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Effect of laser power on porosity and mechanical properties of GH4169 fabricated by laser melting deposition

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The porosity and mechanical properties of GH4169 (a precipitation strengthened nickel-base superalloy) specimens fabricated using the laser melting deposition technique at different laser powers were investigated. The results showed that the dendritic structure with the Laves phase and carbides embedded in the Ni-γ matrix formed in the as-fabricated GH4169 due to the strong temperature gradient and the high cooling rates. Porosity remarkably decreased first and slightly increased subsequently as the laser power increased from 300 to 800 W. The lowest porosity of the specimens characterized by 3D X-ray tomography is 0.28%. The specimens fabricated at 600 W tensiled along the direction perpendicular to the building direction exhibit the average yield strength of 587 MPa, the ultimate tensile strength of 903 MPa, and the elongation at fracture of 13.6%. Furthermore, the fatigue limit of the 600 W fabricated specimens is 173.7 MPa corresponding to the fatigue ratio of 0.1. And the relationship among the porosity, laser power and mechanical properties is discussed.


Being the same precipitation strengthened nickel-base superalloys, GH4169 and Inconel 718 (IN718) are only different in trace elements, such as aluminum and titanium. They were widely used in gas turbine engines in the aerospace and power generations due to their excellent mechanical properties, oxidation and corrosion resistance at elevated temperatures [1,2,3,4,5]. Additive manufacturing (AM) is particularly attractive as a novel manufacturing technique for IN718 and GH4169, owing to their possibility of meeting the demands of complex part geometries and shortening the production cycle in industrial component production. Laser melting deposition (LMD), one of the AM technologies, is developed on the basis of rapid prototyping and laser cladding technology, utilizing the laser beam to melt powders fed by a coaxial powder feeding system, then solidifying rapidly layer by layer. The most important feature of LMD is the high cooling rate caused by the strong temperature gradient, which leads to the solid–liquid interface in the molten pool deviating from the equilibrium state. Meanwhile, macro-segregation of the alloy appears in traditional castings and forgings are prevented, and micro-segregation plays an important role in the formation of dendrites. Therefore, the cooling rate and the temperature gradient have important effects on microstructural evolution and the mechanical properties of the LMD-fabricated components.

What is more, there are many great challenges in the LMD-fabricated components, including cracking, porosity and residual stress. The pores in the LMD-fabricated IN718 and GH4169 components, which degrade mechanical properties, are caused by the incomplete melting of the powder, the residual shielding gas or incomplete bonding between the deposition layers [6]. It should be mentioned that the thermal gradient, solidification velocity and porosity of the components depend on the LMD processing parameters, such as the laser power, scanning velocity, and powder feeding rate. Choi et al. [7] investigated the effect of the applied laser scanning speed on the densification behavior and microstructural development of IN718 components fabricated by selective laser melting (SLM), implying that the minimum porosity (< 1%) could be achieved at the laser scanning speed of 800 mm s−1. The laser power and carrier gas flow have an influence on the porosity of the as-fabricated specimens. When the laser power decreases or the carrier gas flow is increased, voids appear in the as-fabricated specimens [8].

To achieve the desired microstructure and improve mechanical properties of the LMD-fabricated components, it is necessary to optimize the processing parameters. Much work so far has been focused on the relationship among processing parameters, including the pores and mechanical properties of the as-fabricated components. Zhu et al. [9] showed the relatively high yield strength and the ultimate tensile strength using a laser power of 1000 W due to the high volume fraction of equiaxed grains in laser solid forming in IN718. Similar researches [10, 11] concentrated on the relationship between laser powers and microstructures. Voisin et al. [12] found that strain-to-failure of selective laser-melted Ti–6Al–4V was strongly influenced by the presence of pores. However, few studies are devoted to the effects of the laser power on porosity and microstructural evolution during the LMD process. In this work, GH4169 specimens were fabricated using LMD at different laser powers. The effects of laser power on the microstructure, porosity and mechanical properties of the specimens were investigated systematically.


Specimen fabrication

GH4169 specimens were fabricated using an IPG Photonics fiber laser system with a coaxial powder feeder, with a beam spot size of 200 µm and the maximum power of 1000 W. High-purity argon was selected as both the shielding gas and carrier gas in the process. The spherical GH4169 powders with a size range of 40–200 μm were selected as the deposition material, and the substrate was a 304 stainless steel plate. During the LMD process, the powder feeding rate was 24 g min−1. The layers were scanned according to a zigzag pattern that was rotated 180° after each layer (see the laser scanning strategy shown in Fig. 1a). The as-fabricated GH4169 specimens with the size of 30 mm × 15 mm × 10 mm were prepared at different energy densities through varying laser powers, as presented in Table 1. Here, the energy density (E) is defined as the laser energy per area in a molten pool and can be calculated by Eq. (1) [13]:

$$ E = P / ( {V \times d} )$$

where P is the laser power, V is the laser scanning velocity, and d is the beam spot diameter perpendicular to the scanning direction (SD).

Fig. 1

a Schematic illustration of laser scanning strategy; b microstructure in three mutually perpendicular planes; c EBSD map showing the grain structure and grain boundary distribution; d inverse pole figures of c of the as-fabricated GH 4169 specimen at P = 600 W

Table 1 LMD manufacturing parameters used in this work

Characterization of microstructures and mechanical properties

The as-fabricated GH4169 specimens were ground and polished first using a standard procedure, and then electroetched for 10 s at 5 V in a 10% phosphoric acid solution for the dendrite observation, and electroetched for 15 s at 5 V in a 10% oxalic acid solution for the grain observation. Optical microscope (OM, Olympus DP71) and scanning electron microscope (SEM, LEO Supra 35) were used for the microstructure observation. The grain morphology, grain orientation, and grain size distribution were characterized by SEM equipped with an electron backscattering diffraction (EBSD) detector (HKL channel 5). The EBSD specimens were mechanically polished, and then electropolished for 50 s at 12 V in a 10% perchloric acid solution. The porosity of the as-fabricated specimens was quantified based on OM images and characterized by three-dimensional X-ray tomography (3D-XRT, Versa XRM-500). Energy-dispersive X-ray spectrometer (EDS) and transmission electron microscope (TEM, JEM-2100F) were used to investigate precipitated phases. Tensile and fatigue specimens were taken along the XY plane normal to the building direction (BD), as shown in Fig. 1a. Tensile and fatigue tests were carried out along SD at room temperature using the Micro-Tester (INSTRON 5848) and Dynamic Tester (INSTRON E1000), respectively. The fatigue load frequency is 50 Hz and stress ratio (R) is 0.1.

Results and discussion


Figure 1b presents a typical OM image of a three-dimensional microstructure of a specimen fabricated at P = 600 W. The laminar molten pools with a width ranging from 500 to 800 µm in XY and XZ planes, and the scaly molten pools in the YZ plane can be observed. It is worth mentioning that the width of the molten pool gradually increases with the increase of the laser power for different specimens by the measurement of OM images. With the increase of the laser power, the energy density input into the specimens increased by fixing the laser scanning velocity and the beam spot diameter according to Eq. (1). The lower laser energy input, such as in the conditions of 300 W and 400 W, enables the number and size of the pores to increase due to the existence of unmolten powder, the shielding gas residual, and the dendrites grow smaller and lose their uniformity. On the contrary, when the laser energy input is increased, such as in the conditions of 700 W and 800 W, respectively, the dendrites grow larger and become more uniform. In YZ and XZ planes of the as-fabricated specimens, the epitaxial growth of columnar dendrites with a length of several millimeters throughout several deposition layers can be observed clearly. The dendrite growth direction is determined by both the heat flux direction and crystallographically favored orientation [14]. Figure 1c shows that most of the columnar grains have a diameter up to ~ 780 µm, and few equiaxial grains have a diameter ranging from 50 to 100 µm. Inverse pole figures from this region shown in Fig. 1d indicate that there is a preferred orientation of <101>//Z (BD) in the columnar grains; however, the texture is relatively weak due to the low multiples of the uniform density, and the growth direction of dendrites is partly affected by the heat flow whose direction varies with the position of deposition layers [15].

Figure 2a shows an SEM image of a specimen fabricated at the laser power of 600 W, in which a number of precipitates and dendrites can be found. The element concentration of the matrix and precipitates examined by EDS is shown in Table 2. Point 1 in Fig. 2b is a γ matrix with relatively high contents of Ni, Fe and Cr, respectively; while the contents of C, Ti and Nb in point 2 are much higher than those in the matrix. Thus, the white island precipitates are MC (M = metal elements)-type carbides of (Ti, Nb)C. The random shape particles shown in Point 3 have high contents of Nb, Mo and Ti according to the EDS result, indicating the presence of the Laves phase. TEM observations indicate that two types of precipitate phases exist in the as-fabricated specimens: the larger random shaped Laves phase shown in Fig. 2c, as observed by Xiao et al. [16], and the smaller square MC-type precipitate with a size of 0.6 µm shown in Fig. 2d, as also found by Sui et al. [17]. Figure 3a–c shows the dendrite morphology (on the XY plane) in the specimens fabricated at the laser power of 400 W, 600 W and 800 W, respectively. The relation between the dendrite size and laser power shown in Fig. 3d indicates that the dendrites in the specimens fabricated at the lower laser power are smaller, but the size tends to be non-uniform with the decrease of the laser power.

Fig. 2

a Dendritic morphology and b amplification of an as-fabricated (P = 600 W) specimen in XY planes by SEM observation; TEM images of the c Laves phase and d carbides

Table 2 Element concentration of different test points (mass fraction %)
Fig. 3

SEM images of dendrites in XY planes fabricated at laser powers of a 400 W, b 600 W and c 800 W; d relation between the dendrite size and fabricating laser power

The porosity of the specimens fabricated at different laser powers was quantified from OM images according to the proportion of the pore area accounted for the total area of the specimens. Figure 4 shows that the porosity remarkably decreases first and slightly increases subsequently with the increase of the laser power. The as-fabricated specimens (P = 600 W) achieved the minimum porosity of 0.132%. This might be attributed to the insufficient melting of powders and incomplete bonding between the deposited layers at the laser powers less than 600 W which cannot provide enough energy input to penetrate into the powder. At the same time, the shielding gas was not able to obtain enough time to escape from the molten pool at the low laser powers. The size and distribution of pores in the specimens varied with the laser power. The pore sizes of the as-fabricated specimens are mainly concentrated in the range of 0–15 µm, but notably, the largest pore diameter is 935 µm in the P = 300 W specimens, presenting a wide distribution of pore sizes. Comparing with others, 94.5% of the pores in the P = 600 W specimens are less than 15 µm in diameter. The distribution and morphology of pores in the P = 600 W specimen were further checked using 3D-XRT (Fig. 4b). One can find that the spherical pores extensively distributed in the whole specimen. The pore sizes mainly range from 3 to 30 µm in diameter, and the maximum diameter of pores is 76.6 µm. The volume of the pores is ~ 0.28% of the total volume of the as-fabricated specimens at P = 600 W, which is much higher than that of the specimen fabricated by blown powder direct layer deposition (0.095%) [18]. The porosity of P = 600 W specimens determined by 3D-XRT is much higher than that determined by OM, which might be related to the fact that 3D-XRT can detect much smaller pores close to ~ 3 μm, but hardly discriminated by OM.

Fig. 4

a Porosity of as-fabricated specimens as a function of the laser power; b pore morphology and distribution of an as-fabricated GH4169 specimen deposited at the laser power of 600 W detected by 3D XRT

Tensile properties

Figure 5a presents tensile stress–strain curves of the as-fabricated GH4169 specimens loading along SD. At least three specimens were tested for each sort of specimens. It was noted that tensile properties of the P = 300 W specimens were not able to be obtained due to the high porosity of the specimens. The P = 800 W specimens exhibit the maximum elongation at fracture of 14.3%. Figure 5b reveals the strength variation of the specimens with the laser power. Generally, the specimens fabricated at 600 W exhibit the highest strength (σys = 587 MPa, σuts = 903 MPa) among the five sorts of specimens and elongation at fracture (εf = 13.6%). The appropriate laser power generates fine dendrites with the relatively uniform size and few pores during the LMD process. It is demonstrated that the P = 400 W and 500 W specimens show lower strength and elongation than that of the P = 600 W specimen. One reasonable explanation is that with the laser power decreasing from 600 to 400 W, the amount of unmelted powder and shielding gas residue increases, resulting in larger pores and higher porosity and finally lower mechanical properties of specimens. What is more, for the yield strength of the specimen, it is not only related to the porosity, but the dendrite size and the uniformity as well. Based on the Hall–Petch relation, an appropriate laser power, such as 600 W, generated refine dendrites with a relatively uniform size, and also few pores enable the specimen to possess high yield strength [2]. However, compared with those of the P = 600 W specimens, the lower strength and higher fracture elongation of the P = 700 W and 800 W specimens may result from the increase of the dendrite size and uniformity (Fig. 3d) due to the higher laser energy input. Notably, from the observation of Fig. 3a–c, the Laves phases are continuous in the 400 W specimen, but become fragmentary in the 800 W one. It is known that the Laves phase and carbides normally form in the inter-dendritic liquid with the concentration of elements C, Al, Ti, Nb and Mo at the last stage, after the formation of the primary γ phase. With the increase of the laser power, the energy density input into the specimens increased, and the heat accumulation in the specimens increased accordingly. Thus, the segregation elements mentioned above enable to be dissolved more in the primary γ phase. As a result, the inter-dendritic Laves phases grow immaturely and become fragmentary.

Fig. 5

a Tensile stress–strain curves of as-fabricated specimens tested in perpendicular to the building direction, showing the different tensile properties of specimens deposited at different laser powers; b effect of the laser power on strength of specimens

Figure 6 shows SEM observations of a tensile fracture surface of a P = 600 W specimen. There existed a large pore and a small one close to the specimen surface, as shown in Fig. 6a. Figure 6b shows that deep and fine dimples distributed in the dendrites, and plenty of bright particles are identified as Laves phases near the dimples. Figure 6c shows a macro-fracture path of the specimen without evident necking. A close observation of the specimen surface reveals that the distinct slip traces appeared, and most of small cracks initiated around Laves phases, as indicated by arrows shown in Fig. 6d, implying the stress concentration between the Laves phase and matrix.

Fig. 6

SEM images of tensile fracture surfaces of an as-fabricated P = 600 W GH4169 specimen: a fracture surface; b amplification observation in a; c surface observation of the tensile fractured specimen; d amplification observation of the specimen surface in c, arrows indicating the elongation of the micro-pore

Fatigue properties

The relation between applied stress amplitude (σa) and fatigue life (Nf) of the P = 600 W specimens at room temperature was obtained, and shown in Fig. 7. The fatigue limit of the specimens at R = 0.1 is 173.7 MPa, and 227.1 MPa at R = −1 converted by the Goodman equation. A ratio of the fatigue limit to the tensile strength (σ−1/σUTS) is 0.25, which is lower than that of the wrought IN 718 specimens (σ−1/σUTS = 0.36) [19]. Compared with specimens fabricated using other conditions such as forging [19,20,21], direct laser sintering [22] and laser AM [21], the fatigue limit of the present as-fabricated P = 600 W specimens is quite low. This may result from the existence of micro-pores and the brittle Laves phases.

Fig. 7

Relationship between fatigue life of GH4169 specimens fabricated at 600 W and applied stress amplitude (R = 0.1)

The fracture surface of an as-fabricated P = 600 W specimen after fatigue failure was observed by SEM. Figure 8a shows that fatigue cracks initiated from pores. Some fine striations appeared in the fatigue crack propagation zone, as shown in Fig. 8b. Figure 8c displays a close observation on the striation region in Fig. 8b. The further observation of the specimen surface reveals that microcracks formed either within/around the Laves phase or along the slip bands (Fig. 8d). Combining with the morphology of the fatigued specimen surface, the crack propagated along hard and brittle Laves phases in the inter-dendritic regions. During fatigue loading, the stress concentration was formed due to the uncoordinated mechanical properties of Laves phases and the matrix phase, thereby causing the brittle Laves phases breaking or separating from the matrix. As a result, micro-pores or cracks were formed [21].

Fig. 8

SEM images of fracture surfaces of as-deposited GH4169 specimens at P = 600 W at different stress amplitudes (σa) at room temperature: aσa = 175.5 MPa; bσa = 202.5 MPa; c amplification observation of fatigue striations; d cracking behavior in b, arrows indicating fragmentation of the Laves phase

The above results clearly reveal a close relationship between the porosity caused by different laser powers and mechanical properties, such as the tensile strength and ductility as well as fatigue reliability. The appropriate laser power (here P = 600 W) can generate low porosity in the as-fabricated specimens, leading to comparative high mechanical properties. Although the porosity and mechanical properties of the present GH4169 fabricated through the LMD process is not well satisfied, and even not comparable with that fabricated using conventional methods, such as castings and forgings, the trend of the parameter-controlled microstructure and porosity is still intuitive. A good performance of the mechanical properties, especially fatigue reliability is still needed to be enhanced through optimizing the LMD process parameters.


The relationship between the porosity caused by laser manufacturing powers and mechanical properties of the LMD-fabricated GH4169 was investigated. The following conclusions can be drawn.

  1. 1.

    Porosity of the LMD-fabricated specimens remarkably decreases first and slightly increases subsequently with increasing the laser power from 300 to 800 W.

  2. 2.

    The GH4169 specimens fabricated at the laser power of 600 W possess the minimum porosity of 0.284%, exhibiting the comparatively excellent mechanical properties. For example, the yield strength is 587 MPa, the ultimate tensile strength is 903 MPa and elongation at fracture is 13.6%.

  3. 3.

    The fatigue limit of the GH4169 specimens fabricated by LMD at P = 600 W is 227.1 MPa. Tensile and fatigue properties are mainly affected by pores, the Laves phase and dendrite size.


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This work was financially supported by the National Key R&D Program of China (Grant No. 2017YFB0305800), the National Natural Science Foundation of China (NSFC, Grant No. 51771207) and the Joint Founds of NSFC Liaoning (Grant No. U1508213).

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Correspondence to B. Zhang.

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Shao, W.W., Zhang, B., Liu, Y. et al. Effect of laser power on porosity and mechanical properties of GH4169 fabricated by laser melting deposition. Tungsten 1, 297–305 (2019). https://doi.org/10.1007/s42864-020-00034-w

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  • Laser melting deposition
  • GH4169
  • Microstructure
  • Porosity
  • Mechanical properties