aging and in-situ heat treatment on the mechanical properties and microstructures of the deposits were analyzed. Kumara et al. [10] studied the homoge nization effect of in-situ heat treatments during electron-beam AM processing. Overall, heat treatments, aging effects and precipitate be haviors of the Inconel 718 alloy deposits produced using metal powder with laser or electron beams have been widely studied [9,11]. However, AM based on laser and electron beams have limitations for the commercial implementation of the production of large deposits, as initial investment cost is high and a relatively low deposition rate is generally expected. The wire-arc additive manufacturing (WAAM) process [12] has been developed to address these issues. The method utilizes plasma arcs and wires, where the plasma arc is typically generated using a conventional welding machine and the consumable electrodes produced for welding are used as the feedstock wires. Only a
* Corresponding author. E-mail address: [email protected] (S.H. Lee). https://doi.org/10.1016/j.msea.2020.139005 Received 7 January 2020; Received in revised form 21 January 2020; Accepted 23 January 2020 Available online 25 January 2020 0921-5093/© 2020 Elsevier B.V. All rights reserved.
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few studies have focused on the WAAM process, but have demonstrated promise for the production of simple designs using a tungsten inert gas (TIG) welding machine and wire. Xiangfang et al. [13] produced Inconel 718 components using the WAAM process and interpass rolling, with TIG as the WAMM heat source. The Inconel 718 components produced by the WAAM process alone were compared with components produced using both WAAM and interpass rolling. The effects of both methods on the mechanical properties and microstructure of the components was reported. Asala et al. [14] manufactured deposits using WAAM with a TIG heat source and ATI 718 Plus feedstock material. The study iden tified a deposited metal heat affected zone (DMHAZ), which is a softened zone created during the deposition process. The microstructures and strengthening phases, including the γ00 and γ0 phases of the Inconel 718 deposits, were also investigated. An issue related to the use of a conventional TIG heat source is that the heat input cannot be as precisely controlled as a laser heat source. This can result in a non-uniform melt pool and unstable deposition conditions, which contribute to defects (e.g. liquation cracking) that deteriorate the mechanical properties of the deposit [15,16]. Control of the arc heat source is essential to addressing this issue and ensuring the desired quality of the products. Cold metal transfer (CMT) [17] provides controllable heat input using a high speed digital control and a high frequency wire feeding system. CMT may be implemented in the WAAM process to provide accurate heat input control. Unlike the AM processes that use a metal powder, applications of the CMT process using the Inconel 718 alloy are very limited. Most of the research involving CMT has focused on the welding process and materials other than the Inconel 718 alloy. In-situ heat treatment using the Inconel 718 alloy and the CMT method has not yet been investigated. Before CMT is applied in the WAAM process, the microstructure characteristics of the deposits after accumulated heat input during solidification must be studied. This study aimed to investigate the changes in microstructure in Inconel 718 multi-layer deposits after successive deposition and continuous heat input by CMT. Ten thin layers were produced with various interpass times as a process parameter. The thermal profile of the deposition process was measured. The effects of the interpass time on the shape of the deposit and the microstructural changes per layer were investigated. Furthermore, the cooling rates of the precipitations, the corresponding microstructures and the behavior of the precipitate were evaluated according to thermal profiles measured at a fixed point on the deposit.
XY stage and Z stage, respectively. The scanning and building directions of the deposit are shown in Fig. 1b. 100% Ar gas with a flow rate of 20 ℓ/min was applied to shield the plasma arc and the deposit and the contact tip to workspace distance (CTWD) was fixed at 10 mm. The experimental conditions are listed in Table 1. The temperature was measured in real-time by using a pyrometer (OPTRIS, CTlaser 3MH3) positioned at the center of the deposit 4 mm from the surface in the z-direction. Different emissivity values were applied for the various materials when measuring temperatures using the pyrometer. A preliminary experiment was conducted to compare the K-type thermocouple temperature with the pyrometer temperature to confirm the surface emissivity of the Inconel 718 deposit. The surface emissivity of the pyrometer was 0.92. Thin walled samples comprising 10 layers were deposited using three interpass times (0, 30 and 60 s), which means that the deposit was aircooled for either 0, 30 and 60 s between layers. Deposition was per formed in one direction along the X axis. When the interpass times are 0, 30, and 60 s, the corresponding deposition rates are about 0.6, 1.7 and 3.5 kg/h, respectively. Deposits were cut along the center of the YZ plane to reveal the sample’s microstructure. The test samples were mounted and polished using standard metallographic methods. The microstructure of the Inconel 718 deposit was observed using a field emission scanning elec tron microscope (FE-SEM, Hitachi, S-4800) and an optical microscope (OM, Olympus, BX51RF) after corroding the surface with Keller’s etchant (190 ml water, 5 ml nitric acid, 3 ml hydrochloric acid and 2 ml hydrofluoric acid). A FE-SEM with an energy dispersive X-ray spec trometer (EDS) was used to analyze the chemical compositions of the samples. Hardness measurements were performed from the bottom to the top of the deposit using a Vickers hardness tester (Matsuzawa, PMTX7) by pressing the test sample for a holding time of 10 s with a force of 1000 gf. The hardness value was calculated as the mean value of three indentations. Lattice parameters of test samples were examined using an X-ray diffraction (XRD) system (Malvern Panalytical, X’Pert PRO MPD). The XRD measurements were conducted on the polished surface of the deposit samples. Cu Kα radiation with a wavelength of 1.5406 Å was Table 1 Experimental parameters for the wire-arc additive manufacturing (WAAM) process with cold metal transfer (CMT). Parameter Wire (diameter) Current/voltage Wire feed speed Welding speed CTWD Shielding gas Welding direction Length (layer) Substrate Interpass time
2. Materials and methods The deposition material was Inconel 718 alloy, and AH36 was chosen the substrate due to its low cost. The deposit was produced using a 1.2 mm diameter Inconel 718 wire on a substrate with the dimensions 250 x 80 x 15 mm. A CMT machine (FRONIUS, TPS 4000) was used during deposition, and a XYZ 3-axis automatic stage (Six Degrees Inc.) with a 1μm resolution was used for transition of the sample and welding torch (Fig. 1a). The deposition sample and the welding torch were fixed to the
Fig. 1. Schematic of (a) the experiment, (b) the deposit showing the scanning and building direction. 2
Inconel 718 (1.2 mm) 120 A/13.4 V 4.8 m/min 0.5 m/min 10 mm Ar 100%, 20 ℓ/min One direction 100 mm (10 layers) AH36 (250 x 80 x 15 mm) 0 s, 30 s, 60 s
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applied and measured in the range 30–100� at intervals of 0.09� with an acquisition time of 2 s per increment. The XRD patterns were analyzed using the HighScope Plus v3.0 software package. Phase analysis of mi crostructures was conducted using an electron backscatter diffraction (EBSD) detector (Bruker, CrystAlign 200).
near top sections. The top of the 10th layer was not subjected to the heat associated with depositing the next layer and exhibited a faster cooling rate and smaller thermal gradient. This resulted in the formation of an equiaxed dendrite structure in the uppermost layer of the deposits. When the first layer was deposited to form the base of the deposit, no heat had yet accumulated. Thus, this first layer had a faster cooling rate than the deposits in the middle and top of the deposit. This was asso ciated to the formation of fine primary dendrite structures in the bottom section. As deposition proceeded, the cooling rate slowed down due to continuous arc heat accumulation. The spacing between the primary dendrites formed during deposition of the middle and top sections became greater than the spacing between the primary dendrites at the bottom. A secondary dendrite arm began to appear, and a welldeveloped secondary dendrite arm was observed in the near top part of the deposit due to the low cooling rate. SEM-EDS analysis was conducted at the dendritic and inter-dendritic regions in the middle section of the deposit produced using an interpass time of 30 s (Fig. 4). The weight percentage (wt%) of each chemical component was investigated at points 1, 2, 3 and 4 indicated in Fig. 4. Point 1 was taken during SEM-EDS analysis of the dendritic region (Fig. 4a), where the wt% of Nb and Mo was substantially lower than at the other points with values of 2.51 and 2.08%, respectively. At point 2, taken in the inter-dendritic region, this increased to 6.30 and 3.41% for Nb and Mo, respectively. This observation was associated with the segregation of chemical components, resulting in a higher concentration of Nb and Mo in the inter-dendritic region than the dendritic region. Elements such as Nb, Mo and Ti tend to form segregation zones during the last stages of solidification in the boundary and inter-dendritic re gions [10,18]. Nb has a particularly small diffusion coefficient and tends to remain in the liquid phase until the final stage of solidification, allowing for easy segregation. From the weight percent of elements listed in Table 2, the phases of the point 3 and 4 were further investigated to analyze the precipitates and segregation due to constitutional liquation. The Nb and Mo contents at points 3 and 4 are higher than at point 1, but the content of Ni, Fe and Cr had decreased. Laves phase is a topologically closed-packed (TCP) phase with a composition of (Ni,Fe,Cr)₂(Nb,Ti,Mo) that contains more Cr and Nb than the γ matrix. Antonson et al. [19] found that the Laves phase is formed when the wt% of Nb in the inter-dendritic region ex ceeds 20. In addition, it is known that SEM microstructures of a nickel-based superalloy such as Inconel 718 displayed lamellae in the Laves phase [20]. As the wt% of Nb was 26.43 and the precipitate had a lamellar form, the phase of point 3 shown in Fig. 4c is regarded as Laves. Fig. 4d indicates that the precipitate at point 4 had a polygonal form
3. Results 3.1. Macrostructural features The Inconel 718 alloy was deposited in ten layers using an interpass time of either 0, 30 or 60 s, and an OM macrostructure photograph was taken of the YZ plane cut along the center in the X direction (Fig. 2). The grain growth and remelting zone in the Z direction was identified in the macrostructure photograph of the YZ plane. The remelting zone was formed when the upper part of the solidified air-cooled layer was remelted by the arc heat applied as the following layer was deposited. As deposition proceeds from the bottom to the top direction (Z direction), the arc heat continues to accumulate in the deposit. Coarse grains were observed below ca. 3 mm from the substrate surface due to this heat accumulation. Interpass time was the only process parameter that was varied during depositing. As the interpass time was increased, the height of the deposit tended to increase while the maximum width decreased (Fig. 2). In addition, the inclination angle increased and approached 90� with longer interpass time. These changes were caused by the interlayer temperature control effect, where a shorter interpass time resulted in less air-cooling and the heat accumulation from previous depositions resulted in a higher preheated temperature for the following layer. Increased preheating temperature resulted in a higher temperature and decreased viscosity of the generated melt pool, which increased the velocity of the fluid. This caused the deposit with the shortest interpass time to be wider and shorter than the others. 3.2. Microstructural features The SEM micrographs of the deposits revealed that a columnar dendrite had grown from the bottom to the top in all layers except the top of the 10th layer (Fig. 3). Solidification of the deposit was initiated with the formation of a primary dendrite. The grain grew in the opposite direction of the heat flow, and crystal growth proceeded in the Z di rection when deposited from the bottom to the near top point. The columnar dendrite at the top of the 10th layer was observed to grow in a different direction than the crystal growth in the bottom, middle and
Fig. 2. OM macrographs of the Inconel 718 deposit samples produced using interpass times of (a) 0 s, (b) 30 s and (c) 60 s. 3
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Fig. 3. SEM micrographs of the Inconel 718 deposit deposits produced using interpass times of 0, 30 and 60 s.
on these weight percentages and geometrical shape, the main pre cipitates observed in the inter-dendritic region in the deposited SEM microstructures given in Figs. 4c and 4d were the Laves phase and MC-type carbide, respectively. Differences in the shape of the Laves phase was observed along the height of the deposit produced using an interpass time of 30 s were observed (Fig. 5). The SEM microstructure of the Inconel 718 exhibited lamellar morphology in the Laves phase. Discrete Laves phases were observed in the bottom sections of the deposit, which becomes more continuous in shape towards the middle section. This was attributed to the change in cooling rate, which was highest at the bottom of the de posit. The cooling rate at the very top of the deposit was also high due to convective air-cooling, thus a similar discrete form was observed. The Laves shape changed from a discrete to a lamella shape from the bottom to the top of the deposit and changed back to discrete at the very top. 3.3. Hardness measurement results The Vickers hardness test found that the hardness of the deposits formed using interpass times of 0, 30 and 60 s varied from 250 to 306 HV (Fig. 6). A previous study found that the Vickers hardness of a deposit produced using laser beams and metal powder was ca. 240 to 250 HV, and thus lower than those in the current study [21]. The maximum hardness was measured at the bottom of the deposits in the previous study, and decreased along the height of the deposit [8,21]. In this study, when the interpass time was 0, 30, and 60 s, the maximum hardness was
Fig. 4. SEM-EDS analysis points of the (a) dendritic region, (b) inter-dendritic region, (c) Laves phase in the inter-dendritic region, and (d) MC carbide in the inter-dendritic region.
which the MC-type carbide is [20]. The wt% of Nb at point 4 was 25.48, the second highest chemical component after Ni. In addition, the wt% of C is 3.33, which was considerably higher than at any other point. Based Table 2 Weight percent of elements at points 1, 2, 3 and 4 (see Fig. 4). Position
Ni
Fe
Cr
Nb
Mo
Ti
C
Ta
Others
Remark
Point 1 Point 2 Point 3 Point 4
53.69 52.81 38.76 36.69
21.02 18.08 12.89 13.27
18.17 17.26 12.03 11.44
2.51 6.30 26.43 25.48
2.80 3.41 7.74 6.98
0.89 1.18 1.01 1.08
– – – 3.33
– – – 1.49
0.92 0.96 1.14 0.24
Wt%
4
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Fig. 5. Shape changes of the Laves phase depending on the height in the z-direction of deposit (interpass time ¼ 30 s).
Fig. 6. Vickers hardness measurement of the Inconel 718 deposit produced using an interpass time of (a) 0 s, (b) 30 s, and (c) 60 s.
291, 306, and 300 HV, respectively, which were within the range of measurement error. An increase of hardness value from the bottom was expected for the deposits in the current study due to precipitation of the strengthening phases such as the γ00 and γ0 phases. γ00 is a strengthening phase of the Inconel 718 alloy, and is a disc-shaped precipitate with a chemical composition of Ni3Nb and has a DO₂₂ ordered body-centered tetragonal (BCT) crystal structure. γ0 is a spherical fine precipitate with a chemical composition of Ni₃(Al,Ti) that also acts as a strength ening phase, and has a L1₂ ordered FCC crystal structure [14,22]. The increased hardness in each sample may be attributed to precipitates such as γ00 and γ0 formed due to aging through successive deposition because Inconel 718 is a typical precipitation hardening alloy and strengthening is usually achieved through post-heat treatment, tempering and aging. The details of the aging effect related with the profile of hardness will be discussed in section 3.5 and 4. After passing the maximum hardness, the decrease of the hardness values was observed from 4 mm to 14 mm in Fig. 6a, from 8 mm and 14 mm in Fig. 6b and from 8 mm and 14 mm in
Fig. 6c. This is caused by the coarse grain induced by the low cooling rate during the solidification. The low cooling rate was caused by the heat accumulation during the multi-layer depositions. The increase in hardness values at the top of the three samples was attributed to convective heat transfer that allowed for rapid cooling and quenching. 3.4. XRD analysis results XRD analysis was conducted at the positions where the maximum hardness was measured in each deposit (Fig. 7). Peaks were mainly attributed to the γ-matrix of the deposits. It is difficult to confirm the precipitation of the strengthening phases γ00 and γ0 using XRD without prolonged annealing or coarsening of the precipitates. However, the precipitation can be indirectly confirmed according to the reduced lat tice parameter [23]. The lattice parameters of the deposits produced in this study were lower than the lattice parameter of the γ-matrix of Inconel 718 (360.05 p.m.) (Table 3) [11]. The lattice parameter of the 5
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Fig. 7. The XRD patterns of the Inconel 718 deposits produced using interpass times of 0, 30 and 60 s.
sample produced using an interpass time 30 s was the lowest, and is indicative of precipitation [23]. These results were consistent with the hardness measurement results in Fig. 6. Strengthening phases such as the γ00 and γ0 dissolved in the γ-matrix that was partially formed and precipitated during deposition.
Fig. 8. (a) EBSD orientation map, (b) EBSD phase map of the Inconel 718 deposit sample (interpass time ¼ 30 s) in middle part, which shows precipita tion morphology (8 mm from the substrate surface).
3.5. EBSD analysis results
Table 4 Phase area fraction in the bottom, middle and top sections of an Inconel 718 deposit (interpass time ¼ 30 s).
EBSD analysis was performed to examine the distribution of the microstructure and precipitation in the deposits [24]. The EBSD orien tation map and the EBSD phase map ca. 8 mm from the substrate surface (z-direction) of the deposit produced using an interpass time of 30 s was evaluated, as this point exhibited the maximum hardness (Fig. 8). The EBSD orientation map was shaded according to the main growth di rection of the columnar grains, with red shading indicating the z-di rection and purple shading indicating precipitates. The growth direction of the precipitation phase was different to that of the γ-phase (Fig. 8a). The EBSD phase map analysis of the distribution of precipitates such as γ00 , γ0 and δ revealed that the area fraction of the strengthening phase γ00 was larger in the middle section (8 mm from the surface) than the bot tom and top sections (Fig. 8b). As shown in Table 4, the area fraction of the strengthening phases (sum of γ00 and γ0 fraction) in the middle section was larger than the other areas. The previous hardness profile of the deposit is related to the content of strengthening phases in the micro structure, and these findings confirmed the hardness findings. As shown in Table 4, the area fraction of the δ phase is listed. In the Inconel 718 alloy, the δ phase is known to precipitate at the interface of the existing Laves phase or NbC in the temperature range from 850 to 1000 � C [25]. In addition, the precipitation of δ phase is also dependent on the residual strain and the amount of formed γ00 precipitates [26,27]. These are the reasons why the small area fraction of the δ phase might exist compared with γ00 precipitates in Table 4.
Phase
γ(Ni–Cr–Fe) γ’(Ni₃(Al, Ti)) γ’’(Ni₃Nb) δ(Ni₃Nb)
Interpass time (s)
Lattice parameter (pm) 359.75 359.22 359.69
Bottom (2 mm from substrate surface)
Middle (8 mm from substrate surface)
Top (14 mm from substrate surface)
90.44 8.49
89.30 8.52
97.22 1.89
0.90 0.17
1.91 0.27
0.61 0.28
4. Discussion The temperature profiles of the deposition process were greatly varied due to the layer by layer nature of the AM method (Fig. 9a). The cooling rates were calculated using the temperature profile of the de posit measured at a fixed point using a pyrometer, and ranged 5.4 to 50.0 � C/s. The minimum value of the cooling rate was observed when an interpass time of 0 s was used, and the maximum value was associated with an interpass time of 60 s. The cooling rate was 16.3 to 20.7 � C/s using an interpass time of 30 s. As shown in the thermal profile of Fig. 9a, the deposited area which stay at the aging temperature band between 600 and 1000 � C produces strengthening phases γ’’ and γ’ [18]. The maximum and minimum cooling rates are indicated in the continuous cooling transformation (CCT) phase diagram of the Inconel 718 sample in Fig. 9b. The CCT phase diagram is referenced from the simulation results of Kumara et al. [18]. The solid lines in Fig. 9b represent 0.5% precipitation in the dendrite core and the dashed lines represent the 0.5% precipitation close to the Laves phase. The changes in the pre cipitation kinetics caused by the segregation of chemical components such as Nb and Mo during deposition and solidification could be simu lated [18,21]. The CCT phase diagram indicated that the precipitation of the strengthening phases γ00 and γ0 were altered by the chemical
Table 3 Lattice parameters of the samples where the Inconel 718 deposit have maximum hardness at interpass times of 0, 30 and 60 s. 0 30 60
Area fraction (%)
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Fig. 9. (a) The temperature profile 4 mm from deposit substrate surface (b) continuous cooling transformation phase diagram for Inconel 718 alloy created from simulations [18]. Solid lines represent 0.5% precipitation in the dendrite core and dashed lines represent the 0.5% precipitation near the Laves phase. Cooling curves were based on the thermal profile in (a).
compositions of the Laves phase and dendrite core regions (Fig. 9b). In the simulation, the level of segregation near the Laves phase was Nb 40.85 wt% and Mo 0.88 wt%, and Nb 2.41 wt% and Mo 2.49 wt% in the dendrite core. In other words, the CCT diagram can shift depending on the level of segregation. Therefore, precipitates such as γ00 , γ0 were formed in the Laves phase or MC-type carbides of the inter-dendritic region at lower cooling rates due to the elemental segregation of ele ments including Nb and Mo. This mechanism may account for the different fractions of strengthening phases throughout the deposit (Table 4). SEM micrographs of the Laves phase taken using high-resolution FESEM of the bottom and top sections of the deposit revealed that the Laves phase had a typical lamellar form (Figs. 10a and 10b). It was difficult to identify the strengthening phases or the δ phase around the Laves phase in Figs. 10a and 10b, but the strengthening phases γ ’’ and γ ’ were observed close to the Laves phase in Fig. 10c [10,28]. The δ phase had precipitated and grown around the Laves phase in the inter-dendritic region, and was observed to be needle-shaped in a 2-D cross-section [13]. A high-resolution FE-SEM micrograph of inter-dendritic region in the middle section (8 mm from the substrate surface) exhibited blocky precipitates with a needle-shaped δ phase (Fig. 11a). Further magnifi cation of the high-resolution FE-SEM micrograph revealed disc-shaped
γ00 and spherical γ0 near to the blocky precipitate (Fig. 11b). The γ00 and γ0 phases in these micrographs were similar to those from near the Laves phase (Fig. 10c). The use of a plasma arc as a heat source in the WAAM process instead of a laser or electron beam allowed for the application of high heat to the deposit [28]. Sustained thermal effects by the successive deposition steps resulted in an aging effect on the deposit similar to heat treatment. The SEM micrographs given in Figs. 10c and 11b showed the formation of γ00 , γ0 and δ precipitates due to the aging effect. This aging effect cannot be produced in Inconel 718 products manufactured using con ventional casting or forging without the use of an additional process (e. g. heat treatment), and the phenomenon has only been observed in Inconel 718 deposits using the CMT method. The cooling rates using the CMT process were slower than that of AM processes that used laser or electron beams, and by slowing this cooling rate, the aging effect was maximized. The strengthening phases such as the γ00 , γ0 and δ phases were generated more quickly using the CMT-based WAMM method than AM methods that use laser or electron beams as heat sources. The method is expected to perform well in in-situ heat treatment applied simultaneously to deposition.
Fig. 10. SEM micrographs of the Laves phase in the Inconel 718 deposit sample (interpass time ¼ 30 s) in the (a) bottom section (height ¼ 2 mm), (b) top section (height ¼ 14 mm), and (c) middle section (height ¼ 8 mm). 7
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Fig. 11. SEM micrographs of the middle section of the Inconel 718 deposit (interpass time ¼ 30 s) showing (a) blocky precipitates with needle-shaped δ phase, and (b) a magnified view of the inter-dendritic region containing the precipitate.
5. Conclusions
References
Inconel 718 deposits were produced using a CMT-based WAAM method. The microstructures of the deposit samples were observed using SEM. Analysis was performed using SEM-EDS, EBSD and XRD. Columnar dendrite microstructures were observed, which were formed along the entire length of the sample (z-direction), with the exception of the top layer. Chemical components such as Nb and Mo were segregated in the inter-dendritic region, and the Laves phase, MC-type carbide, and γ00 , γ0 and δ phase precipitates were identified by SEM, SEM-EDS and EBSD. Successive deposition resulted in an aging effect that was observed for the first time using deposition with CMT. Strengthening phases such as γ00 and γ’ were formed, which contributed to a microhardness of 250 to 306 HV. The highest hardness value of each deposit was observed in the middle sections due to the high precipitation of these strengthening phases. The XRD analysis showed a reduced value of the lattice parameter, which indicated the formation of strengthening phases such as the γ00 , γ0 and δ phases. The thermal profile at a fixed point during deposition was used to estimate the cooling rate that resulted in the formation of strengthening phases. The cooling rate was 5.4 to 50.0 � C/s, which is slower than the cooling rate of deposits produced using an AM process with laser or electron beams. The aging effect was maximized in the WAAM process using CMT when the cooling rate was controlled. The study demon strated that in-situ heat treatment during deposition is possible. WAAM offers the industry with a means to produce large deposits without the need for expensive equipment, specifically the use of a conventional welding machine instead of a laser or electron beam. The addition of CMT will further enhance the strength of the products. This combined approach has great potential for future practical implementation.
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CRediT authorship contribution statement Donghyun Van: Data curation, Investigation, Writing - original draft. G.P. Dinda: Data curation, Conceptualization. Jaewoong Park: Resources, Investigation. Jyoti Mazumder: Supervision, Investigation. Seung Hwan Lee: Data curation, Supervision, Conceptualization, Project administration, Writing - review & editing. Acknowledgement This work was supported by Korea Aerospace University (No. 201901-006) and Korea Evaluation Institute of Industrial Technology (KEIT) grant, which was funded by the Korean government (MOTIE) (No. 20000258).
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