1. Introduction
Due to rising material costs and worsening climatic conditions, there is an increasing worldwide demand for alternative technologies that would significantly reduce the amount of materials consumed, which would have a positive impact on both product prices and the environment. One suitable technology for processing such materials is Wire Arc Additive Manufacturing (WAAM) [
1]. The WAAM process is based on depositing solid wire layer by layer until the component is finished, while the motion mechanism is typically provided by an industrial robot controlled by a computer to ensure movement along the desired path [
2,
3]. The resulting component is composed of multiple layers with a constant or adaptive thickness, which are laid down along the outline trajectory of the future product [
4,
5].
Compared to other additive technologies, WAAM stands out with numerous advantages. WAAM can be considered as an advantageous economic process due to its low acquisition cost and the possibility of utilizing a wide range of feedstock, making it suitable for the production of large-scale components [
6]. With this technology, it is possible to achieve a deposition rate of 10 kg/h for steel, which is up to 16 times higher compared to metal-powder-based additive manufacturing techniques [
7]. Additionally, by using metal powder, the process becomes more susceptible to defect formations, such as pores, that can compromise structural integrity, especially under dynamic loading conditions [
8]. Moreover, WAAM components exhibit superior mechanical properties compared to components produced through traditional casting methods [
9,
10,
11]. Materials with a high cost and limited availability, such as various types of nickel alloy Inconel, find application in the WAAM technology across different industries. The Inconel is known for its high mechanical properties, strong material adhesion to cutting edges, and the presence of abrasive carbide particles. These factors make these alloys difficult to machine, and as a result, their application is found in WAAM technology to prevent significant losses associated with their processing [
12,
13,
14].
Despite its many advantages, WAAM has several drawbacks, such as a higher heat input, which can lead to decreased geometric accuracy, higher surface waviness values, and lower mechanical properties of deposited components [
15,
16,
17,
18]. High heat input in WAAM technology can be reduced by the CMT process, which also results in less spattering, reduced emissions, and high process stability [
19,
20,
21]. Hackenhaar et al. examined the influence of deposition mode selection on the resulting mechanical properties of a component made from AISI H13 tool steel. In their study, they applied CMT, LSC, and PS modes. The authors did not observe significant differences in terms of mechanical properties among the components. However, the component deposited with the CMT mode achieved the greatest height [
22]. Another way to reduce the heat input and improve mechanical properties is by maintaining the interpass temperature or using active cooling.
Wu et al. investigated the influence of a cooling strategy on the resulting mechanical properties of a component made from Ti6Al4V alloy. In their experiment, they applied interpass temperatures of 100, 200, and 300 °C along with forced interpass cooling using compressed CO
2. The authors found that decreasing the interpass temperature led to a slight improvement in the mechanical properties of deposited components. However, the sample cooled with the compressed CO
2 showed the best ultimate tensile strength values [
23]. Similar results were also observed by other authors who studied the influence of an interpass temperature on the resulting mechanical properties [
24,
25,
26]. On the other hand, Shen et al. discovered that applying a higher interpass temperature during the deposition of a component made of Fe
3Al alloy led to reduced weld pool asymmetry, attributed to the different thermal conductivities of the substrate and filler metal. This, in turn, resulted in a higher yield strength in the middle section of the component [
27]. Dash et al. observed the impact of active cooling on the resulting microhardness of the fabricated components from mild steel. The authors found that the application of active cooling led to an increase in microhardness compared to samples deposited without active cooling [
28]. Tonelli et al. investigated the effect of interlayer active cooling on the final mechanical properties of WAAM-produced components of austenitic AISI 304 L steel. The authors concluded that there was no difference in the mechanical properties between components fabricated with and without active cooling and that mechanical properties were more affected by microstructural anisotropy [
29].
The aim of the research is to assess the influence of cooling strategy choice on the resulting mechanical properties of Inconel 718 alloy components deposited using the WAAM-CMT process.
2. Materials and Methods
The overlay weld walls were deposited on an S235JR low carbon steel substrate with dimensions of 200 mm × 100 mm × 10 mm. In WAAM, the base material is typically a more cost-effective option since it is eventually removed after the deposition. WAAM is efficient for manufacturing larger components using expensive alloys due to its favorable buy-to-fly (BTF) ratio. It represents the proportion of input material used in production to the final component, and this ratio is notably smaller compared to conventional machining [
30,
31]. Therefore, the corrosion-resistant nickel alloy Inconel 718 in the form of NiCro 718 MIG wire with a diameter of 1.2 mm was used as a filler material. This precipitation-hardened alloy was developed in the 1960s and is still considered the first choice for most components of aircraft engines operating at temperatures below 650 °C due to the dissolution of the γ″ phase and the formation of a stable δ phase at higher temperatures [
32].
Table 1 shows the chemical composition of the welding wire based on the 3.1 inspection certificate delivered by the producer according to EN10204 [
33].
The fully digitized inventor welding source TransPuls Synergic 3200 CMT from Fronius company (Wels, Austria), which is specially designed for the CMT mode, was used for the production of multilayer components. The movement of the welding torch was ensured by the portal device Multiweld from the company First Welding Company (Bratislava, Slovakia), Inc. Argon 4.6 from Messer Tatragas company (Bratislava, Slovakia) was used as shielding gas.
Two components consisting of 53 overlay beads were made. The first component was made without applying the interpass temperature. The overlaying began according to
Figure 1 by moving the welding torch from reference point 1 to point 2, which was 150 mm away from the reference point in the direction of the Y axis. At point 2, the welding process was finished, and the welding torch was subsequently moved by 3 mm in the direction of the Z axis. Then, the welding torch moved 150 mm from point 2 to point 1 along the direction of the Y axis. In point 1, the torch was moved again by 3 mm, and the whole cycle was repeated until the component was finished.
In the case of applying the interpass temperature, the next overlay bead was made when the temperature of the previous layer dropped below 100 °C. The interpass temperature was monitored by a K-type thermocouple and HH-507 R thermometer from Omega Engineering Inc. (Norwalk, CA, USA). The deposition parameters are listed in
Table 2.
The edges of the components (
Figure 2) with a length of 15 mm were cut by an electric discharge machine, Sharmilles Robofil 310 (Charmilles technologies, Geneva, Switzerland), from both sides. Cross-sections acquired this way were used for macro- and micro-analysis as well as microhardness evaluation. The remaining part was milled from both sides to a thickness of 7.5 mm on the DMU 85 MonoBlock device (DMG Mori, Leonberg, Germany). Tensile and notch test specimens were subsequently cut by water abrasive cutting using Aquacut 6001.20 device from Microstep (Bratislava, Slovakia) in order to prevent thermal influence. The location of samples on the deposit is documented in
Figure 3. The tensile test specimens were oriented vertically at the same height. The Charpy impact test samples were oriented vertically, with the notch located at 1/3, ½, and 2/3 of the height of the component.
The final dimensions of the samples for the tensile test and Charpy pendulum impact test were finished to the final size in accordance with ISO 6892-1 and ISO 148-1 (
Figure 4 and
Figure 5) by milling on the DMU 85 MonoBlock device (DMG Mori, Leonberg, Germany) [
34,
35].
The obtained data were statistically evaluated using GraphPad Prism 8. All data showed normal distribution. The data from the static tensile test, Charpy impact test, and microhardness test of components made with and without interpass temperature were compared with a parametric T-test. Used levels of significance were * p < 0.05, *** p < 0.001, **** p < 0.0001. All error bars shown in the figures represent standard deviation (SD).
Samples for macroscopic and microscopic analysis were prepared using a standard procedure on a PHOENIX 4000 device from Buehler company (Lake Bluff, IL, USA), which consisted of wet grinding on 240, 600, and 1200 grit silicon carbide papers. Subsequently, the samples were polished on 9, 3, and 1 µm diamond suspensions.
Electrolytic etching in a 10% solution of oxalic acid (C2H2O4) at 10 V for 10 s was used to highlight the microstructure of Inconel 718. The microstructure was observed on a NEOPHOT 32 metallographic light microscope from Carl Zeiss (Jena, Germany).
A more detailed examination of the microstructure was performed using the transmission electron microscope JEOL 200 CX (TEM) from Jeol producer (Nieuw-Vennep, The Netherlands) operating at 200 KV. Thin foils were used to identify secondary phases. The preparation of thin metal foils consisted of cutting a 1 mm thick sample with a Isomet 5000 linear precision saw from Buehler company (Lake Bluff, IL, USA). This was followed by mechanical grinding on silicon carbide paper (1200 grit) to a thickness of approximately 150 μm. Final thinning was carried out using an electro-polisher Struers-Tenupol5, which was done in a solution of 10% perchloric acid (HClO4) in 90% methanol (CH3OH) at −30 °C temperature and with a voltage of 25 V.
The microhardness measurement, static tensile test, and Charpy impact test were used to evaluate the mechanical properties. The microhardness HV0.1 was measured on a Indentamet 1105 from Buehler company (Lake Bluff, IL, USA) upwards from the substrate in the wall axis with a step of 5 mm. The static tensile test was performed on a LabTech Labtest 5.250 SP1 device at a temperature of 22 °C with a loading speed of 5 mm/s. The ultimate tensile strength (UTS) and offset yield strength of 0.2% (YS0.2) were evaluated. The Charpy impact test was performed on a CHK-300 instrumented pendulum hammer at a temperature of 22 °C with a hammer energy of 300 J.
4. Discussion
The microstructure of the components was formed by a columnar dendritic structure, while the orientation of the dendrites was in the direction of the height of the component. Due to the application of individual layers, it was possible to observe annealed areas characterized by softened dendrites. At the edges of the manufactured components, a change in the orientation of the dendrites was observed, which were no longer oriented in the direction of the height of the component but in the direction of the greatest heat dissipation—to the edge parts of the components. The same character of the microstructure was also observed in the studies of other authors [
38,
39,
40].
Components made without applying the interpass temperature were characterized by the highest UTS and YS
0.2. During overlaying, these components reached higher temperatures, which led to longer cooling in the range of temperatures necessary for the precipitation of strengthening phases. This process increases the content of strengthening phases, which subsequently contributes to the improvement of the strength characteristics of the manufactured components. Conversely, applying the interpass temperature resulted in a lower UTS and YS
0.2. This finding does not correspond to the results of the authors [
23,
25], who, in their studies, investigated the effect of heat input and interpass temperature on the final mechanical properties of the produced deposits. The authors [
25] found that the application of an interpass temperature of 100 °C and the reduction in the heat input led to an increase in the mechanical properties of the components produced by WAAM technology. The reason for the different results may be the different types of Inconel 625 filler material investigated.
TEM analysis confirmed the presence of strengthening phases γ′ and γ″ that precipitated in the matrix. The γ′ phase was characterized by a quasi-globular morphology and was observed only in the component with the applied interpass temperature. The γ″ phase was observed in the component made without applying the interpass temperature. From the point of view of morphology, this phase was characterized by a disk-like shape, which could contribute to the achievement of higher mechanical properties of the component. On the other hand, the increased density of dislocations in the dendrites, where precipitation of strengthening phases did not occur, could have contributed to the increase in mechanical properties.
The interdendritic space of both components was mainly formed by the δ phase and carbides. M6C, M23C6, and MC carbides (TiC, NbC) were identified in the component made with the application of interpass temperature. The exact determination of MC carbides is problematic due to their almost identical interplanar spaces dhkl. EDX analysis would be necessary to accurately determine the MC carbides. The Laves phase was observed sporadically at dendrite boundaries.
Bhujangrao et al. [
41] studied the microstructure of a component made of Inconel 718 alloy by WAAM using Plasma Arc Welding (PAW) technology. They concluded that the interdendritic space is formed by MC carbides and the Laves phase, which are formed in the material under non-equilibrium cooling conditions. The presence of the delta phase was also identified by X-ray diffraction. The authors [
42] also used the same technology (PAW) and material, while X-ray diffraction identified the presence of Laves phase and MC carbides. The discrepancy with our results could be caused by the use of the MIG CMT welding mode, which is characterized by a lower heat input.
The presence of secondary phases in the microstructure depends on the height (1, 3, 6, 9, and 12 mm) of the deposit investigated by Jang et al. [
43]. In the experiment, they used the CMT welding method and IN718 filler wire. The Laves phase was confirmed in the interdendritic space at all heights. At the height of 3, 6, 9, and 12 mm, an δ phase with needle-like morphology was confirmed in the interdendritic space together with the Laves phase. The strengthening phases γ′ and γ″ were observed only at distances of 3 and 6 mm above the substrate. The reason for the absence of strengthening phases just above the substrate was the large heat dissipation into the substrate. On the other hand, the strengthening phase was not observed in the upper parts of the deposit. Also, in our case, the strengthening phases were observed to a greater extent in the lower parts of both manufactured components.
Author Contributions
Conceptualization M.M. and F.S.; methodology, M.M.; validation, M.M., F.S. and J.B.; formal analysis, K.B., M.P. and J.U.; investigation, M.M. and M.P.; data analysis, F.S., K.B., M.D. and M.P.; typing—initial draft preparation, M.M. and F.S.; typing—review and editing, F.S.; visualization, J.B. and J.U.; supervision, M.M.; project administration, F.S. All authors have read and agreed to the published version of the manuscript.
Funding
This work was supported by the Slovak Research and Development Agency under the contract No. APVV-20-0259.
Data Availability Statement
The data presented in this study are available upon request from the appropriate author.
Conflicts of Interest
The authors declare no conflicts of interest.
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Figure 1.
The experiment WAAM setup.
Figure 2.
Components deposited: (a) without interpass temperature, (b) by applying interpass temperature of 100 °C.
Figure 3.
Samples’ locations on the deposited components after abrasive water jet cutting.
Figure 4.
Geometric dimensions of specimens for tensile test.
Figure 5.
Geometric dimensions of specimen for Charpy impact test.
Figure 6.
Microstructural analysis of a component deposited without applying interpass temperature.
Figure 7.
Microstructural analysis of a component deposited by applying interpass temperature.
Figure 8.
Microstructure of transition between annealed and non-annealed areas: (a) w/o interpass, (b) interpass.
Figure 9.
Boundary regions of the overlay component produced w/o interpass temperature: (a) left side, (b) right side.
Figure 10.
Microhardness course in the direction of the height of the component along the axis passing through the center of cross-section.
Figure 11.
T-test of the microhardness of deposited components regarding cooling strategy ns—non-significant.
Figure 12.
T-Test of the UTS of the component with respect to the cooling strategy; **** p ≤ 0.0001.
Figure 13.
T-Test of the YS0.2 of the component with respect to the cooling strategy; * p ≤ 0.05.
Figure 14.
T-test of the impact energy of the component with respect to the cooling strategy; ns—non-significant.
Figure 15.
Interdendritic regions (a) interpass temperature 100 °C (b) w/o interpass temperature.
Figure 16.
Detail of δ-phase in interdendritic region (interpass 100 °C) (a) bright field (b) dark field—reflection (201) was used—(c) diffraction pattern.
Figure 17.
Detail of M6C carbide in interdendritic region (interpass 100 °C) (a) bright field (b) dark field—reflection (044) was used—(c) diffraction pattern.
Figure 18.
Detail of Laves phase at dendritic boundaries (interpass 100 °C) (a) bright field (b) dark field—reflection (110) was used—(c) diffraction pattern.
Figure 19.
Detail of M23C6 carbide on dendrite boundaries (interpass 100 °C) (a) bright field (b) dark field—reflection (242) was used—(c) diffraction pattern.
Figure 20.
Detail of MC carbide inside the dendrite (interpass 100 °C) (a) bright field (b) diffraction pattern.
Figure 21.
Detail of strengthening phases γ′ + matrix γ phase (interpass 100 °C) (a) bright field (b) dark field—reflection (100) was used—(c) diffraction pattern.
Figure 22.
Strengthening γ″ phase + matrix γ phase (w/o interpass) (a) bright field (b) dark field—reflection () was used—(c) diffraction pattern.
Figure 23.
Dislocation in the matrix (sample w/o interpass).
Table 1.
Chemical composition of NiCro 718 wire in %.
Element | Composition [wt %] |
---|
C | 0.07 |
Mn | 0.1 |
Si | 0.15 |
S | 0.001 |
P | 0.008 |
Cr | 17.5 |
Ni | 52 |
Mo | 3 |
Al | 0.4 |
Ti | 0.9 |
Co | 0.05 |
Cu | 0.05 |
Nb + Ta | 5 |
B | 0.005 |
Fe | Balance |
Table 2.
Deposition parameters.
Current [A] | Voltage [V] | Deposition Speed [mm/s] | Wire Feed Rate [m/min] | Gas Flow Rate [L/min] | Heat Input [J/mm] |
---|
205 | 15 | 5 | 8 | 15 | 492 |
Table 3.
Tensile properties of samples with and their average values and standard deviations.
Cooling Strategy | Sample | UTS [MPa] | AVG UTS [MPa] | SD [MPa] | YS0.2 [MPa] | AVG YS0.2 [MPa] | SD [MPa] |
---|
w/o interpass | 1 | 832 | 829.7 | | 518 | 510.3 | |
2 | 830 | 2.5 | 502 | 8 |
3 | 827 | | 511 | |
interpass temperature | 1 | 772 | 764.8 | | 461 | 428.3 | |
2 | 760 | 6.3 | 445 | 44 |
3 | 763 | | 378 | |
Table 4.
Impact energy, average impact energy and standard deviation of samples.
Cooling Strategy | Sample | KV [J] | AVG KV [J] | SD [J] |
---|
w/o interpass | 1 | 28.1 | 26.8 | |
2 | 26.0 | 0.9 |
3 | 26.4 | |
interpass temperature | 1 | 21.5 | 26.4 | |
2 | 25.6 | 4.4 |
3 | 32.2 | |
Table 5.
Confirmed phases by TEM.
Phase | Interpass | w/o Interpass |
---|
γ | ✓ | ✓ |
γ′ | ✓ | – |
γ″ | ✓ | ✓ |
δ | ✓ | ✓ |
M6C | ✓ | – |
M23C6 | ✓ | ✓ |
MC | ✓ | – |
σ | – | – |
Laves | ✓ | – |
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