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Microstructural characteristics and crack formation in additively manufactured bimetal material of 316L stainless steel and Inconel 625
Journal Pre-proof Microstructural characteristics and crack formation in additively manufactured bimetal material of 316L stainless steel and Inconel 625 Nannan Chen (Formal analysis) (Investigation) (Writing - original draft), Haris Ali Khan (Data curation) (Writing - review and editing), Zixuan Wan (Software) (Data curation) (Investigation) (Writing original draft), John Lippert (Data curation) (Writing - original draft), Hui Sun (Software) (Data curation), Shun-Li Shang (Formal analysis) (Writing - original draft), Zi-Kui Liu (Funding acquisition) (Supervision), Jingjing Li (Supervision) (Project administration) (Funding acquisition) (Writing - review and editing)
PII:
S2214-8604(19)31472-1
DOI:
https://doi.org/10.1016/j.addma.2020.101037
Reference:
ADDMA 101037
To appear in:
Additive Manufacturing
Received Date:
2 September 2019
Revised Date:
4 December 2019
Accepted Date:
1 January 2020
Please cite this article as: Chen N, Khan HA, Wan Z, Lippert J, Sun H, Shang S-Li, Liu Z-Kui, Li J, Microstructural characteristics and crack formation in additively manufactured bimetal material of 316L stainless steel and Inconel 625, Additive Manufacturing (2020), doi: https://doi.org/10.1016/j.addma.2020.101037
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Microstructural characteristics and crack formation in additively manufactured bimetal material of 316L stainless steel and Inconel 625 Nannan Chen,1 Haris Ali Khan,1 Zixuan Wan,1 John Lippert,1 Hui Sun,2 Shun-Li Shang,2 Zi-Kui Liu,2 Jingjing Li 1* 1
Department of Industrial and Manufacturing Engineering, Pennsylvania State University, University Park, PA 16802, USA
Department of Materials Science and Engineering, Pennsylvania State University, University
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Park, PA 16802, USA
* To whom correspondence should be addressed. E-mail: [email protected], Tel: 1-814-863-1300, Fax: 1-814-863-4745. 1
Highlights • Changing printing sequence generated two types of interfaces with unique morphology • IN625 grains grew epitaxially on the fine grains of SS316L forming Type-I interface • Bidirectional nucleation from IN625 and mushy zone at SS316L formed Type-II interface • Cracking was formed at Type-II interface and in the SS316L tracks • Cracking mechanisms include solidification, liquidation, and ductility dip cracking
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Abstract This research illustrates the rationale of adopting a preferred printing sequence by examining crack
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generation predominated by resultant interfaces and microstructural inhomogeneity, through underlying governing mechanisms in directed energy deposition of 316L stainless steel/Inconel 625 (SS316L/IN625) bimetals. For this purpose, microstructural and crystallographic
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characterizations augmented by numerical simulations were employed on additively manufactured two distinct interfaces, i.e. Type-I (IN625 deposition on SS316L) and Type-II (SS316L deposition
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on IN625). Changing the printing sequence generated these two types of interfaces with unique morphologies, which was found attributable to the compositional variations and mismatch in
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thermal properties. Type-I interface was typified by gradual-change composition in the transition zone, causing the IN625 grains to grow epitaxially on the grains of SS316L. Type-II interface was characterized as a compositional sudden-change zone (CSCZ) adjacent to SS316L, leading to
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merging bidirectional nucleation and grain growth from the bottom IN625 and upper CSCZ, and lack of epitaxial growth. Additionally, high cracking susceptibility occurred near the Type-II interface rather than the Type-I interface, which was related to solidification and liquidation cracking, and further promoted ductility dip cracking. This research will provide a guideline for the additive manufacturing of bimetals with the consideration of printing sequence to control
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interface formation for a crack-free structure.
Keywords: directed energy deposition; bimetal; microstructure; crack formation; interfacial behavior
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1. Introduction Extreme environment applications, for instance, nuclear plants and aerospace vehicles involve components with requisite of varied properties at different locations [1, 2]. However, it is seldom or impossible to leverage such characteristics from a single material, and this need drives the unremitting pursuit of tailored properties by developing dissimilar materials structures. Different manufacturing techniques have been employed to date for their fabrication. Additive manufacturing (AM) processes, particularly directed energy deposition (DED) technique, emerge
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as a promising opportunity for manufacturing dissimilar materials structures by a layered approach with careful selection of process parameters [3] or by utilizing composition as a spatial design
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parameter [4]. Zhang and Bandyopadhyay [3] demonstrated the feasibility of fabricating metalceramic multi-layer structures with a laser-based AM method. Carroll et al. [5] obtained a bimetal
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structure of 304L stainless steel to Inconel 625 (IN625) using DED through a gradient zone to avoid sharp compositional or microstructural changes at the interface. Shah et al. [6] analyzed the influences of process parameters (e.g. laser power levels and powder mass flow rates) on
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microstructure and mechanical properties in DED processed Inconel-steel bimetal graded structure, where the authors demonstrated that desired mechanical properties can be achieved by controlling
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the input parameters. Onuike et al. [7] utilized three different building strategies (namely direct deposition, compositional gradation, and use of an intermediate bond layer) to produce DED processed bimetal structure of Ti-64 and IN718 alloys. Based on the microstructure and
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mechanical properties, the authors defined an intermediate bond layer as the most suitable technique. DED processed Inconel and copper alloy bimetal structure was produced for improved thermal conductivity behavior in high-temperature aerospace applications [8]. The DED process was also utilized for obtaining intricate bimetal structures with varied properties. For instance, bimetallic stainless tube structure with transitions directly from 316L stainless steel (SS316L) to
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SS430L was produced to effectively obtain transition from non-magnetic to magnetic material [9]. The above discussion manifests that AM processes are favorable in the fabrication of multimaterial into a single structure. Nickel-based superalloys (e.g. IN718 and IN625) and different grades of stainless steel (e.g. SS304L and SS316L) are widely used in nuclear reactors, aerospace, automotive and petrochemical components for their improved corrosion and oxidation resistance, and mechanical
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strength [10-12]. These two alloys offer a promising opportunity in dissimilar materials structures, where an increased strength at elevated temperature (where IN625 is suitable) and corrosion resistance (where SS316L is suitable) is achievable at different locations simultaneously [13]. Nonetheless, utilizing the two materials together bring forth perplexing challenges, including inhomogeneity of the microstructure introduced by lack of fusion, nonuniform melting and rapid solidification, which can cause deterioration of mechanical properties and microsegregation of the alloying elements during the solidification process [14]. In addition, deposition of a dissimilar
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material on these metals is susceptible to problems such as interface cracking and carbon diffusion from the low chromium to high chromium alloy [15,16].
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It is observed that altering the printing sequence changed the interfacial characteristics in bimetal structures. For instance, when using DED to print IN718 and SS316L structure, Li et al.
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[17] found different interfacial characteristics with alternated printing schemes. The authors observed a clear interface devoid of any diffusion layer when IN718 was treated as a substrate while the same parameters yielded a diffusion layer with greater microhardness than SS316L when
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IN718 was deposited on to SS316L. This phenomenon may be related to the nature of microstructural evolution particularly at and near the interface region during the heating and
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cooling stages of the layer. In another research [18], different defects were observed with the different printing sequence (i.e. cracks close to the IN718/316SS interface, and holes at the 316SS/IN718 interface, where the first material was on the top). In essence, it has been reported
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that the printing or deposition scheme is a dominant factor that not only influences the interfacial microstructure but also if appropriately selected, can help in mitigating the undesirable inhomogeneity of the microstructure and other defects (e.g. pores and cracks). Current publications focus on reporting the phenomena; however, it is scarce for scientific understanding that can explain the rationale of selecting a particular printing scheme based on the resultant interfacial
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microstructure and subsequent defects formation. This research targets at ascertaining the formation mechanisms of the distinct interfacial
characteristics and microstructural evolution during the DED process of IN625 and SS316L alloys. Particular emphasis has been laid on establishing the relationship between different printing sequences and the subsequent generated microstructure and cracks. Well-integrated materials characterization techniques were employed to comprehend the microstructural features (i.e. grain
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morphology, texture, and phases), and their formation mechanisms and further influences on the microstructural inhomogeneity and cracks. The experimental results are also augmented by numerical analysis for ascertaining material flow. The remainder of the paper follows the defined sequence. Section 2 and section 3 describe the experimental procedure and basis of numerical analysis, respectively. Section 4 relates to the obtained results consisting of structure morphology, crystallographic analysis, and defects characterization. The paper is culminated by discussing key conclusions in section 5.
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2. Experimental procedure This section provides details of the DED process and explanations of the procedures for
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preparing metallographic samples and microstructural characterization adopted in this research.
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2.1. DED process
A rectangular bimetal structure was additively manufactured using SS316L and IN625
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powders with the DED method in an argon atmosphere. The dimensions of the built part were 110 mm (Length) x 10 mm (Width) x 65 mm (Height) as illustrated in Fig. 1(a), where the X, Y, and Z axes are the step-over, laser scanning, and building directions, respectively. The build started
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with 8 layers of SS316L depositions, followed by 8 layers of Inconel 625 depositions, and continuously alternated with those respective layers until the build finished. The laser scanning pattern is schematically illustrated in Fig. 1(b). Sandvik SS316L powder (particle sizes of 45-150
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µm) and Carpenter Technology’s IN625 alloy powder (particle sizes of 120-270 µm) were used and their chemical compositions are provided in Table 1.
The additive build was manufactured with the Penn State Applied Research Laboratory’s
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high-power high-deposition system which houses a custom cladding head with water-cooled optics and an inert gas enclosure, and the process parameters are listed in Table 2. 2.2. Microstructural characterization The bimetal structure was sectioned from XZ and YZ planes to characterize microstructures of the printed bulk materials (i.e. SS316L and IN625) and the interfaces between these two materials. The sectioned samples were first mounted in epoxy, then polished with silicon 5
papers following a sequence of #320, #600, and #800, and finally polished with diamond suspensions down to 1 m. The samples were further electrolytically etched with a 4% oxalic acid reagent for 5-10 seconds. Vibratory polishing was applied to prepare the samples for Electron Back-Scattered Diffraction (EBSD) analysis. Macrostructures were acquired under a Nikon optical microscope. A Nova nano scanning electron microscope (SEM) 630 field emission analysis equipped with an Oxford X-Max EDS (Energy Dispersive Spectroscopy) was employed to obtain the microstructures and elemental distributions. Finally, the crystallographic features were
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revealed by EBSD analysis performed on a Helios NanoLab 660 SEM equipped with an Oxford EBSD detector.
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3. Numerical analysis
DED process was simulated by commercial computational fluid mechanics (CFD) software
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with customized heat and focal subroutines. Two distinct interfaces, i.e. IN625 deposition on SS316L and SS316L deposition on IN625, were modeled separately. For simplification, a partial
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cylinder was added on the bottom metal to represent the deposited particles during the deposition process, as depicted in Fig. 2(a). A total number of 416,000 cells with 0.1 mm cell size were
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employed in the simulation. The dimensions of the initial fluid cross-sections are indicated in Fig. 2(b). The power density of the laser spot was assumed to follow the Gaussian distribution (Fig. 2c), where the peak power density in the central region was approximately 283 W/mm2. Free
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surfaces of both materials were tracked by volume-of-fluid, which can be represented by Equation 1, where F is fluid volume fraction; u, v and w represent velocity components in x, y and z directions, respectively. The mixture of the two different materials was completely driven by the flow field, i.e. no inter-diffusion was considered in the model. The material properties after the mixture were averaged over the volume fraction of the material as indicated in Equation 2, where
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is the averaged material property after mixture, f1 , 1 and f 2 , 2 are volume fractions and
material properties of SS316L and IN625, respectively.
F F F F u v w 0 t z y x
(1)
f11 f 22
(2)
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Temperature-dependent material properties were considered for both SS316L and IN625, which are listed in the Appendix (Fig. A-1 and Table A-1). Additionally, other material properties such as liquidus, solidus, viscosity, are also given in the appendix (Table A-2).
4. Results and discussion
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This section presents a detailed analysis of microstructures and crystallographic features. CFD analysis is performed to explain the material flow and differences in re-melting. The
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calculations of phase diagram (CALPHAD) approach is used to predict liquidus temperatures to help explain the solidification behavior at these two interfaces during the cooling stage [19]. Based
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on the analysis, formation mechanisms of these two types of interfaces are identified. Crack behavior at and near the Type-II interface is also elucidated.
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4.1. Morphology of the DED dual-material structure
Figs. 3(a) and (b) present the optical macrostructures of the etched cross-sections on the
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XZ and YZ planes, respectively. The printed structure can be easily demarcated into the top (IN625), middle (SS316L), and bottom (IN625) regions. Two distinct interfaces are formed as a result of switching printing sequences, i.e. the SS316L/IN625 interface (the upper interface) by
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depositing the IN625 on the SS316L, and the IN625/316L interface (the lower interface) formed by the deposition of SS316L on the IN625. In this research, the SS316L/IN625 interface and the IN625/316L interface hereinafter are referred to as the Type-I interface and the Type-II interface, respectively. A multilayer structure with regular and continuous molten pool tracks is evident in the printed two bulk materials, which is formed due to layer-wise depositions with constant hatch
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spacing. A comparison of the two interfaces on the XZ plane (Fig. 3a) reveals a wavy feature (high degree of molten pool banding) for Type–II interface in contrast to a relatively flat Type–I interface. This difference suggests a deeper penetration of SS316L into IN625 (Type-I interface) than that of IN625 into SS316L (Type-II interface). Figs. 3(c) and (d) present the magnified images of molten pool tracks of SS316L and IN625 on the XZ plane, respectively. The molten pool track of SS316L reveals two distinct microstructures, where the upper portion depicts an equiaxed dendrite structure while the bottom portion is dominated by a cellular structure combined with cellular 7
dendrites. In contrast, the molten tracks of IN625, presented in Fig. 3(d), are predominated by columnar dendrites. Figs. 3(e) and (f) present the magnified images of the Type-I interface and Type-II interface, respectively. Spherical particles, as marked by arrows in Figs. 3(c) and (f), with a diameter ranging from 30 to 100 μm are observed in the upper portion of the molten tracks of SS316L. As the spherical particle size is similar to that of the virgin SS316L metal powder, the spherical particles are rendered to be the un-melted or partially-melted SS316L powders, which is a typical observation associated with DED or SLM parts [20,21].
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CFD analysis was performed to understand the molten pool formation and the material flow near and at these two types of interfaces. The analysis considered the fact that the flow field
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of the molten pool determines the morphology of the faying interface and the distributions of the elements within the molten track. The simulated cross-sections of the XZ plane at these two types
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of interfaces are shown in Fig. 4. The contour color is presented by melting ratio, a ratio between the volume of the melted region and the total volume considered. From the simulation, it can be seen that the Type-I interface depicts a shallower penetration (red part) than the Type-II interface,
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which agrees with the observations in Figs. 3(e) and (f). , Therefore, the re-melting area of SS316L above the Type-I interface is significantly larger than the re-melting area of IN625 above the Type-
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II interface (Figs. 4a and b). This observation indicates that SS316L acting as a substrate in TypeI interface undergoes more re-melting than IN625 acting as a substrate in the Type-II interface. This difference can be explained by the lower energy absorption rate and lower thermal
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conductivity of IN625 relative to SS316L (the values can be found in Table A-2 and Fig. A-1 in the Appendix).
4.2. Crystallographic features of printed bulk materials EBSD analysis was performed to characterize the grain morphology and orientation for the
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printed base materials. Figs. 5(a) and (b) illustrate the inverse pole figures (IPFs) in the printing direction (Z-direction) for SS316L and IN625, respectively. It is believed that during solidification, decreasing temperature gradient and increasing solidification rate from the bottom to the top within a single molten pool track of SS316L is responsible for the gradient grain morphologies (i.e. columnar grains in the bottom and equiaxed grains formation in the top, as illustrated in Fig. 5a). The growth direction of columnar grains tends to be normal to the interfaces between tracks, i.e. parallel to the temperature gradient direction during solidification. Epitaxial grain growth in re8
melting is responsible for the grain morphology near the interfaces [22], where grains tend to maintain the growth orientation of their adjacent grains from the previous track. The generation of semicircular curved grains can be observed across the track boundaries, as marked by arrows in Fig. 5(a), which is caused by side-branch growth from the existing grains induced by epitaxial growth [23]. For the IN625 printed base metal, coarse columnar grains emerge as the dominant structure. The long columnar grains in red color, indicating [100] direction, are presented at the centerline
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of tracks. This preferred growth orientation is led by the epitaxial grain growth along the centerline with a consistent thermal gradient direction (Z-direction) during continuous re-melting with the
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deposition of additional layers [23]. It is imperative to note that such continuity of grain growth on the centerline is absent in SS316L (Fig. 5a), which is attributed to the clustering of the fine
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equiaxed grains in the top region and further inhibits the columnar growth.
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Texture evolutions in two printed base materials are examined on the YZ plane. Figs. 6(a) and (b) show the IPFs of SS316L and IN625, respectively. From Fig. 6(a), slanted grains with a
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similar inclination angle, approximately 30°, can be seen in a single layer of SS316L while the inclination direction changes between adjacent layers. For the IN625, the grain orientation is along the [100] direction (Fig. 6b). The YZ plane is cross-sectioned along the centerline of IN625 tracks,
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employing that the grain growth direction and crystal orientation are essentially consistent for different layers, i.e. parallel to the Z-direction. To further evaluate the grain orientation and texture for SS316L, pole figures (PFs) for the three layers in Fig. 6(a) are plotted in Figs. 6(c), (d) and (e), respectively. A fiber texture in the [100] direction is observed for the three layers, while the angle between the texture direction and Z-axis is about 30°, which is consistent with the observed growth
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direction. Additionally, Fig. 6(f) plots the PF in the region D in Fig. 6(b), which confirms the fiber texture in the [100] direction. The texture features in SS316L and IN625 are attributed to the γ phase with the preferential growth direction of [100]. 4.3. Characterization of the two types of interfaces 4.3.1. Compositional variation of interfaces
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SEM coupled with EDS was implemented to reveal the compositional variation near the two types of interfaces on the XZ and YZ planes (Fig. 7). For the Type-I interface, a composite zone is evident on the IN625 side with a decreasing concentration level of Fe along the build direction (Z-direction) (Figs. 7a and b). For the Type-II interface, as shown in Figs. 7(c) and (d), the streaks (i.e. unmixed zones) in the SS316L exhibit significant accumulations of Ni, Nb, and Mo elements. In addition, the pattern of the unmixed zones (streaks) on the YZ plane, shown in Fig. 7(d), indicates that the directions of flow velocity vectors (marked by black arrows) of the re-
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melted IN625 at the Type-II interface are roughly opposite to the laser scanning direction. In order to gain insights into the different mixing behaviors, the flow fields of the molten
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pools of the Type-I interface and Type-II interface are both analyzed on the YZ plane by numerical simulation (Fig. 8). The black arrows indicate the flow velocity in the molten pool, which also
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means the area without arrows is solid status at this time. As shown in Fig. 8(a), an elongated anticlockwise flow field (Region A) is evident above the Type-I interface. The flow at the rear of molten pool (the area with back arrows), marked by a black rectangle, brings the re-melted SS316L
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adjacent to the interface forward into the front area with a higher temperature and better fluidity, which promotes a fully mixed zone (composite zone). Fig. 8(b) provides the flow field of the Type-
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II interface. Two distinct zones can be discerned, i.e. one is at the front of the molten pool with an anticlockwise flow (Region B), and the other is at the rear of the molten pool with a clockwise flow (Region C). The front anticlockwise flow tends to produce the composite zone; however, the
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clockwise flow at the rear of the molten pool drives the re-melted IN625 at the interface to flow in the opposite direction of laser scanning, then turn upwards into the molten pool and finally be captured by the subsequent solidification.
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Fig. 9 presents the magnified SEM images and EDS line scanning plots for the Type-I and
Type-II interfaces. For the Type-I interface (Fig. 9a), a distinct boundary is seen between the SS316L and IN625. However, the elemental distribution plot across the Type-I interface (Fig. 9b), indicates a transition zone with a width of ~ 50 µm, where the main alloying elements, Fe and Ni, exhibit a gradual transition. The grains within the transition zone tends to be cellular dendrites and further evolve into dendrites as they enter in the IN625. In contrast, the EDS line scanning plot in Fig. 9(d) confirms a high elemental ratio of Ni in the transition zone and presents a narrow 10
compositional sudden-change zone (CSCZ) in the transition zone adjacent to the SS316L. As shown in Fig. 9(c), a primary dendritic arm can be observed in the CSCZ, while its secondary arms develop into a cellular structure on both sides. Dev et al. [24] also observed a similar feature at the boundary between the filler metal and Inconel 718 in gas tungsten arc welding of stainless steel AISI 416 and Inconel 718. The authors concluded two key underlying formation mechanisms of the transition zone, which are the difference in the melting temperatures of the two materials as a
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consequence of the difference in chemical composition, and the sluggish nature of molten IN625.
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4.3.2 Crystallographic features of interfaces
Fig. 10 presents the IPFs of the Type-I and Type-II interfaces on the XZ and YZ planes. For the Type-I interface (Figs. 10a and b), the IN625 grains grow epitaxially on the fine columnar
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and equiaxed grains of SS316L, leading to grains with different shapes and crystal orientations across the interface. As the grains grow further, the grains with the [001] orientation closer to the
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local temperature gradient direction tend to be more competitive, thus evolve into coarse columnar grains with a preferential growth direction. Figs. 10(c) and (d) present the grain morphologies for
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the Type-II interface. Unlike the Type-I interface, it is seen that limited grains are across the TypeII interface, which implies that the epitaxial grain growth is inhibited as the SS316L solidifies on the IN625. Further, irregular-shaped coarse grains are observed in the composite zone. It is also
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noticed that composition plays an important role in solidification and grain growth. Referring to the EDS mapping in Figs. 7(c) and (d), the grains in the Ni-enrichment region (unmixed zone) solidify lastly as Ni has a lower melting point compared to Fe, and thus display a unique curved
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geometry shown in Figs. 10(c) and (d).
4.3.3. Formation mechanisms of interfaces Based on the EDS and EBSD analyses at the interfaces, it is realized that the Type-I (IN625
deposition on SS316L) interface is featured by a gradual compositional transition and a typical epitaxial growth, while the Type-II (SS316L deposition on IN625) interface has an abrupt compositional transition and is absent of epitaxial growth.
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The gradual-change of composition at the Type-I interface indicates a mixure of IN625 and SS316L, which is mainly attributed to the lower density and lower viscosity of SS316L compared to IN625. The lower density and viscosity boost the re-melted SS316L to flow upwards and mix into the IN625. On the contrary, for the Type-II interface, these factors hinder the mixing of the IN625 into the SS316L. To illuminate the solidification behavior at these two interfaces during the cooling stage, liquidus temperatures across the interfaces were calculated based on the elemental distributions
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software
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(Fig. 9) and the CALPHAD. Fig. 11 shows the calculated liquidus temperatures using the Thermothermodynamic
database
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(https://www.thermocalc.com). Besides, the schematic diagrams for illustrating the formation behaviors of the Type-I and Type-II interfaces during solidification are presented in Figs. 12 and
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As shown in Fig. 11(a), the Type-I interface presents a gradual decrease of liquidus
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temperature in the transition zone from the SS316L to IN625, while the actual temperature across the interface increases from SS316L to IN625 due to the laser heating source above, which inhibits
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protrusions of SS316L into the transition zone and thus produces a relatively smooth interface (Fig. 12a). As illustrated in Fig. 12(b), when the cooling begins, nucleation starts from the interface between solid SS316L and liquid transition zone (the bottom dashed line in Fig. 12) due to its
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higher undercooling than the top region. A constitutional undercooling region forms subsequently ahead of the solidification front as the presence of partitioned solutes promotes the cellular dendrites to grow in the transition zone (Fig. 12c). Epitaxial growth is seen from the interface between the transition zone and IN625 (the upper dashed line in Fig. 12) and grains evolve into the dendritic structures in the IN625 due to the compositional change (Fig. 12d).
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As shown in Fig. 11(b), because of the uneven mixing, a slight increase of the liquidus
temperature in the transition zone is observed, causing a typical mushy zone (Fig. 13a). A sudden rise of liquidus temperature is seen in the CSCZ located between the transition zone and SS316L (Fig. 11b). The rise in liquidus temperature increases the local undercooling, and thus generates a potential nucleation region. When the cooling begins, nucleation occurs at both the solid fronts of IN625 and CSCZ (Fig. 13b). As shown in Fig. 13(c), the grain growth at the solid front of IN625 follows a typical epitaxial growth. However, driven by the higher undercooling in CSCZ compared 12
to the adjacent transition zone and liquid SS316L, horizontal to CSCZ is the preferred direction to form the primary dendritic arm, while the resultant partitioned solutes further induce a constitutional undercooling and promote the growth of secondary dendritic arms. The grains in the CSCZ continue to grow towards the IN625 and form a cellular structure together with the dendrites growing from the IN625 side and finally develop the transition zone (Fig. 13d). This structure in the absence of epitaxial growth at the Type-II interface is detrimental for interfacial mechanical properties due to the interrupted grain growth. Besides, this solidification behavior increases the
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possibility of solidification cracking as a result of the intersection of grains growing from the two
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sides of the transition zone.
4.4. Cracks formation and its mechanisms
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Forescatter diodes (FSD) equipped on an EBSD detector were employed to detect the defects (i.e. porosity and cracks) as shown in Fig. 14. The images revealed no obvious defects in
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the printed base metals and the Type-I interface (Figs. 14a-c). However, cracks are evident at both the Type-II interface and SS316L tracks adjacent to the Type-II interface (Fig. 14d). Moreover,
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limited porosities (black dots) can be seen in the bulk materials and interfaces.
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Further examinations on probing the crack formation mechanism were conducted through SEM and EDS. Fig. 15 presents the magnified SEM image of the cracks located at the Type–II interface. The cracks are predominantly located at the boundary between the cellular structure grew from the CSCZ and the IN625 dendrites. The EDS mapping indicates that there is no obvious elemental enrichment in the crack confirming solidification cracking that happened at the Type–II
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interface. As mentioned in Section 4.3.3, during the solidification of the Type-II interface, the grain growth in the transition zone from the two sides (i.e. the CSCZ and IN625) is responsible for the solidification cracking due to lack of liquid metal to fill the gap between solidifying metals.
Fig. 15. SEM images and EDS mapping for the cracks at the Type-II interface.
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Fig. 16(a) shows the zigzag cracks in SS316L, where the dash lines indicate the track boundaries. Three segments of the zigzag crack are named as crack #1, crack #2 and crack #3. The corresponding Kernel average misorientation (KAM) mapping is extracted by EBSD to characterize the local strain distributions around the cracks, as indicated in Fig. 16(b). Three corresponding locations from each crack marked as A, B and C are chosen to perform the EDS mapping, and the corresponding results are plotted in Figs. 16, 17 and 18, respectively. The crack at location A is resultant from ductility dip cracking (DDC), where the cracks happened along the
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grain boundaries, as seen in Fig. 17(a). This can also be confirmed from Figs. 17(b-e) that no liquidation or elemental accumulation occurred at this location. DDC is induced during solidification by the strain accumulation at grain boundaries and a sharp drop in ductility during
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the temperature range from one-half of the alloy solidus to the solidus temperature [25]. Qian et al. [26] further explained the formation mechanism of DDC from the perspective of dislocation
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migration. As the thermal tension stress introduced by temperature gradient exceeds the yield stress, the dislocation slip starts within the grains followed by their pile up at the high angle grain
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boundaries, enhancing the strain and stress accumulations and eventually causes DDC.
Crack #2 has a wavy shape (Fig. 16a) where elemental enrichment is barely visible (Fig. 18), which indicates a solidification crack formation. In addition, bidirectionally grown columnar
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dendrites can be found adjacent to the crack as marked by the arrow in Fig. 18(a), which further confirms the cracking occurred during the solidification. The formation mechanism is similar to the cracks at the Type-II interface, where the insufficient supply of liquid metal to fill the space between solidifying metals from opposite sides and space opens further under shrinkage strain. As exhibited in Fig. 19, for the location C of crack #3, the EDS mapping reveals noticeable
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elemental accumulation of Nb and Mo in the crack, indicating the formation of liquidation cracking. From Fig. 16(a) and EDS mapping in Fig. 7(c), it is found that the crack is located at the unmixed zone with a higher concentration of Nb and Mo, resulting in increased susceptibility of microsegregation of the Nb and Mo which are deposited as long-chain Laves phase at the grain boundaries under a relatively low cooling rate [27,28]. With the deposition of subsequent layers, the Laves phase and its surrounding γ phase undergo eutectic remelting at the grain boundary under
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reheating, resulting in crack initiation, and the long-chain configuration promotes the crack propagation.
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Author statement Nannan Chen: Formal analysis, Investigation, Writing- Original draft preparation. Haris Ali Khan.: Data curation, Writing- Review & Editing. Zixuan Wan: Software, Data curation. Investigation. Writing- Original draft preparation. John Lippert: Data curation, Writing- Original draft preparation.: Hui Sun: Software, Data curation. Shun-Li Shang: Formal analysis, Writing- Original draft preparation. Zi-Kui Liu: Funding acquisition, Supervision. Jingjing Li: Supervision, Project administration, Funding acquisition, Writing - Review & Editing.
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5. Conclusions
The importance of printing sequence in a DED processed bimetal structure of SS316L and
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IN625 was interpreted by probing interfacial characteristics and crack generation, and their formation mechanisms were further unearthed by microstructural and crystallographic
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characterizations together with numerical simulation of molten pool and flow behavior. The main conclusions are as follows:
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1. The continuous growth of large columnar grains at the centerlines was observed in IN625 because of the epitaxial growth between the layers. In the SS316L, the fine equiaxed dendrites
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were formed at the top of each track which prevented the continual growth of columnar grains. 2. Compositional variation at the two interfaces was explained aided by numerical simulation. For the Type-I interface, the anticlockwise flow at the rare of the molten pool facilitated a fully mixed zone by driving the re-melted SS316L move forward into the front area with higher temperature and better fluidity. For the Type-II interface, the unmixed streaks enriched by the IN625 was
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responsible for the clockwise flow at the rear of the molten pool, which boosts the re-melted IN625 to move upwards into the molten pool and be captured by the solidification. 3. The formation mechanisms of the Type-I and Type-II interfaces were explained through the integration of experimental results and calculations of liquidus temperature. For the Type-I interface, a gradual-change compositional transition zone was observed because the lower density and lower viscosity of SS316L compared to IN625 promotes a good mixing. Epitaxial growth was formed at the Type-I interface due to the gradual decrease in the liquidus temperature from SS316L 15
to IN625 as a result of the compositional variation. For the Type-II interface, the higher density and higher viscosity of IN625 hindered the compositional mixing and thus generated a suddenchange zone. The resultant sudden rise of liquidus temperature finally created a unique interface without epitaxial growth. 4. Type-II interface and the SS316L tracks near the Type-II interface were found susceptible to crack formation. The formation of solidification cracks at the interface and track boundaries were owing to bidirectional solidification. Moreover, the unmixed zone from the re-melted IN625
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caused microsegregation of Nb and Mo, leading to liquidation cracking. The solidification and liquidation cracking further increased the susceptibility of ductility dip cracking which were
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generated under solidus temperature. Acknowledgments
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The authors acknowledge the US National Science Foundation Civil, Mechanical and Manufacturing Innovation Grants No. 1651024. H.S., S.L.S., and Z.K.L. acknowledge the
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financial support from the Office of Naval Research (ONR) under contract No. N00014-17-1-2567.
Appendix Temperature-dependent material properties applied in the numerical model in Section
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3.
This appendix reprensts a complilatin of important properties used in the CFD simulation [29-31]. The temperature-dependent thermal properties, including specific heat capacity, thermal conductivity, and density of SS316L and IN625 are presented in Fig. A-1. Table A-1 summerizes the temperature-dependent energy absorption rates of these two materials at room temperature,
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solidus and liquidus temperatures; and other material properties, such as liquidus and solidus temperatures, latent heat of fusion, viscosity, surface tension and latent heat of vaporization, are listed in Table A-2. Note that the viscosity of IN625 is referred from the information of IN718 [31].
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properties, Archives of Civil and Mechanical Engineering 19(3) (2019) 820-831. [18] X. Mei, X. Wang, Y. Peng, H. Gu, G. Zhong, S. Yang, Interfacial characterization and mechanical properties of 316L stainless steel/Inconel 718 manufactured by selective laser melting, Mat. Sci. Eng. A 758 (2019) 185-191. [19] Z.-K. Liu, First-principles calculations and CALPHAD modeling of thermodynamics, J. Phase Equilibria Diffus. 30 (2009) 517–534.
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steel, AISI 416. Journal of Manufacturing Processes 32 (2018) 685-698. [25] J. Q. Chen, H. Lu, C. Yu, J. M. Chen, M. L. Zhang, Ductility dip cracking mechanism of Ni–
18(4) (2013) 346-353.
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Cr–Fe alloy based on grain boundary energy. Science and Technology of Welding and Joining
[26] D. Qian, J. Xue, A. Zhang, Y. Li, N. Tamura, Z. Song, K. Chen, Statistical study of ductility-
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dip cracking induced plastic deformation in polycrystalline laser 3D printed Ni-based superalloy. Scientific Reports 7(1) (2017) 2859.
[27] Y. Chen, K. Zhang, J. Huang, S. R. E. Hosseini, Z. Li, Characterization of heat affected zone liquation cracking in laser additive manufacturing of Inconel 718. Materials & Design, 90, (2016)
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586-594.
[28] Y. C. Zhang, Z. G. Li, P. L. Nie, Y. X. Wu, Effect of ultrarapid cooling on microstructure of laser cladding IN718 coating. Surf. Eng. 29 (2013) 414–418. [29] C. S. Kim, Thermophysical properties of stainless steels (No. ANL-75-55). Argonne National Lab., Ill. (USA) (1975).
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[30]
Special
Metals
datasheet
Inconel
alloy
625,
http://www.specialme
tals.com/assets/smc/documents/alloys/inconel/inconel-alloy-625.pdf. [31] R. A. Overfelt, C. A. Matlock, M. E. Wells, Viscosity of superalloy 718 by the oscillating
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vessel technique. Metallurgical Transactions, B 27(4) (1996).
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Fig. 1. Illustrations of (a) DED sample structure, and (b) laser scanning pattern.
Fig. 2. The setup of numerical simulation: (a) mesh block of the model, (b) dimensions of the initial fluid cross-section, and (c) laser power density followed the Gaussian distribution.
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Fig. 3. Optical macrostructure of cross-sections on (a) XZ plane, (b) YZ plane; magnified optical
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images with inset SEM micrographs on the XZ cross-section: (c) SS316L, (d) IN625, (e) Type-I
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interface, and (f) Type-II interface.
Fig. 4. Simulated cross-sections of (a) Type-I interface (IN365 is on the top of SS316L), and (b) Type-II interface (SS316L in on the top of IN365).
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Fig. 5. IPFs of (a) SS316L and (b) IN625 on XZ plane.
Fig. 6. IPFs of (a) SS316L and (b) IN625 on YZ plane, {100} pole figures extracted from: (c) region A, (d) region B, (e) region C of SS316L, and (f) region D of IN625 (SD refers to the scanning direction).
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Fig. 7. SEM images and EDS mapping for Type-I interface (IN625 on the top of SS316L) on (a)
planes.
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XZ and (b) YZ planes, and Type-II interface (SS316L on the top of IN625) on (c) XZ and (d) YZ
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Fig. 8. The simulated flow field of the (a) Type-I interface and (b) Type-II interface on the YZ
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plane.
Fig. 9. Magnified morphologies and elemental distributions: (a) SEM image and (b) EDS line scanning plot of Type-I interface, (c) SEM image and (d) EDS line scanning plot of Type-II
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interface.
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Fig. 10. IPFs for the Type-I interface of (a) XZ and (b) YZ planes, and the Type-II interface of
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(c) XZ and (d) YZ planes.
Fig. 11. Calculated liquidus temperatures across (a) the Type-I interface and (b) the Type-II interface based on the elemental distributions in Fig. 9 using the CALPHAD approach. 26
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Fig. 12. Schematic illustration of the formation mechanism of the Type-I interface.
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Fig. 13. Schematic illustration of the formation mechanism of the Type-II interface.
Fig. 14. FSD images on XZ cross-section: (a) SS316L, (b) IN625, (c) Type-I interface, and (d)
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Type-II interface.
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Fig. 15. SEM images and EDS mapping for the cracks at the Type-II interface.
Fig. 16. (a) SEM image and (b) corresponding KAM image for the zigzag cracks in SS316L
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tracks near the Type-II interface (i.e. the upper crack in Fig. 14d).
Fig. 17. (a) SEM image of Crack #1 (Location A in Fig. 16), and (b)-(e) corresponding EDS mapping for elements Fe, Ni, Cr, and Mo, respectively.
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Fig. 18. (a) SEM image of crack #2 (Location B in Fig. 16), and (b)-(e) corresponding EDS
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mapping for elements Fe, Ni, Cr, and Mo, respectively.
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Fig. 19. (a) SEM image of crack #3 (Location C in Fig. 16), and (b)-(f) corresponding EDS
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mapping for elements Fe, Ni, Cr, Nb, and Mo, respectively.
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Fig. A-1. Temperature-dependent thermal properties: specific heat capacity of (a) SS316L and (b) IN625, thermal conductivity of (c) SS316L and (d) IN625, and density of (e) SS316L and (f) IN625 [29,30].
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Table 1. Chemical composition of SS316L and IN625 powders. Powder
Fe
Ni 1014
SS316L Rem. 5
Rem.
C
Si
0.03
1
2
0.1
0.5
0.5
0.045
0.03
0.015 0.015
Mo
-
2-3
3.154.15
8-10
Cr 1618 2023
Ti
Al
-
-
0.4
0.4
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IN625
Elements (Wt.%) Mn P S Nb
Table 2. DED process parameters. Powder Flow Rate (g/min) 20.5 20.5
Powder Feeder Vibration (%) 60 60
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Travel Speed (IPM)
SS316L IN625
2 2
25 25
Step Over Per Pass (mm) 2.5 2.5
Focus Head Standoff (mm) 10 10
Spot Size (mm) 4 4
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Powder
Laser power (kW)
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Table A-1. Temperature-dependent energy absorption rate (A) of SS316L and IN625 for a 1060 nm laser beam [29,30]. SS316L A (%)
Temperature (K)
IN625 A (%)
293
38%
293
30%
1673 (Solidus)
40%
1563 (Solidus)
30%
1698 (Liquidus)
45%
1623 (Liquidus)
35%
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Temperature (K)
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IN625 1623 1563 3.1e5 7.40 1.88 6.35e6
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Liquidus (K) Solidus (K) Latent heat of fusion (J/kg) Viscosity (mPa·s) Surface tension (N/m) Latent heat of vaporization (J/kg)
SS316L 1698 1673 2.6e5 6.42 1.80 6.25e6
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Table A-2. Other material properties of SS316L and IN625 [29-31].
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PII:
S2214-8604(19)31472-1
DOI:
https://doi.org/10.1016/j.addma.2020.101037
Reference:
ADDMA 101037
To appear in:
Additive Manufacturing
Received Date:
2 September 2019
Revised Date:
4 December 2019
Accepted Date:
1 January 2020
Please cite this article as: Chen N, Khan HA, Wan Z, Lippert J, Sun H, Shang S-Li, Liu Z-Kui, Li J, Microstructural characteristics and crack formation in additively manufactured bimetal material of 316L stainless steel and Inconel 625, Additive Manufacturing (2020), doi: https://doi.org/10.1016/j.addma.2020.101037
This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2020 Published by Elsevier.
Microstructural characteristics and crack formation in additively manufactured bimetal material of 316L stainless steel and Inconel 625 Nannan Chen,1 Haris Ali Khan,1 Zixuan Wan,1 John Lippert,1 Hui Sun,2 Shun-Li Shang,2 Zi-Kui Liu,2 Jingjing Li 1* 1
Department of Industrial and Manufacturing Engineering, Pennsylvania State University, University Park, PA 16802, USA
Department of Materials Science and Engineering, Pennsylvania State University, University
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2
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Park, PA 16802, USA
* To whom correspondence should be addressed. E-mail: [email protected], Tel: 1-814-863-1300, Fax: 1-814-863-4745. 1
Highlights • Changing printing sequence generated two types of interfaces with unique morphology • IN625 grains grew epitaxially on the fine grains of SS316L forming Type-I interface • Bidirectional nucleation from IN625 and mushy zone at SS316L formed Type-II interface • Cracking was formed at Type-II interface and in the SS316L tracks • Cracking mechanisms include solidification, liquidation, and ductility dip cracking
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Abstract This research illustrates the rationale of adopting a preferred printing sequence by examining crack
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generation predominated by resultant interfaces and microstructural inhomogeneity, through underlying governing mechanisms in directed energy deposition of 316L stainless steel/Inconel 625 (SS316L/IN625) bimetals. For this purpose, microstructural and crystallographic
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characterizations augmented by numerical simulations were employed on additively manufactured two distinct interfaces, i.e. Type-I (IN625 deposition on SS316L) and Type-II (SS316L deposition
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on IN625). Changing the printing sequence generated these two types of interfaces with unique morphologies, which was found attributable to the compositional variations and mismatch in
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thermal properties. Type-I interface was typified by gradual-change composition in the transition zone, causing the IN625 grains to grow epitaxially on the grains of SS316L. Type-II interface was characterized as a compositional sudden-change zone (CSCZ) adjacent to SS316L, leading to
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merging bidirectional nucleation and grain growth from the bottom IN625 and upper CSCZ, and lack of epitaxial growth. Additionally, high cracking susceptibility occurred near the Type-II interface rather than the Type-I interface, which was related to solidification and liquidation cracking, and further promoted ductility dip cracking. This research will provide a guideline for the additive manufacturing of bimetals with the consideration of printing sequence to control
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interface formation for a crack-free structure.
Keywords: directed energy deposition; bimetal; microstructure; crack formation; interfacial behavior
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1. Introduction Extreme environment applications, for instance, nuclear plants and aerospace vehicles involve components with requisite of varied properties at different locations [1, 2]. However, it is seldom or impossible to leverage such characteristics from a single material, and this need drives the unremitting pursuit of tailored properties by developing dissimilar materials structures. Different manufacturing techniques have been employed to date for their fabrication. Additive manufacturing (AM) processes, particularly directed energy deposition (DED) technique, emerge
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as a promising opportunity for manufacturing dissimilar materials structures by a layered approach with careful selection of process parameters [3] or by utilizing composition as a spatial design
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parameter [4]. Zhang and Bandyopadhyay [3] demonstrated the feasibility of fabricating metalceramic multi-layer structures with a laser-based AM method. Carroll et al. [5] obtained a bimetal
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structure of 304L stainless steel to Inconel 625 (IN625) using DED through a gradient zone to avoid sharp compositional or microstructural changes at the interface. Shah et al. [6] analyzed the influences of process parameters (e.g. laser power levels and powder mass flow rates) on
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microstructure and mechanical properties in DED processed Inconel-steel bimetal graded structure, where the authors demonstrated that desired mechanical properties can be achieved by controlling
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the input parameters. Onuike et al. [7] utilized three different building strategies (namely direct deposition, compositional gradation, and use of an intermediate bond layer) to produce DED processed bimetal structure of Ti-64 and IN718 alloys. Based on the microstructure and
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mechanical properties, the authors defined an intermediate bond layer as the most suitable technique. DED processed Inconel and copper alloy bimetal structure was produced for improved thermal conductivity behavior in high-temperature aerospace applications [8]. The DED process was also utilized for obtaining intricate bimetal structures with varied properties. For instance, bimetallic stainless tube structure with transitions directly from 316L stainless steel (SS316L) to
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SS430L was produced to effectively obtain transition from non-magnetic to magnetic material [9]. The above discussion manifests that AM processes are favorable in the fabrication of multimaterial into a single structure. Nickel-based superalloys (e.g. IN718 and IN625) and different grades of stainless steel (e.g. SS304L and SS316L) are widely used in nuclear reactors, aerospace, automotive and petrochemical components for their improved corrosion and oxidation resistance, and mechanical
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strength [10-12]. These two alloys offer a promising opportunity in dissimilar materials structures, where an increased strength at elevated temperature (where IN625 is suitable) and corrosion resistance (where SS316L is suitable) is achievable at different locations simultaneously [13]. Nonetheless, utilizing the two materials together bring forth perplexing challenges, including inhomogeneity of the microstructure introduced by lack of fusion, nonuniform melting and rapid solidification, which can cause deterioration of mechanical properties and microsegregation of the alloying elements during the solidification process [14]. In addition, deposition of a dissimilar
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material on these metals is susceptible to problems such as interface cracking and carbon diffusion from the low chromium to high chromium alloy [15,16].
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It is observed that altering the printing sequence changed the interfacial characteristics in bimetal structures. For instance, when using DED to print IN718 and SS316L structure, Li et al.
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[17] found different interfacial characteristics with alternated printing schemes. The authors observed a clear interface devoid of any diffusion layer when IN718 was treated as a substrate while the same parameters yielded a diffusion layer with greater microhardness than SS316L when
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IN718 was deposited on to SS316L. This phenomenon may be related to the nature of microstructural evolution particularly at and near the interface region during the heating and
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cooling stages of the layer. In another research [18], different defects were observed with the different printing sequence (i.e. cracks close to the IN718/316SS interface, and holes at the 316SS/IN718 interface, where the first material was on the top). In essence, it has been reported
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that the printing or deposition scheme is a dominant factor that not only influences the interfacial microstructure but also if appropriately selected, can help in mitigating the undesirable inhomogeneity of the microstructure and other defects (e.g. pores and cracks). Current publications focus on reporting the phenomena; however, it is scarce for scientific understanding that can explain the rationale of selecting a particular printing scheme based on the resultant interfacial
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microstructure and subsequent defects formation. This research targets at ascertaining the formation mechanisms of the distinct interfacial
characteristics and microstructural evolution during the DED process of IN625 and SS316L alloys. Particular emphasis has been laid on establishing the relationship between different printing sequences and the subsequent generated microstructure and cracks. Well-integrated materials characterization techniques were employed to comprehend the microstructural features (i.e. grain
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morphology, texture, and phases), and their formation mechanisms and further influences on the microstructural inhomogeneity and cracks. The experimental results are also augmented by numerical analysis for ascertaining material flow. The remainder of the paper follows the defined sequence. Section 2 and section 3 describe the experimental procedure and basis of numerical analysis, respectively. Section 4 relates to the obtained results consisting of structure morphology, crystallographic analysis, and defects characterization. The paper is culminated by discussing key conclusions in section 5.
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2. Experimental procedure This section provides details of the DED process and explanations of the procedures for
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preparing metallographic samples and microstructural characterization adopted in this research.
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2.1. DED process
A rectangular bimetal structure was additively manufactured using SS316L and IN625
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powders with the DED method in an argon atmosphere. The dimensions of the built part were 110 mm (Length) x 10 mm (Width) x 65 mm (Height) as illustrated in Fig. 1(a), where the X, Y, and Z axes are the step-over, laser scanning, and building directions, respectively. The build started
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with 8 layers of SS316L depositions, followed by 8 layers of Inconel 625 depositions, and continuously alternated with those respective layers until the build finished. The laser scanning pattern is schematically illustrated in Fig. 1(b). Sandvik SS316L powder (particle sizes of 45-150
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µm) and Carpenter Technology’s IN625 alloy powder (particle sizes of 120-270 µm) were used and their chemical compositions are provided in Table 1.
The additive build was manufactured with the Penn State Applied Research Laboratory’s
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high-power high-deposition system which houses a custom cladding head with water-cooled optics and an inert gas enclosure, and the process parameters are listed in Table 2. 2.2. Microstructural characterization The bimetal structure was sectioned from XZ and YZ planes to characterize microstructures of the printed bulk materials (i.e. SS316L and IN625) and the interfaces between these two materials. The sectioned samples were first mounted in epoxy, then polished with silicon 5
papers following a sequence of #320, #600, and #800, and finally polished with diamond suspensions down to 1 m. The samples were further electrolytically etched with a 4% oxalic acid reagent for 5-10 seconds. Vibratory polishing was applied to prepare the samples for Electron Back-Scattered Diffraction (EBSD) analysis. Macrostructures were acquired under a Nikon optical microscope. A Nova nano scanning electron microscope (SEM) 630 field emission analysis equipped with an Oxford X-Max EDS (Energy Dispersive Spectroscopy) was employed to obtain the microstructures and elemental distributions. Finally, the crystallographic features were
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revealed by EBSD analysis performed on a Helios NanoLab 660 SEM equipped with an Oxford EBSD detector.
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3. Numerical analysis
DED process was simulated by commercial computational fluid mechanics (CFD) software
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with customized heat and focal subroutines. Two distinct interfaces, i.e. IN625 deposition on SS316L and SS316L deposition on IN625, were modeled separately. For simplification, a partial
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cylinder was added on the bottom metal to represent the deposited particles during the deposition process, as depicted in Fig. 2(a). A total number of 416,000 cells with 0.1 mm cell size were
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employed in the simulation. The dimensions of the initial fluid cross-sections are indicated in Fig. 2(b). The power density of the laser spot was assumed to follow the Gaussian distribution (Fig. 2c), where the peak power density in the central region was approximately 283 W/mm2. Free
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surfaces of both materials were tracked by volume-of-fluid, which can be represented by Equation 1, where F is fluid volume fraction; u, v and w represent velocity components in x, y and z directions, respectively. The mixture of the two different materials was completely driven by the flow field, i.e. no inter-diffusion was considered in the model. The material properties after the mixture were averaged over the volume fraction of the material as indicated in Equation 2, where
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is the averaged material property after mixture, f1 , 1 and f 2 , 2 are volume fractions and
material properties of SS316L and IN625, respectively.
F F F F u v w 0 t z y x
(1)
f11 f 22
(2)
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Temperature-dependent material properties were considered for both SS316L and IN625, which are listed in the Appendix (Fig. A-1 and Table A-1). Additionally, other material properties such as liquidus, solidus, viscosity, are also given in the appendix (Table A-2).
4. Results and discussion
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This section presents a detailed analysis of microstructures and crystallographic features. CFD analysis is performed to explain the material flow and differences in re-melting. The
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calculations of phase diagram (CALPHAD) approach is used to predict liquidus temperatures to help explain the solidification behavior at these two interfaces during the cooling stage [19]. Based
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on the analysis, formation mechanisms of these two types of interfaces are identified. Crack behavior at and near the Type-II interface is also elucidated.
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4.1. Morphology of the DED dual-material structure
Figs. 3(a) and (b) present the optical macrostructures of the etched cross-sections on the
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XZ and YZ planes, respectively. The printed structure can be easily demarcated into the top (IN625), middle (SS316L), and bottom (IN625) regions. Two distinct interfaces are formed as a result of switching printing sequences, i.e. the SS316L/IN625 interface (the upper interface) by
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depositing the IN625 on the SS316L, and the IN625/316L interface (the lower interface) formed by the deposition of SS316L on the IN625. In this research, the SS316L/IN625 interface and the IN625/316L interface hereinafter are referred to as the Type-I interface and the Type-II interface, respectively. A multilayer structure with regular and continuous molten pool tracks is evident in the printed two bulk materials, which is formed due to layer-wise depositions with constant hatch
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spacing. A comparison of the two interfaces on the XZ plane (Fig. 3a) reveals a wavy feature (high degree of molten pool banding) for Type–II interface in contrast to a relatively flat Type–I interface. This difference suggests a deeper penetration of SS316L into IN625 (Type-I interface) than that of IN625 into SS316L (Type-II interface). Figs. 3(c) and (d) present the magnified images of molten pool tracks of SS316L and IN625 on the XZ plane, respectively. The molten pool track of SS316L reveals two distinct microstructures, where the upper portion depicts an equiaxed dendrite structure while the bottom portion is dominated by a cellular structure combined with cellular 7
dendrites. In contrast, the molten tracks of IN625, presented in Fig. 3(d), are predominated by columnar dendrites. Figs. 3(e) and (f) present the magnified images of the Type-I interface and Type-II interface, respectively. Spherical particles, as marked by arrows in Figs. 3(c) and (f), with a diameter ranging from 30 to 100 μm are observed in the upper portion of the molten tracks of SS316L. As the spherical particle size is similar to that of the virgin SS316L metal powder, the spherical particles are rendered to be the un-melted or partially-melted SS316L powders, which is a typical observation associated with DED or SLM parts [20,21].
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CFD analysis was performed to understand the molten pool formation and the material flow near and at these two types of interfaces. The analysis considered the fact that the flow field
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of the molten pool determines the morphology of the faying interface and the distributions of the elements within the molten track. The simulated cross-sections of the XZ plane at these two types
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of interfaces are shown in Fig. 4. The contour color is presented by melting ratio, a ratio between the volume of the melted region and the total volume considered. From the simulation, it can be seen that the Type-I interface depicts a shallower penetration (red part) than the Type-II interface,
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which agrees with the observations in Figs. 3(e) and (f). , Therefore, the re-melting area of SS316L above the Type-I interface is significantly larger than the re-melting area of IN625 above the Type-
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II interface (Figs. 4a and b). This observation indicates that SS316L acting as a substrate in TypeI interface undergoes more re-melting than IN625 acting as a substrate in the Type-II interface. This difference can be explained by the lower energy absorption rate and lower thermal
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conductivity of IN625 relative to SS316L (the values can be found in Table A-2 and Fig. A-1 in the Appendix).
4.2. Crystallographic features of printed bulk materials EBSD analysis was performed to characterize the grain morphology and orientation for the
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printed base materials. Figs. 5(a) and (b) illustrate the inverse pole figures (IPFs) in the printing direction (Z-direction) for SS316L and IN625, respectively. It is believed that during solidification, decreasing temperature gradient and increasing solidification rate from the bottom to the top within a single molten pool track of SS316L is responsible for the gradient grain morphologies (i.e. columnar grains in the bottom and equiaxed grains formation in the top, as illustrated in Fig. 5a). The growth direction of columnar grains tends to be normal to the interfaces between tracks, i.e. parallel to the temperature gradient direction during solidification. Epitaxial grain growth in re8
melting is responsible for the grain morphology near the interfaces [22], where grains tend to maintain the growth orientation of their adjacent grains from the previous track. The generation of semicircular curved grains can be observed across the track boundaries, as marked by arrows in Fig. 5(a), which is caused by side-branch growth from the existing grains induced by epitaxial growth [23]. For the IN625 printed base metal, coarse columnar grains emerge as the dominant structure. The long columnar grains in red color, indicating [100] direction, are presented at the centerline
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of tracks. This preferred growth orientation is led by the epitaxial grain growth along the centerline with a consistent thermal gradient direction (Z-direction) during continuous re-melting with the
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deposition of additional layers [23]. It is imperative to note that such continuity of grain growth on the centerline is absent in SS316L (Fig. 5a), which is attributed to the clustering of the fine
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equiaxed grains in the top region and further inhibits the columnar growth.
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Texture evolutions in two printed base materials are examined on the YZ plane. Figs. 6(a) and (b) show the IPFs of SS316L and IN625, respectively. From Fig. 6(a), slanted grains with a
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similar inclination angle, approximately 30°, can be seen in a single layer of SS316L while the inclination direction changes between adjacent layers. For the IN625, the grain orientation is along the [100] direction (Fig. 6b). The YZ plane is cross-sectioned along the centerline of IN625 tracks,
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employing that the grain growth direction and crystal orientation are essentially consistent for different layers, i.e. parallel to the Z-direction. To further evaluate the grain orientation and texture for SS316L, pole figures (PFs) for the three layers in Fig. 6(a) are plotted in Figs. 6(c), (d) and (e), respectively. A fiber texture in the [100] direction is observed for the three layers, while the angle between the texture direction and Z-axis is about 30°, which is consistent with the observed growth
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direction. Additionally, Fig. 6(f) plots the PF in the region D in Fig. 6(b), which confirms the fiber texture in the [100] direction. The texture features in SS316L and IN625 are attributed to the γ phase with the preferential growth direction of [100]. 4.3. Characterization of the two types of interfaces 4.3.1. Compositional variation of interfaces
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SEM coupled with EDS was implemented to reveal the compositional variation near the two types of interfaces on the XZ and YZ planes (Fig. 7). For the Type-I interface, a composite zone is evident on the IN625 side with a decreasing concentration level of Fe along the build direction (Z-direction) (Figs. 7a and b). For the Type-II interface, as shown in Figs. 7(c) and (d), the streaks (i.e. unmixed zones) in the SS316L exhibit significant accumulations of Ni, Nb, and Mo elements. In addition, the pattern of the unmixed zones (streaks) on the YZ plane, shown in Fig. 7(d), indicates that the directions of flow velocity vectors (marked by black arrows) of the re-
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melted IN625 at the Type-II interface are roughly opposite to the laser scanning direction. In order to gain insights into the different mixing behaviors, the flow fields of the molten
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pools of the Type-I interface and Type-II interface are both analyzed on the YZ plane by numerical simulation (Fig. 8). The black arrows indicate the flow velocity in the molten pool, which also
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means the area without arrows is solid status at this time. As shown in Fig. 8(a), an elongated anticlockwise flow field (Region A) is evident above the Type-I interface. The flow at the rear of molten pool (the area with back arrows), marked by a black rectangle, brings the re-melted SS316L
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adjacent to the interface forward into the front area with a higher temperature and better fluidity, which promotes a fully mixed zone (composite zone). Fig. 8(b) provides the flow field of the Type-
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II interface. Two distinct zones can be discerned, i.e. one is at the front of the molten pool with an anticlockwise flow (Region B), and the other is at the rear of the molten pool with a clockwise flow (Region C). The front anticlockwise flow tends to produce the composite zone; however, the
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clockwise flow at the rear of the molten pool drives the re-melted IN625 at the interface to flow in the opposite direction of laser scanning, then turn upwards into the molten pool and finally be captured by the subsequent solidification.
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Fig. 9 presents the magnified SEM images and EDS line scanning plots for the Type-I and
Type-II interfaces. For the Type-I interface (Fig. 9a), a distinct boundary is seen between the SS316L and IN625. However, the elemental distribution plot across the Type-I interface (Fig. 9b), indicates a transition zone with a width of ~ 50 µm, where the main alloying elements, Fe and Ni, exhibit a gradual transition. The grains within the transition zone tends to be cellular dendrites and further evolve into dendrites as they enter in the IN625. In contrast, the EDS line scanning plot in Fig. 9(d) confirms a high elemental ratio of Ni in the transition zone and presents a narrow 10
compositional sudden-change zone (CSCZ) in the transition zone adjacent to the SS316L. As shown in Fig. 9(c), a primary dendritic arm can be observed in the CSCZ, while its secondary arms develop into a cellular structure on both sides. Dev et al. [24] also observed a similar feature at the boundary between the filler metal and Inconel 718 in gas tungsten arc welding of stainless steel AISI 416 and Inconel 718. The authors concluded two key underlying formation mechanisms of the transition zone, which are the difference in the melting temperatures of the two materials as a
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consequence of the difference in chemical composition, and the sluggish nature of molten IN625.
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4.3.2 Crystallographic features of interfaces
Fig. 10 presents the IPFs of the Type-I and Type-II interfaces on the XZ and YZ planes. For the Type-I interface (Figs. 10a and b), the IN625 grains grow epitaxially on the fine columnar
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and equiaxed grains of SS316L, leading to grains with different shapes and crystal orientations across the interface. As the grains grow further, the grains with the [001] orientation closer to the
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local temperature gradient direction tend to be more competitive, thus evolve into coarse columnar grains with a preferential growth direction. Figs. 10(c) and (d) present the grain morphologies for
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the Type-II interface. Unlike the Type-I interface, it is seen that limited grains are across the TypeII interface, which implies that the epitaxial grain growth is inhibited as the SS316L solidifies on the IN625. Further, irregular-shaped coarse grains are observed in the composite zone. It is also
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noticed that composition plays an important role in solidification and grain growth. Referring to the EDS mapping in Figs. 7(c) and (d), the grains in the Ni-enrichment region (unmixed zone) solidify lastly as Ni has a lower melting point compared to Fe, and thus display a unique curved
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geometry shown in Figs. 10(c) and (d).
4.3.3. Formation mechanisms of interfaces Based on the EDS and EBSD analyses at the interfaces, it is realized that the Type-I (IN625
deposition on SS316L) interface is featured by a gradual compositional transition and a typical epitaxial growth, while the Type-II (SS316L deposition on IN625) interface has an abrupt compositional transition and is absent of epitaxial growth.
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The gradual-change of composition at the Type-I interface indicates a mixure of IN625 and SS316L, which is mainly attributed to the lower density and lower viscosity of SS316L compared to IN625. The lower density and viscosity boost the re-melted SS316L to flow upwards and mix into the IN625. On the contrary, for the Type-II interface, these factors hinder the mixing of the IN625 into the SS316L. To illuminate the solidification behavior at these two interfaces during the cooling stage, liquidus temperatures across the interfaces were calculated based on the elemental distributions
Calc
software
and
the
TCFE8
Steels/Fe-alloys
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(Fig. 9) and the CALPHAD. Fig. 11 shows the calculated liquidus temperatures using the Thermothermodynamic
database
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(https://www.thermocalc.com). Besides, the schematic diagrams for illustrating the formation behaviors of the Type-I and Type-II interfaces during solidification are presented in Figs. 12 and
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13, respectively.
As shown in Fig. 11(a), the Type-I interface presents a gradual decrease of liquidus
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temperature in the transition zone from the SS316L to IN625, while the actual temperature across the interface increases from SS316L to IN625 due to the laser heating source above, which inhibits
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protrusions of SS316L into the transition zone and thus produces a relatively smooth interface (Fig. 12a). As illustrated in Fig. 12(b), when the cooling begins, nucleation starts from the interface between solid SS316L and liquid transition zone (the bottom dashed line in Fig. 12) due to its
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higher undercooling than the top region. A constitutional undercooling region forms subsequently ahead of the solidification front as the presence of partitioned solutes promotes the cellular dendrites to grow in the transition zone (Fig. 12c). Epitaxial growth is seen from the interface between the transition zone and IN625 (the upper dashed line in Fig. 12) and grains evolve into the dendritic structures in the IN625 due to the compositional change (Fig. 12d).
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As shown in Fig. 11(b), because of the uneven mixing, a slight increase of the liquidus
temperature in the transition zone is observed, causing a typical mushy zone (Fig. 13a). A sudden rise of liquidus temperature is seen in the CSCZ located between the transition zone and SS316L (Fig. 11b). The rise in liquidus temperature increases the local undercooling, and thus generates a potential nucleation region. When the cooling begins, nucleation occurs at both the solid fronts of IN625 and CSCZ (Fig. 13b). As shown in Fig. 13(c), the grain growth at the solid front of IN625 follows a typical epitaxial growth. However, driven by the higher undercooling in CSCZ compared 12
to the adjacent transition zone and liquid SS316L, horizontal to CSCZ is the preferred direction to form the primary dendritic arm, while the resultant partitioned solutes further induce a constitutional undercooling and promote the growth of secondary dendritic arms. The grains in the CSCZ continue to grow towards the IN625 and form a cellular structure together with the dendrites growing from the IN625 side and finally develop the transition zone (Fig. 13d). This structure in the absence of epitaxial growth at the Type-II interface is detrimental for interfacial mechanical properties due to the interrupted grain growth. Besides, this solidification behavior increases the
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possibility of solidification cracking as a result of the intersection of grains growing from the two
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sides of the transition zone.
4.4. Cracks formation and its mechanisms
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Forescatter diodes (FSD) equipped on an EBSD detector were employed to detect the defects (i.e. porosity and cracks) as shown in Fig. 14. The images revealed no obvious defects in
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the printed base metals and the Type-I interface (Figs. 14a-c). However, cracks are evident at both the Type-II interface and SS316L tracks adjacent to the Type-II interface (Fig. 14d). Moreover,
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limited porosities (black dots) can be seen in the bulk materials and interfaces.
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Further examinations on probing the crack formation mechanism were conducted through SEM and EDS. Fig. 15 presents the magnified SEM image of the cracks located at the Type–II interface. The cracks are predominantly located at the boundary between the cellular structure grew from the CSCZ and the IN625 dendrites. The EDS mapping indicates that there is no obvious elemental enrichment in the crack confirming solidification cracking that happened at the Type–II
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interface. As mentioned in Section 4.3.3, during the solidification of the Type-II interface, the grain growth in the transition zone from the two sides (i.e. the CSCZ and IN625) is responsible for the solidification cracking due to lack of liquid metal to fill the gap between solidifying metals.
Fig. 15. SEM images and EDS mapping for the cracks at the Type-II interface.
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Fig. 16(a) shows the zigzag cracks in SS316L, where the dash lines indicate the track boundaries. Three segments of the zigzag crack are named as crack #1, crack #2 and crack #3. The corresponding Kernel average misorientation (KAM) mapping is extracted by EBSD to characterize the local strain distributions around the cracks, as indicated in Fig. 16(b). Three corresponding locations from each crack marked as A, B and C are chosen to perform the EDS mapping, and the corresponding results are plotted in Figs. 16, 17 and 18, respectively. The crack at location A is resultant from ductility dip cracking (DDC), where the cracks happened along the
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grain boundaries, as seen in Fig. 17(a). This can also be confirmed from Figs. 17(b-e) that no liquidation or elemental accumulation occurred at this location. DDC is induced during solidification by the strain accumulation at grain boundaries and a sharp drop in ductility during
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the temperature range from one-half of the alloy solidus to the solidus temperature [25]. Qian et al. [26] further explained the formation mechanism of DDC from the perspective of dislocation
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migration. As the thermal tension stress introduced by temperature gradient exceeds the yield stress, the dislocation slip starts within the grains followed by their pile up at the high angle grain
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boundaries, enhancing the strain and stress accumulations and eventually causes DDC.
Crack #2 has a wavy shape (Fig. 16a) where elemental enrichment is barely visible (Fig. 18), which indicates a solidification crack formation. In addition, bidirectionally grown columnar
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dendrites can be found adjacent to the crack as marked by the arrow in Fig. 18(a), which further confirms the cracking occurred during the solidification. The formation mechanism is similar to the cracks at the Type-II interface, where the insufficient supply of liquid metal to fill the space between solidifying metals from opposite sides and space opens further under shrinkage strain. As exhibited in Fig. 19, for the location C of crack #3, the EDS mapping reveals noticeable
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elemental accumulation of Nb and Mo in the crack, indicating the formation of liquidation cracking. From Fig. 16(a) and EDS mapping in Fig. 7(c), it is found that the crack is located at the unmixed zone with a higher concentration of Nb and Mo, resulting in increased susceptibility of microsegregation of the Nb and Mo which are deposited as long-chain Laves phase at the grain boundaries under a relatively low cooling rate [27,28]. With the deposition of subsequent layers, the Laves phase and its surrounding γ phase undergo eutectic remelting at the grain boundary under
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reheating, resulting in crack initiation, and the long-chain configuration promotes the crack propagation.
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Author statement Nannan Chen: Formal analysis, Investigation, Writing- Original draft preparation. Haris Ali Khan.: Data curation, Writing- Review & Editing. Zixuan Wan: Software, Data curation. Investigation. Writing- Original draft preparation. John Lippert: Data curation, Writing- Original draft preparation.: Hui Sun: Software, Data curation. Shun-Li Shang: Formal analysis, Writing- Original draft preparation. Zi-Kui Liu: Funding acquisition, Supervision. Jingjing Li: Supervision, Project administration, Funding acquisition, Writing - Review & Editing.
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5. Conclusions
The importance of printing sequence in a DED processed bimetal structure of SS316L and
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IN625 was interpreted by probing interfacial characteristics and crack generation, and their formation mechanisms were further unearthed by microstructural and crystallographic
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characterizations together with numerical simulation of molten pool and flow behavior. The main conclusions are as follows:
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1. The continuous growth of large columnar grains at the centerlines was observed in IN625 because of the epitaxial growth between the layers. In the SS316L, the fine equiaxed dendrites
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were formed at the top of each track which prevented the continual growth of columnar grains. 2. Compositional variation at the two interfaces was explained aided by numerical simulation. For the Type-I interface, the anticlockwise flow at the rare of the molten pool facilitated a fully mixed zone by driving the re-melted SS316L move forward into the front area with higher temperature and better fluidity. For the Type-II interface, the unmixed streaks enriched by the IN625 was
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responsible for the clockwise flow at the rear of the molten pool, which boosts the re-melted IN625 to move upwards into the molten pool and be captured by the solidification. 3. The formation mechanisms of the Type-I and Type-II interfaces were explained through the integration of experimental results and calculations of liquidus temperature. For the Type-I interface, a gradual-change compositional transition zone was observed because the lower density and lower viscosity of SS316L compared to IN625 promotes a good mixing. Epitaxial growth was formed at the Type-I interface due to the gradual decrease in the liquidus temperature from SS316L 15
to IN625 as a result of the compositional variation. For the Type-II interface, the higher density and higher viscosity of IN625 hindered the compositional mixing and thus generated a suddenchange zone. The resultant sudden rise of liquidus temperature finally created a unique interface without epitaxial growth. 4. Type-II interface and the SS316L tracks near the Type-II interface were found susceptible to crack formation. The formation of solidification cracks at the interface and track boundaries were owing to bidirectional solidification. Moreover, the unmixed zone from the re-melted IN625
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caused microsegregation of Nb and Mo, leading to liquidation cracking. The solidification and liquidation cracking further increased the susceptibility of ductility dip cracking which were
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generated under solidus temperature. Acknowledgments
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The authors acknowledge the US National Science Foundation Civil, Mechanical and Manufacturing Innovation Grants No. 1651024. H.S., S.L.S., and Z.K.L. acknowledge the
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financial support from the Office of Naval Research (ONR) under contract No. N00014-17-1-2567.
Appendix Temperature-dependent material properties applied in the numerical model in Section
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3.
This appendix reprensts a complilatin of important properties used in the CFD simulation [29-31]. The temperature-dependent thermal properties, including specific heat capacity, thermal conductivity, and density of SS316L and IN625 are presented in Fig. A-1. Table A-1 summerizes the temperature-dependent energy absorption rates of these two materials at room temperature,
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solidus and liquidus temperatures; and other material properties, such as liquidus and solidus temperatures, latent heat of fusion, viscosity, surface tension and latent heat of vaporization, are listed in Table A-2. Note that the viscosity of IN625 is referred from the information of IN718 [31].
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[30]
Special
Metals
datasheet
Inconel
alloy
625,
http://www.specialme
tals.com/assets/smc/documents/alloys/inconel/inconel-alloy-625.pdf. [31] R. A. Overfelt, C. A. Matlock, M. E. Wells, Viscosity of superalloy 718 by the oscillating
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Fig. 1. Illustrations of (a) DED sample structure, and (b) laser scanning pattern.
Fig. 2. The setup of numerical simulation: (a) mesh block of the model, (b) dimensions of the initial fluid cross-section, and (c) laser power density followed the Gaussian distribution.
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Fig. 3. Optical macrostructure of cross-sections on (a) XZ plane, (b) YZ plane; magnified optical
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images with inset SEM micrographs on the XZ cross-section: (c) SS316L, (d) IN625, (e) Type-I
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interface, and (f) Type-II interface.
Fig. 4. Simulated cross-sections of (a) Type-I interface (IN365 is on the top of SS316L), and (b) Type-II interface (SS316L in on the top of IN365).
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Fig. 5. IPFs of (a) SS316L and (b) IN625 on XZ plane.
Fig. 6. IPFs of (a) SS316L and (b) IN625 on YZ plane, {100} pole figures extracted from: (c) region A, (d) region B, (e) region C of SS316L, and (f) region D of IN625 (SD refers to the scanning direction).
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Fig. 7. SEM images and EDS mapping for Type-I interface (IN625 on the top of SS316L) on (a)
planes.
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XZ and (b) YZ planes, and Type-II interface (SS316L on the top of IN625) on (c) XZ and (d) YZ
24
Fig. 8. The simulated flow field of the (a) Type-I interface and (b) Type-II interface on the YZ
ur na
lP
re
-p
ro
of
plane.
Fig. 9. Magnified morphologies and elemental distributions: (a) SEM image and (b) EDS line scanning plot of Type-I interface, (c) SEM image and (d) EDS line scanning plot of Type-II
Jo
interface.
25
of ro -p re
lP
Fig. 10. IPFs for the Type-I interface of (a) XZ and (b) YZ planes, and the Type-II interface of
Jo
ur na
(c) XZ and (d) YZ planes.
Fig. 11. Calculated liquidus temperatures across (a) the Type-I interface and (b) the Type-II interface based on the elemental distributions in Fig. 9 using the CALPHAD approach. 26
of ro -p re
Jo
ur na
lP
Fig. 12. Schematic illustration of the formation mechanism of the Type-I interface.
27
re
-p
ro
of
Fig. 13. Schematic illustration of the formation mechanism of the Type-II interface.
Fig. 14. FSD images on XZ cross-section: (a) SS316L, (b) IN625, (c) Type-I interface, and (d)
Jo
ur na
lP
Type-II interface.
28
ro
of
Fig. 15. SEM images and EDS mapping for the cracks at the Type-II interface.
Fig. 16. (a) SEM image and (b) corresponding KAM image for the zigzag cracks in SS316L
Jo
ur na
lP
re
-p
tracks near the Type-II interface (i.e. the upper crack in Fig. 14d).
Fig. 17. (a) SEM image of Crack #1 (Location A in Fig. 16), and (b)-(e) corresponding EDS mapping for elements Fe, Ni, Cr, and Mo, respectively.
29
of ro -p
re
Fig. 18. (a) SEM image of crack #2 (Location B in Fig. 16), and (b)-(e) corresponding EDS
Jo
ur na
lP
mapping for elements Fe, Ni, Cr, and Mo, respectively.
30
of ro -p
re
Fig. 19. (a) SEM image of crack #3 (Location C in Fig. 16), and (b)-(f) corresponding EDS
Jo
ur na
lP
mapping for elements Fe, Ni, Cr, Nb, and Mo, respectively.
31
of ro -p re lP ur na
Jo
Fig. A-1. Temperature-dependent thermal properties: specific heat capacity of (a) SS316L and (b) IN625, thermal conductivity of (c) SS316L and (d) IN625, and density of (e) SS316L and (f) IN625 [29,30].
32
Table 1. Chemical composition of SS316L and IN625 powders. Powder
Fe
Ni 1014
SS316L Rem. 5
Rem.
C
Si
0.03
1
2
0.1
0.5
0.5
0.045
0.03
0.015 0.015
Mo
-
2-3
3.154.15
8-10
Cr 1618 2023
Ti
Al
-
-
0.4
0.4
re
-p
ro
of
IN625
Elements (Wt.%) Mn P S Nb
Table 2. DED process parameters. Powder Flow Rate (g/min) 20.5 20.5
Powder Feeder Vibration (%) 60 60
lP
Travel Speed (IPM)
SS316L IN625
2 2
25 25
Step Over Per Pass (mm) 2.5 2.5
Focus Head Standoff (mm) 10 10
Spot Size (mm) 4 4
Jo
ur na
Powder
Laser power (kW)
33
Table A-1. Temperature-dependent energy absorption rate (A) of SS316L and IN625 for a 1060 nm laser beam [29,30]. SS316L A (%)
Temperature (K)
IN625 A (%)
293
38%
293
30%
1673 (Solidus)
40%
1563 (Solidus)
30%
1698 (Liquidus)
45%
1623 (Liquidus)
35%
of
Temperature (K)
-p
IN625 1623 1563 3.1e5 7.40 1.88 6.35e6
Jo
ur na
lP
re
Liquidus (K) Solidus (K) Latent heat of fusion (J/kg) Viscosity (mPa·s) Surface tension (N/m) Latent heat of vaporization (J/kg)
SS316L 1698 1673 2.6e5 6.42 1.80 6.25e6
ro
Table A-2. Other material properties of SS316L and IN625 [29-31].
34