Surface macrodefects formed during vertical wall growth do not affect the internal macrostructure of the sample, except for the disturbance of layer thicknesses. Due to the difference in melting temperature of the materials, a disturbance of the uniform thickness of the copper layer in the longitudinal section can be observed in Figure 4 . The average value of the thickness of the copper layers is 4.1 ± 0.5 mm. Typical characteristics for additively manufactured products are observed, where overlapping molten pools with a semi-ellipse shape are formed. A middle-ground approach uses geometry of the melt pool and simulates a series of overlapping semi-circles or semi-ovals. Lack-of-fusion porosity depends on both the melt pool size as well as the scanning pattern. For a given melt pool size, the minimum depth of overlap is defined [ 28 ]. This quantity is affected by the hatch distance of the melt pool, which then dictates the maximum layer thickness. The formation of a molten pool with a semi-ellipse shape is characteristic of various methods of additive manufacturing, both for ferrous [ 29 ] and for non-ferrous metals [ 30 ], due to the peculiarities of the process. In terms of their size, melt pools are statistically 1.1–1.3 mm in width and 0.5–0.7 mm in depth. The molten pool boundaries in the sample are visible as “fish-scales” [ 31 ], filled with a cellular structure [ 32 ].
A gradual decrease from 0.42 kJ/mm to 0.38 kJ/mm occurs during the printing process as one moves away from the cooled substrate. Due to the heat accumulation in the sample, the creation of the molten pool and the melting of the filament occur with a reduced electron beam power. After printing the first five steel layers, the heat input values of zones II–V occurred in the range of 0.33–0.30 kJ/mm. This printing mode was chosen to ensure the contact of layers of dissimilar materials. The printing parameters used were aimed at adjusting the values of thermal investment, influencing the temperature gradient and taking into account the difference of chemical and thermophysical properties of dissimilar materials.
For convenience, the bimetallic layered macrocomposite was divided into five areas: area I—printing the first five layers of stainless steel onto the substrate; II and IV—printing five layers of copper onto the already deposited steel; III and V—printing stainless steel layers onto the already deposited copper layers.
In this work, the value of the heat input was calculated using the formula [ 27 ]:where U and I are voltage and current of the electron beam, and v is the printing speed. Using the data on printing parameters ( Table 2 ) and Formula (1), the dependence of layer-by-layer heat deposition is obtained ( Figure 3 ).
In the course of growing the vertical wall, macrodefects are formed. This occurs due to insufficient heating or overheating of the fed filament, which entails a difference in the size of the formed layer. In the case of overheating, material spreads, droplets are formed and the shape of the vertical wall is distorted across the width. In the case of insufficient heating, unmelted parts of the wire remain, which leads to non-melting of the first layers with the substrate and layers with each other, and pores are formed in the areas not filled with unmelted filament. In the present work, these factors were taken into account, and optimal parameters were selected for vertical wall printing from heterogeneous materials in order to avoid the influence of insufficient heating or overheating. This is affected by a significant difference, of 21–24 times, in the thermal conductivity values of the used materials. The alternation of copper and steel wire material during printing creates conditions of high heat dissipation of previously applied copper layers, in which the steel wire does not have time to melt. At the same time, during the application of copper, it spills out due to the insufficiently high thermal conductivity of steel. For printing with steel wire, in this case, the selection of parameters was based on already known data [ 22 ]. When using heterogeneous materials with different thermal physical properties as filaments for printing, it is necessary to vary the low and high values of heat inputs depending on the material.
During the melting of fed wire by the electron beam with the formation of a molten pool, there is a local increase in the substrate temperature near the molten pool. As a result, a temperature gradient occurs, which leads to the presence of internal stresses. After a number of thermal cycles, the substrate deforms under the influence of temperature stresses.
When C11000 copper layers are deposited onto already consolidated 321 stainless steel layers, some grains grow, due to the adjacent ones, by migration of high-angular boundaries, i.e., collective recrystallization takes place. In this case, grain boundaries migrate, and an equiaxial structure with minimal surface energy and grains of equal size and shape is formed. One of the factors contributing to the formation of an equiaxed fine-grained copper structure can be the alloying of copper by iron, as well as elements present in the composition of austenitic steel. Particles of more refractory metals will contribute to the formation of a large number of crystallization nuclei during solidification of the product. At the same time, impurity atoms and second phase inclusions (in the transition zone) are the factors that prevent the migration of grain boundaries. As the driving force of collective recrystallization decreases as it develops, the grain growth stops when it reaches a certain value. However, first of all, the difference between directional solidification (constrained growth) results in the formation of oriented cells or oriented dendrites and free solidification (unconstrained growth), which results in equiaxed dendrite (equiaxed grain) formation [ 34 ]. The columnar structure is formed due to the so-called constrained solidification whereas the equiaxed grains appear when the unconstrained solidification takes place. Since only five layers of copper have been deposited, the influence of these factors is not reduced and there is no directed grain growth of the elongated shape. Directed grain growth is facilitated by the slow heat dissipation and high temperature gradient generated by the additive manufacturing process [ 35 ].
In a previous study, when C11000 copper was deposited on an AISI 321 steel substrate, a region with a fine-grained structure, of not more than 1 mm in length, was observed by Osipovich et al. [ 33 ]. Since the growth process of the sample in zones II and IV was continuous, a homogeneous growth of coarse grains over the entire wall height would be expected. However, microstructural analysis reveals only equiaxed grains with an average size of 28.5 ± 0.2 μm.
Copper interlayers of the sample are of particular scientific interest due to the effect of grain size, which is described by the Hall–Petch ratio. In this case ( Figure 5 ), a fine-grained structure is observed in the copper interlayers (zone II and IV).
A common feature of additive steel macroimages is the fine-grained microstructure, typical for AM [ 32 ]. Elongated and oriented grains depending on manufacturing process parameters have been discussed previously [ 37 ]. It has been noted that the evolution of crystallization cells is related to the segregation of alloying elements at the interfaces, leading to microsegregation. It should be noted that solute microsegregation transforms into solute redistribution after back-diffusion into the solid [ 38 ]. Equiaxial cells can be either small or large with sizes of 0.2–0.6 μm and 1–2 μm, respectively. In this case, ferrite is located between the primary arms of this cell structure (dendrites) [ 39 ]. Heating ferrite to 912 °C leads to the formation of tiny austenite grains at the boundaries of the ferrite grains. Further heating leads to the growth of these new austenitic grains with complete replacement of the ferrite grains by austenitic grains—iron transformations occur. Since the electron beam process of additive manufacturing is a layer-by-layer deposition of material during sample manufacturing, there is a constant re-melting of the already deposited layers, i.e., thermal cycling. It seems that solidification which appears after re-melting is, in reality, rapid solidification. In the case of the rapid solidification, the phenomenon of microsegregation (as well as redistribution) is slightly modified by an increase in the partition ratio, k, as explained previously [ 40 ]. However, the proper behavior of the partition ratio during rapid solidification has been defined [ 41 ]. Thermal cycling leads to an evolution of the structure by height in individual zones—I, III and V—from small to large dendrites. Optical microscopy showed that the structures of stainless steel in zones I and V are similar ( Figure 6 a,c). This means an increase in dendrite grain size. The explanation for this is the thermal energy provided by the powerful electron beam, which to some extent can reduce the rate of the solidification process, thereby giving the dendrite more time to become coarse-grained during the solidification process. The dendritic structure mainly developed with a deviation from the building direction (z) due to a change in the temperature gradient during solidification.
The microstructures for EBAM-manufactured steel 321 in different parts of zone I, III and V are shown in Figure 7
TEM images clearly show that the steel microstructure contains two phases—Austenite and acicular δ-ferrite. The morphology of δ-ferrite dendrites does not depend on the position in the steel blank ( Figure 7 ). According to the EDS analysis, δ-ferrite lamellae are enriched in chromium and depleted in nickel. The typical composition of the ferrite phase is Fe-(27.2–29.3)Cr-(3.2–3.8)Ni (wt. %). The elemental composition of the austenitic phase is proportional to the initial wire composition, but the chromium and nickel content is somewhat excessive—Fe-(18.3–19.1)Cr-(9.4–10.3)Ni (wt. %). δ-ferrite in the form of lamellae 300–500 nm wide is uniformly distributed in the steel structure. The boundaries between phases are partially curved ( Figure 7 ) and partially faceted by the Kurdjumov–Sachs orientation relationship between austenite and ferrite. However, δ-ferrite is a high-temperature phase and cannot be observed as it undergoes phase transformations at room temperature. However, based on the literature analysis [ 42 ], the above features indicate that δ-ferrite formed at high temperature was fixed by quenching, and it became possible to see it at room temperature.
Thus, the morphology of zones I, III and V is characterized by a grain–dendrite structure, with the growth of elongated dendrites, with ferrite located between the primary dendrite arm spacing, associated with thermal cycling, and the direction of dendrite growth inherited by the direction of the fabrication process.
R δ = ( I δ ∑ I A ) × 100 %
(2)
Iδ
is the integral intensity of the peaks corresponding to the δ-Fe phase; ∑IA
is the sum of all analyzed integral intensities.To better understand the formation of δ-Fe and its different volume fraction in zones I, III and V of the sample, quantitative studies were carried out using several methods: X-ray, metallographic and local-type ferritometer measurements ( Figure 8 and Figure 9 ). The relative intensities of the γ-Fe and δ-Fe phases were calculated using Formula (2) [ 43 ]:whereis the integral intensity of the peaks corresponding to the δ-Fe phase; ∑is the sum of all analyzed integral intensities.
Moreover, the volume fraction of δ-Fe in different parts of the sample is different: for zones I and V of steel sections, the volume fraction of δ-Fe varies within the range fI,V(δ-Fe) = (16.3–17.2); for zone III, a decrease in fII(δ-Fe) of 1.7–1.9 times as compared to the volume fraction for zones I and V of steel sections is typical.
It is seen that all experimental methods of determining the volume fraction of the δ-Fe phase are consistent with each other. There is a general tendency for the volume fraction to decrease in the middle of the sample and when approaching the copper interlayers, zones II and IV. This is due to the fact that copper, like nickel, is a stabilizer of γ-Fe, in this case, the ferrite that occurs during the last stages of solidification [ 44 ]. The liquid projection begins in the peritectic region in the Fe–Ni system and proceeds to the eutectic reaction in the Cr–Ni system in the Fe–Cr–Ni ternary system. The eutectic ferrite is then located along the boundaries of the austenite grains, in contrast to Widmanstatten pattern morphologies, in which the ferrite is contained mainly in the core grains. The microstructure, as shown in Figure 6 , in this alloy, can be characterized by regions of both eutectic ferrite and skeletal ferrite. The characteristics of Widmanstatten pattern morphologies of alloy 321 were also revealed. Values for the alloying elements to calculate the Ni–Cr equivalents ratio were obtained by SEM. Figure 10 shows the variation in the basic and alloying elements of the material.
eq/Nieq ratios [Creq = Cr + 1.37 × Mo + 1.5 × Si +2 × Nb + 3 × Ti
(3)
Nieq = Ni + 0.31 × Mn + 22 × C + 14.2 × N + Cu
(4)
The element composition in the 2 to 4 mm and 10 to 14 mm areas corresponds to the chemistry of C11000 copper. In the remaining height areas, the chemical composition corresponds to the composition of stainless steel 321. In the graph of the dependence of elemental composition on the distance to the zones of material supply change from stainless steel to copper and vice versa, there is an area in which the composition of austenitic steel is different from that shown in Table 1 and there is a fairly high concentration of copper. Analysis of the data allows us to conclude that in the transition zone, the austenitic steel is alloyed with copper and there is the formation of a two-phase region—copper and austenitic stainless steel alloyed with copper. This does not contradict the available data on the Fe–Cu phase diagram [ 45 ], according to which intermetallic compounds based on iron and copper are not formed, but a limited amount of copper (≈5.8% at 1083 °C) can dissolve in γ-Fe. As the number of spherical inclusions of copper decreases with distance from the interface, the concentration of copper also becomes lower. The presence of a small concentration of iron in the copper part of the bimetal sample near the interface may be due to the formation of a solid solution of iron in copper, although the solubility of iron in copper is rather low (≈2.8% at 1083 °C). Cr/Niratios [ 46 ] were used for the theoretical calculation, using formulas for equivalents ((3) and (4)):
The graph of the equivalent dependence is shown in Figure 11
eq/Nieq value is less than 1.5, then the solidification mainly starts from the γ-phase. It is seen that the value of the equivalent Creq and Nieq ratio in zone III is essentially lower, which agrees with the results of the measurement of the δ-Fe volume fraction presented above. In other words, the studied zone III is provided by the process conditions under which δ-Fe at the crystallization stage is formed to a lesser extent than in zones I and V. Just as with further cooling, the least amount of residual δ-Fe will remain in this zone.The results of studies [ 47 ] showed that if the content of Cr and Ni in stainless steel is the same and the Cr/Nivalue is less than 1.5, then the solidification mainly starts from the γ-phase. It is seen that the value of the equivalent Crand Niratio in zone III is essentially lower, which agrees with the results of the measurement of the δ-Fe volume fraction presented above. In other words, the studied zone III is provided by the process conditions under which δ-Fe at the crystallization stage is formed to a lesser extent than in zones I and V. Just as with further cooling, the least amount of residual δ-Fe will remain in this zone.
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