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Effect of heat treatment on the microstructure and mechanical properties of maraging steel by selective laser melting
Materials Science & Engineering A 760 (2019) 105–117
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Effect of heat treatment on the microstructure and mechanical properties of maraging steel by selective laser melting
T
Yuchao Baia,b, Di Wanga, Yongqiang Yanga,*, Hao Wangb,** a b
School of Mechanical and Automotive Engineering, South China University of Technology, Guangzhou, 510640, China Department of Mechanical Engineering, Faculty of Engineering, National University of Singapore, 9 Engineering Drive 1, Singapore, 117575, Singapore
ARTICLE INFO
ABSTRACT
Keywords: Maraging steel Selective laser melting Heat treatment Microstructure Mechanical properties
This work investigates the solution phenomenon, aging behavior and room-temperature mechanical properties of maraging steel manufactured by selective laser melting (SLM). Different heat treatment experiments, including solution treatment (ST), direct aging treatment (DAT) and solution + aging treatment (SAT) are designed. Microstructure analysis indicates that ST and SAT will eliminate the cellular and lath structures, but DAT has little effect on these. The content of austenite increases with the addition of DAT temperature and holding time. While austenite is almost undetectable in ST and SAT samples. Meanwhile, both the elongation and toughness of the samples with DAT gain a slight improvement with the temperature increasing. Importantly, DAT yields similar microhardness, tensile strength and impact toughness to SAT, although the resultant microstructures are completely different. The results demonstrate that DAT can achieve the similar mechanical properties to SAT samples. Samples with high mechanical properties (microhardness of 653.93 HV and ultimate strength of 2126.30 MPa) have been obtained by DAT at 520 °C for 6 h as well as solution treatment at 900 °C for 1 h and aging treatment at 520 °C for 6 h. This investigation reveals the evolution regularity of microstructure, microhardness, tensile performance and impact toughness of maraging steel manufactured by SLM after different heat treatments.
1. Introduction Selective laser melting (SLM) can manufacture high relative density metal parts through a layer-by-layer method using laser beam [1-4], which is absolutely different from traditional casting and subtractive manufacturing. The moving laser offers energy to melt metal powders fully to form a metal track and then the powders besides and on the track will also be melted to form adjacent tracks. Parts will be eventually achieved by joining these tracks. Due to the layer-by-layer manufacturing method, functional parts can be manufactured directly by SLM. Moreover, because the melt pool is very small, the solidification rate (104-106) of melt pool is very high. Equiaxed and columnar crystal structure with grain spacing less than 1 μm is obtained due to the lack of enough time for grain growth. As such, the microstructure and mechanical properties are improved significantly. Because of the advantages of SLM technology, it has broad application prospects in the fields of aerospace, automobile, medical and industry [5-8]. Many researchers have been focusing on the manufacture of metal parts using SLM, such as titanium alloy [9,10], 316L stainless steel [11], Al–Si alloy
*
[12], IN718 alloy [13] and iron-based materials [14]. In the meantime, maraging steel parts manufactured by SLM also attracted the attention of many scholars [15]. Maraging steel is one kind of ultra-high strength steels, especially after SAT, which is widely used in aircraft, aerospace, precision gear and molds [16,17]. The reason for the high strength of maraging steel is that the nanometer-sized Ni3(Mo, Ti) and Fe2Mo intermetallic particles precipitated during the aging [18-20]. These particles will significantly enhance the microhardness and strength by inhibiting the mobility of dislocations. Currently, there are some researchers studying the properties of maraging steel by SLM, such as residual stress, process parameters, microstructure, tensile strengthen, and microhardness. Kempen et al. [21] studied the tensile and impact properties under different process parameters and aging treatment. Suryawanshi et al. [22] investigated the tensile, fracture, and fatigue crack growth properties of a 3D printed maraging steel through selective laser melting in both before- and after-aging conditions. Charpy impact of maraging steel 300 was tested by Yasa et al. [23]. Some other scholars obtained the maraging steel parts with a density of more than 99% by optimizing the
Corresponding author. Corresponding author. E-mail addresses: [email protected] (Y. Yang), [email protected] (H. Wang).
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https://doi.org/10.1016/j.msea.2019.05.115 Received 25 February 2019; Received in revised form 29 May 2019; Accepted 31 May 2019 Available online 01 June 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
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process parameters of SLM, and the maximum tensile strength was more than 2000 MPa after SAT [24,25]. Because of the positive contribution of SAT to the improvements of mechanical prosperities, heat treatment is essential for maraging steel after SLM process. Currently, most researchers focus on the effect of traditional heat treatments (e.g. solution treatment for 1 h at 840 °C and aging treatment for 6 h at 480 °C) which were suitable for casting and wrought maraging steel samples. However, SLM is a very complex manufacturing technology, which leads to different microstructures and properties. Thus, the heat treatment for SLM maraging steel may differ from traditional ones. Based on the literature review, it is known that there are few papers investigating the heat treatment process based on the fact that the microstructure, composition uniformity and microsegregation of the SLM maraging steel parts were different from the casting or wrought ones. This paper aims to investigate the effect of heat treatment process, which is different from traditional method, on the martensitic transformation, the residual austenite and the precipitation behavior of maraging steel manufactured by SLM. Moreover, the evolution mechanism of microstructure and mechanical properties subject to different heat treatment methods are studied and discussed.
Table 1 Chemical composition of the 18Ni-300 maraging steel powder (in wt. %). Element
Ni
Co
Mo
Ti
Al
C
Fe
Content,%
18.2
9
5.3
0.8
0.1
< 0.2
Bal.
2. Experiments 2.1. Material and process The powder of 18Ni-300 maraging steel, provided by Carpenter Technology Corporation, was manufactured by gas atomization method, and the particle size, approximate spherical shape, ranged from 15 μm to 45 μm. The substrate was a 316L stainless steel plate. The micrograph and particle size distribution of the 18Ni-300 powder are shown in Fig. 1, and the chemical compositions are listed in Table 1. The experiments were carried out on South China University of Technology Dimetal-100 selective laser melting machine. The key components of this system include a 200 W fiber laser (Yb: YAG, wavelength: 1075 nm, SPI), a high-speed and high-precision galvanometer scanning unit, and an f-θ lens. High-purity argon was introduced during the manufacturing process to isolate oxygen. Optimized processing parameters with 99.19% relative density were adopted to manufacture samples. The laser power was 160 W, laser space was 0.07 mm, laser scanning speed was 400 mm/s, powder layer thickness was 0.03 mm and laser scanning strategy was S-cross orthogonal scanning (as shown in Fig. 2(a)). Square samples, tensile samples and impact samples were manufactured to study the microstructure, phase, microhardness, tensile strength and impact toughness. The size of square sample is 10 mm × 10 mm × 10 mm and the top surface is etched to observe the microstructure. The size of tensile sample is show in Fig. 2(b). The impact toughness is achieved by standard V-type Charpy Impact Test according to ISO 148–3:2016(en): Sample size is
Fig. 2. (a) Schematic of S-cross orthogonal scanning strategy and (b) size of tensile test samples.
55 mm(L) × 10 mm(W) × 10 mm(H); V-notch size is 2 mm deep with 45° and 0.25 mm radius along the base. 2.2. Heat treatment The heat treatment was performed in a box-type furnace MXQ160040. In order to prevent oxidation of the samples during heating, the samples were placed in sealed glass tubes filled with high-purity argon gas. The sample was placed in the furnace at room temperature. Because of the high cooling rate of the SLM process, the grain size is very small and the microsegregation is obvious, which leads to a result that the microstructure is significantly different from the casting and wrought ones. Thus, the heat treatment process should be different from traditional SAT, and a new scheme is needed. To study the heat treatment process, the differential scanning calorimeter (DSC) analysis of the as-built and solution-treated samples was conducted using STA 449C synchronous thermal analyzer at a rate of 10 °C/min in an argon atmosphere, from which the heat treatment schemes were set as Tables
Fig. 1. 18Ni-300 maraging steel powder (a) and particle size distribution (b).
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Table 2 Heat treatment scheme of maraging steel based on SLM process. Experiment No.
1 2 3 4 5
ST
DAT
SAT
Heating rate°C/min
Temperature/°C
Holding time/h
Temperature/°C
Holding time/h
Temperature/°C
Holding time/h
780 840 900 960 1020
1 1 1 1 1
400 440 480 520 560
6 6 6 6 6
900 900 900 900 900
6 6 6 6 6
2 and 3. In order to ensure the reliability of the data, three samples were deposited for each heat treatment process.
+ + + + +
400 440 480 520 560
10 10 10 10 10
peaks #5 and #6 indicate that the powders have been treated through SAT. According to the DSC curve, it can be concluded that the aging temperature range for the SLM maraging steel is 400–600 °C, and the solution temperature is 600–800 °C. However, it is known that the microsegregation of samples of SLM parts is pronounced, and the elements need high temperature to diffuse. Thus, in order to ensure a more effective solution treatment, the solution temperature is elevated to 780–1020 °C in this paper.
2.3. Characterization of microstructure and mechanical properties The square samples were etched for 2 min using dilute aqua regia (H2O: HNO3: HCl = 6: 1: 3), and then rinsed with alcohol and dried. The microstructure was observed by Quanta-650 Scanning Electron Microscopy (SEM). The micro-Vickers hardness of samples was measured under the load of 1.96 N (200 g) using DHV-10 00Z according to the ASTM E384-17 standard. The tensile strength, yield strength and elongation were tested by CMT5105. The impact test of samples was tested by JBW-750YD impact test machine. The fracture surface analysis was conducted using SEM.
3.2. Microstructure characterization The scanning electron microscopy micrographs of the samples under different heat treatment conditions are presented in Fig. 4(a)-(r). Fig. 4(a) shows the microstructure of the as-built sample, in which melt boundary can be clearly seen (red circle). There are different morphologies on two sides of the boundary with small angle bunches consisting of strips (blue arrows) and fine cellular structures (orange arrow). After ST, the boundaries, strips and cells disappear gradually with the incense of temperature and holding time, which is shown in Fig. 4(b)-(g). At low ST temperature (ST780-1 h), the bigger size grain boundaries still exist (red circle) and some white particles (orange arrows) can be seen in the boundaries. With the addition of temperature, the microstructure shows intertwined large slats which is shown by orange and white arrows in Fig. 4(c). This is because that the high temperature during ST causes the growth of austenite grain, which leads to bigger martensite laths after ST. Moreover, improving the holding time during ST shows little effect on the microstructure (Fig. 4(e)-(g)). Fig. 4(h)-(m) present the microstructure of DAT samples at different temperatures and holding time, respectively. The results are distinct from those obtained by ST. The strips (blue arrows), melt boundary (shown by red circles and arrows) and cellular structures (orange arrows) do not completely disappear, but become blurred with the temperature and holding time rising. In Fig. 4(h), the morphology similar to that of the as-built sample could be seen clearly at DAT400-6 h. However, the melt boundary starts to dissolve at DAT520-6 h, and then the long strips and cell walls are broken into short strips and spherical particles when the temperature rises to 560 °C. The similar result also appears in the morphologies of samples at 520 °C with the holding time rising from 1 h to 12 h. Another significant change is that the residual cell walls and strips become much thinner. Fig. 4(n)-(r) display other morphologies which are different from
3. Results 3.1. Thermodynamic characteristics DSC characteristic shows the endothermic and exothermic phenomena of materials during heating, which is related to the phase transformation. The DSC analysis of the as-built and solution-treated samples was investigated first, as shown in Fig. 3. It is worth noting that the difference between the two curves is different from the report by other researchers [26]. The DSC curve for the as-built sample exhibits four peaks comparing to that for the powder with only two peaks. The first exothermic peak (peak #1) can be attributed to the formation of carbide, precipitation of coherent precipitation zones or recovery of martensite [27,28], and the next exothermic peak #2 is usually associated with grain growth of retained austenite and the formation of the main intermetallic precipitates, such as Ni3(Ti, Mo) phases followed by the decomposition of Fe2Mo [29]. Tewari et al. [30] found that with the temperature and holding time rising, ε, ω, Fe2Mo and Ni3 (Ti, Mo) appeared in sequence from 675 to 825 K. And the latter two peaks are corresponding to the solution process. The endothermic peaks #3 and #4 are in the high-temperature section. The former may be related to the phase transition from α-phase (martensite, BCC-body cubic centered) to γ-phase (austenite, FCC-face cubic centered) and the latter is considered to be caused by the decomposition of precipitates [31]. Some disturbances occur at high temperatures in the DSC curve for the as-built sample (peak #7), which may be related to the inclusion of impurities like oxides during the molding process. The endothermic Table 3 Heat treatment scheme of maraging steel based on SLM process. Experiment No.
1 2 3 4 5
ST
DAT
SAT
Heating rate°C/min
Temperature/°C
Holding time/h
Temperature/°C
Holding time/h
Temperature/°C
Holding time/h
900 900 900 900 900
0.25 0.5 1 2 4
520 520 520 520 520
1 3 6 9 12
900 900 900 900 900
6 6 6 6 6
107
+ + + + +
520 520 520 520 520
10 10 10 10 10
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Fig. 3. DSC curves for maraging steel powder and sample manufactured by SLM.
Fig. 4. SEM images showing the microstructure of SLM-formed maraging steel under different heat treatments: (a) as-built; (b) ST780-1 h; (c) ST900-1 h; (d) ST10201 h; (e) ST900-0.5 h; (f) ST900-2 h; (g) ST900-4 h; (h) DAT400-6 h; (i) DAT480-6 h; (j) DAT520-6 h; (k) DAT520-1 h; (l) DAT520-3 h; (m) DAT520-12 h; (n)SAT4006 h; (o) SAT480-6 h; (p) SAT560-6 h; (q) SAT520-1 h; (r) SAT520-12 h.
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Fig. 4. (continued)
the ones of ST and DAT. They consist of packets of martensites at SAT400-6 h within prior-austenite grains as shown by green circles and arrows in Fig. 4(n). The grain boundaries indicate that the martensite packets do not stretch out of the prior-austenite grain boundary. The microstructure of other metal alloys like Ti6Al4V also consists of martensite at 780 °C for 2 h, but no boundaries are visible [32]. Fig. 4(o) and (q) present micrographs of the samples at SAT480-6 h and SAT520-1 h, respectively. It can be seen that the martensite laths become longer and wider (shown by orange arrows). However, the boundary is still faintly visible. But when the temperature rises to 560 °C, or the holding time increases to 12 h, the martensitic laths disappear. Instead, many irregular bright bars are embedded in the dark matrix which are shown by white lines in Fig. (p) and (r). Fig. 5 shows the low magnification light optical micrographs ( × 100) of ST900-1 h, SAT440-6 h and SAT560-6 h, respectively. It can
be seen that the microstructure is lath martensite (in red circle) with small size after solution treatment at 900 °C for 1 h (Fig. 5(a)). After aging treatment at 440 °C for 6 h, there are little changes in the microstructure, as shown by red circle in Fig. 5(b). But with the aging temperature increasing, the size of martensite becomes much bigger and comparable to the spacing of laths (red circle in Fig. 5(c)). 3.3. XRD analysis The XRD patterns of the SAT samples are shown in Fig. 6. It is evident that α-phase (BCC) and γ-phase (FCC) coexist in the as-built sample, which indicates that the FCC structure has been produced during SLM. With the solution treatment temperature increasing, the (200) peak of γ-phase disappears above 840 °C and the intensity of (200) peak of α-phase decreases. On the contrary, the intensity of (211)
Fig. 5. Light optical micrographs: (a) ST900-1 h; (b) SAT440-6 h; (c) SAT560-6 h. 109
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Fig. 6. The XRD spectra of maraging steel samples: (a) ST at different temperatures; (b) ST at different holding time; (c) DAT at different temperatures; (d) DAT at different holding time; (e) SAT at different temperatures; (f) SAT at different holding time.
peak of α-phase increases slightly. The same result comes from ST at 900 °C with different holding time, as shown in Fig. 6(b). It can be concluded that ST makes γ phase transfer to α phase to achieve full martensite matrix and the martensite phase preferentially grows in the (211) direction with the temperature and holding time increasing. Fig. 6(c) and (d) show the patterns of DAT at different temperatures and holding time, respectively. The most apparent feature suggests an increase in the intensity of (200) peak of γ-phase. Especially, when the temperature reaches to 560 °C, the (111) peak and (220) peak of γphase appear. This indicates that the BCC-matrix changes into FCCmatrix during DAT and high temperature induces austenite phase
growth in other directions. Following the increasing of holding time, only the intensity of (200) peak of γ-phase increases, and there exist no other peaks. The result indicates that prolonging holding time can only promote the γ phase increasing in the direction of (200). The XRD patterns of SAT samples are also investigated, as shown in Fig. 6 (e) and (f). They are similar to the ones of ST, but different from the ones of DAT. No peaks of γ phase exist in the SAT samples, which indicates aging treatment will not lead to the γ-phase transformation after solution treatment, including increasing temperature (560 °C) and prolonging holding time (12 h) [24]. In order to identify the austenite content and its changes, Eq. (1) is 110
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Fig. 7. The content of austenite of DAT maraging steel samples: (a) DAT at different holding time; (b) DAT at different temperatures.
adopted [33], where the Ifcc represents the (110) peak area integration of α phase, the Ibcc represents the (200) peak area integration of γ phase, and the RIRfcc and RIRbcc are collected from PDF-2004 card.
f fcc =
time. However, the growth rate is relatively stable comparing to aging temperature with content of γ phase increasing to 12.3% at 12 h. 3.4. Mechanical properties
Ifcc / RIRfcc Ifcc / RIRfcc + Ibcc / RIRbcc
(1)
3.4.1. Microhardness An investigation on microhardness of the as-built, ST, DAT and SAT samples has been conducted. From Fig. 8(a) and (b), it can be seen that the microhardness of solution treatment at different temperature (ST-A) samples decreases gradually and levels up slightly after reaching its minimum with the temperature increasing when the holding time is kept at 1 h. However, the microhardness of solution treatment with different holding time (ST-B) samples drop sharply comparing with the as-built one with the holding time increasing when temperature is kept at 900 °C. The above results indicate that ST could reduce the microhardness of maraging steel manufactured by SLM by up to 15.7%. Fig. 8(c) and (d) present the microhardness of DAT and SAT samples, in which the DAT-A, DAT-B, SAT-A and SAT-B mean direct aging
Fig. 7(a) presents the content of austenite after DAT at different temperatures. According to Eq. (1), the content of γ phase is 6.2% comparing with α phase (93.8%) in the as-built sample. After direct aging at 400 °C for 6 h, the content of γ phase increases to 6.9%, and the content of α phase decreases correspondingly. As the aging temperature increases, the content of γ phase increases gradually, which means aging temperature promotes austenite transformation. When the temperature rises to 560 °C, the content of γ phase increases to 17.9% sharply. This indicates high aging temperature may lead to the rapid growth of γ phase beside transformation. Fig. 7(b) shows the content of austenite at 520 °C for holding time from 0 h (as-built) to 12 h. It can be seen that the content of γ phase increases with the extension of holding
Fig. 8. Micro-Vickers hardness of maraging steel at different heat treatments: (a) (b) ST; (c) (d) DAT and SAT. 111
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Fig. 9. Stress-strain curves at different heat treatments: (a) (b) ST; (c) (d) DAT; (e) (f) SAT.
treatment at different temperature, direct aging treatment with different holding time, solution + aging treatment at different temperature and solution + aging treatment with different holding time respectively. The microhardness can be significantly improved after DAT and SAT. And the microhardness of DAT samples increases first and then drop slightly with the aging temperature and holding time rising. The hardness reaches the maximum value 653.94 HV at 520 °C for 6 h. The similar results also appear in the SAT samples, which indicates that the solution treatment conducted before aging treatment has little effect on microhardness.
3.5. Temperature-dependent tensile property Fig. 9 shows the room-temperature tensile curves under ST, DAT and SAT to highlight the difference in tensile behavior under different heat treatments. Comparing with the as-built sample, the tensile strengths decrease slightly under ST (Fig. 9(a) and (b)). However, the tensile strengths increase sharply under DAT and SAT, as shown in Fig. 9(c)-(f). In detail, the tensile strength drops first and then rises slightly with the increase of ST temperature, but the fracture elongations are nearly the same. Conversely, the tensile strengths are nearly the same and the fracture elongation rises first and then drops with the 112
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Fig. 10. Fracture morphology of tensile samples: (a) as-built; (b) ST900-2 h; (c) DAT520-6 h; (d) SAT520-6 h.
holding temperature increasing at 900 °C. For the DAT samples, the tensile strength and elongation both appear to increase first and then decrease with the temperature and holding time rising. Especially, the aging temperature shows significant influence. This is caused by the comprehensive effect of austenite and precipitated particles. The increase of austenite content can improve the elongation, but reduce the tensile strength. With the generation of precipitated particles, the tensile strength increase. But the tensile strength will drop when the particles become too big. And the big particles not only have bad effect on the tensile strength but also result in the reduction of elongation (The effect is not significant when particle is small). As for the SAT samples, the tensile strength increases and drops drastically with the increase of temperature after ST at 900 °C for 1 h, but the elongation is almost constant. However, the elongation increases first and then drops with the holding time prolonging, and the change in strength is relatively small. The above results show that the effect of heat treatment on the tensile property is different. Both DAT and SAT could significantly improve the tensile strength to the same extent, which means that aging treatment alone could enhance the tensile strength of parts manufactured by SLM, and no additional solution treatment is needed. The excellent comprehensive performance of the work material can be obtained at 520 °C for 6 h. The fracture morphologies obtained from the as-built, ST900-2 h, DAT520-6 h and SAT520-6 h samples are shown in Fig. 10(a)-(d), respectively. Apparent dimples (orange arrows) are displayed in Fig. 10(a), which indicates good plasticity corresponding to the high elongation. Meanwhile, the co-existing deep holes (blue arrow) are resulted from the SLM process caused by molten liquid shrinkage or vaporization. Fig. 10(b) presents a sizable number of dimples in the as-built sample, which demonstrates the plastic fracture and explains the reason why elongations of as-built and ST are similar. In the meantime, the particle (green arrow) embedded in a big dimple implies the presence of the
undissolved second phase. Fig. 10(c) shows the SEM image of the fracture surface of DAT520-6 h sample, in which lath packets (brown arrow) and dimples appear. It can be confirmed that the crack extends along the martensitic laths in a martensitic lath packet, which is defined by apparent boundaries. This further proves the existence of austenite (with good plasticity) in addition to martensite. Fig. 10(d) displays the fracture surface of SAT520-6 h sample. In addition to the tearing edges (yellow arrow) and small platforms (purple arrow), there are a small number of dimples, which reveals that the development of cracks was mainly ruled by a quasi-cleavage mechanism [34]. For the DAT and SAT samples, the fractures are obviously brittle with little plastic deformation. 3.6. Temperature-dependent impact toughness The toughness test results are shown in Fig. 11. It can be seen that ST can improve the toughness to a certain level of elevated temperature followed by a dip when the temperature is above 960 °C and/or the holding time exceeds 1 h, as shown in Fig. 11(a) and (b). Fig. 11(c)-(f) show that both DAT and SAT lead to the reduction of impact roughness. Though the toughness of DAT-A samples increases after 480 °C for 6 h, all the others drop drastically to a low-toughness state. Fig. 11 displays the fracture surfaces after Charpy impact testing at room temperature. It is obvious that the shear lips and fibrous zone co-exist in Fig. 12(a) and (b), while the radical zone is almost non-existent here, which indicates the good toughness of the as-built and ST900-1 h samples. The shear lips of latter are larger showing the better toughness. In Fig. 12(c) and (d), the fracture surfaces are both flat, and the shear lip and fibrous zone are not found but some oxidized particles and void defects. The flat radical zone demonstrates that the cracks grow rapidly from the Vport in impact progress. This is mainly due to the formation of fine precipitates during aging treatment in both DAT and SAT samples which limit the ductile deformation of the samples [35,36]. 113
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Fig. 11. Charpy impact test results of 18Ni-300 maraging steel: (a) (b) ST; (c) (d) DAT and SAT.
Fig. 12. Fracture surfaces of Charpy samples: (a) as built; (b) ST900; (c) DAT520-6 h; (d) SAT480-6 h.
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Fig. 13. Element map by EDS mapping of the as-built sample: (a) SEM image; (b) spectrum for Fe; (c) spectrum for Ni; (d) spectrum for Co; (e) spectrum for Mo; (f) spectrum for Ti.
much higher than that of the dark area, as shown in Fig. 13 and Table 4. The reason is that the solubility of Mo and Ti in solid is lower than that in liquid. They are forced to move to the liquid zone when the grain formation is initiated during rapid cooling process. As the neighboring grains encounter each other, the Mo-rich and Ti-rich regions are formed in the boundary. High temperature during ST makes the segregated elements (such as Mo and Ti) gain enough energy to diffuse into the interior of the grains. Thus, the white boundaries begin to break up and dissolve. The higher and longer the temperature and holding time, the faster the white boundaries disappear. This can be found in the SAT samples subject to the ST process before aging treatment. During heat treatment of AlSi10Mg prepared by SLM, similar phenomenon is also observed [39]. According to the EDS analysis of white boundaries in Fig. 13 and Table 4, it can be concluded that the white particles in Fig. 4(b) exhibit higher Mo and Ti content which could not be dissolved completely due to low temperature [40]. However, in our experiments, the white boundaries could still be seen in the DAT samples, even
Table 4 The chemical composition of different areas in Fig. 13 (a) (in wt. %). Element
Ni
Co
Mo
Ti
Al
Fe
Spectrum 1, % Spectrum 2, %
18.16 18.52
8.81 9.15
7.14 5.80
1.65 1.11
0.27 0.27
63.97 65.15
4. Discussions From the SEM images in Fig. 4, it can be seen that the microstructure changes significantly with the temperature and holding time. The main micro-features of the as-built sample consist of cells and laths structure with white boundary. This is because high cooling rate results in a large number of nucleation points to produce small grains. The feature is also observed in 316L [37], CoCrMo alloy [33] and Al–Si alloy [38]. And EDS analysis shows that the white boundary is composed of Fe, Ni, Co, Mo, Ti and Al, and the contents of Mo and Ti are 115
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though the temperature increases to 560 °C or the holding time increases to 12 h. That is because the diffusion of Mo and Ti needs much higher energy that cannot be provided by DAT. Furthermore, it is noted that the discontinuous boundaries are due to the intermetallic phases (e.g., Ni3Ti, Ni3Mo, and Ni3Al) precipitated during aging treatment [41,42]. Even though the high magnification SEM micrographs present significant difference between ST900-1 h and SAT560-6 h in Fig. 4, the low magnification optical micrographs show similar lath martensitic lath features. The main difference is that the size of martensite becomes bigger after aging treatment for 6 h. Figs. 8 and 9 show a similar change trend in the microhardness and tensile strength, including ST, DAT and SAT samples. For example, both microhardness and tensile strength decrease obviously after ST. That is because the high cooling rate leads to the rapid solidification of the molten pool to produce martensite tissue, fine grains and large residual stress, which can significantly prevent the movement of dislocation to hinder the formation and expansion of cracks. As such, the microhardness and tensile strength are improved. After ST, the microstructure begins to transform into martensite lathes, and the higher the temperature and holding time, the more the martensite formation. Martensite can improve the tensile strength and microhardness, but the results show a reduction in them. The reason is that the fine grains and residual stress disappear after ST. Their strengthening effects decrease at the same time. And the improvement from martensite transformation is lower than the deduction of disappearance of fine gains and residual stress. As a result, the microhardness and tensile strength decrease eventually. These indicate that the as-built and ST samples are mainly exhibiting ductile fractures [34]. A ductile fracture is always transgranular [43]. For maraging steel manufactured by casting and forging methods, solution treatment is essential before aging treatment to prepare the martensite matrix for precipitation strengthening with the particles, such as Fe2Mo, Ni3Ti, and Ni3Mo, which can increase the resistance of dislocation movement. However, there is a significant difference in maraging steel by SLM. Due to the rapid cooling, the martensite matrix has been achieved during SLM manufacturing process, which can be identified by the XRD analysis (Fig. 6). These provide the tissue base for the aging treatment. Therefore, theoretically speaking, it is possible to achieve high strength and microhardness after conducting aging treatment on the as-built samples without solution treatment [44]. DAT samples retain the similar microstructure to the as-built ones. But the microhardness and tensile strength of the DAT samples are very high, which are similar to SAT. From the DSC curve, it can be observed that the aging behavior occurs at relatively low temperatures comparing with ST. The cells and laths from microsegregation during SLM are reserved due to low energy density. But the Ni, Mo and Ti elements dissolved in the martensite matrix can be precipitated normally to form strengthening particles Ni3(Mo, Ti), which improve the mechanical properties significantly. Therefore, even though the microstructure is almost constant after aging treatment, DAT samples can still achieve high microhardness and tensile strength. Another thing should be noticed that almost no peaks from Fe2Mo and Ni3(Ti,Mo) intermetallic compounds can been found in the XRD test (Fig. 6) in the as-built, DAT and SAT samples. This is mainly because of the extremely small size of precipitates, low phase content, and highly dispersed distribution [24,45,46]. In summary, the microstructures and mechanical properties are very different between the as-built, ST, DAT and SAT samples. But the DAT samples show the similar microhardness, tensile strength and impact toughness to the SAT ones, despite the distinctions in microstructure. Since DAT could obtain the same mechanical properties with less postprocessing comparing to SAT, it will be very meaningful to reduce the cost of heat treatment for the maraging steel by SLM. However, more research is still needed to explore and understand the strengthening mechanism of DAT. The future work will focus on the texture, precipitation behavior and reverse austenite nucleation using transmission electron microscope (TEM) and electron backscattered diffraction (EBSD).
5. Conclusions This paper investigates the heat treatment processes associated with the selective laser melting (SLM) technology, which is different from traditional methods. A series of heat treatment experiments, including solution treatment (ST), direct aging treatment (DAT) and solution + aging treatment (SAT), have been conducted on maraging steel samples manufactured by SLM. The evolution mechanisms of microstructure, microhardness, tensile property and toughness are investigated with different heat treatment conditions. The following conclusions can be drawn from this study: (1) ST and SAT have significant effect on the microstructure evolution. When the ST temperature is at 780 °C, the austenite grain boundary and white particles can be observed. When the temperature reaches 840 °C, the cells, lathes and melt track boundary all disappear. The microstructure of DAT samples is similar to the as-built samples. With the temperature and holding time increasing, the melt track boundary dissolves gradually. As for the SAT samples, the microstructure consists of lath martensite. But when the temperature reaches 560 °C or holding time reaches to 12 h, the lath martensite disappears and is replaced by white and dark tissues. (2) Comparing with the as-built samples, ST reduces the microhardness and tensile strength, and the fracture elongation does not increase because of the disappearance of fine grains obtained in SLM. But the impact toughness is improved. Both DAT and SAT can improve the microhardness and tensile strength significantly. However, both fracture elongation and impact toughness are reduced. The best heat treatment process for high hardness and strength is DAT at 520 °C for 6 h with an achievable microhardness up to 653.93 HV and tensile strength up to 2126.30 MPa. (3) DSC curves show the phase transformation temperature of the asbuilt sample at 456 °C, 572 °C, 670 °C and 715 °C, which indicates the solution and aging temperature range. The microhardness, strength and impact properties of DAT are similar to the ones of SAT. But the microstructure is obviously different between these samples. The content of austenite increases with the temperature and holding time rising in the DAT samples. However, ST and SAT result in the elimination of austenite in the maraging steel manufactured by SLM. Acknowledgments The authors gratefully appreciate the financial support from the Natural Science Foundation of China (NSFC, No.: 51875215), Guangdong Province Science and Technology Project (No.: 2015B090921002 and 2016B09092500), Guangzhou Science and Technology Project (No.: 201704020118), Singapore Ministry of Education AcRF Tier 2 (Project No.: MOE2018-T2-1-140) and A*STAR AME IAF-PP (Grant No.: A1893a0031). References [1] W.E. Frazier, Metal additive manufacturing: a review, J. Mater. Eng. Perform. 23 (6) (2014) 1917–1928. [2] D. Herzog, V. Seyda, E. Wycisk, C. Emmelmann, Additive manufacturing of metals, Acta Mater. 117 (2016) 371–392. [3] E.O. Olakanmi, R.F. Cochrane, K.W. Dalgarno, A review on selective laser sintering/ melting (SLS/SLM) of aluminium alloy powders: processing, microstructure, and properties, Prog. Mater. Sci. 74 (2015) 401–477. [4] B. Vayre, F. Vignat, F. Villeneuve, Metallic additive manufacturing: state-of-the-art review and prospects, Mechanics & Industry 13 (2) (2012) 89–96. [5] D. Dai, D. Gu, Thermal behavior and densification mechanism during selective laser melting of copper matrix composites: simulation and experiments, Mater. Des. 55 (2014) 482–491. [6] Q. Zhang, J. Chen, P. Guo, et al., Texture and microstructure characterization in laser additive manufactured Ti–6Al–2Zr–2Sn–3Mo–1.5 Cr–2Nb titanium alloy, Mater. Des. 88 (2015) 550–557. [7] C.Y. Yap, C.K. Chua, Z.L. Dong, et al., Review of selective laser melting: materials
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Effect of heat treatment on the microstructure and mechanical properties of maraging steel by selective laser melting
T
Yuchao Baia,b, Di Wanga, Yongqiang Yanga,*, Hao Wangb,** a b
School of Mechanical and Automotive Engineering, South China University of Technology, Guangzhou, 510640, China Department of Mechanical Engineering, Faculty of Engineering, National University of Singapore, 9 Engineering Drive 1, Singapore, 117575, Singapore
ARTICLE INFO
ABSTRACT
Keywords: Maraging steel Selective laser melting Heat treatment Microstructure Mechanical properties
This work investigates the solution phenomenon, aging behavior and room-temperature mechanical properties of maraging steel manufactured by selective laser melting (SLM). Different heat treatment experiments, including solution treatment (ST), direct aging treatment (DAT) and solution + aging treatment (SAT) are designed. Microstructure analysis indicates that ST and SAT will eliminate the cellular and lath structures, but DAT has little effect on these. The content of austenite increases with the addition of DAT temperature and holding time. While austenite is almost undetectable in ST and SAT samples. Meanwhile, both the elongation and toughness of the samples with DAT gain a slight improvement with the temperature increasing. Importantly, DAT yields similar microhardness, tensile strength and impact toughness to SAT, although the resultant microstructures are completely different. The results demonstrate that DAT can achieve the similar mechanical properties to SAT samples. Samples with high mechanical properties (microhardness of 653.93 HV and ultimate strength of 2126.30 MPa) have been obtained by DAT at 520 °C for 6 h as well as solution treatment at 900 °C for 1 h and aging treatment at 520 °C for 6 h. This investigation reveals the evolution regularity of microstructure, microhardness, tensile performance and impact toughness of maraging steel manufactured by SLM after different heat treatments.
1. Introduction Selective laser melting (SLM) can manufacture high relative density metal parts through a layer-by-layer method using laser beam [1-4], which is absolutely different from traditional casting and subtractive manufacturing. The moving laser offers energy to melt metal powders fully to form a metal track and then the powders besides and on the track will also be melted to form adjacent tracks. Parts will be eventually achieved by joining these tracks. Due to the layer-by-layer manufacturing method, functional parts can be manufactured directly by SLM. Moreover, because the melt pool is very small, the solidification rate (104-106) of melt pool is very high. Equiaxed and columnar crystal structure with grain spacing less than 1 μm is obtained due to the lack of enough time for grain growth. As such, the microstructure and mechanical properties are improved significantly. Because of the advantages of SLM technology, it has broad application prospects in the fields of aerospace, automobile, medical and industry [5-8]. Many researchers have been focusing on the manufacture of metal parts using SLM, such as titanium alloy [9,10], 316L stainless steel [11], Al–Si alloy
*
[12], IN718 alloy [13] and iron-based materials [14]. In the meantime, maraging steel parts manufactured by SLM also attracted the attention of many scholars [15]. Maraging steel is one kind of ultra-high strength steels, especially after SAT, which is widely used in aircraft, aerospace, precision gear and molds [16,17]. The reason for the high strength of maraging steel is that the nanometer-sized Ni3(Mo, Ti) and Fe2Mo intermetallic particles precipitated during the aging [18-20]. These particles will significantly enhance the microhardness and strength by inhibiting the mobility of dislocations. Currently, there are some researchers studying the properties of maraging steel by SLM, such as residual stress, process parameters, microstructure, tensile strengthen, and microhardness. Kempen et al. [21] studied the tensile and impact properties under different process parameters and aging treatment. Suryawanshi et al. [22] investigated the tensile, fracture, and fatigue crack growth properties of a 3D printed maraging steel through selective laser melting in both before- and after-aging conditions. Charpy impact of maraging steel 300 was tested by Yasa et al. [23]. Some other scholars obtained the maraging steel parts with a density of more than 99% by optimizing the
Corresponding author. Corresponding author. E-mail addresses: [email protected] (Y. Yang), [email protected] (H. Wang).
**
https://doi.org/10.1016/j.msea.2019.05.115 Received 25 February 2019; Received in revised form 29 May 2019; Accepted 31 May 2019 Available online 01 June 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
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process parameters of SLM, and the maximum tensile strength was more than 2000 MPa after SAT [24,25]. Because of the positive contribution of SAT to the improvements of mechanical prosperities, heat treatment is essential for maraging steel after SLM process. Currently, most researchers focus on the effect of traditional heat treatments (e.g. solution treatment for 1 h at 840 °C and aging treatment for 6 h at 480 °C) which were suitable for casting and wrought maraging steel samples. However, SLM is a very complex manufacturing technology, which leads to different microstructures and properties. Thus, the heat treatment for SLM maraging steel may differ from traditional ones. Based on the literature review, it is known that there are few papers investigating the heat treatment process based on the fact that the microstructure, composition uniformity and microsegregation of the SLM maraging steel parts were different from the casting or wrought ones. This paper aims to investigate the effect of heat treatment process, which is different from traditional method, on the martensitic transformation, the residual austenite and the precipitation behavior of maraging steel manufactured by SLM. Moreover, the evolution mechanism of microstructure and mechanical properties subject to different heat treatment methods are studied and discussed.
Table 1 Chemical composition of the 18Ni-300 maraging steel powder (in wt. %). Element
Ni
Co
Mo
Ti
Al
C
Fe
Content,%
18.2
9
5.3
0.8
0.1
< 0.2
Bal.
2. Experiments 2.1. Material and process The powder of 18Ni-300 maraging steel, provided by Carpenter Technology Corporation, was manufactured by gas atomization method, and the particle size, approximate spherical shape, ranged from 15 μm to 45 μm. The substrate was a 316L stainless steel plate. The micrograph and particle size distribution of the 18Ni-300 powder are shown in Fig. 1, and the chemical compositions are listed in Table 1. The experiments were carried out on South China University of Technology Dimetal-100 selective laser melting machine. The key components of this system include a 200 W fiber laser (Yb: YAG, wavelength: 1075 nm, SPI), a high-speed and high-precision galvanometer scanning unit, and an f-θ lens. High-purity argon was introduced during the manufacturing process to isolate oxygen. Optimized processing parameters with 99.19% relative density were adopted to manufacture samples. The laser power was 160 W, laser space was 0.07 mm, laser scanning speed was 400 mm/s, powder layer thickness was 0.03 mm and laser scanning strategy was S-cross orthogonal scanning (as shown in Fig. 2(a)). Square samples, tensile samples and impact samples were manufactured to study the microstructure, phase, microhardness, tensile strength and impact toughness. The size of square sample is 10 mm × 10 mm × 10 mm and the top surface is etched to observe the microstructure. The size of tensile sample is show in Fig. 2(b). The impact toughness is achieved by standard V-type Charpy Impact Test according to ISO 148–3:2016(en): Sample size is
Fig. 2. (a) Schematic of S-cross orthogonal scanning strategy and (b) size of tensile test samples.
55 mm(L) × 10 mm(W) × 10 mm(H); V-notch size is 2 mm deep with 45° and 0.25 mm radius along the base. 2.2. Heat treatment The heat treatment was performed in a box-type furnace MXQ160040. In order to prevent oxidation of the samples during heating, the samples were placed in sealed glass tubes filled with high-purity argon gas. The sample was placed in the furnace at room temperature. Because of the high cooling rate of the SLM process, the grain size is very small and the microsegregation is obvious, which leads to a result that the microstructure is significantly different from the casting and wrought ones. Thus, the heat treatment process should be different from traditional SAT, and a new scheme is needed. To study the heat treatment process, the differential scanning calorimeter (DSC) analysis of the as-built and solution-treated samples was conducted using STA 449C synchronous thermal analyzer at a rate of 10 °C/min in an argon atmosphere, from which the heat treatment schemes were set as Tables
Fig. 1. 18Ni-300 maraging steel powder (a) and particle size distribution (b).
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Table 2 Heat treatment scheme of maraging steel based on SLM process. Experiment No.
1 2 3 4 5
ST
DAT
SAT
Heating rate°C/min
Temperature/°C
Holding time/h
Temperature/°C
Holding time/h
Temperature/°C
Holding time/h
780 840 900 960 1020
1 1 1 1 1
400 440 480 520 560
6 6 6 6 6
900 900 900 900 900
6 6 6 6 6
2 and 3. In order to ensure the reliability of the data, three samples were deposited for each heat treatment process.
+ + + + +
400 440 480 520 560
10 10 10 10 10
peaks #5 and #6 indicate that the powders have been treated through SAT. According to the DSC curve, it can be concluded that the aging temperature range for the SLM maraging steel is 400–600 °C, and the solution temperature is 600–800 °C. However, it is known that the microsegregation of samples of SLM parts is pronounced, and the elements need high temperature to diffuse. Thus, in order to ensure a more effective solution treatment, the solution temperature is elevated to 780–1020 °C in this paper.
2.3. Characterization of microstructure and mechanical properties The square samples were etched for 2 min using dilute aqua regia (H2O: HNO3: HCl = 6: 1: 3), and then rinsed with alcohol and dried. The microstructure was observed by Quanta-650 Scanning Electron Microscopy (SEM). The micro-Vickers hardness of samples was measured under the load of 1.96 N (200 g) using DHV-10 00Z according to the ASTM E384-17 standard. The tensile strength, yield strength and elongation were tested by CMT5105. The impact test of samples was tested by JBW-750YD impact test machine. The fracture surface analysis was conducted using SEM.
3.2. Microstructure characterization The scanning electron microscopy micrographs of the samples under different heat treatment conditions are presented in Fig. 4(a)-(r). Fig. 4(a) shows the microstructure of the as-built sample, in which melt boundary can be clearly seen (red circle). There are different morphologies on two sides of the boundary with small angle bunches consisting of strips (blue arrows) and fine cellular structures (orange arrow). After ST, the boundaries, strips and cells disappear gradually with the incense of temperature and holding time, which is shown in Fig. 4(b)-(g). At low ST temperature (ST780-1 h), the bigger size grain boundaries still exist (red circle) and some white particles (orange arrows) can be seen in the boundaries. With the addition of temperature, the microstructure shows intertwined large slats which is shown by orange and white arrows in Fig. 4(c). This is because that the high temperature during ST causes the growth of austenite grain, which leads to bigger martensite laths after ST. Moreover, improving the holding time during ST shows little effect on the microstructure (Fig. 4(e)-(g)). Fig. 4(h)-(m) present the microstructure of DAT samples at different temperatures and holding time, respectively. The results are distinct from those obtained by ST. The strips (blue arrows), melt boundary (shown by red circles and arrows) and cellular structures (orange arrows) do not completely disappear, but become blurred with the temperature and holding time rising. In Fig. 4(h), the morphology similar to that of the as-built sample could be seen clearly at DAT400-6 h. However, the melt boundary starts to dissolve at DAT520-6 h, and then the long strips and cell walls are broken into short strips and spherical particles when the temperature rises to 560 °C. The similar result also appears in the morphologies of samples at 520 °C with the holding time rising from 1 h to 12 h. Another significant change is that the residual cell walls and strips become much thinner. Fig. 4(n)-(r) display other morphologies which are different from
3. Results 3.1. Thermodynamic characteristics DSC characteristic shows the endothermic and exothermic phenomena of materials during heating, which is related to the phase transformation. The DSC analysis of the as-built and solution-treated samples was investigated first, as shown in Fig. 3. It is worth noting that the difference between the two curves is different from the report by other researchers [26]. The DSC curve for the as-built sample exhibits four peaks comparing to that for the powder with only two peaks. The first exothermic peak (peak #1) can be attributed to the formation of carbide, precipitation of coherent precipitation zones or recovery of martensite [27,28], and the next exothermic peak #2 is usually associated with grain growth of retained austenite and the formation of the main intermetallic precipitates, such as Ni3(Ti, Mo) phases followed by the decomposition of Fe2Mo [29]. Tewari et al. [30] found that with the temperature and holding time rising, ε, ω, Fe2Mo and Ni3 (Ti, Mo) appeared in sequence from 675 to 825 K. And the latter two peaks are corresponding to the solution process. The endothermic peaks #3 and #4 are in the high-temperature section. The former may be related to the phase transition from α-phase (martensite, BCC-body cubic centered) to γ-phase (austenite, FCC-face cubic centered) and the latter is considered to be caused by the decomposition of precipitates [31]. Some disturbances occur at high temperatures in the DSC curve for the as-built sample (peak #7), which may be related to the inclusion of impurities like oxides during the molding process. The endothermic Table 3 Heat treatment scheme of maraging steel based on SLM process. Experiment No.
1 2 3 4 5
ST
DAT
SAT
Heating rate°C/min
Temperature/°C
Holding time/h
Temperature/°C
Holding time/h
Temperature/°C
Holding time/h
900 900 900 900 900
0.25 0.5 1 2 4
520 520 520 520 520
1 3 6 9 12
900 900 900 900 900
6 6 6 6 6
107
+ + + + +
520 520 520 520 520
10 10 10 10 10
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Fig. 3. DSC curves for maraging steel powder and sample manufactured by SLM.
Fig. 4. SEM images showing the microstructure of SLM-formed maraging steel under different heat treatments: (a) as-built; (b) ST780-1 h; (c) ST900-1 h; (d) ST10201 h; (e) ST900-0.5 h; (f) ST900-2 h; (g) ST900-4 h; (h) DAT400-6 h; (i) DAT480-6 h; (j) DAT520-6 h; (k) DAT520-1 h; (l) DAT520-3 h; (m) DAT520-12 h; (n)SAT4006 h; (o) SAT480-6 h; (p) SAT560-6 h; (q) SAT520-1 h; (r) SAT520-12 h.
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Fig. 4. (continued)
the ones of ST and DAT. They consist of packets of martensites at SAT400-6 h within prior-austenite grains as shown by green circles and arrows in Fig. 4(n). The grain boundaries indicate that the martensite packets do not stretch out of the prior-austenite grain boundary. The microstructure of other metal alloys like Ti6Al4V also consists of martensite at 780 °C for 2 h, but no boundaries are visible [32]. Fig. 4(o) and (q) present micrographs of the samples at SAT480-6 h and SAT520-1 h, respectively. It can be seen that the martensite laths become longer and wider (shown by orange arrows). However, the boundary is still faintly visible. But when the temperature rises to 560 °C, or the holding time increases to 12 h, the martensitic laths disappear. Instead, many irregular bright bars are embedded in the dark matrix which are shown by white lines in Fig. (p) and (r). Fig. 5 shows the low magnification light optical micrographs ( × 100) of ST900-1 h, SAT440-6 h and SAT560-6 h, respectively. It can
be seen that the microstructure is lath martensite (in red circle) with small size after solution treatment at 900 °C for 1 h (Fig. 5(a)). After aging treatment at 440 °C for 6 h, there are little changes in the microstructure, as shown by red circle in Fig. 5(b). But with the aging temperature increasing, the size of martensite becomes much bigger and comparable to the spacing of laths (red circle in Fig. 5(c)). 3.3. XRD analysis The XRD patterns of the SAT samples are shown in Fig. 6. It is evident that α-phase (BCC) and γ-phase (FCC) coexist in the as-built sample, which indicates that the FCC structure has been produced during SLM. With the solution treatment temperature increasing, the (200) peak of γ-phase disappears above 840 °C and the intensity of (200) peak of α-phase decreases. On the contrary, the intensity of (211)
Fig. 5. Light optical micrographs: (a) ST900-1 h; (b) SAT440-6 h; (c) SAT560-6 h. 109
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Fig. 6. The XRD spectra of maraging steel samples: (a) ST at different temperatures; (b) ST at different holding time; (c) DAT at different temperatures; (d) DAT at different holding time; (e) SAT at different temperatures; (f) SAT at different holding time.
peak of α-phase increases slightly. The same result comes from ST at 900 °C with different holding time, as shown in Fig. 6(b). It can be concluded that ST makes γ phase transfer to α phase to achieve full martensite matrix and the martensite phase preferentially grows in the (211) direction with the temperature and holding time increasing. Fig. 6(c) and (d) show the patterns of DAT at different temperatures and holding time, respectively. The most apparent feature suggests an increase in the intensity of (200) peak of γ-phase. Especially, when the temperature reaches to 560 °C, the (111) peak and (220) peak of γphase appear. This indicates that the BCC-matrix changes into FCCmatrix during DAT and high temperature induces austenite phase
growth in other directions. Following the increasing of holding time, only the intensity of (200) peak of γ-phase increases, and there exist no other peaks. The result indicates that prolonging holding time can only promote the γ phase increasing in the direction of (200). The XRD patterns of SAT samples are also investigated, as shown in Fig. 6 (e) and (f). They are similar to the ones of ST, but different from the ones of DAT. No peaks of γ phase exist in the SAT samples, which indicates aging treatment will not lead to the γ-phase transformation after solution treatment, including increasing temperature (560 °C) and prolonging holding time (12 h) [24]. In order to identify the austenite content and its changes, Eq. (1) is 110
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Fig. 7. The content of austenite of DAT maraging steel samples: (a) DAT at different holding time; (b) DAT at different temperatures.
adopted [33], where the Ifcc represents the (110) peak area integration of α phase, the Ibcc represents the (200) peak area integration of γ phase, and the RIRfcc and RIRbcc are collected from PDF-2004 card.
f fcc =
time. However, the growth rate is relatively stable comparing to aging temperature with content of γ phase increasing to 12.3% at 12 h. 3.4. Mechanical properties
Ifcc / RIRfcc Ifcc / RIRfcc + Ibcc / RIRbcc
(1)
3.4.1. Microhardness An investigation on microhardness of the as-built, ST, DAT and SAT samples has been conducted. From Fig. 8(a) and (b), it can be seen that the microhardness of solution treatment at different temperature (ST-A) samples decreases gradually and levels up slightly after reaching its minimum with the temperature increasing when the holding time is kept at 1 h. However, the microhardness of solution treatment with different holding time (ST-B) samples drop sharply comparing with the as-built one with the holding time increasing when temperature is kept at 900 °C. The above results indicate that ST could reduce the microhardness of maraging steel manufactured by SLM by up to 15.7%. Fig. 8(c) and (d) present the microhardness of DAT and SAT samples, in which the DAT-A, DAT-B, SAT-A and SAT-B mean direct aging
Fig. 7(a) presents the content of austenite after DAT at different temperatures. According to Eq. (1), the content of γ phase is 6.2% comparing with α phase (93.8%) in the as-built sample. After direct aging at 400 °C for 6 h, the content of γ phase increases to 6.9%, and the content of α phase decreases correspondingly. As the aging temperature increases, the content of γ phase increases gradually, which means aging temperature promotes austenite transformation. When the temperature rises to 560 °C, the content of γ phase increases to 17.9% sharply. This indicates high aging temperature may lead to the rapid growth of γ phase beside transformation. Fig. 7(b) shows the content of austenite at 520 °C for holding time from 0 h (as-built) to 12 h. It can be seen that the content of γ phase increases with the extension of holding
Fig. 8. Micro-Vickers hardness of maraging steel at different heat treatments: (a) (b) ST; (c) (d) DAT and SAT. 111
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Fig. 9. Stress-strain curves at different heat treatments: (a) (b) ST; (c) (d) DAT; (e) (f) SAT.
treatment at different temperature, direct aging treatment with different holding time, solution + aging treatment at different temperature and solution + aging treatment with different holding time respectively. The microhardness can be significantly improved after DAT and SAT. And the microhardness of DAT samples increases first and then drop slightly with the aging temperature and holding time rising. The hardness reaches the maximum value 653.94 HV at 520 °C for 6 h. The similar results also appear in the SAT samples, which indicates that the solution treatment conducted before aging treatment has little effect on microhardness.
3.5. Temperature-dependent tensile property Fig. 9 shows the room-temperature tensile curves under ST, DAT and SAT to highlight the difference in tensile behavior under different heat treatments. Comparing with the as-built sample, the tensile strengths decrease slightly under ST (Fig. 9(a) and (b)). However, the tensile strengths increase sharply under DAT and SAT, as shown in Fig. 9(c)-(f). In detail, the tensile strength drops first and then rises slightly with the increase of ST temperature, but the fracture elongations are nearly the same. Conversely, the tensile strengths are nearly the same and the fracture elongation rises first and then drops with the 112
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Fig. 10. Fracture morphology of tensile samples: (a) as-built; (b) ST900-2 h; (c) DAT520-6 h; (d) SAT520-6 h.
holding temperature increasing at 900 °C. For the DAT samples, the tensile strength and elongation both appear to increase first and then decrease with the temperature and holding time rising. Especially, the aging temperature shows significant influence. This is caused by the comprehensive effect of austenite and precipitated particles. The increase of austenite content can improve the elongation, but reduce the tensile strength. With the generation of precipitated particles, the tensile strength increase. But the tensile strength will drop when the particles become too big. And the big particles not only have bad effect on the tensile strength but also result in the reduction of elongation (The effect is not significant when particle is small). As for the SAT samples, the tensile strength increases and drops drastically with the increase of temperature after ST at 900 °C for 1 h, but the elongation is almost constant. However, the elongation increases first and then drops with the holding time prolonging, and the change in strength is relatively small. The above results show that the effect of heat treatment on the tensile property is different. Both DAT and SAT could significantly improve the tensile strength to the same extent, which means that aging treatment alone could enhance the tensile strength of parts manufactured by SLM, and no additional solution treatment is needed. The excellent comprehensive performance of the work material can be obtained at 520 °C for 6 h. The fracture morphologies obtained from the as-built, ST900-2 h, DAT520-6 h and SAT520-6 h samples are shown in Fig. 10(a)-(d), respectively. Apparent dimples (orange arrows) are displayed in Fig. 10(a), which indicates good plasticity corresponding to the high elongation. Meanwhile, the co-existing deep holes (blue arrow) are resulted from the SLM process caused by molten liquid shrinkage or vaporization. Fig. 10(b) presents a sizable number of dimples in the as-built sample, which demonstrates the plastic fracture and explains the reason why elongations of as-built and ST are similar. In the meantime, the particle (green arrow) embedded in a big dimple implies the presence of the
undissolved second phase. Fig. 10(c) shows the SEM image of the fracture surface of DAT520-6 h sample, in which lath packets (brown arrow) and dimples appear. It can be confirmed that the crack extends along the martensitic laths in a martensitic lath packet, which is defined by apparent boundaries. This further proves the existence of austenite (with good plasticity) in addition to martensite. Fig. 10(d) displays the fracture surface of SAT520-6 h sample. In addition to the tearing edges (yellow arrow) and small platforms (purple arrow), there are a small number of dimples, which reveals that the development of cracks was mainly ruled by a quasi-cleavage mechanism [34]. For the DAT and SAT samples, the fractures are obviously brittle with little plastic deformation. 3.6. Temperature-dependent impact toughness The toughness test results are shown in Fig. 11. It can be seen that ST can improve the toughness to a certain level of elevated temperature followed by a dip when the temperature is above 960 °C and/or the holding time exceeds 1 h, as shown in Fig. 11(a) and (b). Fig. 11(c)-(f) show that both DAT and SAT lead to the reduction of impact roughness. Though the toughness of DAT-A samples increases after 480 °C for 6 h, all the others drop drastically to a low-toughness state. Fig. 11 displays the fracture surfaces after Charpy impact testing at room temperature. It is obvious that the shear lips and fibrous zone co-exist in Fig. 12(a) and (b), while the radical zone is almost non-existent here, which indicates the good toughness of the as-built and ST900-1 h samples. The shear lips of latter are larger showing the better toughness. In Fig. 12(c) and (d), the fracture surfaces are both flat, and the shear lip and fibrous zone are not found but some oxidized particles and void defects. The flat radical zone demonstrates that the cracks grow rapidly from the Vport in impact progress. This is mainly due to the formation of fine precipitates during aging treatment in both DAT and SAT samples which limit the ductile deformation of the samples [35,36]. 113
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Fig. 11. Charpy impact test results of 18Ni-300 maraging steel: (a) (b) ST; (c) (d) DAT and SAT.
Fig. 12. Fracture surfaces of Charpy samples: (a) as built; (b) ST900; (c) DAT520-6 h; (d) SAT480-6 h.
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Fig. 13. Element map by EDS mapping of the as-built sample: (a) SEM image; (b) spectrum for Fe; (c) spectrum for Ni; (d) spectrum for Co; (e) spectrum for Mo; (f) spectrum for Ti.
much higher than that of the dark area, as shown in Fig. 13 and Table 4. The reason is that the solubility of Mo and Ti in solid is lower than that in liquid. They are forced to move to the liquid zone when the grain formation is initiated during rapid cooling process. As the neighboring grains encounter each other, the Mo-rich and Ti-rich regions are formed in the boundary. High temperature during ST makes the segregated elements (such as Mo and Ti) gain enough energy to diffuse into the interior of the grains. Thus, the white boundaries begin to break up and dissolve. The higher and longer the temperature and holding time, the faster the white boundaries disappear. This can be found in the SAT samples subject to the ST process before aging treatment. During heat treatment of AlSi10Mg prepared by SLM, similar phenomenon is also observed [39]. According to the EDS analysis of white boundaries in Fig. 13 and Table 4, it can be concluded that the white particles in Fig. 4(b) exhibit higher Mo and Ti content which could not be dissolved completely due to low temperature [40]. However, in our experiments, the white boundaries could still be seen in the DAT samples, even
Table 4 The chemical composition of different areas in Fig. 13 (a) (in wt. %). Element
Ni
Co
Mo
Ti
Al
Fe
Spectrum 1, % Spectrum 2, %
18.16 18.52
8.81 9.15
7.14 5.80
1.65 1.11
0.27 0.27
63.97 65.15
4. Discussions From the SEM images in Fig. 4, it can be seen that the microstructure changes significantly with the temperature and holding time. The main micro-features of the as-built sample consist of cells and laths structure with white boundary. This is because high cooling rate results in a large number of nucleation points to produce small grains. The feature is also observed in 316L [37], CoCrMo alloy [33] and Al–Si alloy [38]. And EDS analysis shows that the white boundary is composed of Fe, Ni, Co, Mo, Ti and Al, and the contents of Mo and Ti are 115
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though the temperature increases to 560 °C or the holding time increases to 12 h. That is because the diffusion of Mo and Ti needs much higher energy that cannot be provided by DAT. Furthermore, it is noted that the discontinuous boundaries are due to the intermetallic phases (e.g., Ni3Ti, Ni3Mo, and Ni3Al) precipitated during aging treatment [41,42]. Even though the high magnification SEM micrographs present significant difference between ST900-1 h and SAT560-6 h in Fig. 4, the low magnification optical micrographs show similar lath martensitic lath features. The main difference is that the size of martensite becomes bigger after aging treatment for 6 h. Figs. 8 and 9 show a similar change trend in the microhardness and tensile strength, including ST, DAT and SAT samples. For example, both microhardness and tensile strength decrease obviously after ST. That is because the high cooling rate leads to the rapid solidification of the molten pool to produce martensite tissue, fine grains and large residual stress, which can significantly prevent the movement of dislocation to hinder the formation and expansion of cracks. As such, the microhardness and tensile strength are improved. After ST, the microstructure begins to transform into martensite lathes, and the higher the temperature and holding time, the more the martensite formation. Martensite can improve the tensile strength and microhardness, but the results show a reduction in them. The reason is that the fine grains and residual stress disappear after ST. Their strengthening effects decrease at the same time. And the improvement from martensite transformation is lower than the deduction of disappearance of fine gains and residual stress. As a result, the microhardness and tensile strength decrease eventually. These indicate that the as-built and ST samples are mainly exhibiting ductile fractures [34]. A ductile fracture is always transgranular [43]. For maraging steel manufactured by casting and forging methods, solution treatment is essential before aging treatment to prepare the martensite matrix for precipitation strengthening with the particles, such as Fe2Mo, Ni3Ti, and Ni3Mo, which can increase the resistance of dislocation movement. However, there is a significant difference in maraging steel by SLM. Due to the rapid cooling, the martensite matrix has been achieved during SLM manufacturing process, which can be identified by the XRD analysis (Fig. 6). These provide the tissue base for the aging treatment. Therefore, theoretically speaking, it is possible to achieve high strength and microhardness after conducting aging treatment on the as-built samples without solution treatment [44]. DAT samples retain the similar microstructure to the as-built ones. But the microhardness and tensile strength of the DAT samples are very high, which are similar to SAT. From the DSC curve, it can be observed that the aging behavior occurs at relatively low temperatures comparing with ST. The cells and laths from microsegregation during SLM are reserved due to low energy density. But the Ni, Mo and Ti elements dissolved in the martensite matrix can be precipitated normally to form strengthening particles Ni3(Mo, Ti), which improve the mechanical properties significantly. Therefore, even though the microstructure is almost constant after aging treatment, DAT samples can still achieve high microhardness and tensile strength. Another thing should be noticed that almost no peaks from Fe2Mo and Ni3(Ti,Mo) intermetallic compounds can been found in the XRD test (Fig. 6) in the as-built, DAT and SAT samples. This is mainly because of the extremely small size of precipitates, low phase content, and highly dispersed distribution [24,45,46]. In summary, the microstructures and mechanical properties are very different between the as-built, ST, DAT and SAT samples. But the DAT samples show the similar microhardness, tensile strength and impact toughness to the SAT ones, despite the distinctions in microstructure. Since DAT could obtain the same mechanical properties with less postprocessing comparing to SAT, it will be very meaningful to reduce the cost of heat treatment for the maraging steel by SLM. However, more research is still needed to explore and understand the strengthening mechanism of DAT. The future work will focus on the texture, precipitation behavior and reverse austenite nucleation using transmission electron microscope (TEM) and electron backscattered diffraction (EBSD).
5. Conclusions This paper investigates the heat treatment processes associated with the selective laser melting (SLM) technology, which is different from traditional methods. A series of heat treatment experiments, including solution treatment (ST), direct aging treatment (DAT) and solution + aging treatment (SAT), have been conducted on maraging steel samples manufactured by SLM. The evolution mechanisms of microstructure, microhardness, tensile property and toughness are investigated with different heat treatment conditions. The following conclusions can be drawn from this study: (1) ST and SAT have significant effect on the microstructure evolution. When the ST temperature is at 780 °C, the austenite grain boundary and white particles can be observed. When the temperature reaches 840 °C, the cells, lathes and melt track boundary all disappear. The microstructure of DAT samples is similar to the as-built samples. With the temperature and holding time increasing, the melt track boundary dissolves gradually. As for the SAT samples, the microstructure consists of lath martensite. But when the temperature reaches 560 °C or holding time reaches to 12 h, the lath martensite disappears and is replaced by white and dark tissues. (2) Comparing with the as-built samples, ST reduces the microhardness and tensile strength, and the fracture elongation does not increase because of the disappearance of fine grains obtained in SLM. But the impact toughness is improved. Both DAT and SAT can improve the microhardness and tensile strength significantly. However, both fracture elongation and impact toughness are reduced. The best heat treatment process for high hardness and strength is DAT at 520 °C for 6 h with an achievable microhardness up to 653.93 HV and tensile strength up to 2126.30 MPa. (3) DSC curves show the phase transformation temperature of the asbuilt sample at 456 °C, 572 °C, 670 °C and 715 °C, which indicates the solution and aging temperature range. The microhardness, strength and impact properties of DAT are similar to the ones of SAT. But the microstructure is obviously different between these samples. The content of austenite increases with the temperature and holding time rising in the DAT samples. However, ST and SAT result in the elimination of austenite in the maraging steel manufactured by SLM. Acknowledgments The authors gratefully appreciate the financial support from the Natural Science Foundation of China (NSFC, No.: 51875215), Guangdong Province Science and Technology Project (No.: 2015B090921002 and 2016B09092500), Guangzhou Science and Technology Project (No.: 201704020118), Singapore Ministry of Education AcRF Tier 2 (Project No.: MOE2018-T2-1-140) and A*STAR AME IAF-PP (Grant No.: A1893a0031). References [1] W.E. Frazier, Metal additive manufacturing: a review, J. Mater. Eng. Perform. 23 (6) (2014) 1917–1928. [2] D. Herzog, V. Seyda, E. Wycisk, C. Emmelmann, Additive manufacturing of metals, Acta Mater. 117 (2016) 371–392. [3] E.O. Olakanmi, R.F. Cochrane, K.W. Dalgarno, A review on selective laser sintering/ melting (SLS/SLM) of aluminium alloy powders: processing, microstructure, and properties, Prog. Mater. Sci. 74 (2015) 401–477. [4] B. Vayre, F. Vignat, F. Villeneuve, Metallic additive manufacturing: state-of-the-art review and prospects, Mechanics & Industry 13 (2) (2012) 89–96. [5] D. Dai, D. Gu, Thermal behavior and densification mechanism during selective laser melting of copper matrix composites: simulation and experiments, Mater. Des. 55 (2014) 482–491. [6] Q. Zhang, J. Chen, P. Guo, et al., Texture and microstructure characterization in laser additive manufactured Ti–6Al–2Zr–2Sn–3Mo–1.5 Cr–2Nb titanium alloy, Mater. Des. 88 (2015) 550–557. [7] C.Y. Yap, C.K. Chua, Z.L. Dong, et al., Review of selective laser melting: materials
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