Laser melting deposition of fine-grained 90W-7Ni-3Fe alloys using pre-sintered granulated powder
Release time:
2025-06-17
1. Introduction
Tungsten heavy alloy (WHA) is a two-phase composite material consisting of W reinforcement particles in a Ni-Fe-based or Ni-Cu-based solid solution matrix [1,2]. Due to their characteristics of high strength, high hardness, and high ductility [3,4], etc., WHAs are widely used in the fields of national defense, nuclear industry, and aerospace, for which the demanding of WHAs with higher strength is an eternal pursuit. One of the main methods to increase the strength of WHAs is to reduce the W grain sizes to promote the Hall-Petch strengthening. Typically, WHAs are usually prepared by traditional powder metallurgy liquid phase sintering (LPS) method. Due to the long duration at high sintering temperatures and the low cooling rate of LPS, obvious grain growth of W grains usually takes place, leading to final W grain sizes larger than 15 μm. In order to mitigate this growth, several new sintering techniques have been explored, such as microwave sintering [5–8], and spark plasma sintering [9–12], both aimed at lowering sintering temperatures. More recently, a laser melting deposition (LMD) additive manufacturing (AM) method, known for its rapid cooling rate, has attracted extensive attention to the preparations of WHAs [13–22].
Due to the marked difference in the melting point of the W phase and the Ni-based matrix, it is challenging to obtain WHA powders using common atomization techniques, such as gas atomization and centrifugal atomization. In previous research, all the reported studies have simply used mixed W, Ni, and Fe elemental powders as raw materials. Using these mixed powders, WHAs with a W content of up to 90 wt% and a sample density larger than 99% have been successfully prepared by LMD [13]. The fabricated 90WNiFe alloys exhibit a microstructure that is similar to those prepared by LPS, characterized by W grains embedded in a Ni-based matrix. Encouraging results have been reported on the positive effects of the high cooling rate of LMD in hindering grain growth, with observed W grain sizes mirroring or smaller than those of the raw W powder particles [14] [16]. Intuitively, utilizing finer W powder particles as the raw powder would seem ideal to refine W grains. However, when the W content is as high as 90 wt%, the fine W powder particles usually lead to poor flowability, which hinders continuous deposition. The finest W powder particle size reported for LMD application is about 7 μm, resulting in a final W grain size of about 8 μm [16]. Thus, the pursuit of methods to secure even finer grains using LMD remains an attractive avenue.
It has been shown that different elemental powder particles can be adhered together by a spray granulation method. A single spray granulated powder particle contains numerous fine elemental powder particles, maintaining an average composition similar to the designated level [23]. Typically, these spray granulated powder particles are spherical, with sizes up to 100 μm. They can be sintered at a relatively low temperature to remove the resin, resulting in what is referred to as “pre-sintered powder” subsequently. The substantial large size of the pre-sintered (mixed) powder particles guarantees a good flowability for LMD, while also supplying fine powder particles of the desired element.
The aim of this study is to investigate the feasibility of utilizing a presintered powder, comprising fine W powder particles with sizes of 1–3 μm, to manufacture fine-grained 0W-7Ni-3Fe WHAs by LMD. In an initial attempt, thin plate samples were prepared with varying laser powers. The microstructural evolutions and mechanical properties of the resulting WHAs were characterized to understand the processing microstructure-property relationship, thereby offering guidelines for manufacturing WHAs with enhanced properties.
2. Experimental
2.1. Powder
The pre-sintered 90W-7Ni-3Fe powder was provided by Chengdu Macro-Micro New Materials Co., Ltd., China. Figs. 1(a-c) show the morphologies, element distributions, and particle size distribution of the pre-sintered powder, respectively. The pre-sintered particles had spherical-like hollow shapes, with an average size of 74 μm. Each pre sintered particle is comprised of fine W, Ni, and Fe particles with sizes of 1–3 μm. Before the LMD processing, the pre-sintered powder was subjected to drying in a vacuum oven at 120 ◦C for 3 h.
2.2. LMD processing
An in-house developed LMD system equipped with an ytterbiumdoped fiber laser with a wavelength of 1070 nm and a spot diameter of approx. 2 mm was used to deposit the pre-sintered powder on a 316L steel substrate. To prevent oxidation, the deposition was carried out in a chamber filled with constantly flowing high purity argon. The oxygen content in the chamber was maintained below 20 ppm. The powder was fed into the chamber through a coaxial nozzle at a feeding rate of 4–5 g/ min. Thin plate samples with width of 55 mm, height of 40 mm, and thickness of 3.8–4.9 mm were prepared layer-by-layer using laser powers of 500 W, 600 W, 700 W, 800 W and 900 W. These samples are denoted as samples S1-S5 as the laser power increases from 500 W to 900 W, respectively. All the samples were prepared at a consistent laser scanning speed of 150 mm/min, and an identical height increment of 0.15 mm.
2.3. Characterization of microstructure and mechanical properties
The microstructure of the LMD samples was characterized using optical microscopy (OM) and scanning electron microscopy (SEM) on the cross sections perpendicular to the laser scanning direction (SD). Before observation, the samples were cut using spark erosion, then ground with a non-crystallizing colloidal silica suspension, and finally subjected to vibration-assisted polishing.
The mechanical properties were assessed using micro-hardness tests. The micro-hardness tests were carried out using an HVS-1000A Vickers hardness instrument with a diamond indenter under a load of 1 kg and a dwell time of 15 s. The indentation tests were conducted at 80 random locations away from large pores for each sample.
3. Results and discussion
3.1. Microstructure
The fabricated samples have rather rough surfaces, due to the adhesion of powder particles on the surfaces. The microstructures on the cross sections perpendicular to the SDs for all samples are shown in Fig. 2. Large amounts of pores with varied shapes and sizes are observed. These pores can be divided into two categories: smaller ones with size in the range of tens of micrometers and larger ones spanning several hundreds of micrometers. The smaller pores exhibit near spherical shapes, and are distributed nearly randomly across the entire cross sections of all samples. These pores are likely caused by the trapped gas in the pre-sintered powder. In contrast, the larger pores feature irregular shapes and are typically regarded as lack-of-fusion defects. For samples S1-S5, the larger pores are primarily found near the surface regions, where the lower molten pool temperatures lead to a high probability of lack-of-fusion defects. In sample S5, larger pores are also observed in the center regions, likely a result of intense evaporation of elements due to high input laser power. The thinner sample thickness and fewer amounts of smaller pores in sample S5 suggest a higher molten pool temperature, further substantiating the intense evaporation phenomena.
The fabricated LMD samples exhibit two-phase composite structures, similar to those in traditional LPS samples, with near spherical W grains embedded in the matrix, as illustrated in Figs. 2(a3-e3). In the focused areas away from the obvious pores seen in the OM images, the structures look rather uniform and dense. Only a few pores with sizes of several micrometers can be occasionally observed in the matrix and inside W grains. This indicates that if the residual gas can be eliminated, it is possible to fabricate dense, fine-grained WHAs using optimized moderate laser powers avoiding intense evaporation.
The volume fractions of W grains of all samples were determined from the SEM images, and Fig. 3 shows the results. As the input laser power increases from 500 W to 900 W, the volume fraction of W grains increases from 83 vol% to 95 vol%. For 90WNiFe, 93WNiFe and 95WNiFe alloys, the volume fractions of W grains typically fall within the ranges of 78–83 vol% [14,24,25], 83–86 vol% [24–27], and 90–96 vol% [25,28,29], respectively. As shown in Fig. 3, only when the input laser powers are below 700 W, do the volume fractions of W grains descend below 83 vol%, which corresponds with the specifications for the 90WNiFe alloy. At a laser power of 700 W or higher, the volume fraction of W grains increases to 89 vol% and above, matching the 95WNiFe alloy parameters. Thus, to suppress the intense evaporation of Ni and Fe elements, the input laser power should remain below 700 W.
Figure 4 shows the variation of the average W grain size, determined based on the SEM images in Fig. 3, with respect to the applied laser power. For laser power below 700 W, the LMD samples (S1 and S2) exhibit an average W grain sizes of approximately 5 μm. As the laser power exceeds 700 W, this size begins to grow. At 900 W, the average W grain size reaches 17 μm, which is comparable to that of traditional LPS samples.
It is noted that the finest W grain sizes obtained is slightly larger than the initial W powder particles, suggesting W grain growth at the present process parameter and the corresponding cooling rate. Nevertheless, the W grain sizes of samples S1 and S2 are significantly smaller than those of the traditional LPS samples and represent the finest grain size ever reported for the laser additive manufactured WHAs [16]. The increase of W grain sizes with the increase of laser power is likely caused by the corresponding increase in molten pool temperatures, leading to comparatively slower cooling rates and thus more precipitation of W atoms from the matrix during cooling. Moreover, the intense evaporation of Ni and Fe elements at high input laser powers leads to a decreased volume fraction of the matrix and thus shorter diffusion distances between W grains, which may also contribute to the growth of W grains.
Figure 5 shows the variation of micro-hardness with respect to the input laser power. As the laser power increases from 500 W to 900 W, the micro-hardness increases from 357 HV to 406 HV. Note that the hardness measurements were carried out in regions away from the large pores, and the quantity of micrometer-sized pores is similar across all samples (see Fig. 2). Therefore, micro-porosity is unlikely to be an important factor influencing the observed increase in micro-hardness.
It is noticed that even with the presence of micro-porosities, the micro-hardness of the samples prepared in this study is much higher than that of the traditionally prepared LPS 90W-7Ni-3Fe alloy samples with a W grain size of 16 μm [14], the micro-hardness of which is given as a dashed line in Fig. 5. For the samples S1 and S2 with nominal compositions of 90WNiFe, the micro-hardness increase is presumed to primarily stem from the presence of fine W grains. For samples S3-S5, the higher volume fractions of W grains likely enhance the observed increase in micro-hardness, considering the fact that the W grain size of S5 is comparable to the reference LPS sample. Indeed, higher hardness of 397–420 HV has been reported for 93W-5.6Ni-1.4Fe alloy prepared using spark plasma sintering with a W grain size similar to the samples S3-S5 [30].
However, the porosities do affect the micro-hardness. Compared to the 90W-7Ni-3Fe alloys prepared by LMD using mixed elemental powder [14], of which the micro-hardness values fall in the green box in Fig. 5, most samples in this study have lower micro-hardness, despite their finer W grains. Only the sample with a 95% volume fraction of W grains (S5) achieves micro-hardness comparable to the LMD 90W-7Ni- 3Fe samples prepared using mixed elemental powder. 
3.3. LMD processing using pre-sintered powder
To ensure optimal flowability and continuous deposition, LMD usually requires particle sizes larger than 45 μm. Commonly, powders are pre-alloyed, ensuring uniform composition in each particle analogous to the final product. However, manufacturing pre-alloyed powders for WHAs poses challenges, leading contemporary research to primarily employ mixed elemental powders. Since the densities of W, Ni, and Fe are significantly different, separation or delamination of particles often occurs during mixing and delivery of these mixed elemental powders, which can lead to inhomogeneous microstructure in the LMD processed WHAs. Utilizing pre-sintered powder can mitigate this issue, resulting in a more homogeneous microstructure.
Compared to the pre-alloyed and mixed powder, the individual elemental powders within the pre-sintered powder are much finer. This implies swifter melting of these particles due to their increased surface fractions, leading to higher laser absorption rates. Such efficient melting on one hand reduces the required input energy. Conversely, it can intensify the adhesion of powder particles on the surfaces of the fabricated alloys, which deteriorates the surface quality. It is worth mentioning that the hollow nature of the pre-sintered powder particles means that significant amounts of gas are introduced into the molten pool, which leads to the high porosity as shown in Fig. 2. The porosity cannot be reduced efficiently by just adjusting the process parameters. Further modification of the powder particles is needed to densify the pre-sintered powder. Despite the evident porosity, this work demonstrates the possibility of fabricating fine-grained WHAs by LMD using pre-sintered powder. With suitable powder processing [31], as well as optimized deposition parameters to minimize porosity, the production of WHAs with both high strength and high ductility becomes a foreseeable prospect.
4. Conclusions
In this study, by using a unique pre-sintered 90W-7Ni-3Fe powder, fine grained WHAs have been successfully prepared by the LMD technique. The microstructure and mechanical properties of these samples have been investigated, and the main conclusions are as follows: (1) WHAs with a W grain size as fine as 5 μm can be obtained at laser power below 600 W. Beyond 700 W, the W grain size starts to increase with the laser power. At a laser power of 900 W, the W grain size reaches to 17 μm, which is comparable to traditional
LPS samples.
(2) The volume fraction of W grains increases with increasing laser power, attributed to the evaporation of Ni and Fe elements. At laser power higher than 700 W, the volume fraction of W grains of the LMD samples falls into the range observed in 95WNiFe alloys.
(3) All the samples fabricated at different laser powers have high porosities, which is predominately a result of the hollow nature of the pre-sintered powder, introducing high amounts of trapped gas into the molten pool.
(4) All the LMD samples display higher micro-hardness than traditional LPS 90W-7Ni-3Fe alloy samples, a result of either finer grains or an increased volume fractions of W grains. However, due to the presence of porosities, their micro-hardness is lower than the 90W-7Ni-3Fe alloy samples prepared using LMD withmixed elemental powder.
Paper citation information:
International Journal of Refractory Metals and Hard Materials 119 (2024) 106507
Tungsten-nickel-iron alloy (W-Ni-Fe)
Tungsten-nickel-iron alloy is a tungsten-based heavy alloy, which belongs to powder metallurgy materials and is made of three metals: tungsten (W), nickel (Ni), and iron (Fe) through a specific process. It combines the high density and high strength of tungsten with the toughness of the nickel-iron bonding phase, and is an important structural material in the industrial field.
The spherical W-Ni-Fe alloy powder produced by Stardust Technology is prepared by radio frequency plasma spheroidization process, which has the characteristics of high purity and low oxygen, high sphericity, smooth surface, no satellite, uniform particle size distribution, excellent flowability, high bulk density and high tap density.
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