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High Temperature Materials and Processes

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Volume 38, Issue 2019

Issues

Microstructural Evolution and Phase Transformation on the X-Y Surface of Inconel 718 Ni-Based Alloys Fabricated by Selective Laser Melting under Different Heat Treatment

Peng Liu
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  • School of Materials Science and Engineering, Shandong Jianzhu University, Jinan 250101, P. R.China
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/ Siyu Sun / Meiqing Cao
  • School of Materials Science and Engineering, Shandong University of Science and Technology, Qingdao 266590, P. R. China
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Published Online: 2018-09-19 | DOI: https://doi.org/10.1515/htmp-2017-0182

Abstract

Microstructure, hardness, precipitates and phase transformation on the X-Y surface of Inconel 718 Ni-based alloys fabricated by selective laser melting (SLM) were studied before and after a suitable heat treatment. The test results show that the obvious weld beads structure was observed on the surface of the as-built alloys, and the microstructure shows seriously inhomogeneous with distributing columnar crystals and fine dendrites. When the 720°C, 3 h/furnace cooling+620°C and 3 h/air cooling was used, the columnar and small dendritic crystals begun to transform into bulk crystals, and the weld beads structure disappeared instantly. The X-Y surface hardness of alloys is about HV490-540, which is higher than the one of other heat treatment processes. With the increase of solid solution temperature, the microstructure on the surface of alloys shows an obvious refining characterization. The hardness value on the surface was also reduced gradually. The typical columnar crystals, dendrites, intermetallic compounds and precipitated hardening phase (Cr-Fe-Mo-Ni, FeNi and Ni8Nb) were also reduced or could not be found. Under the solid solution process of 950°C, the fine homogeneous δ phase and γ′ phase in the grain or near the grain boundary had an important effect on the suitable surface hardness value (HV476-500) of alloys.

Keywords: Ni-based alloys; selective laser melting; additive manufacturing; heat treatment; microstructure

Introduction

Inconel 718, a Ni-based alloy, has the ability of retaining mechanical stability at elevated temperature up to 650°C, excellent resistant to oxidation, high temperature corrosion, creep resistance and good fatigue life [1, 2]. Hence, it was widely used in aeronautics and in energy industries for various applications, such as aircraft turbines, jet engines and steam turbine power plants [3, 4]. However, the present cast components have coarse grain size, dendritic segregation and some solidification defects, which lead to poor mechanical properties of Inconel 718 alloy. In addition, it was difficult to control the performance of this alloy and produce complex geometries for the special wrought form [5]. Selective laser melting (SLM) with the characters that can produce more homogeneous and finer microstructures [6] is a novel processing method originated from the laser welding and processing [7], which consequently could be a prospective candidate to replace conventional production processes. This technology will offer several advantages compared to conventional technologies, such as reduction of production steps, high flexibility, low material consumption and, the most importantly, the possibility to manufacture parts with high geometrical complexity and dimensional accuracy.

However, some problems, such as residual stresses, porosity, directional grain growth, segregation, creation of non-equilibrium phases and other process induced defects have to be systematically considered and studied since the complexity of the non-equilibrium SLM process [8, 9]. In order to achieve the desired mechanical and microstructure properties, the post-heat treatment were commonly required for relieving residual stresses and for facilitating the precipitates of strengthening phases [10, 11]. The primary strengthening mechanism of Inconel 718 alloy was used by heat treatment that involves solid solution annealing and ageing treatments [12, 13]. However, the post-heat treatment still have no a uniform standard for Inconel 718 alloys fabricated by SLM technology till now. Some typical industrial standards heat treatments for Inconel 718 castings and forgings were used to observe these changes in microstructure, phase precipitate and mechanical properties of Inconel 718 alloys fabricated by SLM [14, 15, 16]. Therefore, it is quite important to optimize the post-heat treatment process for facilitating the precipitates of strengthening phases and enhancing the mechanical properties of Inconel 718 alloys fabricated by SLM.

In this study, according to the field assembly requirement of SLM-ed Inconel 718 alloys components, samples produced by SLM experienced a heat treated process using different solid solution annealing and ageing treatments to observe these changes in microstructure, precipitates and mechanical properties on the X-Y surface, which provides a reference for the heat treatment process of Inconel 718 alloy fabricated by SLM.

Experimental

In this experiment, the rectangular bar of Inconel 718 alloy was fabricated by a SLM technology. The dimension of the rectangular bar is 250 mm×60 mm×12 mm. Inconel 718 spherical powder was used during the process of the SLM, and the power size is 45 ~ 65 μm. The laser power was 350 W, and the scanning speed was 700 mm/s. In addition, the Ar gas was used as the protective gas in the melting process. The chemical composition of Inconel 718 alloys fabricated is shown in Table 1.

Table 1:

The chemical composition of Inconel 718 alloys fabricated by selective laser melting.

Inconel 718 alloys fabricated were made into a series of samples with dimensions 10 mm×10 mm×10 mm by a lining cutting machine. Then, three heat treatment processes are adopted, and the detailed process is shown in Table 2. And then, the microstructure, microhardness distribution and phase constituent on the X-Y surface of alloys (see Figure 1) were performed before and after the heat treatment of samples by means of optical microscopy (OM), microhardness tester, scanning electron microscopy (SEM) and X-ray diffraction (XRD). For a microstructural analysis, a ratio of eroding solution is 5 g CuCl2: 100 ml HCl: 100 ml CH3CH2OH. Other experimental equipments used are as follow: Horizontal metallurgical microscope of Nikon Epiphot 300 U/200, SEM equipment of JSM-6380LA and X-ray diffractometer of D/max-rc.

Theschematic of Inconel 718 Ni-based alloys fabricated by SLM.
Figure 1:

Theschematic of Inconel 718 Ni-based alloys fabricated by SLM.

Table 2:

The different heat treatment process.

Results and discussions

Microstructure on the X-Y surface of Inconel 718 alloys

Figure 2 shows the microstructure on the X-Y surface of Inconel 718 alloys under optimized different heat treatment processes. According to Figure 2(a), the obvious weld beads structure can be observed on the X-Y surface of Inconel 718 alloys with no heat treatment process. This main reason is due to the deposition of metal between the weld beads in the SLM process. The obvious and large weld beads structure existed could induce the intermetallic compounds formed easily. Moreover, the microstructure shows the seriously inhomogeneous on the X-Y surface with lots of distributing columnar crystals and fine dendrites (see Figure 2(a)).

Microstructure on the X-Y surface of Inconel-718 alloys: (a) Microstructure of the as-built alloy, (b) Microstructure of A process, (c) Microstructure of B process and (d) Microstructure of C process.
Figure 2:

Microstructure on the X-Y surface of Inconel-718 alloys: (a) Microstructure of the as-built alloy, (b) Microstructure of A process, (c) Microstructure of B process and (d) Microstructure of C process.

However, when the A process was used, a large of columnar crystals and small dendritic crystals begun to transform into the larger bulk crystals. Then, the obvious weld beads structure begun to fuse and disappear gradually (see Figure 2(b)). When the B process, a high temperature solution treatment at temperature of 950°C, was used, the weld beads structure completely disappeared. Only some relatively large columnar crystals and dendrites could be observed (see Figure 2(c)). Finally, after the C process, the microstructure of alloys showed the homogeneous bulk crystal structure, and the columnar crystals and dendrites also completely disappeared (see Figure 2(d)).

Precipitates and microstructural transition on the X-Y surface of Inconel-718 alloys

In order to further know the influence of the heat treatment process on the microstructural transformation on the X-Y surface of Inconel-718 alloys, SEM with energy dispersive spectrometer was used to observe and analyze the alloys under different conditions. The phase constituents of alloys were analyzed by means of XRD. Figure 3 shows the microstructure of alloys by OM and SEM. The EDS results of some typical precipitates on the X-Y surface of alloys are shown in Table 3. In addition, the XRD results under different conditions are shown in Figure 4.

SEM morphology of Inconel 718 alloys under different conditions: (a) Optical microscope of the as-built alloy, (b) SEM morphology of the as-built alloy, (c) SEM morphology of A process, (d) SEM morphology of B process and (e) SEM morphology of C process.
Figure 3:

SEM morphology of Inconel 718 alloys under different conditions: (a) Optical microscope of the as-built alloy, (b) SEM morphology of the as-built alloy, (c) SEM morphology of A process, (d) SEM morphology of B process and (e) SEM morphology of C process.

XRD curve of Inconel-718 Ni-based alloys under different conditions: (a) As-built alloy, (b) A process, (c) B process and (d) C process.
Figure 4:

XRD curve of Inconel-718 Ni-based alloys under different conditions: (a) As-built alloy, (b) A process, (c) B process and (d) C process.

Table 3:

The EDS analysis of segregations and precipitates on the X-Y surface of alloys(wt.%).

The test results indicate that the main phase of Inconel 718 alloys with no heat treatment is γ phase (see Figure 3(a)), and elements Cr, Al and Mo could be dissolved in γ phase. Some researches indicated that the properties of main phase would be improved by means of lattice distortion [17]. Moreover, the chemical composition among dendrites on the X-Y surface is composed of high temperature elements, such as Fe, Cr, Ni and Nb, and the precipitates Cr-Ni-Fe-C can be also found in the region (see Figure 3(b), Table 3 and Figure 4). The carbide structure is brittle and hard, which leads to cracks by the separation of matrix interface, and cracks are easily broken at a high temperature. Therefore, the comprehensive mechanical properties of alloys will be seriously affected. According to the XRD analysis, It is obvious that a γ′ (Ni3 (Al, Ti, Nb)) phase that has higher structure stability (γ′ phase is precipitated at 600°C) could be also obtained in this region [11, 14].

When the A process was used, the partially precipitated phases and intermetallic compounds were dissolved, and the obvious carbides Cr-Ni-Fe-C also disappeared (see Table 3 and Figure 4). On the contrary, a mass needle-like γ′ phase with rich Nb, Ni3 (Al, Ti,Nb), was precipitated in crystal grain (see Figure 3(c)). Therefore, the strength and thermal stability of alloys will be further improved under the effect of γ′ phase with rich Nb. When the high temperature solution treatment at 950°C was finished, some intermetallic compounds and partial segregated hardening phases were also dissolved (see Figure 4). At this time, the massive white precipitates were observed at the grain boundaries (see Figure 3(d)). The EDS analysis indicated the element Nb was released to form NbNi3 phase (δ phase of solution) at this high temperature. Moreover, there is no obvious carbide precipitates on the X-Y surface of alloys. As a result, δ phase of solution precipitated at the grain boundaries will play a certain role for grain refinement based on dislocation pinning effect [18].

When the C process was used, some intermetallic compounds and segregated hardening phases were also dissolved. However, a large number of δ phases of solution were further promoted to precipitate. According to analysis of SEM and EDS, the massive white δ phases of solution were observed at the grain boundaries and grain interiors, and the content of element Nb increased further. This is according to the analysis of SEM. At this time, a large number of δ phases of solution were dispersed in grain interiors and grain boundaries (see Figure 3e), and some carbide (such as Cr-Ni-Fe-C) was also precipitated (see Figure 4).

Hardness distribution on the X-Y surface of Inconel-718 Ni-based alloys

According to the above analysis of OM, SEM and XRD, the microstructure of alloys had an obvious change under the different heat treatment process. Therefore, it is necessary to reflect the relation between microstructure change and heat treatment process through analyzing the hardness distribution of alloys. The hardness distribution of alloys under the different heat treatment process is shown in Figure 5.

The curve of hardness distribution on the X-Y surface of Inconel 718 alloys.
Figure 5:

The curve of hardness distribution on the X-Y surface of Inconel 718 alloys.

When the A process was used, the surface hardness of alloys was about HV490-540. It is higher than the surface hardness of any other states. Then, some partially segregated phases and intermetallic compounds were dissolved, and elements Cr, Al and Ti were also released. These elements were dissolved in γ phases to improve the strength of alloys. Therefore, the needle-like γ′ phase (Ni3(Al, Ti, Nb)) is quite important for the ordered strengthening and coherent strengthening of alloys. When the B process was used, the alloys showed the high hardness distribution (about HV476-500). At this time, no carbide could be observed, and element Nb was released to form NbNi3 phase (δ phase of solution). As a result, the δ phase of solution precipitated at the grain boundaries will be helpful for grain refinement under the dislocation pinning effect. Compared with the A process, some intermetallic compounds and partial segregated hardening phases were dissolved, and γ phase Ni3(Al, Ti,Nb) also decreased obviously. Therefore, the surface hardness of alloys was reduced obviously. When the C process was used, the surface hardness of alloys was close to the one of the as-built alloys (about HV400-450). During this process, some intermetallic compounds and segregated hardening phases were further dissolved. Element Nb was released further to form NbNi3 phase (δ phase of solution). Moreover, a large number of δ phases of solution were also dispersed in grain interiors and grain boundaries. At this time, the grain growth will be limited by the obvious dislocation pinning effect [19]. As a result, the alloys showed relatively fine structure. Therefore, it has a certain effect on the impact resistance and low-temperature brittleness resistance of alloys. On the contrary, the hardness distribution of the as-built alloys is lower relatively than the one of other states, which is about HV300-350. The reason is that the obvious brittle precipitated phases and the consumption of element Nb.

Discussion

Precipitates under B and C process

SEM morphology and EDS analysis of precipitates on the X-Y surface of alloys after B process and C process are shown in Figure 6. After B process, the particle-like precipitates were observed at the grain boundaries, and the needle-like γ’ precipitates was dispersed and distributed in grains (see Figure 6(a)). When the C process was used, a mass particle-like precipitates were observed at the grain boundaries and in the grain interiors. At this time, some precipitates were gathered with different size at the grain boundaries. According to the EDS analysis (see Table 3), the particle-like precipitates was mainly composed of element Nb, Ni and Fe. In addition, compared with the chemical composition of precipitates under different solution treatment temperature, the content of element Nb was improved with the increase of solid solution temperature. Under the same solid solution temperature, the content of element Nb was also improved with the increase of the volume of particle-like precipitates.

SEM morphology of typical precipitates under different solution temperature: (a) Local SEM morphology of B process and (b) Local SEM morphology of C process.
Figure 6:

SEM morphology of typical precipitates under different solution temperature: (a) Local SEM morphology of B process and (b) Local SEM morphology of C process.

However, after 950°C solution treatment, some intermetallic compounds and partial segregated hardening phases, such as Cr-Fe-Mo-Ni, FeNi and Ni8Nb, were further dissolved. Compared with the A process, no carbide could be found (Cr-Ni-Fe-C). The element Nb will be precipitated in the form of δ phase of solution. As a result, a substitutional solid solution of NbNi3 (δ phase of solution) was formed. During the process, Fe in FeNi3 was displaced. In addition, the atomic radius of Nb is slightly larger than the one of Fe. Compared with FeNi3, the lattice structure and interplanar spacing of δ phase are basically the same as those of FeNi3. Therefore, δ phase of solution could not be found by means of the XRD test.

After the C process, the most remarkable change of alloys is that the content of Nb increases greatly. A large number of δ phases of solution were dispersed in grain interiors, and the δ phase of solution was gathered in grain boundaries with different size (see Figure 6(b)). In addition, some carbide could be found (such as Cr-Ni-Fe-C). However, some intermetallic compounds and segregated hardening phase (such as Cr-Fe-Mo-Ni, FeNi and Ni8Nb) were dissolved. At this time, the grain growth process was limited by obviously dislocation pinning. Therefore, the structure of alloys shows fine relatively. This has a certain effect on the impact resistance and low-temperature brittleness resistance.

Effect of solid solution temperature on hardness and phase transformation

According to the above heat treatment analysis including the direct double ageing and solid solution+double ageing, the microstructure on the surface of Inconel 718 alloys had an obvious change in some respects of structure characters, hardness, precipitates and phase constituent. Based on OM and SEM analysis, the obvious weld bead structure has been got off by means of a direct double ageing process. However, the large substrate grain and dendritic structures still exited in the alloys, and the hardness value was higher than the one of other processes. According to the XRD result, the reason could be relation to be existence of Cr-Ni-Fe-C and Ni-Cr-Co-Mo (see Figure 4(a)). When the solid+double ageing process was added, the surface structure showed a typical grain refining characteristic with the increase of solid solution temperature. At present, the above high hardness phases decreased, and even were disappeared. Thus, the surface hardness shows an obvious decreasing trend (see Figure 5).

In addition, with the dissolve and solution of some elements, the phase constituents on the surface of alloys are also uncomplicated. However, it is clear that the surface hardness rapidly decreased and closed to the hardness of alloys with no heat treatment when the 1050°C solid solution was used. On the contrary, the grain structure of alloys shows the obvious refining and homogenized characteristic. This is contradictory with the Hall-Petch theory [20]. This is to say, the hardness shows an obvious decreasing characteristic when the grain obtained an obvious refining. According to Figure 6(b), the gather of bulk δ phase near the grain boundary could be an important effect factor. The disappearance of high density dislocations and changes of surface stress on the surface of alloys both were also the possible influencing factor under high temperature process [21].When the 950°C solid solution process was used, the hardness on the surface of alloys was in middle of the as-built and 1050°C solid solution alloys. Moreover, the δ phase near the grain boundary and γ′ phase with rich Nb in the grain show uniform and fine characteristic (see Figure 6(a)). In this sense, the surface structure will be favorable to enhance the hardness. Unfortunately, the surface grain size is still not uniform. Therefore, the heat treatment process still need to be further optimized to obtain a perfect surface structure of alloys. More detailed researches will be reported in the future.

Conclusions

(1) The obvious weld beads structure was observed on the X-Y surface of the as-built Inconel-718 alloys. The microstructure of alloys shows seriously inhomogeneous with distributing lots of columnar crystals and fine dendrites. After the direct double ageing, the columnar crystals and small dendritic crystals were transformed into bulk crystals, and the weld beads structure begun to fuse. After the 950°C solid solution+double ageing, the weld beads structure disappeared, and the alloys showed relatively large columnar crystals and dendrites. When the 1050°C solid solution+double ageing were used, the microstructure showed a typical homogeneous bulk crystal structure, and the columnar crystals and dendrites disappeared.

(2) The X-Y surface hardness of alloys after the direct double ageing is about HV490-540, and it is higher than the surface hardness of any other state. After the 950°C solid solution+double ageing, the alloys have a high hardness, and it is about HV476-500. δ phase (NbNi3) phase of solution precipitated at the grain boundaries could play a certain role in grain refinement. After the 1050°C solid solution+double ageing, the surface hardness of alloys was close to the surface hardness of the as-built alloys (HV400-450). Some intermetallic compounds and segregated hardening phases were further dissolved. Element Nb was released further to form NbNi3 phase. In addition, the hardness of the as-built alloys is relatively low since the increase of high brittle precipitated phase and the consumption of Nb.

(3) XRD and SEM analysis indicated that a mass needle-like γ′ phase with rich Nb, Ni3 (Al, Ti,Nb), was precipitated in crystal grain during the direct double ageing. During the 950°C solid solution+double ageing, element Nb was precipitated in the form of δ phase of solution. The substitutional solid solution of NbNi3 was formed through displacing Fe in FeNi3. After the above two processes, no carbide was found. However, after the 1050°C solid solution+double ageing process, the content of Nb increased greatly. A large number of δ phases of solution were dispersed in grain interiors, and the δ phase of solution was gathered in grain boundaries with different size. The grain growth will be limited, which shows relatively fine structure.

Acknowledgements

This research was financially supported through the Shandong Provincial Natural Science Foundation, China (Grant No. ZR2016JL017) and National Natural Science Foundation of China (Grant No. 51305240).

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About the article

Received: 2018-02-15

Accepted: 2018-06-24

Published Online: 2018-09-19

Published in Print: 2019-02-25


Citation Information: High Temperature Materials and Processes, Volume 38, Issue 2019, Pages 229–236, ISSN (Online) 2191-0324, ISSN (Print) 0334-6455, DOI: https://doi.org/10.1515/htmp-2017-0182.

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