The microstructure evolution and tensile properties of Inconel 718 fabricated by high-deposition-rate laser directed energy deposition

Journal Pre-proof The microstructure evolution and tensile properties of Inconel 718 fabricated by high-deposition-rate laser directed energy deposition Zuo Li, Jing Chen, Shang Sui, Chongliang Zhong, Xufei Lu, Xin Lin

PII:

S2214-8604(19)30845-0

DOI:

https://doi.org/10.1016/j.addma.2019.100941

Reference:

ADDMA 100941

To appear in:

Additive Manufacturing

Received Date:

25 June 2019

Revised Date:

4 October 2019

Accepted Date:

8 November 2019

Please cite this article as: Li Z, Chen J, Sui S, Zhong C, Lu X, Lin X, The microstructure evolution and tensile properties of Inconel 718 fabricated by high-deposition-rate laser directed energy deposition, Additive Manufacturing (2019), doi: https://doi.org/10.1016/j.addma.2019.100941

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The microstructure evolution and tensile properties of Inconel 718 fabricated by high-deposition-rate laser directed energy deposition

Zuo Li a, b, Jing Chen a, b, *, Shang Sui a, b, Chongliang Zhong c, Xufei Lu a, b, Xin Lin a, b

State Key Laboratory of Solidification Processing, Northwestern Polytechnical

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a

University, 127 Youyixilu, Xi’an, Shanxi 710072, PR China

Key Laboratory of Metal High Performance Additive Manufacturing and Innovative

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b

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Design, MIIT China, Northwestern Polytechnical University, 127 Youyixilu, Xi’an,

Fraunhofer Institute for Laser Technology ILT, Steinbachstr. 15, 52074, Aachen,

Corresponding author: Jing Chen ([email protected])

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*

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Germany

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c

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Shanxi 710072, PR China

Tel: +8629 8849 4001 Fax: +8629 8849 4001

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Graphical Abstract

Highlights

The IN718 sample with deposition rate of 2.2 kg/h and height 75 mm was prepared.

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δ, γ" and γ' phase are precipitated in bottom and middle region due to thermal cycle.

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The microhardness and room temperature tensile properties exhibit a high value.

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Abstract:

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In order to meet the requirements for rapid manufacturing of large-scale highperformance metal components, the unique advantages of high-deposition-rate laser

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directed energy deposition (HDR-LDED, deposition rate ≥ 1 kg/h) technology have

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been attracted great attention. HDR-LDED technology significantly improves the efficiency by simultaneously increasing the mass and energy input on basis of conventional laser directed energy deposition (C-LDED, deposition rate ≤ 0.3 kg/h), which dramatically changes the solidification condition and thermal cycling effect compared to C-LDED processes. Based on this, Inconel 718 bulk samples were fabricated with a deposition rate of 2.2 kg/h and a height of 75 mm. Through

experimental observation combined with finite element simulation, the precipitation morphology, thermal cycling effect and tensile properties at room temperature of the block samples at heights of 6 mm (bottom region), 37 mm (middle region) and 69 mm (top region) from the substrate were investigated. The results show that both temperature interval and incubation time satisfy the precipitation conditions of the second phases because of the intense thermal cycling effect so that δ, γ" and γ' phase

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are precipitated in the bottom and middle region of the as-deposited sample during the

HDR-LDED process. As a result, the micro-hardness and the yield strength of the bottom region (385 HV; 745.1±5.2 MPa) are similar to those of the middle region (381

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HV; 752.2±12.1 MPa), respectively. And they are both higher than those of the top

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region (298 HV; 464.7±44.2 MPa). The tensile fracture mechanism is shown in both fracture and debonding of the Laves phase. The inhomogeneous microstructures and

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corresponding mechanical property differences of Inconel 718 fabricated by HDR-

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LDED along the deposition direction suggest the necessity to conduct further research of the post heat treatment in the future.

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Key words: high-deposition-rate laser directed energy deposition, Inconel 718, thermal cycle, second phase, mechanical properties

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1. Introduction

As one of the most representative nickel-base superalloys, Inconel 718 is widely

used in large-scale integrated equipment such as blades (Blade-Integrated-Disks) and turbine engines due to its excellent mechanical properties and oxidation resistance. To realize the manufacturing of large-scale complex Inconel 718 structural components,

laser directed energy deposition (LDED) technology has gradually become widely accepted. However, for the manufacturing of integral structural parts, the production efficiency of conventional laser directed energy deposition (C-LDED), with a deposition rate of approximately 0.5 kg/h, has been not meet the rapid production requirements [1-3]. As of today, the deposition rates of HDR-LDED technology is at least ten times as high as of C-LDED [4-6], which greatly reduces manufacturing time

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and costs. Therefore, HDR-LDED technology, which provides an effective way for rapid production of large-scale metal structural parts, is expected to become one of the most important technologies in this field.

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Compared with C-LDED, HDR-LDED technology is achieved by significantly

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increases both energy and mass input at the same time, which changes the process conditions during HDR-LDED compared to C-LDED process. Firstly, the change in

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process conditions focused on the solidification process in HDR-LDED technology.

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M.M. Ma proposed that a lower energy input will lead to a refined microstructure and less Laves phase in the as-deposited Inconel 718 sample since the cooling rate is

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increased [7]. C.L. Zhong investigated the microstructure of Inconel 718 fabricated by HDR-LDED and found that Laves phase is coarser in HDR-LDED compared to C-

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LDED, which is the result of the decrease in cooling rate [8]. Additionally, there are differences in Nb concentrations between C-LDED (4.92 wt. %) and HDR-LDED (5.2 wt. %) Inconel 718 powder, which was used by C.L. Zhong. He found that the concentration of Nb (15.82 wt. %) in the Laves phase of high-deposition-rate laser direct energy deposited (HDR-LDEDed) Inconel 718 was significantly lower than that

of conventional laser direct energy deposited (C-LDEDed) Inconel 718 (25.44 wt. %) [4, 9]. This research indicates that the changes of solidification conditions will affect micro-segregation regardless of alloy composition. Secondly, the change in the process conditions focused on the thermal cycling effect. Currently, many scholars have studied the effects of thermal cycling on the microstructure [10-13]. W. Sames et al. found that during electron beam melting (EBM) process, the intense thermal cycling effect causes

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a large amount of the strengthening phases γ" and γ' to precipitate in the as-deposited Inconel 718 sample and the corresponding micro-hardness is up to 400 HV [13]. Y. Tian

et al. also observed this phenomenon in C-LDED technology, which is considered to be

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the result of thermal cycle as well [11]. According to the above research results, it can

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be said that the change of thermal cycling effect significantly affects the precipitation phase of the as-deposited sample. However, F.C. Liu et al. also chose the C-LDED

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technology to prepare the as-deposited Inconel 718, and no precipitation of the

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strengthening phase was observed in the samples, although there is the thermal cycling effect in C-LDED technology. X.Q. Wang hasn't observed the precipitation of

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strengthening phase in selective laser melting (SLM) technology when investigating the effects of the thermal cycling on the microstructures [10]. In HDR-LDED, the

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increase in energy input will intensify the thermal cycling effect, and conversely, the increase in mass input will weaken this effect. It remains unclear how the combined factors of the two will affect the thermal cycle. Based on literature, the question whether the change in thermal cycling effect influences the microstructures of the as-deposited Inconel 718 remains unanswered. This needs to be researched especially in HDR-

LDED. In terms of the microstructure of HDR-LDEDed Inconel 718, C.L. Zhong initially deposited the single-track cladding layer, and it was found that there was no significant difference between the microstructure of HDR-LDEDed samples and that of C-LDED [14], although the mass and energy input were increased during the forming process. However, it is worth mentioning that the dilution area of the single-track cladding layer

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is very large according to the results of the research. This means that the re-melting depth of the forming layer to the solidified layer will be significantly increased when

the bulk sample is deposited. Correspondingly, the average temperature inside the bulk

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sample and the duration time of the high temperature range will be improved. As a

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result, the thermal cycling effect inside the deposited sample will be particularly significant. The intense thermal cycling effects may cause the temperature to rise to the

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phase transition point or even induce solid phase transition during HDR-LDED

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technology since Inconel 718 is a precipitation-strengthening alloy. Therefore it is estimated that when the sample transitions from a single-track to bulk one, the thermal

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cycling effect may play an important role in the microstructure and further affect the type of precipitated phases. Obviously, the changes of the precipitation phases will

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eventually affect the tensile properties and fracture mechanism of the as-deposited sample. Based on the above analysis, whether the thermal cycling effect will affect the microstructure and tensile properties of the bulk samples in HDR-LDED technology still needs to be further explored. In this paper, the influences of the thermal cycling effect on the precipitation phase,

micro-hardness and tensile properties of HDR-LDEDed Inconel 718 alloy have been investigated. In addition, the tensile fracture mechanism of HDR-LDEDed Inconel 718 also has been analyzed. 2. Experimental details

2.1 Experimental materials and procedures

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The experimental equipment is established by State Key Laboratory of Solidification Processing. This system consists of a 6 kW semiconductor laser, a threedimensional numerical controlled working table, an inert atmosphere processing

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chamber (oxygen content ≤50 ppm), a powder feeder with a coaxial nozzle and an adjustable automatic feeding device with high precision (type DPSF-2).

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A set of HDR-LDED parameters (diode laser power, 4000 W; scanning speed,

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1600 mm/min; beam diameter, 5 mm; overlap, 50%; layer thickness: 1.3 mm) was utilized to fabricate an Inconel 718 cube with the dimension of 45 mm×45 mm×75 mm

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(length × width × height) on a C45E4 stainless steel substrate with the dimension of 140 mm×50 mm×5 mm (length × width × height). The deposition rate was about 2.2

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kg/h. The HDR-LDEDed Inconel 718 cube is shown in Fig. 1 (a). The substrate surface

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was polished to remove the oxide film by abrasive paper and cleaned with acetone beforehand. Spherical powder Inconel 718, which has been manufactured by the Plasma-Rotating Electrode Process (PREP), was used as the deposition material. The chemical composition of the Inconel 718 powder is as follows (wt. %): 18.48Cr, 0.54Al, 5.30Nb, 3.04Mo, 0.95Ti, 0.03C, 52.76Ni and balance Fe. The diameters of the used PREP powder particles range from 45 μm to 90 μm.

The microstructures were revealed by using the etchant of 1 g CrO3+2 ml H2O+6 ml HCl and the microstructure of the sample was examined by optical microscopy (OM), transmission electron microscopy (TEM) and scanning electron microscope (SEM) equipped with energy disperse spectrometer (EDS). The Vickers micro-hardness distributions along the deposition direction in the bottom, middle and top region were determined using a Struers Duramin-A300 Vickers micro-hardness tester at a load of

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500 g for 15 s. The interval between two adjacent points is 1000 μm. JMatPro software was used to calculate the time-temperature-transformation (TTT) curves of δ, γ" and γ' phase, which is based on new computer models for materials characteristics simulation and

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displays the results via a user friendly interface. JMatPro has been evaluated in different

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studies and good agreement with predicted material characteristics (alloy composition, phases, phase composition, physical and mechanical properties, etc.) and experimental

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results have been achieved. The calculation of the TTT curve depends on the specific

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composition of the alloy. Element chemical composition (wt. %) of Inconel 718 is as follows: Fe: 4.91%, Cr: 21.82%, Ni: 59.437%, Nb: 3.52%, Mo: 9.28%, Ti: 0.34%, C:

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0.033%, Al: 0.29%, Si: 0.27%, Mn: 0.10%. ANSYS software was used to calculate the thermal cycles of the bottom region

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(Z=~6 mm), the middle region (Z=~37 mm) and the top region (Z=~69 mm) during HDR-LDED process. Room temperature tensile tests were performed at a constant crosshead displacement rate of 1 mm/min on sheets with gage dimensions of 24 mm × 6 mm × 2 mm by an INSTRON3382 machine. As shown in Fig. 1(a), three plate-shaped tensile specimens were cut from the bottom, middle and top region of the bulk sample

along the x-axis direction and were sequentially labeled as 1#, 2# and 3#, respectively. The tensile specimen size is shown in Fig. 1(b). Three tensile specimens at different

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regions were prepared and the results were averaged out.

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Fig. 1 The HDR-LDEDed Inconel 718 alloy: (a) block sample; (b) specimen size for tensile test

2.2 Thermal cycle simulation

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The thermal history is calculated by performing a three-dimensional transient

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thermal analysis using the finite element model (FEM). The establishment of the thermal model is mainly based on the following equations and follow the law of

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conservation of energy. That is, the heat stored inside the entire material volume is equal to the external input sum minus the heat flow loss caused by radiation and

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convection. The detailed description of the model is given in the reference [15-17].

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The transient heat transfer energy balance is controlled in the entire volume of the material as follows: 𝜌𝐶𝜌

𝑑𝑇 𝐷𝑡

= −𝛻 · 𝑞(𝑟, 𝑡) + 𝑄(𝑟, 𝑡)

(1)

where ρ is the material density, Cp is the specific heat capacity, T is the temperature, t is the time, Q is the internal heat generation rate, r is the relative reference coordinate, and q is the heat flux.

The Fourier heat flux constitutive relation is followed by: 𝑞 = −𝑘 ∙ 𝛻𝑇

(2)

The heat flux is mainly determined by the thermal conductivity k depending on the temperature. Thermal radiation qrad is calculated using Stefan-Boltzmann's law: 𝑞𝑟𝑎𝑑 = 𝜉𝜎(𝑇𝑠 4 − 𝑇∞ 4 )

(3)

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where ξ is the surface radiance, σ is the Boltzmann constant, Ts is the surface temperature of the sample, and T∞ is the ambient temperature.

Newton's law of cooling describes the heat loss qconv due to convection:

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𝑞𝑐𝑜𝑛𝑣 = ℎ(𝑇𝑠 − 𝑇∞ )

(4)

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where h is the convective heat transfer coefficient.

When the finite element method is applied to the laser directed energy deposition,

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the total input heat is mainly characterized by three parameters: laser power P, scanning

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speed V and laser absorption rate η. The heat loss during deposition is characterized by the radiation coefficient α. The high deposition rate is characterized by the layer

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thickness δ in the process. Following the law of conservation of energy, the thermal cycle curve of HDR-LDEDed Inconel 718 superalloy is finally realized.

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The boundary conditions for simulating the thermal cycling process are as follows:

both heat radiation and heat convection conditions are applied to all the external surfaces of the base-plate and the metal deposition. The heat transfer coefficient by convection for HDR-DED process of Inconel 718 is set to 18 [W/m2·ºC], while the emissivity parameter is set to 0.285. The reference does state that the value is only an

approximation, as the emissivity will vary with surface finish. The environment temperature is 25 °C. Finally, the energy absorption coefficient, η, is set to 0.3.

3. Results and discussion

3.1 Microstructure

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Fig. 2 shows the microstructures of HDR-LDEDed Inconel 718 from the bottom region to the top region. Columnar grains growing epitaxially along the deposition direction in the bottom, middle and top region can be seen in the optical micrographs

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(Fig. 2(a), (b) and (c)). Therefore the heat flow direction during HDR-LDED process is

approximately perpendicular to the surface of the substrate or the pre-deposited layers

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[18]. The columnar grains in these three regions are large and some even extend over

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several deposited layers. The widths of the columnar grains range from 100 μm to 500 μm. Clearly visible layer bands are shown in Fig. 2(a)-(c). It is induced by the different

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microstructures between the bottom areas (planar interface growth) and other areas (dendrites growth) of a layer [9].

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A lot of irregular and light-grey phases are noted in the inter-dendritic region.

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According to previous research, these inter-dendritic light grey phases are Laves phases [4]. In addition, the transmission electron experiment was performed and the selected area electron diffraction (SAED) results further indicate that the light gray phase is Laves phase as shown in Fig.4 (i). The EDS element analysis on the Laves phase at different regions is carried out in Fig. 3 and the results of chemical compositions are shown in Table 1. It can be seen that the Nb concentration in the Laves phase in the

bottom, middle and top region are 24.23 wt. %, 25.39 wt. % and 24.02 wt. % respectively, which is much higher than the one in the matrix (max. 2.50 wt. %). In addition, a certain amount of sub-micron needle-shaped phases and nano-scale phases precipitate around the inter-dendritic Laves phases in both the bottom region and the middle region in Fig. 2(d) and (e). The SAED results indicate that the needleshaped precipitation is δ phase and the nano-scale precipitations are γ" and γ' phase as

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shown in Fig. 4. Ding figured out that the surrounding δ phase has an orientation relationship with the Laves phase [19]: (010)𝛿 ∥ (1̅010)𝐿𝑎𝑣𝑒𝑠

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[100]𝛿 ∥ [1̅21̅6]𝐿𝑎𝑣𝑒𝑠

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The distribution of γ" and γ' phase is non-uniform as shown in Fig. 2(d) and (e). It can be seen that γ" and γ' phase is mainly precipitated around the Laves phase, as indicated

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by the red dotted line. This is mainly due to the non-uniform distribution of elements

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such as Nb, Al and Ti [1]. After the long-term aging treatment, the principal strengthening precipitate in alloy 718 is γ" phase, a disc shaped coherent precipitate

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with ordered D022 crystal structure represented as Ni3Nb. A small amount of a secondary strengthening precipitate γ' phase, a spherical coherent precipitate with L12

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crystal structure represented as Ni3(Al,Ti), also forms. However, under the short-term thermal cycle in this paper, γ" phase and γ' phase are not completely precipitated, and it is difficult to distinguish the two. In the top region, there is no strengthening phases observed as revealed in Fig. 2(f) and Fig. 4(g).

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Fig. 2 The microstructures of the HDR-LDEDed Inconel 718 sample from the bottom region to

the top region: a), d) and g) the bottom region; b), e) and h) the middle region; c), f) and i) the top

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region

Fig. 3 The EDS element analysis on the different measuring positions:

a) the matrix; b) the irregular precipitation; c) the area around the irregular precipitation

Table 1 EDS analysis results of as-deposited HDR-LDEDed Inconel 718 superalloy Al

Nb

Mo

Ti

Cr

Fe

Ni

Spectrum 1

0.46

2.34

2.23

0.91

20.26

20.78

53.01

Spectrum 2

0.25

13.51

2.92

2.20

15.20

14.68

51.25

Spectrum 3

0.42

24.23

4.81

1.94

13.54

13.33

41.72

Spectrum 4

0.44

2.37

2.04

1.00

20.07

20.89

53.19

Spectrum 5

0.47

9.41

2.60

1.72

17.58

16.53

51.69

Spectrum 6

0.34

25.39

5.49

1.69

13.52

13.37

40.21

Spectrum 7

0.42

2.50

2.05

0.92

20.07

20.40

53.63

Spectrum 8

0.28

11.55

2.56

2.07

16.23

15.71

51.60

Spectrum 9

0.31

24.02

5.55

1.53

13.70

41.20

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Location

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13.70

Fig. 4 TEM figures of the sub-micron needle-shaped phases and the nano-scale phases: a) the bright field image of the needle-shaped phases; b) the dark field image of the needle-shaped phases; c) the SAED image of the needle-shaped phases; d) the bright field image of the nano-

scale phases; e) the dark field image of the nano-scale phases; f) the SAED image of the nanoscale phases; g) the bright field image of the light grey phases; h) the dark field image of the light grey phases; i) the SAED image of the light grey phases;

The formation of δ phase mainly occurs in two ways. Firstly, it forms in the boundary, such as grain boundary, twinning boundary or phase boundary. Secondly, it forms inside gains through the reaction γ"→δ [20]. Obviously, the precipitation of δ phase in this study is carried out in the first way. Azadian pointed out that the nucleation

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and growth of δ phase needes a certain amount of Nb [21]. Radavich indicated that the Nb concentration required for the formation of δ phase is between 6 wt. % and 8 wt. % [22]. In HDR-LDEDed Inconel 718 samples, Nb segregation regions exist around the

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inter-dendritic Laves phases. The Nb content in this region is higher than 10 wt. % as shown in Fig. 3(c). This provides favorable conditions for the precipitation of δ phase.

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Besides, the formation of δ phase also needs a certain temperature and time. Fig.

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5 shows the TTT curve of δ phase. It is calculated by JMatPro software based on the chemical composition of the Nb segregation region. It can be seen that when the

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temperature is between 890 ℃ and 1040 ℃, the minimum incubation time required for

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δ phase precipitation is 540 s.

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Fig. 5 The Time-Temperature-Transformation (TTT) curves of γ", γ' and δ phase

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The time that the thermal cycle curve stays in the fastest temperature range of the δ phase is shown in the Fig. 6. The heights of 6 mm, 37 mm and 69 mm from the

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substrate along the deposition direction were selected to represent the bottom, middle and top region of the deposited sample, respectively. According to the thermal cycle

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curve, it can be found that the holding time of the top region between 890 ℃ and 1040 ℃ is only about 85 s, which is far less than the minimum time required for the formation

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of δ phase (223 s). So there is no δ phase forming in the top region. However, in the

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bottom and middle region, the holding times between 890 ℃ and 1040 ℃ are 1132 s and 1013 s, respectively. Therefore, the formation of δ phase occurs in these regions.

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Fig. 6 The time that the thermal cycle curve stays in the fastest temperature range of the δ phase

γ" phase is the main strengthening precipitation and γ' phase is the auxiliary

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strengthening precipitation in Inconel 718 superalloy. It has been reported that the volume fraction of γ" phase is four times that of γ' phase [23]. Therefore, most of the

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nano-scale phases observed in Fig. 2(g) and (h) are γ" phases. As can be seen in Fig. 5, in terms of γ" phase, the nose of the TTT curve occurs at about 960 ℃ for about 16 s

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and it precipitates rapidly in the temperature range between 850 ℃ and 1010 ℃. The γ"

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phase stays in its fastest temperature range for only 57 s as shown in Fig. 5. In the bottom and middle region, the holding times between 850 ℃ and 1010 ℃ are 1281 s and 903 s, respectively as shown in Fig. 7. They are far greater than the needed holding time of 57 s. So γ" phase forms in both the bottom and the middle region. At the top, the holding time between 850℃ and 1010℃ is only 88 s, which is slightly larger than the

minimum holding time (57 s). The precipitation and the growth of γ" phase are not

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sufficient and therefore are difficult to be detected.

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Fig. 7 The time that the thermal cycle curve stays in the fastest temperature range of the γ" phase

Although the volume fraction of γ' phase is smaller than that of γ" phase, its

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formation is still detected in the bottom and middle region. As shown in Fig. 5, the nose of the TTT curve for γ' phase occurs at about 920 ℃ for about 36 s. When the

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temperature ranges from 810 ℃ to 960 ℃, the precipitation of γ' phase is rapid. In this

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case, the holding time is about 124 s. Fig. 8 shows the holding times between 810 ℃ and 960 ℃ are 1252 s and 610 s in the bottom and middle region. Therefore γ' phase is generated in these regions. In the top region, the holding time between 810 ℃ and 960 ℃ is only about 94 s, which is shorter than the required holding time. Therefore, no γ' phase can be found in this region.

In the precipitation-strengthening Inconel 718 superalloy, the proper heat treatment system is essential for improving its mechanical properties. The heat treatment system should be formulated according to the as-deposited microstructure. When the as-deposited microstructure changes, the heat treatment system selected will change too. Thus, exploring post heat treatment system for HDR-LDED technology

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will become one of the important issues that needs to be further studied in the future.

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Fig. 8 The time that the thermal cycle curve stays in the fastest temperature range of the γ' phase

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3.2 Micro-hardness

Different microstructures can lead to different mechanical properties. Fig. 9 shows

the micro-hardness of HDR-LDEDed Inconel 718 from the bottom to the top along the deposition direction. It can be seen that the average micro-hardness of the bottom region (385 HV) is similar to that of the middle region (381 HV). And they are both larger than

the average micro-hardness of the top region (298 HV). This phenomenon is mainly due to the precipitation of γ" and γ' phase in the bottom and middle region. Meanwhile, their holding time in the fastest temperature range are short (Fig. 5(c) and (d)) so that the precipitations of γ" and γ' phase are not sufficient. As a result, the micro-hardness is still lower than that of the directly aged Inconel 718 sample (438 HV) [7]. It means that post heat treatment is still necessary for HDR-LDEDed Inconel 718. The

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phenomenon that the micro-hardness of the as-deposited sample is non-uniformly

distributed has been mentioned earlier before and the micro-hardness in the bottom

region of the bulk sample fabricated by LMD technology has reached 400 HV [11]. The

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author also shows that this phenomenon is due to the presence of non-uniformly

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distributed strengthening phases in the deposited samples.

Fig. 9 The micro-hardness of the bottom region, the middle region and the top region

Subsequently, the micro-hardness of the as-deposited Inconel 718 fabricated by different additive manufacturing (AM) technologies was determined [9, 11, 18, 20, 2435]. As can be seen in Fig. 10, different colors represent the mirco-hardness in different additive manufacturing methods. The green numbers in the right column represent the corresponding references. The micro-hardness of the bulk sample used for this paper is basically comparable to that of the as-deposited sample by selective laser melting

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(SLM). Nevertheless, the reasons for the increase in micro-hardness are significantly different. From the statistical results, the micro-hardness in the middle and bottom

region of HDR-LDEDed material used for this paper has far exceeded that of generally

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laser metal deposited Inconel 718. In the whole, the hardness of middle and bottom

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regions are comparable to the relatively higher values for the as-fabricated Inconel 718

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in literatures, and that of the top region is comparable to the relatively lower values.

Fig. 10 The micro-hardness for as-built Inconel 718 superalloy fabricated by different AM

3.3 Mechanical properties The nominal stress-strain curves of HDR-LDEDed Inconel 718 superalloy is shown in Fig. 11. The red, blue, and green curves represent the bottom, middle and top regions of the deposited sample respectively. Three repetitive tests were performed in every region to obtain the average and standard deviation as shown in Fig. 12. In the nominal stress-strain curves, it can be seen that the yield strength and tensile strength

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of the top region are lower than those of the middle and bottom. The elongation has a

large dispersion, and the reason will be analyzed later. Fig. 12 shows the average strength and plasticity in different regions of the deposited samples. It can be seen that

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the yield strength at the bottom, middle and top region are 745.1±5.2 MPa, 752.2±

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12.1 MPa and 464.7±44.2 MPa, respectively. The difference in elongation is not significant, but it is higher than aerospace material specifications (AMS) 5662/5663

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(solution treatment at 980°C/1 h/water quenching followed by two step aging at

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720°C/8 h/ furnace cooling to 620°C/8 h/air cooling to room temperature) [36] for Inconel 718 fabricated by conventional manufacturing processes such as casting and

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forging.

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superalloy

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Fig. 11 Room temperature tensile strain-strain curves of as-deposited HDR-LDEDed Inconel 718

Fig. 12 Tensile test results of HDR-LDEDed Inconel 718 alloy at room temperature

In order to clarify the level of the tensile test values in as-deposited Inconel 718 with additive manufacturing (AM) in this paper, the σ0.2 vs. total elongation to failure for as-deposited Inconel 718 made by SLM, EBM and LMD are displayed in Fig. 13.

Similarly, different colors represent different σ0.2 and elongation in different additive manufacturing methods and the green numbers also represent the corresponding references [9, 14, 30, 32, 33, 35, 37-52]. It can be seen that σ0.2 and elongation of this sample in the bottom and middle region are comparable to those of the as-deposited sample by SLM. It is well known that the grain size of as-deposited Inconel 718 by SLM is finer compared to LMDed Inconel 718. In terms of SLM, the increase in yield

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strength is mainly determined by the fine grains. Although the grain size is coarser

during laser directed energy deposition, it is commonly known that Inconel 718 also is a precipitation strengthened superalloy. The interaction between the precipitation and

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the dislocation is the essence of the precipitation strengthening. In this paper, the

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strengthening phases (γ" and γ') are precipitated in the bottom and middle region, which increases the yield strength of the as-deposited sample. According to Fig. 13, it seems

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that the σ0.2 and elongation of the sample have almost exceeded those of the as-built

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sample by LMD and EBM considering the literature listed so far. But it is worth mentioning that compared with the C-LDED and SLM, the microstructure of HDR-

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LDEDed sample is still relatively rough, which is mainly reflected in the nonuniformity of the strengthening phase and inconsistency of the microstructure from

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bottom to top region.

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Fig. 13 The yield strength vs. total elongation to failure for as-built Inconel 718 superalloy fabricated by different AM

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Fig. 14 shows the tensile fracture morphologies in the different regions of asdeposited sample. As shown in a), b) and c), it can be said that the dendritic morphology

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is still preserved on the top, middle and bottom region after the room temperature tensile

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test, respectively. There is a considerable number of dimples with obvious directional orientation along the columnar dendrites in the fibrous zone, which is the typical

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characteristic of ductile fracture (in (a'), (b') and (c')). It is well known that the hard and brittle Laves phase is precipitated between the dendrites in laser direct energy deposited

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Inconel 718 alloy and dislocations tend to cause stress concentration around the Laves phase, which provides a favorable condition for crack initiation and expansion. Therefore, when the material is subjected to tensile load, micro-pores easily formed around the Laves phase and cause the final fracture by growing and interconnecting under the stress loading. Since the Laves phase is precipitated from the inter-dendrites

(shown in (a')), the dimple morphology within the inter-dendritic region as the center of the dimple and the dendritic region as the tearing edge is formed. Because the growth direction of the dendrites is perpendicular to the tensile direction and the inter-dendritic region is the center of the dimple, the dimples on the fracture will exhibit a distinct

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feature along the columnar dendrites.

Fig. 14 Tensile fracture morphology along the tensile direction: a) and a ') top region; b) and b ') middle region; c) and c ') bottom region

Fig. 15 shows the longitudinal sections of the fractures. Fig 15(a) and Fig 15(a') represent the top region of the as-deposited sample. The green arrows point to the fracture of Laves phases and the separation of Laves phases which can be detected in the matrix. It is difficult to conclude the dominant role in breaking and debonding of the Laves phase, by the results givens. It is worth noting that the average yield strength of the top region of the deposited sample is about 464.7 MPa, but the elongation can be

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up to 52.08 %. In particular, the fracture morphology of longitudinal section of the

specimen was observed as shown in Fig. 16. Compared with (a) and (a'), it can be seen

Laves phase in this region is finer and more granular. The slip line is easier to go through

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the granular Laves phase than the large Laves phase, which facilitates plastic

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deformation of the matrix [53]. In addition, the fracture is mainly based on the fracture

phase is observed.

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of the Laves phase from the longitudinal section, and almost no debonding of the Laves

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Fig 15(b) and Fig 15(b') show the fracture morphology in the middle region of the as-deposited sample. It can be seen that both debonding and fracture of the Laves phase

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are existent. Compared with the top region, although the elongation in the middle region is slightly reduced, the yield strength and tensile strength are greatly improved. This is

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mainly because the dislocation is started and moves along the slip surface under stress loading when it encounters the γ" and γ' phase. On the one hand, the resistance to dislocation motion is generated, that is, the strengthening effect. On the other hand, the dislocations around the strengthening phase will cause stress concentration, which makes the deformation between the strengthening phase and the Laves phase

uncoordinated, leading to preferential fracture of the Laves phase. In addition, the elongation of the 1# sample in the middle region is significantly decreased compared with the 2# and 3# sample. It can be seen from the fracture morphology that the fracture is also caused by the breaking (green arrows) and debonding (blue arrows) of the Laves phase in Fig. 17. It is worth to note that the size and volume fraction of the δ phase are larger. The previous studies have shown although the δ phase does not influence the

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tensile strength, yield strength and micro-hardness values at room temperature, have only a slight influence on the ductility [50, 54]. Therefore, the increase of δ phase

volume fraction may be one of the reasons for plasticity decline in this paper.

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Additionally, the significant difference in the morphology of the δ phase also fully

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demonstrates the non-uniformity of the microstructure. This further affirms the necessity of a good heat treatment in the following. At last, Fig (c) and (c') are the

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fracture morphology in the bottom region of the as-deposited sample. Here, the fracture

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and debonding of the Laves phase are still dominant. The samples still demonstrate a

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good yield and tensile strength caused by the presence of the strengthening phase.

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Fig. 15 The longitudinal sections of the fractures: a) and a ') top region; b) and b ') middle region;

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c) and c ') bottom region

Fig. 16 The fractures of 3# sample in the top region

Fig. 17 The fractures of 1# sample in the middle region

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Combined with the microstructure of the deposited sample and the tensile test at room temperature, it can be seen that there is a significant inconsistency in the

microstructure of the deposited sample from the bottom to the top region. In addition,

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in the bottom and middle region, due to the limited thermal cycling effect, although there is the precipitation of strengthening phase, its distribution is still uneven and

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insufficient. On one hand, although there is the strengthening phases precipitation in

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the as-deposited sample, the yield strength and tensile strength are obviously lower than those of the heat treated, cast and wrought samples [37, 55-57], which means that the

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heat treatment system is necessary for improving the mechanical properties of the sample. Furthermore, the microstructure by HDR-LDED is significantly different from

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the as-deposited samples fabricated by C-LDED, which means that the typical heat

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treatment regime (homogenization at 1050°C/1 h/water quenching followed by two step aging at 720°C/8 h/ furnace cooling to 620°C/8 h/air cooling to room temperature [53]) may not apply to the HDR-LDEDed samples. It is necessary to explore the suitable heat treatment regime to improve the mechanical properties of the deposited samples in the following work.

Conclusion In summary, the thermal cycle during HDR-LDED process influences the precipitation of δ, γ" and γ' phase in the as-deposited Inconel 718. In the bottom and middle region, δ phases precipitate around the inter-dendritic Laves phases. Furthermore, a certain amount of γ" and γ' phases also form in these two regions. However, there is no δ, γ" and γ' phase existent in the top region. As a result, the micro-

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hardness of the bottom region (385 HV) is similar to that of the middle region (381 HV), and they are both higher than the micro-hardness of the top region (298 HV). The

yield and tensile strength of the top region are lower than those of the middle and bottom

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region. The elongation has a large dispersion. According to the fracture morphology,

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the fine granular Laves phase is beneficial to plasticity, while the increase of the δ phase volume fraction reduces the plasticity to some extent. The precipitation of the

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strengthening phase in the as-deposited sample imply a high importance to conducting

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technology.

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further research in finding a suitable post heat treatment system for HDR-LDED

Conflict of Interest

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We wish to draw the attention of the Editor to the following facts which may be considered as potential conflicts of interest and to significant financial contributions to this work. [OR] We wish to confirm that there are no known conflicts of interest associated with this publication and there has been no significant financial support for this work that could have influenced its outcome. We confirm that the manuscript has been read and approved by all named authors and that there are no other persons who satisfied the criteria for authorship but are not listed. We further confirm that the order of authors listed in the manuscript has been approved by all of us.

We confirm that we have given due consideration to the protection of intellectual property associated with this work and that there are no impediments to publication, including the timing of publication, with respect to intellectual property. In so doing we confirm that we have followed the regulations of our institutions concerning intellectual property.

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We understand that the Corresponding Author is the sole contact for the Editorial process (including Editorial Manager and direct communications with the office). He/she is responsible for communicating with the other authors about progress, submissions of revisions and final approval of proofs. We confirm that we have provided a current, correct email address which is accessible by the Corresponding Author and which has been configured to accept email from (e-mail address: [email protected])

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Acknowledgement

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Funding: This work was supported by the Sino-German Science Foundation (No.

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GZ1267).

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23

List of Figures Fig. 1 The HDR-LDEDed Inconel 718 alloy: (a) block sample; (b) specimen size for tensile test. Three plate-shaped tensile specimens were cut from the bottom, middle and top region of the bulk sample along the x-axis direction and were sequentially labeled as 1#, 2# and 3#, respectively. Three tensile specimens at different regions were

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prepared and the results were averaged out.

Fig. 2 The microstructures of the HDR-LDEDed Inconel 718 sample from the bottom

region to the top region: a), d) and g) the bottom region; b), e) and h) the middle region;

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c), f) and i) the top region.

Fig. 3 The EDS element analysis on the different measuring positions: a) the matrix; b)

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the irregular precipitation; c) the area around the irregular precipitation. Nb

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concentration in the Laves phase in the bottom, middle and top region are 24.23 wt. %, 25.39 wt. % and 24.02 wt. % respectively, which is much higher than the one in the

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matrix (max. 2.50 wt. %). There is a significant micro-segregation in the as-deposited

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sample.

Fig. 4 TEM figures of the sub-micron needle-shaped phases and the nano-scale phases: a) the bright field image of the needle-shaped phases; b) the dark field image of the needle-shaped phases; c) the SAED image of the needle-shaped phases; d) the bright field image of the nano-scale phases; e) the dark field image of the nano-scale phases;

f) the SAED image of the nano-scale phases; g) the bright field image of the light grey phases; h) the dark field image of the light grey phases; i) the SAED image of the light grey phases.

Fig. 5 The Time-Temperature-Transformation (TTT) curves of γ", γ' and δ phase. When the temperature is between 890 ℃ and 1040 ℃, the minimum incubation time required

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for δ phase precipitation is 540 s. When the temperature is between 850 ℃ and 1010 ℃, the minimum incubation time required for γ" phase precipitation is 57 s. When the

temperature is between 810 ℃ and 960 ℃, the minimum incubation time required for γ'

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phase precipitation is 57 s.

Fig. 6 The holding time that the thermal cycle curve stays in the fastest temperature

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range of the δ phase. it can be found that the holding time of the top region between

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890 ℃ and 1040 ℃ is only about 85 s, which is far less than the minimum time required for the formation of δ phase (223 s). So there is no δ phase forming in the top region.

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However, in the bottom and middle region, the holding times between 890 ℃ and 1040 ℃ are 1132 s and 1013 s, respectively. Therefore, the formation of δ phase occurs in these

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regions.

Fig. 7 The time that the thermal cycle curve stays in the fastest temperature range of the γ" phase. In the bottom and middle region, the holding times between 850 ℃ and 1010 ℃ are 1281 s and 903 s, respectively. They are far greater than the needed holding time of

57 s. So γ" phase forms in both the bottom and the middle region. At the top, the holding time between 850℃ and 1010℃ is only 88 s, which is slightly larger than the minimum holding time (57 s). The precipitation and the growth of γ" phase are not sufficient and therefore are difficult to be detected.

Fig. 8 The time that the thermal cycle curve stays in the fastest temperature range of the

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γ' phase. The holding times between 810 ℃ and 960 ℃ are 1252 s and 610 s in the bottom and middle region. Therefore γ' phase is generated in these regions. In the top

region, the holding time between 810 ℃ and 960 ℃ is only about 94 s, which is shorter

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than the required holding time. Therefore, no γ' phase can be found in this region.

Fig. 9 The micro-hardness of the bottom region, the middle region and the top region.

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The average micro-hardness of the bottom region (385 HV) is similar to that of the

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middle region (381 HV). And they are both larger than the average micro-hardness of

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the top region (298 HV).

Fig. 10 The micro-hardness for as-built Inconel 718 superalloy fabricated by different

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AM. From the statistical results, the micro-hardness of HDR-LDEDed material used for this paper is basically in good condition.

Fig. 11 Room temperature tensile strain-strain curves of as-deposited HDR-LDEDed Inconel 718 superalloy. The yield strength and tensile strength of the top region are

lower than those of the middle and bottom.

Fig. 12 Tensile test results of HDR-LDEDed Inconel 718 alloy at room temperature. the yield strength at the bottom, middle and top region are 745.1±5.2 MPa, 752.2±12.1 MPa and 464.7±44.2 MPa, respectively. The difference in elongation is not significant, but it is higher than aerospace material specifications (AMS) for Inconel 718 fabricated

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by conventional manufacturing processes such as casting and forging.

Fig. 13 The yield strength vs. total elongation to failure for as-built Inconel 718

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superalloy fabricated by different AM. It seems that σ0.2 and elongation of the sample

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have almost exceeded those of the as-built sample by LMD and EBM considering the

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literature listed so far.

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Fig. 14 Tensile fracture morphology along the tensile direction: a) and a ') top region; b) and b ') middle region; c) and c ') bottom region. The dendritic morphology is still

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preserved on the top, middle and bottom region after the room temperature tensile test, respectively. There is a considerable number of dimples with obvious directional

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orientation along the columnar dendrites in the fibrous zone, which is the typical characteristic of ductile fracture.

Fig. 15 The longitudinal sections of the fractures: a) and a ') top region; b) and b ') middle region; c) and c ') bottom region. The fracture and debonding of the Laves phase

are the main fracture mechanism.

Fig. 16 The fracture of 3# sample in the top region. The fracture is mainly based on the fracture of the Laves phase from the longitudinal section, and almost no debonding of the Laves phase is observed.

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Fig. 17 The fracture of 1# sample in the middle region. It is worth to note that the size

and volume fraction of the δ phase are larger. The increase of δ phase volume fraction

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may be one of the reasons for plasticity decline.

List of Tables Table 1 - Table 1 EDS analysis results of as-deposited HDR-LDEDed Inconel 718

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superalloy.