A large temperature gradient led to rapid temperature changes in the selective laser melting (SLM) process, resulting in high residual stress and micro-cracks, which were the main challenges faced by the 3D printing of tungsten (W) materials. On the other hand, the volumetric energy density had a significant effect on the density of the pure W structural component fabricated by SLM. It was difficult to eliminate defects in the formed parts solely by optimizing SLM technology parameters. The crystalline W powder with a concentrated grain size was used as the raw material, and substrate preheating was adopted to prepare the W structural component with a lattice structure by selective laser melting (SLM). The processed samples were subsequently subjected to heat treatment at 1100℃, 1400℃, 1700℃and 1900℃ for 2 hours. The mechanical properties and microstructure of the specimens at different heat treatment temperatures were analyzed. The results indicate that the optimal volumetric energy density for obtaining a theoretical density of 96.2% is around 972 J/mm3. After the heat treatment at 1700℃, the samples achieved the best mechanical performance, with an ultimate compressive strength and microhardness reaching 1200 MPa and 457HV, respectively. The fracture morphology shows cleavage fracture characteristics, and microcracks were largely eliminated after the heat treatment at 1700℃, indicating strong metallurgical bonding between layers and particles.
Research Article
The effects of energy density and heat treatment on the properties of 3D printed tungsten
https://doi.org/10.21203/rs.3.rs-3776346/v1
This work is licensed under a CC BY 4.0 License
Journal Publication
published 14 Jun, 2024
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Tungsten
3D printing
Heat treatment
Energy density
Tungsten has multiple unique properties, including a high melting temperature (3420℃), superior thermal conductivity (174 W/(m⋅K) at room temperature), good mechanical performance, low thermal expansion (4 µm⋅(m⋅K)−1 at room temperature), and high chemical stability, which make it effectively applied in various engineering fields. For example, plasma-oriented structures and high-performance rocket nozzles in the future nuclear industry, military, and medical fields. On the other hand, the aforementioned properties of tungsten increase the difficulty of forming complex tungsten structural components by traditional processing methods. Therefore, tungsten parts are usually manufactured using methods such as chemical vapor deposition and powder metallurgy, including spark plasma sintering, etc. [1]. Therefore, compared with the traditional powder metallurgy for W materials, current additive manufacturing (AM) techniques show a more significant potential to achieve the goal of fabricating W parts in a flexible and integrated way, since these processing routes require no pre-set molds [2] .
Selective laser melting (SLM) is a classification of additive manufacturing. The technology uses a laser beam as an energy source and directly forms parts with complex geometric shapes layer by layer according to a sliced model. During the forming process, the metal powder is melted directly by the high-energy laser beam to achieve metallurgical bonding [3]. SLM technology has been studied in the preparation of metals such as stainless steel, aluminum alloys, and nickel-based alloys [4–9], but there are still many challenges for forming refractory metals.
Materials suitable for SLM manufacturing typically exhibit good laser absorption and high thermal conductivity, as well as good surface tension and viscosity. Refractory metal tungsten has been proven to be processable by SLM. During SLM, a high-energy laser beam can generate an extremely high energy density on the powder layer surface. Even tungsten, with a melting point of 3420℃, can be completely melted. However, selective laser melting of pure tungsten involves many of the challenges associated with metal additive manufacturing. The element’s ultra-high melting point results in high cohesive energy, and its high surface tension promotes the formation of spheroidization. The excellent thermal conductivity leads to rapid solidification and cooling, resulting in cracks formed within the part. The mechanical properties of samples prepared by SLM can be affected by heating conditions, including rapid cooling, which may lead to non-equilibrium states, residual impurities, the generation of residual stresses, or the formation of amorphous regions. Unmelted powders may have existed in the manufacturing process when the powder was not properly wetted to the underlying metal due to the influence of the internal oxygen content. The porosity is also generated through powder spattering and metal evaporation [10, 11]. These transitions may have a negative impact on the mechanical properties of the final manufactured part. Therefore, research on pure tungsten additive manufacturing mainly focuses on defects such as thermal stress cracking, oxygen cracking, high residual stress, and droplet spheroidization. Xiong Z et al. [12] obtained pure tungsten samples with an ultimate compressive strength of up to 1200 MPa through heat treatment using selective laser remelting technology. The laser energy density applied by Wang et al. [13] had a significant impact on liquid splashing during the SLM process. An increase in energy density led to an increase in splashing intensity, which in turn led to an increase in splashing and dispersion around the melt. R. K. Enneti et al. [14] discovered that the densification of tungsten treated with SLM rose as the energy density grew. Nevertheless, the topmost compactness acquired in this study amounts to solely 75% of the anticipated compactness. Therefore, the preparation of pure tungsten parts by selective laser melting technology alone still has defects, such as a large grain size and high residual stress. Guo M et al. [15] systematically investigated the impact of volumetric energy density on the densification, microstructure, and mechanical properties of pure tungsten manufactured by SLM. At a volumetric energy density of 1000 J/mm3, they achieved a maximum relative density of 98.4% and an ultimate compressive strength of up to 902 MPa. Ren LYX et al. [16] optimized processing parameters to allow for the production of bulk pure tungsten samples with a high relative density of up to 98.51%. Zheng T et al. [17] found that the introduction of ZrC nanoparticles as a secondary phase into the tungsten matrix offers a promising strategy for additive manufacturing of high-quality tungsten. Typically, reducing cracks requires the addition of secondary phase particles. This article focuses on improving the performance of pure tungsten fabricated through SLM by employing heat treatment.
As the focus of the above research report is on the process parameters of SLM, in our work, we study in detail the effects of heat treatment on the densification behavior, microstructure, and mechanical properties at the optimum volumetric energy density. In this paper, the purpose is to characterize the SLM manufacturing technology from the density and mechanical properties of manufactured and heat-treated samples. The microstructure and mechanical properties of printed tungsten material after heat treatment were analyzed and discussed through SEM and quasi-static uniaxial compression testing. The changes in density and microstructure of the produced samples before and after heat treatment were detected by SEM and metallographic analysis, which provided an effective reference for improving the structural performance of pure tungsten processed by SLM.
Process and sample preparation
The raw material selected was 325-mesh polygonal crystalline W powder, with its morphology shown in Fig. 1. The oxygen content of the W powder reached 300 ppm, the purity was above 99.5%, and the content of other impurities was less than 0.001%. The powder particle size distribution was measured using an Mastersizer 2000 laser particle size analyzer. Figure 1(b) depicts the particle size distribution of pure W powder. It is noted that the pure W powder has a size distribution as D10 = 35.61 µm, D50 = 49.11 µm and D90 = 67.54 µm respectively. Data about the starting physical properties of the powder are included in Table 1. The crystalline W powder exhibited a uniform particle size distribution and intact crystal structure, with no obvious aggregation. The flowability of both the crystalline W powder and common W powder was tested using a powder comprehensive characteristic tester, SYRJX93-DDBT series, and the results are listed in Table 2. The angle of repose has the greatest effect on the fluidity of the powder, and the smaller the angle of repose, the better the powder’s fluidity. The crystalline W powder showed higher flowability.
|
Composition |
W |
O |
C |
CO2 |
Fe |
|---|---|---|---|---|---|
|
Proportion(%) |
≥ 99.5 |
≤ 0.030 |
0.0008 |
0.0075 |
0.0075 |
|
Crystalline W powder |
Common W powder |
|
|---|---|---|
|
Grain size/(um) Angle of repose/(°) |
15 ~ 45 26.40 |
2 ~ 3 45.50 |
|
Angle of rupture/(°) |
21.10 |
32.20 |
|
Differential angle/(°) |
5.30 |
3.30 |
|
Flat angle/(°) |
26.05 |
53.75 |
|
Bulk density/(g/cm3) |
8.9 |
9.7 |
|
Liquidity |
Better |
Ordinary |
In this work, the SLM 3D printer used was the WXL-120T model. The equipment consisted of a laser transmission and control unit, a water cooling circulation system, and a forming control system. It was equipped with a 500W fiber laser in single beam mode, with a beam diameter of 50um. Based on the lightweight and high-strength characteristics, we chose the curving lattice struts, as studied by Long Bai et al. [18], as the printing target. Before printing, the Unigraphics NX modeling software was used to construct a lattice structure model with dimensions of 6×6×6 mm (width x length x height), as shown in Fig. 2a and Fig. 2b. The actual printed model is shown in Fig. 2c. Figure 2d shows a 67° scan policy pattern diagram. Prior to printing, the substrate was preheated to 200℃, with argon as the protective gas and a maximum oxygen limit in the chamber of 100 ppm. The high-energy laser was the main energy source for melting the metal powder, and the scanning power, powder layer thickness, and scanning spacing were important parameters for controlling the quality of the forming process. The volume energy density was defined as [19]:
In the equation, E represents the volumetric energy density, P represents the laser power, v represents the scanning speed, h represents the scanning distance, and d represents the thickness of the powder layer. Appropriate input of laser energy density can effectively improve the densification and mechanical properties of pure W samples. Samples were processed by controlling the optimal volumetric energy density through changing laser power, scanning speed, and scanning spacing.
Heat treatment and mechanical properties test
The microstructure and mechanical properties of the components are related to the thermal behavior generated by any process. In the SLM process, the local rapid scanning of the high-energy laser beam leads to quick melting and solidification of the powder, which complicates and makes it difficult to control the SLM process. Liu et al. [20] showed that the thermal cycle in selective laser melting (SLM) is a unique thermal behavior in additive manufacturing processes, which affects the microstructure and properties of the samples. The mechanical properties of materials can be improved through heat treatment. In this work, the heat treatment conditions were divided into three groups, and the samples were held at each temperature for 2 hours. All three groups of experiments were heated in an argon atmosphere with the aim of improving the mechanical properties of the pure W components. After the heat treatment process, the samples were characterized by polishing with different sizes of sandpaper according to standard metallographic procedures. Then, a mixed solution of 10% potassium hydroxide and potassium ferrocyanide was used to etch the samples for 15 seconds before observing the surface morphology and microstructure of the samples using a scanning electron microscope (SEM, Sigma 500). The compressive strength of the samples was tested using a Tensile Testing Machine, and the compression test parameters for the crystal structure were referenced from the ISO 13314:2011 standard [21]. Before the experiment, sandpaper was used to polish the upper and lower surfaces of the sample to ensure that the sample surface and the pressure plate surface were parallel. During testing, the sample was placed between two circular hardened steel blocks, and the lubricated contact surface was used to reduce the friction coefficient. The sample was loaded with a Z-direction loading rate of 1mm/min at 25℃. When the load plummeted sharply, the pressure was stopped, and the sample was removed after slowly unloading it. The grain structure was observed using an optical microscope. The sample hardness was measured using a Vickers hardness tester.
Effects of SLM processing parameters on microstructure on printed W parts
In the SLM process, the volume energy density has an important influence on the densification behavior of the sample. As shown in Fig. 3, using different laser energy densities during the processing of pure W parts can affect the density and relative density of the sample. Changing the process parameters, as indicated in Table 3, can lead to a change in energy density. When the laser energy density is lower than 972 J/mm3, increasing the energy density can promote the densification process, thereby enhancing the density and relative density of SLM processed W parts. However, when the laser energy density is higher than 972 J/mm3, the density and relative density will decrease as the laser energy density increases. This indicates that there exists an optimal laser energy density (972 J/mm3) in the SLM processing of pure W. By using the optimal energy density, a pure W sample with a relative density of up to 96.2% can be obtained.
| Laser power (W) | Laser speed (mm·s-1) | Layer thickness (um) | Hatch space (um) | Relative Density(%) | Energy density(J/mm3) |
|---|---|---|---|---|---|
| 200 | 1000 | 30 | 30 | 77.4 | 278 |
| 250 | 800 | 30 | 30 | 80.1 | 347 |
| 300 300 | 600 500 | 30 30 | 40 30 | 88.7 91.0 | 417 667 |
| 350 350 | 400 300 | 30 30 | 30 30 | 96.2 92.9 | 972 1296 |
| 350 | 200 | 30 | 40 | 92.5 | 1458 |
The surface morphology of pure W specimens processed by SLM under different volume energy densities is shown in Fig. 4. It is apparent that, with the use of different volume energy densities, the specimen’s surface morphology presents different features under SEM scanning, but the commonality is that the surface is coated with a large amount of powder particles due to the significant temperature difference between the unmelted and melted powders, causing thermal diffusion between the materials and leading to powder bonding on the surface and pores of the support structure [22]. Different sizes of metal particles can also be observed in the figure, which are due to the adhesion of unmelted or partially melted metal particles to the substrate surface. These particles are distributed on the relatively smooth pillar substrate surface due to the phenomenon of remelting and seriously affect the stability and size of the melt pool during the manufacturing process.
When the volume energy densities are 278 J/mm3 and 347 J/mm3, there are more irregular pores and large cracks on the sample’s surface (Fig. 4a and b), and the obtained relative densities are only 77.4% and 80.1%, respectively. With the increase in volume energy density to 417 J/mm3, the number of large cracks decreases (Fig. 4c), but some regularly shaped pores can still be observed, yielding a density of 17.16 g/cm3, which is 88.7% of the theoretical density. The laser power had a significant impact on the porosity and mechanical properties, so we adjusted the variation of laser power within a small range (200W-350W) [23].The scanning speed and hatch space had been changed, which had enabled an increase in energy density to 667 J/mm3. The surface morphology was significantly different (Fig. 4d). The number of pores decreased. The resulting density was 17.61 g/cm3, which is 91.0% of the theoretical density. As the volume energy density continues to increase to 972 J/mm3, the surface pores of the sample almost disappear (Fig. 4e), and the measured relative density is 96.2%, with a density of 18.61 g/cm3, indicating the successful production of high-density W parts via SLM. When the energy density increased to 1296 J/mm3, the obtained density decreased, and small pores appeared on the surface (Fig. 4f). The obtained density was 17.98 g/cm3, which is 92.9% of the theoretical density. However, excessive heat input could reduce forming quality, causing an increase in either porosity or pore size. When the applied energy density reaches 1458 J/mm3, the density of the W sample decreases mainly due to the presence of voids (Fig. 4g), and the forming effect is similar to that at an energy density of 1296 J/mm3, resulting in a density of 17.89 g/cm3, which is 92.5% of the theoretical density.
The surface roughness of pure W was measured using a MITUTOYO surface roughness measuring instrument. Figure 5 shows the variation of surface roughness of pure W parts with the applied energy density. From the figure, it can be observed that the energy density significantly affects the wear performance of W components fabricated by SLM. When the applied energy density is less than 600 J/mm3, the average surface roughness is Ra 18 um. In this case, the applied energy density is insufficient to fuse the powder particles and form a dense component. This phenomenon is due to the presence of large pores in the samples at low densities. When the energy density is above 600 J/mm3 and below 1200 J/mm3, the average surface roughness obtained is the lowest among all the samples, at Ra 7.6 um. Further increasing the energy density above 1200 J/mm3 results in an average surface roughness of Ra 9.5 um, showing an increasing trend. This is similar to the variation in hardness.
In the SLM manufacturing of pure W, W is prone to cracking, which is also the main difficulty in achieving additive manufacturing. In order to fully understand the influence of process parameters on pure W, we further studied the microstructure of W parts.
The main defects of W samples are unmelted and gas pores caused by gas entrapment during the manufacturing process, which is also the reason for the formation of cracks and residual porosity defects. Therefore, controlling the indoor oxygen content to the lowest level is the primary factor in forming high-quality W materials. Thermal shrinkage is the main cause of microcrack formation during SLM metal processing [24]. Reducing the scanning speed or increasing the laser power can increase the thermal shrinkage of the part. Therefore, thermal shrinkage tends to occur in SLM components that use high energy density, causing thermal stress. For common metal parts processed by SLM, thermal tensile stress exists at the top and bottom of the SLM parts. Due to thermal tensile stress, microcracks appear on the surface, as shown in Fig. 6. Another reason for microcracks is the high ductile-to-brittle transition temperature (DBTT) of W, typically between 200°C and 400°C [25]. When the brittle temperature range is encountered, the sensitivity of cracks increases, making it easy to produce cracks. During the manufacturing process, a high scanning speed will result in a short laser exposure time, causing an unstable melt pool. This unstable melt pool leads to the cracks and ruptures shown in the figure. The melted droplets rapidly spread and solidify on the printing substrate. Under extreme temperature gradients, when the bottom thermal stress exceeds the yield stress of W, cracks occur in the W samples. Therefore, at the appropriate volumetric energy density, the molten pool is in a relatively stable state. At low volumetric energy densities, the temperature inside the molten pool is lower, and the surface tension of the melt increases, resulting in higher viscosity and poor fluidity.
At the same time, more unmelted W particles appear at the forefront of the molten spread under low energy, further hindering the flow of the melt and making the spread of the melt more difficult. According to reports, the diffusion time of W droplets is 86.3µs, while their solidification time is only 46µs [26]. The low solidification time and several times longer spread time cause the W melt to solidify without completing the spreading process, resulting in severe balling on the surface. Meanwhile, due to the fast scanning speed, the cooling rate of the melt increases. The faster solidification rate and lower fluidity result in discontinuities in the scanned tracks and lower relative density of the material. With changes in process parameters, the overlap between the melt tracks improves. At an energy density of 972 J/mm3, the spreading process of W powder particles is improved, resulting in an increase in the relative density of the formed W parts and a reduction in internal cracks and voids. However, further increasing the volumetric energy density will lead to overheating of the melt pool, an increase in the overlap rate of adjacent laser tracks, and the occurrence of sintering losses, surface wrinkles, and a small amount of splatters, which will damage the forming quality. The evaporation and combustion of elements will also lead to a further reduction in structural density.
Tungsten grain analysis and XRD phase analysis after heat treatment
Figure 7 presents the XRD patterns of pure W at different heat treatment temperatures, where the diffraction peaks of pure W powder and specimens correspond to BCC W (JCPDS card no. 040806). Due to the protection of the Ar atmosphere during the SLM process, there were no secondary peaks of oxide such as WO3 in the XRD patterns. However, after heat treatment, a noticeable shift in the 2θ angles of the W diffraction peaks occurred. As the heat treatment temperature increased from 1400℃ to 1900℃, the 2θ angles of the diffraction peaks changed from high angles to low angles, with the lowest recorded at 1700℃. The residual stresses in the SLM process were reduced after heat treatment, which explains the reason for the decrease in the 2θ angles.
Figure 8 shows the grain morphology under different heat treatment conditions. Large grain structure is present in W samples that are not heat-treated, which are actually w powder bonding masses structures that have not been fully melted. Due to the fast laser scanning speed, the temperature in the molten pool did not remain stable after reaching the melting temperature of W, and the temperature quickly decreased, causing only the boundaries of W particles to melt and their internal grains to solidify without fully melting. Therefore, there was no process of nucleation and growth of dendritic structures during complete melting and solidification, and no dendritic structure appeared. They are wrapped in Large grains, which also confirms that the internal grains of W were not melted. Compared with the two groups at 1100°C and 1400°C, after heat treating at 1700°C, the residual stress generated by rapid solidification can be eliminated, and the grain shape is closer to spherical. But there is little difference in grain size, the grain size ranges from 40–70 um. After SLM processing, the tungsten samples exhibited low stored deformation energy, and the pres-ence of small amounts of impurity elements in the crystalline tungsten powder hindered dislocation movement and grain boundary migration, As a result, the recrystallization temperature rises compared to normal. However, when the heat treatment temperature increased to 1900℃, the tungsten grains exhibited noticeable growth compared to the other samples, the grain size is more than 100 um. Due to tungsten being a pure metal, there are no second-phase particles anchoring the grain boundaries, making it prone to grain migration and growth.
Figure 9 shows the surface morphology of pure tungsten sample after heat treatment. The results show that both 1700℃ and 1900℃ improve the bonding degree of unfused powder particles on tungsten surface to a certain extent, and the bonding effect of 1700℃ is more obvious. There are no significant changes in the surface morphology of pure tungsten during low-temperature heat treatment (1100°C, 1400°C), and the influence on the grain structure is relatively small. The heat treatment temperature of 1700℃ was sufficient to partially melt and slightly recrystallize the pure tungsten particles. During the recrystallization process, there was a rearrangement of crystal planes and lattice recombination inside the particles, leading to the repositioning of grain boundaries. This phenomenon of particle recrystallization helped improve the lattice structure and grain boundary characteristics of pure tungsten, promoting grain boundary motion and rearrangement, thereby enhancing the strength and stability of the grain boundaries and improving the mechanical properties of the material [27]. In Fig. 9d, we can observe the agglomeration of pure tungsten particles and partial melting flow of some particle structures at high temperatures. This resulted in changes in the surface morphology, with occurrences of particle recrystallization, surface leveling, and other phenomena, which were also the reasons for the improvement in mechanical properties. However, when the heat treatment temperature was increased to 1900℃, the enlargement of grain size led to the accumulation and slip of dislocations, accelerating the formation and propagation rate of fatigue cracks, thereby reducing the mechanical properties of the material.
Mechanical properties under different heat treatment conditions
To release the internal stress of the material, the study investigated the changes in the properties of pure W samples under different heat treatment temperatures. The experiment used a vacuum hot press for annealing. Based on the relative density of the sample, the optimum process parameters were selected with an energy density of 972 J/mm3, which corresponded to a relative density of 96.2%. The sample was annealed at four heat treatment temperatures of 1100℃, 1400℃, 1700℃ and 1900℃ for 2 hours each. The compression test equipment for the sample was a Taisite universal testing machine. The compression performance of the samples after treatment at different temperatures is shown in Fig. 10a.The hardness of the samples at different heat treatment temperatures was measured using a Vickers hardness tester, with a load of 10 kgf and a dwell time of 10s. Figure 10b shows the microhardness of W parts at different heat treatment temperatures.
The residual stresses generated during rapid solidification were eliminated in the samples after heat treatment at 1000℃ and 1400℃. The heat treatment at 1700°C caused remelting of pure tungsten particles and changes in the surface morphology. Simultaneously, the combined effects of particle recrystallization, grain boundary migration, and grain growth enhanced the mechanical properties of pure tungsten, including hardness and compressive strength. At 1900℃, since tungsten had already melted and solidified during SLM fabrication, there was no process of particle rearrangement and formation and growth of sintering necks. Only grain growth occurred, leading to a decrease in the material’s mechanical properties, hardness value included. The failure characteristics of the structure mainly depend on the material composition [28, 29]. The deformation process is the same in all samples. When the lattice structure is deformed under stress, stress concentration is clearly visible at the nodes, and the structure failure is mainly caused by node fracture. Further analysis of the deformation of each structure shows that elastic deformation occurs in the initial stage, and as the load increases, the structure begins to undergo plastic deformation. Plastic deformation mainly occurs on the top column, as shown in Fig. 11, and the interference between the columns tends to saturate. The slope of the compression curve increases rapidly. When the stress of the structure reaches the peak value, the unit column node fractures, the stress rapidly decreases, and the sample collapses layer by layer, resulting in structural failure. The support column completely fractures from the nodes, and then the strength of the lattice structure is significantly weakened, and the cell columns are crushed by contacting the next layer of support columns. The layer-by-layer collapse damage characteristics are shown at the marked position in the figure.
Morphology analysis of sample fracture at different heat treatment temperatures
The image in Fig. 12 displays the fracture morphology of pure W parts after heat treatment. The fracture exhibits typical brittle fracture characteristics with less plastic deformation. Typical fracture modes at different heat treatment temperatures were characterized. In its pristine state, many voids were visible inside the specimen (Fig. 12a). Weak bonding between layers and particles resulted in voids and unmelted particles being the primary causes of early fracture, the crack originates from and propagates toward the selective laser melting (SLM) processing defects [30]. Additionally, the waviness and roughness of the pillars increased local stress concentration, resulting in a lower compressive strength [31]. As the heat treating temperature increased to 1100℃ and 1400℃, the specimens became denser, with fewer unmelted particles. During compression, the grains were squeezed together, and the crack propagated along the grain boundaries due to the reduced intergranular strength from the high ductile-to-brittle transition temperature. These initial cracks continued to extend along the grain boundaries as they slipped, causing the sample to fail along the grain boundaries, leading to overall material failure. Due to tungsten’s inherent properties, microcracks are difficult to avoid, but the compression process is not sensitive to them, and there is almost no difference in ultimate compressive strength compared to the non-heat-treated state. Further increasing the heat treating temperature to 1700℃ showed a significant difference in fracture morphology. After treatment at 1700℃, the fracture surface exhibited typical brittle characteristics with evident cleavage steps and facets. The presence of voids and microcracks was reduced. During the compression process, particles were mutually compressed. Due to the low strength of grain boundaries caused by the high ductile-to-brittle transition temperature, these initial cracks propagated along the grain boundaries as they continuously slipped. The samples fractured along the grain boundaries, and after heat treatment at 1700℃, the grain size was reduced, resulting in an increase in grain boundary strength. As a result, the ultimate compressive strength showed an improvement. As shown in Fig. 12d, most of the pores and microcracks disappeared. The fracture mode indicated that the melting of the unmelted powder particles was sufficient at this temperature, resulting in strong metallurgical bonding between layers and particles. Additionally, cleavage steps and planes were evident at this temperature. In this case, the obtained ultimate compressive strength was the highest among all heat-treated samples. When the heat treatment temperature reached 1900℃, as shown in Fig. 12e, only microcracks were observed. The fracture surface morphology indicated that the low grain boundary strength caused by the high ductile-to-brittle transition temperature led to the initiation of cracks at the grain boundaries. These initial cracks propagated along the grain boundaries. In comparison to 1100℃ and 1400℃, there was almost no difference in mechanical properties.
This study investigated the comprehensive effects of energy density and heat treatment on the performance of 3D printed W, focusing on the variation of energy density during the printing process and how post-printing heat treatment affects the microstructure and mechanical properties of the final product. By comprehensively analyzing these interacting factors, this study provides new insights for the production of high-performance 3D printed W components under optimal conditions. The specific conclusions are as follows:
·The volumetric energy density plays a crucial role in the densification of W parts fabricated by SLM. A lower energy density leads to porosity due to incomplete melting of powder particles, resulting in a decrease in density. However, density further decreases with increasing energy density. In this study, the optimal process parameters were selected with an energy density of 972 J/mm3, corresponding to a theoretical density of 96.2% for pure W specimens.
·After heat treatment at 1100℃, 1400℃, 1700℃ and 1900℃, high relative density (96%) pure W samples fabricated by SLM exhibited a maximum compressive strength of 1261 MPa after treatment at 1700℃, representing a 65% increase from the untreated state. The maximum Vickers hardness was found to be 457 HV, and the average surface roughness was Ra 6 um. The fracture morphology revealed that at this temperature, there was sufficient melting of unmelted powder particles, resulting in strong metallurgical bonding between layers and particles. Additionally, distinct fracture steps and surfaces were observed, indicating cleavage fracture, which provided the highest ultimate compressive strength among all heat treatments. Under conditions of 1000℃ and 1400℃, the microstructure of the samples showed minimal changes, resulting in mechanical properties consistent with their original state.
Acknowledgments This work was financially supported by the Natural Science Foundation of Fujian Province (No.2021J011213), and National Natural Science Foundation of China (52173266).
Conflict of interest The authors declare no interest conflict.
Author Contribution
C.WL：A complete experiment was done and the article was written.Y.YH:Experiment and paper writing guidance.
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No competing interests reported.