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Microstructure Characteristics of Inconel 625 Superalloy Manufactured by Selective Laser Melting
Accepted Manuscript Microstructure Characteristics of Inconel 625 Superalloy Manufactured by Selective Laser Melting Shuai Li, Qingsong Wei, Yusheng Shi, Zicheng Zhu, Danqing Zhang PII:
S1005-0302(15)00066-3
DOI:
10.1016/j.jmst.2014.09.020
Reference:
JMST 492
To appear in:
Journal of Materials Science & Technology
Received Date: 22 April 2014 Revised Date:
24 August 2014
Accepted Date: 18 September 2014
Please cite this article as: S. Li, Q. Wei, Y. Shi, Z. Zhu, D. Zhang, Microstructure Characteristics of Inconel 625 Superalloy Manufactured by Selective Laser Melting, Journal of Materials Science & Technology (2015), doi: 10.1016/j.jmst.2014.09.020. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
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Microstructure Characteristics of Inconel 625 Superalloy Manufactured by Selective Laser Melting Shuai Li1, 2, Qingsong Wei1, *, Yusheng Shi1, Zicheng Zhu2, Danqing Zhang2 State Key Laboratory of Die & Mould Technology, Huazhong University of Science and
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1
Technology, Wuhan 430074, China 2
NTU Additive Manufacture Centre, Nanyang Technological University, 50 Nanyang Avenue,
Singapore 639798, Singapore
*Corresponding author. Prof.; Tel.:+86 13296512995.
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E-mail address:[email protected] (Qingsong Wei).
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[Manuscript received 22 April 2014; in revised from 24 August 2014; Accepted 18 September 2014]
Selective laser melting (SLM), an additive manufacturing process, is capable of manufacturing metallic parts with complex shapes directly from computer-aided design (CAD) models. SLM parts are created on a layer-by-layer manner, making it more flexible than traditional material processing
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techniques. In this paper, Inconel 625 alloy, a widely used material in the aerospace industry, were chosen as the build material. Scanning electron microscopy (SEM), electron back scattering diffraction (EBSD) and X-ray diffraction (XRD) analysis techniques were employed to analyze its
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microstructure. It was observed that the molten pool was composed of elongated columnar crystal. Due to the rapid cooling speed, the primary dendrite arm space was approximately 0.5 µm and the
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hardness of SLM state was very high (343 HV). The inverse pole figure (IPF) indicated that the growing orientation of the most grains was <001> due to the epitaxial growth and heat conduction. The XRD results revealed that the austenite structure with large lattice distortion was fully formed. No carbides or precipitated phases were found. After heat treatment the grains grew into two microstructures with distinct morphological characters, namely, rectangular grains and limited in the molten pool, and equiaxed grains along the molten boundaries. Upon experiencing the heat treatment, MC carbides with triangular shapes gradually precipitated. The results also identified that a large
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ACCEPTED MANUSCRIPT number of zigzag grain boundaries were formed. In this study, the grain formation and microstructure, and the laws of the molten pool evolution were also analyzed and discussed. Key words: Selective laser melting; Nickel based superalloy; Texture; Lattice constant; Zigzag grain
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boundary
1. Introduction
Inconel 625 is a nickel based solid solution strengthening superalloy, which is largely
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strengthened by Mo and Nb elements. Inconel 625 has been widely applied to various areas especially in the aerospace industry since the 1960s. This material has drawn particular attention due perfect combination of good yield, tensile and creep
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to its excellent characteristics including
strengths as well as strong resistance to high temperature corrosion on prolonged exposure to aggressive environments[1,2]. However, it is difficult to control the performance when casting or forging for this material. Selective laser melting (SLM) utilizes a laser beam to melt pure metal or pre-alloyed powders layer-by-layer according to the given computer-aided design (CAD) model,
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directly creating nearly fully dense metal parts with complex geometries[3‒5]. It shows unique advantages in material saving, process control and parts performance. At present, extensive work has been conducted on producing nickel based superalloy parts. Yadroitsev et al.[6,7] studied the influence
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of different scanning interval and scanning strategies on Inconel 625 powders fabricated with porous components. It was identified that the tensile strength of Inconel 625 is much higher than forgings
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standard level. Mumtaz and Hopkinson[8] used a pulsed laser to explore the impact of different laser pulses on Inconel 625 forming ability. Vilaro et al.[3] investigated the microstructure and phase composition of Nimonic 263 nickel based alloy and the mechanical properties were studied under different heat treatment conditions. They found that the resulting microstructure is quite homogeneous but remains out-of-equilibrium[3]. All the above researches obtained nickel based parts with high densities, good surface quality and mechanical properties by designing different experiments and optimizing the process parameters.
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ACCEPTED MANUSCRIPT However, Inconel 625, as a solid solution strengthened superalloy, is normally subjected to varying application conditions, including continuous mechanical and thermal stresses, and other environmental factors such as corrosion. The internal microstructure of the parts has significant influence on its macroscopic properties[9]. The majority of the current studies concentrated on the
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macroscopic properties, paying less attention on the microstructure and the phase composition. G.P Dinda et al.[9] used laser aided direct metal deposition to fabricate Inconel 625 superalloy samples and the microstructure evolution and thermal stability was also studied. Nevertheless, the laser deposition process uses a higher heat input, resulting in a larger single track and lower temperature
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gradient, which is different from the SLM process. For instance, the cooling speed can be as high as 106 K/s for SLM process[10]. As a consequence, the SLM parts exhibit different properties in the
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microstructure and phase composition in comparison to the traditional casting and forging parts. Although the SLM parts presented good mechanical properties[11], they also showed a strong anisotropy due to the directional columnar grain growth caused by directional thermal conduction during SLM process[10]. The texture has significant influence on the mechanical properties.
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Unfortunately, little literature gives a deep analysis about the texture so far. In the meantime, due to the differences in application conditions, the original structure of the SLM parts may not be ideal in terms of long time use. Therefore, it is necessary to explore the influence of heat treatment on the SLM manufactured parts, which is aimed at improving its mechanical properties.
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In this study, several Inconel 625 samples were fabricated using SLM and the microstructures of both the X-Y and Y-Z sections have been observed and analyzed by scanning electron microscopy
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(SEM). By applying the electron back scattering diffraction (EBSD) method, the texture and grain morphology were studied to show the anisotropy of the SLM parts. The effect of annealing temperatures on the microstructures related to the micro-hardness, the lattice constant and the distribution of the carbides was also analyzed.
2. Experiments
2.1. Materials and processes 3
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A commercial Inconel 625 powder (H.C. Starck GmbH, Germany) produced by gas atomized process with the average particle size of 34.63 µm was used. Fig. 1(a) illustrates the SEM morphology of the powder. The shape is approximately spherical and the surface is smooth. The
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powder is composed of 53.5 Ni, 21.5 Cr, 0.96 Fe, 3.71 Nb, 8.8 Mo, 0.47 Mn, 0.41 Si in wt%. The powder was heated at 50 °C for 5 h prior to the SLM process. This procedure was to eliminate the water vapor inside the powder and get a good flow ability of the powder.
The parts were fabricated on an HRPM-II machine developed by Huazhong University of
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Science and Technology, China. The machine is equipped with a 200-W fiber laser, of which the laser spot is 50‒80 µm in diameter and scanning speed ranges from 200 to 1000 mm/s. The chamber
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was evacuated followed by filling in a high-purity argon gas to form an oxygen-free atmosphere. The optimized processing parameters obtained by our group were adopted for manufacturing of almost fully dense Inconel 625 parts, including the laser power of 160 W, the scan speed of 500 mm/s, the hatch space of 0.06 mm, and the layer thickness of 0.02 mm. Bidirectional scan mode was used as
2.2. Microstructures
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shown in Fig. 1(b).
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The samples were then subjected to mechanical polishing with a grit size of 1 µm. Subsequently, the polished samples were etched for 30 s in a mixture solution comprising of 10 ml HNO3, 10 ml
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HCL and 15 ml CH3COOH. The distribution characteristics of the elements were analyzed by the energy dispersive spectrometer (EDS). The grain orientation, grain size and grain boundary characteristics were determined by using the EBSD system mounted on the scanning electron microscopy machine (JEOL 7600F). EBSD-scans were conducted on the Y-Z and X-Y sections of the SLM samples. For this purpose, mechanically pre-polished samples were electro polished for 20 s at 20 V in a 5% perchloric acid solution. The data was analyzed using HKL Channel 5. The crystal structure and the lattice parameters of the samples were tested using XRD Cu K radiation at 40 kV and 100 mA. All samples were scanned in the standard geometry from 30° to 100° with a 0.01° step 4
ACCEPTED MANUSCRIPT size and 1-s dwelling time. In order to evaluate the microstructure evolution after heat treatment, a set of samples were annealed respectively at 700, 1000, 1150 °C for 1 h followed by cooling in still air. The Vickers hardness was measured under 500 g load applied for 30 s. Each hardness value is
3. Results and Discussion
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3.1. Microstructures and characteristics of the SLM samples
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the average of 5 measurements.
Fig. 2(a) shows the surface morphology of the samples prior to polishing. A very clear “V” shape
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morphology is identified, which normally appears in welding processing. This phenomenon has already been found in other materials[10]. The adjacent melting tracks tightly overlap, with a width of approximately 100 µm, generating a nearly full-dense structure devoid of lacuna. Therefore, it can be identified that the Inconel 625 superalloy powder has a good forming performance. The formation of
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this morphology can be attributed to the laser heat source induced by the moving laser. The laser beam can be considered to be a point of intense heat source leading to a spindle temperature field in the molten pool[12]. The powder is continuously melted in the front of the pool and the melted liquid metal rapidly solidifies as soon as the laser moves away. Meanwhile, the uneven temperature
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distribution causes the difference between the density and surface tension. As a result, the molten pool keeps stirring and generally forms convection throughout the entire process, leading to the
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formation of “V” shape morphology. The angle of “V” morphology largely depends on the laser scanning speed. The faster the laser scans, the smaller the “V” angle becomes, and the narrower the molten pool will be. The morphology of the Y-Z section presents a “scales” shape as illustrated in Fig. 2(b). In this experiment, the layer thickness was 0.02 mm. Most of the previous layer was remelted by the following process. The scales shape matched the Gaussian distribution of laser energy. Due to most of the laser energy being focused on the laser beam center, a depression appears at the bottom of the molten pool as indicated by the arrows. If the density of laser energy increases, a key hole is likely to form which will potentially improve the joint between adjacent layers. A higher 5
ACCEPTED MANUSCRIPT magnification of molten pool boundary area is shown in Fig. 2(c), from which three different zones were observed: molten pool, fusion line and heat affect zone (HAZ). The cellular structure sizes change from approximately 0.2 to 1 µm. The microstructures in the molten are complicated due to the convection in the molten pool. In Fig. 2(c), cellular structure was found while in molten pool 2
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the cellular was different. The width of the fusion line is approximately 0.5 µm, which is as thin as one layer of sub-grain. Moreover, the temperature outside the fusion line is still very high when the powders are being melted. This causes the cellular structure boundary, which is primarily composed of Nb element dissolved into the substrate gradually. The sub-grain boundary is thus interrupted in
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this area. Fig. 2(c) also indicates that the columnar dendrite structure is highly degenerated with the second dendrite growth being entirely suppressed. Due to rapid solidification, the average primary
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dendrite arm space is approximately 0.5 µm, which is nearly 2 orders of magnitude smaller as compared with the traditional casting (100‒300 µm)[13]. Rapidly solidified dendrite space can be determined by the following equation[14]:
=
(1)
where d is the primary dendrite arm space; b and a are associated with the material constants (for
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nickel based alloys, b=1/3, a≈50 µm(K/s)[15]). Thus, the solidification cooling rate ε can be calculated, i.e. ε =106 K/s, which is consistent with the numerical simulation of the SLM process[3]. The EBSD clearly shows the texture of the samples. Fig. 3 is the inverse pole figure (IPF) of the
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Y-Z and X-Y section, in which a region with a single color represents the presence of one grain only. Hence, some larger elongation grains of several hundred micrometers in length grow cross a few
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layers. A very strong texture can be identified for the SLM sample. The long axes of most of the grains are aligned in the <001> direction and a cube texture is observed in the Y-Z section, as shown in Fig. 3(a). This phenomenon was also observed by other researchers[16]. Inconel 625 has a facecentered cubic crystal structure. The preferred orientation is <001>[17]. During the SLM process, the laser melts the powder on the top of the sample but the substrate or the layer that has solidified still remains cold. So, the temperature gradient forms from the top to the bottom. The dendrites get the priority to grow; as a result, the angle between <001> and the direction of heat flux is minimum. Meanwhile, the maximum temperature in the molten pool can reach 3000 K[18], leading to superheat. 6
ACCEPTED MANUSCRIPT These further result in the difficulty to form homogeneous nucleation in the molten pool, which can enhance the growth of columnar crystal in the molten pool. Under this condition, the columnar grain will grow from the bottom of the molten pool towards the top. However, the interaction of the laser and powder, liquid metal, surrounding atmosphere, solidified base metal and un-melted metal
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powders can collaboratively form a large temperature gradient, which may result in surface tension gradient and Marangoni convection[10]. The velocity of convection flow in the melt pool can reach tens of meter per second according to the simulation result[19]. The convection in the molten pool can break the original growth law of columnar crystal. Parts of dendrites change growth directions,
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causing the formation of steering dendrite. Thus, some grains change the grow directions rather than always being towards <001>. Recently, some researchers used two yttrium fiber lasers with the
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power of 400 W and 1000 W, respectively to melt the powder. They found that the texture was mostly influenced by laser energy input condition[20]. Under a laser power of 1000 W, even more strong texture is obtained. For the texture that has a significant impact on the mechanical properties,
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this phenomenon should be considered in the future.
3.3. Influence of heat treatment on the microstructure
The hardness variation of the SLM samples after annealing at different temperatures is shown in
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Fig. 4. Since the cooling rate in the SLM process is very high (about 106 K/s), most of the strengthening elements such as Mo, Nb remain in the Ni matrix. A lager lattice distortion is caused
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by these point defects. The hardness of the original SLM samples is 343 HV, which is higher than that of the forged samples (typically 305 HV). Although Inconel 625 belongs to solid solution strengthening superalloy, phase transition also takes place during the heat treatment. The hardness of the sample is slightly lower than that of SLM state due to the release of the residual stress after annealing at 700 °C. The samples show the increase of the hardness when annealing at 800 and 900 °C due to the precipitation of δ (Ni3Nb). δ is an orthorhombic phase with lattice constant a = 0.574 nm, b = 0.422 nm, c = 0.454 nm. The mismatch with the Ni matrix is relatively large; thereby, the hardness increases with the heat treatment at 800 and 900 °C. When the annealing temperature is 7
ACCEPTED MANUSCRIPT above 1000 °C, the dissolution of δ phase into the matrix and the decrease of the lattice distortion cause the decrease of hardness rapidly. The phase transition will be discussed in the following in detail. Fig. 5(a) shows that the microstructure of the SLM formed parts is austenite; no carbides and
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other phases are found. The reason is as follows: as the laser beam moves very fast (600 mm/s), the solidification time is very short (often less than 1 ms). The atomic reconstruction speed in front of the short-range liquid/solid solidification is much higher than the diffusion speed, so solution atoms will be trapped by the rapid solidification, and thus the so-called “solute trapping” is formed. Most of
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solute atoms such as Cr, Mo, Nb are captured in the Ni matrix and phase transition is hard to occur. In the same way, carbides are also difficult to aggregate and precipitate.
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The XRD results under different heat treatment are shown in Fig. 5(c). The microstructures are the same as the SLM state and all of them are the solution of Cr in Ni matrix. The intermetallic phases such as γ′′ (Ni3Nb), δ (Ni3Nb), which can usually be found in casting and forging but are not detected in the XRD results because the sizes of these particles are small[9]. However, the lattice
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constants of the matrix that are calculated from the XRD results can imply the nature of precipitation. The (200) diffraction peak in Fig. 5(c) was used to calculate the lattice parameters according to Bragg diffraction. The equations are given as follows: (2)
= ⁄√ℎ + +
(3)
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2 sin =
Table 1 lists the lattice constants of the samples at different states. For the SLM state, most of the
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strengthening elements such as Nb, Mo are embedded into the matrix, causing large lattice distortion, and the lattice constant reaches the maximum. When the samples are treated at 700 °C, the Nb element precipitates to form γ′′ (Ni3Nb) phase, and the lattice constants decline accordingly. When the temperature rises to 850 °C, the γ′′ phase transforms to a new stable structure of orthorhombic δ (Ni3Nb). Moreover, when the samples are treated at 1000 °C, the strengthening elements are gradually dissolved into the matrix, resulting in the increase in lattice constant. However, the opposite results were obtained. This is mainly caused by a large number of carbides precipitating from the matrix at 1000 °C. An obvious peak representing the MC carbides can be observed in the 8
ACCEPTED MANUSCRIPT XRD results. The carbide element precipitates from the matrix, making the lattice constant increase. Fig. 5(b) shows the diffraction peak positions of the (200) crystal plane in different heat treatment processes. The changes of the integral line-breadth can be detected. The integral line-breadth becomes narrow gradually from SLM state, 700 °C, 1000 to 1150 °C, which means that the crystal
the space of the austenite phase crystal plane has increased.
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size becomes larger whilst the angles of the diffraction peaks become larger, which demonstrates that
Fig. 6 shows the microstructures of the samples that have undergone air cooling for 1 h after being heated at 1000, and 1150 °C, respectively. At 700 °C, the microstructure is the same as that of
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the SLM state. When the temperature reaches 1000 °C, the homogenization of the elements takes place and large amount of fine grains form, causing the original molten pool boundary gradually
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disappears. In addition, a number of particulates appear along the grain boundaries (as shown in Fig. 6(c)). The analysis of energy spectrum demonstrates that the particles, which are pointed by arrows have high content of Nb, Mo, and C. Mo and Nb have strong affinities to form carbides with carbon. The total content of Mo and Nb is approximately 45% and carbon content is nearly 6% in weight.
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The result is in consistent with the carbon content of MC carbide. The NbC, MoC carbide particulates distribute around the grain boundaries. These carbides play a role in pinning the migration of grain boundaries, thereby increasing the strength of the alloy. When the samples are heated at 1150 °C for 1 h, the grains grow obviously, and the boundary carbides still exist and
forming process.
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gradually grow up. However, the grain sizes are still much smaller than those of the traditional
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Two phenomena were identified. First, after heat treatment most of the grains have an obvious preferred orientation. In the middle of the molten pool, the grain grows and finally becomes a rectangular shape and is confined in the molten pool. Almost all of them have the same orientation. This phenomenon is caused by heat conduction in the SLM process. Numerous small grains are formed along the boundary of the molten pool. This is because impurities are mainly concentrated in the pool boundary during the solidification process. These impurities provide heterogeneous cores for the grains; as a result small grains grow up along the boundary. The second phenomenon is the generated zigzag grain boundary after heat treatment as shown in Fig. 6(c). The zigzag grain 9
ACCEPTED MANUSCRIPT boundary is significant for the alloy to improve the ductility, toughness and comprehensive performance[21,22]. Although there are many different theories to explain the forming of the zigzag grain boundary[23,24], there are no explanations for Inconel 625. A new explanation was developed to explain the zigzag grain boundary for Inconel 625, as plotted in Fig. 7. The formation of the zigzag
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largely results from the existence of MC carbides. The bending degree of the grain boundary is affected by the carbide orientation, carbon particle spacing and carbides size. MC carbides with bigger sizes may cause severer bending. If adjacent MC carbides have the same orientation, there will be a straight grain boundary rather than a zigzag grain boundary. This model can be used to
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explain the zigzag formation of the SLM parts after heat treatment.
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4. Conclusions
In summary, Inconel 625 superalloy samples have been successfully fabricated. Various methods were applied to exploring the microstructures of the as-fabricated samples and the samples
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after heat treatment. Conclusions can be summarized below:
(1) SEM results show that the molten pool is composed of elongated columnar crystal. The primary dendrite arm space is about 0.5 µm. The cooling rate is about 106 K/s. Due to the thermal history, a strong texture was formed from the substrate to the top of the sample.
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(2) The original microstructure of the SLM manufactured Inconel 625 superalloy is austenite with high solution distortions. No carbides or phases were precipitated. The hardness of the
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SLM fabricated parts was very high (254±6) due to the very fine microstructure with higher degree of supersaturation of strengthening elements. (3) After heat treatment, the microstructures are similar to the SLM primitive state, although the lattice constant descends slightly. The decrease is primarily caused by a large number of NbC and MoC carbides precipitated from the matrix. (4) The grains have a clear preferred orientation, which has a relationship with the heat conduction after heat treatment. Due to the influence of MC carbides, zigzag grain boundary is formed, which is crucial for the alloy. 10
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Acknowledgements This study was funded by the National Science & Technology Pillar Program of China (Granted No. 2012BAF08B00), the National Natural Science Foundation of China (Granted Nos. 51375189
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and 51375188), the independent R&D subjects of Huazhong University of Science and Technology and the State Key Laboratory of Materials Processing and Die & Mold Technology. This study was also co-funded by China Scholarship Council (CSC).
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[19] D. Dai, D. Gu, Mater. Des. 55 (2014) 482‒491. [20] T. Niendorf, S. Leuders, A. Riemer, H.A. Richard, T. Tröster, D. Schwarze, Metall. Mater. Trans. B 44 (2013) 794‒796. [21] Z. Mu, R. Ye, L. Gao, Mater. Sci. Technol. 4 (1988) 540‒547.
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[23] O. Miyagawa, M. Yamamoto, M. Kobayashi, in: Proceedings of the Third International Symposium on Metallurgy and Manufacturing of Superalloys, 1976, pp. 245‒254.
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[24] R.F. Decker, C.T. Sims, The Superalloys, John Wiley & Sons, Inc., New York, 1972, p. 33.
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ACCEPTED MANUSCRIPT Figure and table captions Table 1 Lattice constants of the specimens in different states
Fig. 1. (a) SEM morphology of Inconel 625 powder, (b) illustration of the bidirectional scan mode.
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The arrows indicate the movement of the laser. Fig. 2. SEM micrographs of SLM samples: (a) top surface morphology of the samples (unpolished), (b) “scales” shape of Y-Z section, (c) high magnification picture of the molten pool boundary. The lines present the three different areas: HAZ, fusion line and molten pool.
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Fig. 3. Microstructure morphology of the SLM formed Inconel 625 superalloy: (a) inverse pole figure (IPF) colored map of the Y-Z section, (b) inverse pole figure (IPF) colored map of the X-Y
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section, (c) (100) pole figure of the IPF image of the SLM formed Inconel 625 part, (d) the index map of the IPF and the reference coordinate.
Fig. 4. Hardness of the samples after annealing at different temperatures. Fig. 5. (a) X-ray diffraction of the SLM deposited sample, (b) diffraction peak positions of (200)
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crystal plane under different heat treatment processes, (c) X-ray diffraction of the SLM deposited sample at different temperatures.
Fig. 6. Microstructure evolutions of X-Y plane after heat treatment: (a) optical metallographic image of the specimen after 1000 °C, (b) optical metallographic image of the specimen after 1150 °C, (c)
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magnified view showing the grain boundary structure and the carbides, (d) inverse pole figure (IPF) colored map of the X-Y section.
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Fig. 7. Model of the segment of grain boundary in MC zone.
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ACCEPTED MANUSCRIPT Table and Figure List Table 1 Lattice constant (nm)
SLM
0.3604±0.0001
SLM + 700 °C, 1 h, AQ*
0.3599±0.0002
SLM + 1000 °C, 1 h, AQ
0.35972±0.00001
SLM + 1150 °C, 1 h, AQ
0.3595±0.00001
*
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Note: AQ: water quenching
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Heat treatment
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Fig. 2
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Fig. 1
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Fig. 4
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Fig. 3
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Fig. 5
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Fig. 7
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Fig. 6
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S1005-0302(15)00066-3
DOI:
10.1016/j.jmst.2014.09.020
Reference:
JMST 492
To appear in:
Journal of Materials Science & Technology
Received Date: 22 April 2014 Revised Date:
24 August 2014
Accepted Date: 18 September 2014
Please cite this article as: S. Li, Q. Wei, Y. Shi, Z. Zhu, D. Zhang, Microstructure Characteristics of Inconel 625 Superalloy Manufactured by Selective Laser Melting, Journal of Materials Science & Technology (2015), doi: 10.1016/j.jmst.2014.09.020. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
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Microstructure Characteristics of Inconel 625 Superalloy Manufactured by Selective Laser Melting Shuai Li1, 2, Qingsong Wei1, *, Yusheng Shi1, Zicheng Zhu2, Danqing Zhang2 State Key Laboratory of Die & Mould Technology, Huazhong University of Science and
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1
Technology, Wuhan 430074, China 2
NTU Additive Manufacture Centre, Nanyang Technological University, 50 Nanyang Avenue,
Singapore 639798, Singapore
*Corresponding author. Prof.; Tel.:+86 13296512995.
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E-mail address:[email protected] (Qingsong Wei).
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[Manuscript received 22 April 2014; in revised from 24 August 2014; Accepted 18 September 2014]
Selective laser melting (SLM), an additive manufacturing process, is capable of manufacturing metallic parts with complex shapes directly from computer-aided design (CAD) models. SLM parts are created on a layer-by-layer manner, making it more flexible than traditional material processing
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techniques. In this paper, Inconel 625 alloy, a widely used material in the aerospace industry, were chosen as the build material. Scanning electron microscopy (SEM), electron back scattering diffraction (EBSD) and X-ray diffraction (XRD) analysis techniques were employed to analyze its
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microstructure. It was observed that the molten pool was composed of elongated columnar crystal. Due to the rapid cooling speed, the primary dendrite arm space was approximately 0.5 µm and the
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hardness of SLM state was very high (343 HV). The inverse pole figure (IPF) indicated that the growing orientation of the most grains was <001> due to the epitaxial growth and heat conduction. The XRD results revealed that the austenite structure with large lattice distortion was fully formed. No carbides or precipitated phases were found. After heat treatment the grains grew into two microstructures with distinct morphological characters, namely, rectangular grains and limited in the molten pool, and equiaxed grains along the molten boundaries. Upon experiencing the heat treatment, MC carbides with triangular shapes gradually precipitated. The results also identified that a large
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ACCEPTED MANUSCRIPT number of zigzag grain boundaries were formed. In this study, the grain formation and microstructure, and the laws of the molten pool evolution were also analyzed and discussed. Key words: Selective laser melting; Nickel based superalloy; Texture; Lattice constant; Zigzag grain
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boundary
1. Introduction
Inconel 625 is a nickel based solid solution strengthening superalloy, which is largely
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strengthened by Mo and Nb elements. Inconel 625 has been widely applied to various areas especially in the aerospace industry since the 1960s. This material has drawn particular attention due perfect combination of good yield, tensile and creep
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to its excellent characteristics including
strengths as well as strong resistance to high temperature corrosion on prolonged exposure to aggressive environments[1,2]. However, it is difficult to control the performance when casting or forging for this material. Selective laser melting (SLM) utilizes a laser beam to melt pure metal or pre-alloyed powders layer-by-layer according to the given computer-aided design (CAD) model,
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directly creating nearly fully dense metal parts with complex geometries[3‒5]. It shows unique advantages in material saving, process control and parts performance. At present, extensive work has been conducted on producing nickel based superalloy parts. Yadroitsev et al.[6,7] studied the influence
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of different scanning interval and scanning strategies on Inconel 625 powders fabricated with porous components. It was identified that the tensile strength of Inconel 625 is much higher than forgings
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standard level. Mumtaz and Hopkinson[8] used a pulsed laser to explore the impact of different laser pulses on Inconel 625 forming ability. Vilaro et al.[3] investigated the microstructure and phase composition of Nimonic 263 nickel based alloy and the mechanical properties were studied under different heat treatment conditions. They found that the resulting microstructure is quite homogeneous but remains out-of-equilibrium[3]. All the above researches obtained nickel based parts with high densities, good surface quality and mechanical properties by designing different experiments and optimizing the process parameters.
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ACCEPTED MANUSCRIPT However, Inconel 625, as a solid solution strengthened superalloy, is normally subjected to varying application conditions, including continuous mechanical and thermal stresses, and other environmental factors such as corrosion. The internal microstructure of the parts has significant influence on its macroscopic properties[9]. The majority of the current studies concentrated on the
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macroscopic properties, paying less attention on the microstructure and the phase composition. G.P Dinda et al.[9] used laser aided direct metal deposition to fabricate Inconel 625 superalloy samples and the microstructure evolution and thermal stability was also studied. Nevertheless, the laser deposition process uses a higher heat input, resulting in a larger single track and lower temperature
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gradient, which is different from the SLM process. For instance, the cooling speed can be as high as 106 K/s for SLM process[10]. As a consequence, the SLM parts exhibit different properties in the
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microstructure and phase composition in comparison to the traditional casting and forging parts. Although the SLM parts presented good mechanical properties[11], they also showed a strong anisotropy due to the directional columnar grain growth caused by directional thermal conduction during SLM process[10]. The texture has significant influence on the mechanical properties.
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Unfortunately, little literature gives a deep analysis about the texture so far. In the meantime, due to the differences in application conditions, the original structure of the SLM parts may not be ideal in terms of long time use. Therefore, it is necessary to explore the influence of heat treatment on the SLM manufactured parts, which is aimed at improving its mechanical properties.
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In this study, several Inconel 625 samples were fabricated using SLM and the microstructures of both the X-Y and Y-Z sections have been observed and analyzed by scanning electron microscopy
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(SEM). By applying the electron back scattering diffraction (EBSD) method, the texture and grain morphology were studied to show the anisotropy of the SLM parts. The effect of annealing temperatures on the microstructures related to the micro-hardness, the lattice constant and the distribution of the carbides was also analyzed.
2. Experiments
2.1. Materials and processes 3
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A commercial Inconel 625 powder (H.C. Starck GmbH, Germany) produced by gas atomized process with the average particle size of 34.63 µm was used. Fig. 1(a) illustrates the SEM morphology of the powder. The shape is approximately spherical and the surface is smooth. The
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powder is composed of 53.5 Ni, 21.5 Cr, 0.96 Fe, 3.71 Nb, 8.8 Mo, 0.47 Mn, 0.41 Si in wt%. The powder was heated at 50 °C for 5 h prior to the SLM process. This procedure was to eliminate the water vapor inside the powder and get a good flow ability of the powder.
The parts were fabricated on an HRPM-II machine developed by Huazhong University of
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Science and Technology, China. The machine is equipped with a 200-W fiber laser, of which the laser spot is 50‒80 µm in diameter and scanning speed ranges from 200 to 1000 mm/s. The chamber
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was evacuated followed by filling in a high-purity argon gas to form an oxygen-free atmosphere. The optimized processing parameters obtained by our group were adopted for manufacturing of almost fully dense Inconel 625 parts, including the laser power of 160 W, the scan speed of 500 mm/s, the hatch space of 0.06 mm, and the layer thickness of 0.02 mm. Bidirectional scan mode was used as
2.2. Microstructures
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shown in Fig. 1(b).
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The samples were then subjected to mechanical polishing with a grit size of 1 µm. Subsequently, the polished samples were etched for 30 s in a mixture solution comprising of 10 ml HNO3, 10 ml
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HCL and 15 ml CH3COOH. The distribution characteristics of the elements were analyzed by the energy dispersive spectrometer (EDS). The grain orientation, grain size and grain boundary characteristics were determined by using the EBSD system mounted on the scanning electron microscopy machine (JEOL 7600F). EBSD-scans were conducted on the Y-Z and X-Y sections of the SLM samples. For this purpose, mechanically pre-polished samples were electro polished for 20 s at 20 V in a 5% perchloric acid solution. The data was analyzed using HKL Channel 5. The crystal structure and the lattice parameters of the samples were tested using XRD Cu K radiation at 40 kV and 100 mA. All samples were scanned in the standard geometry from 30° to 100° with a 0.01° step 4
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3. Results and Discussion
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3.1. Microstructures and characteristics of the SLM samples
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the average of 5 measurements.
Fig. 2(a) shows the surface morphology of the samples prior to polishing. A very clear “V” shape
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morphology is identified, which normally appears in welding processing. This phenomenon has already been found in other materials[10]. The adjacent melting tracks tightly overlap, with a width of approximately 100 µm, generating a nearly full-dense structure devoid of lacuna. Therefore, it can be identified that the Inconel 625 superalloy powder has a good forming performance. The formation of
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this morphology can be attributed to the laser heat source induced by the moving laser. The laser beam can be considered to be a point of intense heat source leading to a spindle temperature field in the molten pool[12]. The powder is continuously melted in the front of the pool and the melted liquid metal rapidly solidifies as soon as the laser moves away. Meanwhile, the uneven temperature
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distribution causes the difference between the density and surface tension. As a result, the molten pool keeps stirring and generally forms convection throughout the entire process, leading to the
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formation of “V” shape morphology. The angle of “V” morphology largely depends on the laser scanning speed. The faster the laser scans, the smaller the “V” angle becomes, and the narrower the molten pool will be. The morphology of the Y-Z section presents a “scales” shape as illustrated in Fig. 2(b). In this experiment, the layer thickness was 0.02 mm. Most of the previous layer was remelted by the following process. The scales shape matched the Gaussian distribution of laser energy. Due to most of the laser energy being focused on the laser beam center, a depression appears at the bottom of the molten pool as indicated by the arrows. If the density of laser energy increases, a key hole is likely to form which will potentially improve the joint between adjacent layers. A higher 5
ACCEPTED MANUSCRIPT magnification of molten pool boundary area is shown in Fig. 2(c), from which three different zones were observed: molten pool, fusion line and heat affect zone (HAZ). The cellular structure sizes change from approximately 0.2 to 1 µm. The microstructures in the molten are complicated due to the convection in the molten pool. In Fig. 2(c), cellular structure was found while in molten pool 2
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the cellular was different. The width of the fusion line is approximately 0.5 µm, which is as thin as one layer of sub-grain. Moreover, the temperature outside the fusion line is still very high when the powders are being melted. This causes the cellular structure boundary, which is primarily composed of Nb element dissolved into the substrate gradually. The sub-grain boundary is thus interrupted in
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this area. Fig. 2(c) also indicates that the columnar dendrite structure is highly degenerated with the second dendrite growth being entirely suppressed. Due to rapid solidification, the average primary
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dendrite arm space is approximately 0.5 µm, which is nearly 2 orders of magnitude smaller as compared with the traditional casting (100‒300 µm)[13]. Rapidly solidified dendrite space can be determined by the following equation[14]:
=
(1)
where d is the primary dendrite arm space; b and a are associated with the material constants (for
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nickel based alloys, b=1/3, a≈50 µm(K/s)[15]). Thus, the solidification cooling rate ε can be calculated, i.e. ε =106 K/s, which is consistent with the numerical simulation of the SLM process[3]. The EBSD clearly shows the texture of the samples. Fig. 3 is the inverse pole figure (IPF) of the
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Y-Z and X-Y section, in which a region with a single color represents the presence of one grain only. Hence, some larger elongation grains of several hundred micrometers in length grow cross a few
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layers. A very strong texture can be identified for the SLM sample. The long axes of most of the grains are aligned in the <001> direction and a cube texture is observed in the Y-Z section, as shown in Fig. 3(a). This phenomenon was also observed by other researchers[16]. Inconel 625 has a facecentered cubic crystal structure. The preferred orientation is <001>[17]. During the SLM process, the laser melts the powder on the top of the sample but the substrate or the layer that has solidified still remains cold. So, the temperature gradient forms from the top to the bottom. The dendrites get the priority to grow; as a result, the angle between <001> and the direction of heat flux is minimum. Meanwhile, the maximum temperature in the molten pool can reach 3000 K[18], leading to superheat. 6
ACCEPTED MANUSCRIPT These further result in the difficulty to form homogeneous nucleation in the molten pool, which can enhance the growth of columnar crystal in the molten pool. Under this condition, the columnar grain will grow from the bottom of the molten pool towards the top. However, the interaction of the laser and powder, liquid metal, surrounding atmosphere, solidified base metal and un-melted metal
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powders can collaboratively form a large temperature gradient, which may result in surface tension gradient and Marangoni convection[10]. The velocity of convection flow in the melt pool can reach tens of meter per second according to the simulation result[19]. The convection in the molten pool can break the original growth law of columnar crystal. Parts of dendrites change growth directions,
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causing the formation of steering dendrite. Thus, some grains change the grow directions rather than always being towards <001>. Recently, some researchers used two yttrium fiber lasers with the
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power of 400 W and 1000 W, respectively to melt the powder. They found that the texture was mostly influenced by laser energy input condition[20]. Under a laser power of 1000 W, even more strong texture is obtained. For the texture that has a significant impact on the mechanical properties,
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this phenomenon should be considered in the future.
3.3. Influence of heat treatment on the microstructure
The hardness variation of the SLM samples after annealing at different temperatures is shown in
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Fig. 4. Since the cooling rate in the SLM process is very high (about 106 K/s), most of the strengthening elements such as Mo, Nb remain in the Ni matrix. A lager lattice distortion is caused
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by these point defects. The hardness of the original SLM samples is 343 HV, which is higher than that of the forged samples (typically 305 HV). Although Inconel 625 belongs to solid solution strengthening superalloy, phase transition also takes place during the heat treatment. The hardness of the sample is slightly lower than that of SLM state due to the release of the residual stress after annealing at 700 °C. The samples show the increase of the hardness when annealing at 800 and 900 °C due to the precipitation of δ (Ni3Nb). δ is an orthorhombic phase with lattice constant a = 0.574 nm, b = 0.422 nm, c = 0.454 nm. The mismatch with the Ni matrix is relatively large; thereby, the hardness increases with the heat treatment at 800 and 900 °C. When the annealing temperature is 7
ACCEPTED MANUSCRIPT above 1000 °C, the dissolution of δ phase into the matrix and the decrease of the lattice distortion cause the decrease of hardness rapidly. The phase transition will be discussed in the following in detail. Fig. 5(a) shows that the microstructure of the SLM formed parts is austenite; no carbides and
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other phases are found. The reason is as follows: as the laser beam moves very fast (600 mm/s), the solidification time is very short (often less than 1 ms). The atomic reconstruction speed in front of the short-range liquid/solid solidification is much higher than the diffusion speed, so solution atoms will be trapped by the rapid solidification, and thus the so-called “solute trapping” is formed. Most of
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solute atoms such as Cr, Mo, Nb are captured in the Ni matrix and phase transition is hard to occur. In the same way, carbides are also difficult to aggregate and precipitate.
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The XRD results under different heat treatment are shown in Fig. 5(c). The microstructures are the same as the SLM state and all of them are the solution of Cr in Ni matrix. The intermetallic phases such as γ′′ (Ni3Nb), δ (Ni3Nb), which can usually be found in casting and forging but are not detected in the XRD results because the sizes of these particles are small[9]. However, the lattice
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constants of the matrix that are calculated from the XRD results can imply the nature of precipitation. The (200) diffraction peak in Fig. 5(c) was used to calculate the lattice parameters according to Bragg diffraction. The equations are given as follows: (2)
= ⁄√ℎ + +
(3)
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2 sin =
Table 1 lists the lattice constants of the samples at different states. For the SLM state, most of the
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strengthening elements such as Nb, Mo are embedded into the matrix, causing large lattice distortion, and the lattice constant reaches the maximum. When the samples are treated at 700 °C, the Nb element precipitates to form γ′′ (Ni3Nb) phase, and the lattice constants decline accordingly. When the temperature rises to 850 °C, the γ′′ phase transforms to a new stable structure of orthorhombic δ (Ni3Nb). Moreover, when the samples are treated at 1000 °C, the strengthening elements are gradually dissolved into the matrix, resulting in the increase in lattice constant. However, the opposite results were obtained. This is mainly caused by a large number of carbides precipitating from the matrix at 1000 °C. An obvious peak representing the MC carbides can be observed in the 8
ACCEPTED MANUSCRIPT XRD results. The carbide element precipitates from the matrix, making the lattice constant increase. Fig. 5(b) shows the diffraction peak positions of the (200) crystal plane in different heat treatment processes. The changes of the integral line-breadth can be detected. The integral line-breadth becomes narrow gradually from SLM state, 700 °C, 1000 to 1150 °C, which means that the crystal
the space of the austenite phase crystal plane has increased.
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size becomes larger whilst the angles of the diffraction peaks become larger, which demonstrates that
Fig. 6 shows the microstructures of the samples that have undergone air cooling for 1 h after being heated at 1000, and 1150 °C, respectively. At 700 °C, the microstructure is the same as that of
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the SLM state. When the temperature reaches 1000 °C, the homogenization of the elements takes place and large amount of fine grains form, causing the original molten pool boundary gradually
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disappears. In addition, a number of particulates appear along the grain boundaries (as shown in Fig. 6(c)). The analysis of energy spectrum demonstrates that the particles, which are pointed by arrows have high content of Nb, Mo, and C. Mo and Nb have strong affinities to form carbides with carbon. The total content of Mo and Nb is approximately 45% and carbon content is nearly 6% in weight.
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The result is in consistent with the carbon content of MC carbide. The NbC, MoC carbide particulates distribute around the grain boundaries. These carbides play a role in pinning the migration of grain boundaries, thereby increasing the strength of the alloy. When the samples are heated at 1150 °C for 1 h, the grains grow obviously, and the boundary carbides still exist and
forming process.
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gradually grow up. However, the grain sizes are still much smaller than those of the traditional
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Two phenomena were identified. First, after heat treatment most of the grains have an obvious preferred orientation. In the middle of the molten pool, the grain grows and finally becomes a rectangular shape and is confined in the molten pool. Almost all of them have the same orientation. This phenomenon is caused by heat conduction in the SLM process. Numerous small grains are formed along the boundary of the molten pool. This is because impurities are mainly concentrated in the pool boundary during the solidification process. These impurities provide heterogeneous cores for the grains; as a result small grains grow up along the boundary. The second phenomenon is the generated zigzag grain boundary after heat treatment as shown in Fig. 6(c). The zigzag grain 9
ACCEPTED MANUSCRIPT boundary is significant for the alloy to improve the ductility, toughness and comprehensive performance[21,22]. Although there are many different theories to explain the forming of the zigzag grain boundary[23,24], there are no explanations for Inconel 625. A new explanation was developed to explain the zigzag grain boundary for Inconel 625, as plotted in Fig. 7. The formation of the zigzag
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largely results from the existence of MC carbides. The bending degree of the grain boundary is affected by the carbide orientation, carbon particle spacing and carbides size. MC carbides with bigger sizes may cause severer bending. If adjacent MC carbides have the same orientation, there will be a straight grain boundary rather than a zigzag grain boundary. This model can be used to
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explain the zigzag formation of the SLM parts after heat treatment.
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4. Conclusions
In summary, Inconel 625 superalloy samples have been successfully fabricated. Various methods were applied to exploring the microstructures of the as-fabricated samples and the samples
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after heat treatment. Conclusions can be summarized below:
(1) SEM results show that the molten pool is composed of elongated columnar crystal. The primary dendrite arm space is about 0.5 µm. The cooling rate is about 106 K/s. Due to the thermal history, a strong texture was formed from the substrate to the top of the sample.
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(2) The original microstructure of the SLM manufactured Inconel 625 superalloy is austenite with high solution distortions. No carbides or phases were precipitated. The hardness of the
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SLM fabricated parts was very high (254±6) due to the very fine microstructure with higher degree of supersaturation of strengthening elements. (3) After heat treatment, the microstructures are similar to the SLM primitive state, although the lattice constant descends slightly. The decrease is primarily caused by a large number of NbC and MoC carbides precipitated from the matrix. (4) The grains have a clear preferred orientation, which has a relationship with the heat conduction after heat treatment. Due to the influence of MC carbides, zigzag grain boundary is formed, which is crucial for the alloy. 10
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Acknowledgements This study was funded by the National Science & Technology Pillar Program of China (Granted No. 2012BAF08B00), the National Natural Science Foundation of China (Granted Nos. 51375189
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and 51375188), the independent R&D subjects of Huazhong University of Science and Technology and the State Key Laboratory of Materials Processing and Die & Mold Technology. This study was also co-funded by China Scholarship Council (CSC).
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[19] D. Dai, D. Gu, Mater. Des. 55 (2014) 482‒491. [20] T. Niendorf, S. Leuders, A. Riemer, H.A. Richard, T. Tröster, D. Schwarze, Metall. Mater. Trans. B 44 (2013) 794‒796. [21] Z. Mu, R. Ye, L. Gao, Mater. Sci. Technol. 4 (1988) 540‒547.
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[23] O. Miyagawa, M. Yamamoto, M. Kobayashi, in: Proceedings of the Third International Symposium on Metallurgy and Manufacturing of Superalloys, 1976, pp. 245‒254.
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[24] R.F. Decker, C.T. Sims, The Superalloys, John Wiley & Sons, Inc., New York, 1972, p. 33.
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ACCEPTED MANUSCRIPT Figure and table captions Table 1 Lattice constants of the specimens in different states
Fig. 1. (a) SEM morphology of Inconel 625 powder, (b) illustration of the bidirectional scan mode.
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The arrows indicate the movement of the laser. Fig. 2. SEM micrographs of SLM samples: (a) top surface morphology of the samples (unpolished), (b) “scales” shape of Y-Z section, (c) high magnification picture of the molten pool boundary. The lines present the three different areas: HAZ, fusion line and molten pool.
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Fig. 3. Microstructure morphology of the SLM formed Inconel 625 superalloy: (a) inverse pole figure (IPF) colored map of the Y-Z section, (b) inverse pole figure (IPF) colored map of the X-Y
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section, (c) (100) pole figure of the IPF image of the SLM formed Inconel 625 part, (d) the index map of the IPF and the reference coordinate.
Fig. 4. Hardness of the samples after annealing at different temperatures. Fig. 5. (a) X-ray diffraction of the SLM deposited sample, (b) diffraction peak positions of (200)
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crystal plane under different heat treatment processes, (c) X-ray diffraction of the SLM deposited sample at different temperatures.
Fig. 6. Microstructure evolutions of X-Y plane after heat treatment: (a) optical metallographic image of the specimen after 1000 °C, (b) optical metallographic image of the specimen after 1150 °C, (c)
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magnified view showing the grain boundary structure and the carbides, (d) inverse pole figure (IPF) colored map of the X-Y section.
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Fig. 7. Model of the segment of grain boundary in MC zone.
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ACCEPTED MANUSCRIPT Table and Figure List Table 1 Lattice constant (nm)
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0.3604±0.0001
SLM + 700 °C, 1 h, AQ*
0.3599±0.0002
SLM + 1000 °C, 1 h, AQ
0.35972±0.00001
SLM + 1150 °C, 1 h, AQ
0.3595±0.00001
*
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Note: AQ: water quenching
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Heat treatment
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