Microstructural evolution, nanoprecipitation behavior

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Materials and Design 134 (2017) 23–34

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Microstructural evolution, nanoprecipitation behavior and mechanical properties of selective laser melted high-performance grade 300 maraging steel Chaolin Tan a,b, Kesong Zhou a,b,⁎, Wenyou Ma b, Panpan Zhang b, Min Liu b, Tongchun Kuang a a

School of Materials Science and Engineering, South China University of Technology, Guangzhou 510640, China Guangdong Institute of New Materials, National Engineering Laboratory for Modern Materials Surface Engineering Technology, Key Lab of Guangdong for Modern Surface Engineering Technology, Guangzhou 510651, China





• Evolutions of the typical SLMed microstructures are illustrated and theoretically explained. • Precipitation behavior and phase transformation of SLMed maraging steel are characterized by TEM and XRD. • Significant improvement of strength after solution and aging treatment was evaluated and explained. • Relationships between massive nanoprecipitates and improved mechanical performances are elucidated.

a r t i c l e

i n f o

Article history: Received 24 May 2017 Received in revised form 8 August 2017 Accepted 10 August 2017 Available online 12 August 2017 Keywords: Selective laser melting Maraging steel Precipitate Age hardening Microstructural evolution Orowan mechanism

a b s t r a c t High-performance grade 300 maraging steels were fabricated by selective laser melting (SLM) and different heat treatments were applied for improving their mechanical properties. The microstructural evolutions, nanoprecipitation behaviors and mechanical properties of the as-fabricated and heat-treated SLM parts were carefully characterized and analysed. The evolutions of the massive submicron sized cellular and elongated acicular microstructures are illustrated and theoretically explained. Nanoprecipitates triggered by intrinsic heat treatment and amorphous phases in as-fabricated specimens are observed by TEM. High-resolution TEM (HRTEM) images of the age hardened specimens clearly exhibit massive nanosized needle-shaped nanoprecipitates Ni3X (X = Ti, Al, Mo) and 50–60 nm sized spherical core-shell structural nanoparticles embedded in amorphous matrix. XRD analyses reveal austenite reversion and probable phase transformations during heat treatments. The hardness and tensile strength of the as-fabricated and age-treated SLM specimens absolutely meet the standard wrought requirements. Furthermore, the lost ductility after aging can be compensated by preposed solution treatments. Relationships between massive nanoprecipitates and dramatically improved mechanical performances of age hardened specimens are elaborately analysed and perfectly explained by Orowan mechanism. This study demonstrates that high-performance grade 300 maraging steels, which is comparable to the standard wrought levels, can be produced by SLM additive manufacturing. © 2017 Elsevier Ltd. All rights reserved.

⁎ Corresponding author at: School of Materials Science and Engineering, South China University of Technology, Guangzhou 510640, China. E-mail address: [email protected] (K. Zhou).

http://dx.doi.org/10.1016/j.matdes.2017.08.026 0264-1275/© 2017 Elsevier Ltd. All rights reserved.


C. Tan et al. / Materials and Design 134 (2017) 23–34

1. Introduction The latest emerging additive manufacturing (AM), also called 3D printing, is an incremental layer-by-layer manufacturing in which materials, such as plastic or metal, are deposited onto one another in layers to produce a 3D objects [1–3]. Selective laser melting (SLM), a typical metal AM technology, integrates multidisciplinary fields including information technology, material technology and manufacturing technology [4]. SLM produces 3-D parts from laser radiation to powder materials on the powder beds through successive layer bonded to already existing layers [5,6]. Due to many irreplaceable superiorities such as high material and resource efficiency, good part design and production flexibility, improved mechanical performances of SLM technology [7–9], SLM has been wildly applied in customized medical and dental application fields, tooling inserts with conformal cooling channels and functional components with high geometrical complexity such as porous and lattice constructs [10–12]. Many researches have investigated SLM fabrication of Ti-based alloys [13,14], Al-based alloys [15,16], Ni-based alloys [17,18] and Febased alloys [19,20], in terms of optimizing the process conditions, proceeding heat treatments and mixing reinforced ceramic powders, in order to minimize the porosity or cracks, adjust the acquisitive performances and endow specialized functions. Maraging steels, combining ultra-high strength with good toughness and ductility, are a kind of special advanced high strength steel (AHSS) being widely applied in the aircraft, aerospace and tooling industries [21,22]. Compared with other Fe-based alloys, such as 316L stainless steels, Fe-Al high strength steels and H13 tooling steels, maraging steels are well suited for the SLM process mainly due to the following three reasons. First and foremost, maraging steels can be age hardened by intermetallic precipitates, such as η-Ni3Ti [23,24], Fe2Mo [24], Ni3Mo [25], Ni3Al [26], NiAl [22,27], Ni3(Al, Ti) [27,28], Ni3(Ti, Mo) [29], Ni(Al, Fe) [30], etc. Some maraging steels can show hardening even after only a few seconds of heat treatments due to solute clustering [31]. In the SLM process, a previously deposited material may experience cyclic reheating with gradually decaying laser intensity during deposition of neighboring tracks and subsequent layers, i.e., the as-deposited material in one layer is in-situ heat-treated by the subsequent tracks and layers. The sharp temperature pulses in this type of intrinsic heat treatment (IHT) almost reach the melting point [23,32,33]. Consequently, this in-situ IHT may trigger clustering or nucleation of hardening precipitates during SLM process without need of succeeding additional heat treatment process [22]. Secondly, martensitic matrix materials need to be quenched rapidly from the austenitic region to the Ms temperature. The maraging steel produced by SLM has a significantly higher yield strength and ultimate tensile strength than conventional processing methods without heat treatment [34]. This can be attributed to the very fine cellular microstructure (diameter ≤ 1 μm) caused by the small sized molten pool and extremely high cooling rate (up to 108 K/s) in the SLM process [32]. Thirdly, maraging steels are mainly applied in the aerospace and tool-manufacturing industries, which often call for geometrically complex components with excellent mechanical properties in relative small quantities, such as modes with conformal cooling channels and components with lattice constructs [32,35]. Consequently, the SLM technology is an efficient method to meet the aforementioned requirements and expand application fields of maraging steels. Nevertheless, relatively few researches concerning SLM fabricated maraging steels have been reported, and merely focused on process optimization and mechanical performances [21,34,36]. The precipitation behaviors and strengthening mechanisms of SLM fabricated maraging steels during age hardening have not been profoundly and systematically investigated. In this paper, high-performance maraging steels were produced by SLM technique, high-resolution TEM (HRTEM) analysis was conducted to investigate the microstructural evolution, precipitation behaviors and strengthening mechanisms of SLM specimens.

Mechanical properties of the SLM produced specimens were characterized and systematically investigated in comparison with traditionally fabricated maraging steels. 2. Experimental details 2.1. Materials and SLM process The maraging steel powder produced by Electro Optical Systems (EOS) GmbH (Germany) was used as the raw material. The composition of the powder was listed in Table 1 and the oxygen content of the raw powder was measured to be 342 ppm. Fig. 1 shows the SEM morphology and size distributions (measured by a HORIBA LA960WET laser scattering particle size analyser) of the powder. The powder particles with a mean diameter of 41.62 μm are almost spherical shapes. The experiments were carried out in an EOS M290 SLM (powder bed) system. Before starting the SLM process, the base plate was preheated to 40 °C by a heater placed inside the building platform, and the oxygen content in the process chamber was maintained under 0.6% by pumping in continuous nitrogen. The process parameters were optimized by orthogonal experiment method, in which the laser power (P), scan speed (V) and hatch space (h) were systematically studied in the range of 200–370 W, 500–1800 mm/s and 50–150 μm, respectively. The schematic diagram of the SLM process and optimal process parameters are shown in Fig. 2. The laser scanning was in zigzag pattern with 67° rotation (as depicted by θ = 67° in Fig. 2) between the adjacent layers. As mentioned above, aging heat treatment can effectively improve the performances of maraging steels, so the as-fabricated specimens were heat treated at 490 °C, 6 h for age-hardening. Besides, the solution treatment at 840 °C for 1 h, followed by aging at 490 °C for 6 h was also conducted for comparison. All the heat treatments were protected in argon atmosphere and air-cooled after duration. 2.2. Characterizations The surface roughness of the fresh vertical and horizontal crosssections for an as-fabricated specimen was characterized by a BMT SMS Expert 3-D model optical profilometer to evaluate the roughness Sa with an area of 2 mm × 2 mm. Relative density ρr of the asfabricated specimens was measured by the formula ρr = (m0ρ1)/ (m0ρ0 − m1ρ0) according to Archimedes' principle, in which m0, ρ0, m1 and ρ1 are the maraging steel specimen's weight in the air, theoretical full-density (8.01 g/cm3), weight when submerged in the water and density of the applied water under normal atmospheric pressure (0.9982 g/cm3), respectively. The polished and etched vertical and horizontal cross-sections were observed by a Leica Dmi5000m optical microscope (OM) and a Zeiss Merlin field emission scanning electron microscope (FE-SEM), fitted with an Oxford X-MaxN20 energy dispersive spectrometer (EDS). A JEOL 2100F transmission electron microscopy (TEM) was used for nanosized structures and precipitates observation operating at 200 kV. Scanning transmission electron microscopy (STEM) observations, selected area electron diffraction (SAED) patterns and energy-dispersive X-ray spectroscopy (EDX) mappings were also carried out on the TEM. X-ray diffraction (XRD) was conducted by using a Bruker D8 Advance Diffractometer with a Cu Kα radiation (wavelength, λ = 0.15418 nm), at 40 kV and 40 mA in a 2θ range of 30–90° using a step size of 0.02°. The mechanical properties were evaluated by hardness and tensile tests. The hardness of the SLM specimens was measured by a TH320 Rockwell hardness tester according to the 150 kg loaded Rockwell C scale (HRC), and estimated by an average value from 10 measured points. The tensile properties were evaluated by an Instron 5900 universal material testing machine (referring to ASTM E8 [38]) with the cross head speed of 1 mm/min. The ultimate tensile strength (UTS), yield strength (YS) and break elongation (El) of the SLM specimens were estimated by an average value from 3

C. Tan et al. / Materials and Design 134 (2017) 23–34


Table 1 Main chemical composition of grade 300 maraging steels. Element (wt%)










ASTM A538 Powder As-fabricated

18–19 18.2 17.79

8.5–9.5 9.02 9.52

4.6–5.2 5.22 5.18

0.5–0.8 0.79 0.67

0.05–0.15 0.15 0.09

≤0.5 0.21 0.15

≤0.5 0.06 0.07

≤0.03 – –

Bal. Bal. Bal.

measured specimens. Fracture morphologies of specimens after tensile tests were investigated by the FE-SEM.

3. Results and discussion 3.1. Surface appearance and densification analysis The characteristic fresh horizontal surface morphologies of the asfabricated SLM parts are provided in Fig. 3. As shown in Fig. 3a, a relatively smooth and dense surface is obtained and the regular liquid fronts are formed, which reveals that all the laser tracks have good metallurgical bonding with each other. Similar phenomenon was also observed in the SLM fabricated Fe-based composites [39]. The element compositions of as-fabricated specimens are listed in Table 1, almost all of them meet the ASTM A538 standard requirements. Furthermore, there are no obvious elemental contents distinctions between the powder and asfabricated specimens, which indicates that the burning loss and alloying elements evaporation during laser irradiation were negligible. Interestingly, massive nanocrystals are observed at the center of laser tracks (Fig. 3b) and near the overlapped regions (Fig. 3c). This would form in response to the rapid solidification with extremely high cooling rate up to 106–108 K/s during the laser process [10,32]. The relatively coarse crystals in the overlapped regions, as provided in Fig. 3c, were triggered by the crystal growth during laser re-irradiation process.

The preparation accuracy of SLM fabricated parts is seriously affected by the surface roughness. Therefore, it is of great significance to characterize and evaluate the surface conditions. The non-polished primary roughness (Sa) of SLM fabricated specimens examined from a 2 mm × 2 mm area are 4.16 μm for the horizontal surfaces and 4.79 μm for the vertical surfaces. The relatively low surface roughness certified the high accuracy of forming process. The density of SLM parts is another important factor that affects the performances of the SLM parts. The relative density of as-fabricated specimens reaches 99.98%, which confirms the rationality and optimization of processing parameters. The formation processes of the micropores can be explained as follows: firstly, the protective atmosphere of nitrogen gas in the building chamber was rolled in the active flow of molten pools incited by the high-energy density laser irradiation; furthermore, the dynamics of the Marangoni flow increases with laser energy input, thus increasing the probability of any gas (either newly formed or originally trapped within the highly packed powder particles) being dragged towards the bottom of the molten pool then the micropores will be enveloped in solidified molten pools [40,41].

3.2. Microstructure analysis 3.2.1. Macro-morphology observation Fig. 4 shows the typical optical micrographs of the etched horizontal and vertical cross-sections of the as-fabricated SLM specimen. There are

Fig. 1. SEM morphology (a) and size distribution (b) of the grade 300 maraging steel powders.

Fig. 2. Schematic diagram and main parameters of the SLM process. (The right picture is adapted from [37])


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Fig. 3. SEM images showing typical surface morphologies of SLM-fabricated specimens: (a) low-magnification observation of laser tracks; high-magnification observation of central track (b) and overlapped area (c) corresponding to the region 1 and region 2 marked in panel (a) respectively.

scarcely any micropores or defects in the specimens, suggesting the achievement of a nearly fully dense part. As shown in the Fig. 4a, the horizontal cross section displays the motion tracks of laser radiation. The ellipsoid braided structures of the horizontal cross section are about 90–110 μm in width, which equals to the laser spot diameter (about 100 μm). By contrast, the structures of the vertical cross section (Fig. 4b) are flaky half-ellipses, similar to the fish scales, with the dimension of 40–50 μm in height and 80–100 μm in base width. The height of half-ellipses approximately equals to the powder layer thickness while the base width approximately equals to laser spot diameter. It clearly reflects the profile of the laser molten pool. Therefore, the structures of vertical cross section were attributed to the Gaussian distribution of laser intensity.

3.2.3. Microstructure evolutions Fig. 6 shows a characteristic morphologies of the horizontal and vertical cross sections of an as-fabricated SLM specimen encompassing the border between laser tracks and molten pools. As revealed in Fig. 6a, in the middle region of a laser track, submicron sized hexagonal cellular grains are uniformly distributed; by contrast, in the margin region of the laser tracks, elongated acicular grains which are perpendicular to the laser moving direction are present. These microstructure characteristics can be depicted and analysed by Fig. 6b, the Gaussian laser energy distribution caused the corresponding heat input (Qs) distribution in the molten pool, as described by the equation below [43]:

QS ¼ 3.2.2. Microstructural observation Fig. 5 shows microstructural observations by OM and SEM for the horizontal cross-sections of the as-fabricated and heat treated specimens. The OM images show that the as-fabricated specimen consists of coarse plate martensites (Fig. 5a) with uneven distribution, and the overlapped trails could be slightly observed. By contrast, the martensites are refined after age treatment (Fig. 5b), and the dense acicular martensites are formed in the solution-aged specimens (Fig. 5c). The interfacial features between laser tracks disappeared after heat treatments. In addition, the spot-like austenite located among the acicular martensites can be observed in the higher magnification SEM image of Fig. 5c. Massive submicron sized fine cellular microstructures with size of approx. 0.2–0.6 μm could be observed in the SEM images in Fig. 5a– c. These characteristic fine cellular microstructures would form in response to the instant melting and rapid solidification with extremely high cooling rate during the laser irradiation [42]. As one can see that, the as-fabricated and heat-treated specimens present distinct microstructural characteristics. Compared with the as-fabricated specimens, the grain boundaries of heat-treated specimens are vague and irregular, it was possibly caused by precipitates squeezing into the boundaries, residual stress releasing and phase transformations during heat treatments.

fP V 2

πd h

exp −3

! R2  d



z h

where Pv is the absorbed laser power, f is the factor of heat distribution that influences the power distribution, d is the beam radius, R is the radial distance from the laser beam center, z is the current depth in the thickness direction and h is the depth of the energy source. In the same horizontal plane, the heat input exponential decreases when the R increases, as such, the temperature of the center in a laser track is much higher than that of the margin area, and the thermal flux will generate from the center to margin due to the thermal dissipation. The high dissipation and cooling rate cause the temperature of liquid metal (Tl) below the melting point (Tm) in the central zone, and the degree of undercooling (△T = Tm − Tl) is highly enough for new grains nucleation. Hence, the liquid metal will experience synchronously nucleation and randomly orientation. Moreover, the growth rates in all directions of crystal nucleus are consistent, so equiaxial crystals are easily formed. As illustrated in Fig. 6b, the equiaxial crystals exhibit hexagonal cellular structures, which can be theoretically explained by low energy boundary theory. As many researches revealed, the polycrystalline materials are normally filled with high-angle boundaries; if the angles of boundaries vary from each other, then the interfacial energies would have significant differences among the boundaries, which will cause unsteadiness of the polycrystalline system. In fact, on the basis of

Fig. 4. Optical macrographs of the horizontal (a) and vertical (b) cross-sections of the SLM fabricated specimens.

C. Tan et al. / Materials and Design 134 (2017) 23–34


Fig. 5. The OM and SEM images taken from horizontal cross-sections of SLM specimen: (a) as-fabricated, (b) aging treated and (c) solution-aging treated specimens.

minimum Gibbs free energy theory, the grain boundaries are prone to have identical angles and three crystals intersecting interface will preferentially form. Because interfacial tensions of ternary grain boundaries can be steadily maintained when interfacial energies (γ) satisfied the equation γ1 = γ2 = γ3. Besides, the intersecting of four or more crystals could not steadily hold and the status of three crystals intersecting will come into being via structural decomposition under the favorable conditions. Therefore, to meet the equilibrium equation of force (as stated in

Fig. 6b), the values of boundary angles shall be equivalent, i.e., α1 = α2 = α3. With regarding to the elongated acicular crystals, the input thermal flux and crystallizing latent heat released from the center decreased the degree of undercooling, which could suppresses the nucleation of new grains in return. As such, the elongated acicular crystals are prevalent due to high crystal growth rate along the direction of thermal flux. A sectional SEM morphology of the molten pool in vertical cross section of an as-fabricated SLM specimen was provided in Fig. 6a. There are

Fig. 6. Microstructural evolutions analyses of the as-fabricated specimens: (a) the characteristic morphologies of the horizontal and vertical cross-sections; (b) the schematics and formation mechanism analyses of the cellular crystals and elongated acicular crystals; (c) the schematics and formation mechanism analyses of the microstructures in the molten pool and overlapped area.


C. Tan et al. / Materials and Design 134 (2017) 23–34

columnar dendritic structures at the bottom of the molten pool, cellular microstructures in the middle of the molten pool and coarse equiaxial crystal at the border between the molten pools. The formation mechanisms of these various crystal morphologies are illustrated in Fig. 6c, the microstructure evolution is mainly determined by the ratio of temperature gradient to solidification rate (G/R) [44]. At the bottom of the molten pool, G has the largest value, due to the decreasing of heat input (Qs). While, R is close to 0, so the ratio of G/R is comparatively large and the solidification structure will be planar extension growth. Similar results were also reported by Acharya in SLM IN718 powders [45]. Ascending from the bottom of the molten pool, R increases gradually and the value of G/R decreases gradually. Therefore, a cellular dendritic structure is formed along the layer stacking direction, this direction is also the direction of heat flow. With the further decrease of G/R in the middle of the molten pool, the microstructure transforms to cellular structure [46]. Besides, the coarse equiaxial crystals form at the borders between the molten pools, it is mainly because these regions are affected by the heat flux of the subsequent laser irradiation. So the solidification rate decreased and coarse crystals developed in this region. 3.3. Nanoprecipitate characteristics 3.3.1. Nanoprecipitate characteristics of as-fabricated specimen Further detailed microstructural observations of as-fabricated specimen by TEM have been carried out and given in Fig. 7. Columnar martensites with width of about 200 nm (Fig. 7a) were parallelly distributed and determined by the corresponding selected area electron diffraction (SAED) pattern from the [011] zone axis. More interestingly, massive nanoprecipitates with size of about 5–20 nm in the matrix were clearly observed in Fig. 7b, the reasons of precipitates formation could be explained as follows: firstly, during powder material deposition, a melting track solidified after laser irradiation. And then, the solidified track will experience cyclic reheating and gradually cooling down during laser deposition of neighboring tracks and subsequent layers. This intrinsic heat treatment (IHT) phenomenon consists of sharp temperature pulses near to the melting point of maraging steels [22]. Meanwhile, as previously mentioned, maraging steel is a martensite age hardening steel which is strengthened by the precipitation of intermetallics at temperature of about 455–510 °C. Therefore, the nanoprecipitates can be triggered by IHT effect, similar discovery was also reported earlier by Kürnsteiner [22] during laser deposition of Fe-19Ni-Al maraging steel, in which high number density (10 25 precipitates per m3 ) of NiAl precipitates was formed and detected. Besides, as clearly exhibited by HRTEM and corresponding SAED pattern in Fig. 7c, the amorphous-nanocrystalline composite microstructures are obtained due to extremely high cooling rate during solidification. Crystalline structures can also be observed in bright-field TEM image in Fig. 7c, and it can be confirmed from the inset SAED pattern in which diffraction spots (crystalline phases) surrounds the amorphous diffraction ring.

3.3.2. Nanoprecipitate characteristics of age hardened specimen However, the number of precipitates in as-fabricated grade 300 maraging steel triggered by IHT during SLM process is limited, due to the relatively sluggish precipitation of the Mo-containing precipitate phases [32]. Therefore, a heat treatment for age hardening is complementary. Scanning transmission electron microscopy (STEM) image and energy-dispersive X-ray spectroscopy (EDX) mappings taken from the horizontal section of an age-treated grade 300 TEM specimen are provided in Fig. 8. Numerously and densely distributed acicular precipitates and dispersively distributed 40–80 nm sized spherical nanoparticles are observed in the STEM image. EDX mappings show that the nanoparticles are rich of Ti and Al elements, it is mainly due to the very fast precipitation kinetics of Ti and Al. The richer elemental contents of Ti and Al, predict that particles may be Ti/Al-based or Ti-Albased intermetallic compounds. TEM bright-field images and corresponding SAED patterns taken from age-treated specimen are given in Fig. 9 for intensive investigation of precipitates. Fig. 9a shows BF-TEM overview of massive nanoprecipitates and dispersive nanoparticles embedded in martensite matrix. Extensive amorphous phases could be clearly observed in the matrix. Besides, a lot of dislocations (arrowed) can be also observed in Fig. 9a, similar dislocations morphology was also presented in Ni(Al, Fe) maraging steel [30]. The needle-shaped nanoprecipitates are about 6–10 nm in diameter and range from 15 nm to 45 nm in length, the interspacings of precipitates are about 25 nm in average. Compared with the 5–20 nm sized nanoprecipitates in the as-fabricated specimens, the increasement of nanoparticles in the age-treated specimens presents not only in quantity, but also in terms of diameter size. This kind of variation indicates that a large quantity of nanoprecipitates may nucleate and grow up during 6-hour age treatment. This phenomenon could be explained in two aspects. On one hand, as a majority of observations revealed [23,47], intermetallics generated from long time aging will heterogeneous precipitate on dislocations, and subsequently grow up by pipe diffusion. The growth behavior obeys a relationship of the type d = ktn, where d is the precipitate diameter, t is the aging time, n responds the growth exponent and k is a constant related to activation energy of growth [24]. So the precipitates generated from aging for several hours are larger than nanoprecipitates triggered by IHT effect in asfabricated specimens. On the other hand, the primary nanoprecipitates, which triggered by IHT effect, would also grow up during aging, so the size of nanoparticles in age-treated specimens increased. Spherical nanoparticles that lie among the nanoprecipitates are about 50–60 nm. Notably, as exhibited in high-magnificational observation of nanoparticles in Fig. 9b, nanoparticles present core-shell structure with crystalline core and amorphous shell. Based on the EDX mapping results in Fig. 8, the cores may be Ti- or Al-based intermetallics. SAED pattern of the particles taken from encircled region in Fig. 9a is shown in Fig. 9c, it reveals phase constitution of hexagonal η-Ni3Ti (a = 5.101 Å, c = 8.306 Å), cubic Ni3Al (a = 3.572 Å) and α-Fe (a = 2.866 Å). Fig. 9d shows HRTEM images and the corresponding SAED pattern (inset) of amorphous region taken from region 1 of Fig. 9b. An amorphous diffraction is clearly observed, it confirmed the white

Fig. 7. TEM bright-field images taken from as-fabricated specimen showing micrographs of (a) columnar martensites with corresponding SAED pattern (inset), (b) nanoprecipitates (arrowed) and (c) amorphous-nanocrystalline composite microstructures with corresponding SAED pattern (inset).

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Fig. 8. STEM image and EDX mapping of the TEM thin foil taken from an age-treated specimen showing distribution (arrowed) and higher Ti, Al contents of the nanoparticles.

contrast morphologies in Fig. 9a are almost complete amorphous structures. The diffraction spots surrounding the amorphous diffraction are the embedded precipitates in amorphous phase. As shown in Fig. 9e, HRTEM images of precipitates region corresponding to the region 2 in Fig. 9b, confirms the existence of η-Ni3Ti precipitates by the dspacings of {0004} reflections. Fig. 9f shows SAED pattern of the precipitates corresponding to the regions marked in Fig. 9a, it reveals that the phase constitutions of precipitates are Ni3Al and orthorhombic Ni3Mo (a = 5.064 Å, b = 4.224 Å and c = 4.448 Å). As many reports revealed, typically, interaction between Ni and Ti is the most rapid interaction during aging, because the η phase is the dominate precipitate which could be quickly generated [23,24,32,48,49].

Moreover, the hexagonal η-Ni3Ti is the main precipitate phase responsible for the strengthening of the maraging steels containing Ti, and Ti might be substituted partially by other elements such as Mo or Al, depending on alloy composition [23,24,32]. In addition, as presented in Fig. 8, the particles consist of abundant Ti and relative less Al, it indicates that the η-Ni3Ti precipitates formed very quickly upon aging, and the ηNi3Ti precipitates could act as the nucleation site for the later precipitation of Ni3Al and Ni3Mo precipitates. Besides, the Ni3Al and Ni3Mo precipitates may well form through substitution of the remaining Ti in the matrix by Mo and Al. HRTEM images and schematic diagram depicting formation process of nanoparticles are provided in Fig. 10. The overall HRTEM image of

Fig. 9. Bright-field TEM images and corresponding SAED patterns taken from age-treated specimen: (a) BF-TEM overview showing massive nanoprecipitates and dispersive nanoparticles embedded in amorphous matrix; (b) megascopic BF-TEM image of nanoparticle taken from the given region in (a); (c) SAED pattern of the nanoparticles corresponding to the SAED-1 region marked in (a); (d) HRTEM image and SAED pattern (inset) of amorphous region corresponding to the region 1 marked in (b); (e) HRTEM images of precipitates region corresponding to the region 2 marked in (b); (f) SAED pattern of the precipitates corresponding to the SAED-2 regions marked in (a).


C. Tan et al. / Materials and Design 134 (2017) 23–34

Fig. 10. HRTEM bright-field images and schematic diagram depicting formation processes of nanoparticles in age-treated specimen.

nanoparticles, as exhibited in Fig. 10b, shows that the central region is a crystalline region, and crystalline phases gradually transform to amorphous phase from the center to the boundary region. Detailed HRTEM observations of nanoparticles taken from left side and right side are shown in Fig. 10a and c, respectively. The η-Ni3Ti phase was determined from Fig. 10a by the d-spacings of {0004} reflections and α-Fe (martensite) phase was determined from Fig. 10c by the d-spacings of {110} reflections. According to the heterogeneous nucleation theory [50], the nonsteady amorphous structure can easily nucleate at nanoprecipitates, because the solid nanoprecipitates can act as “intermetallide nuclei” for nucleation and then lower the surface energy and nucleation energy. As illustrated in Fig. 10d, amorphous structures will primarily nucleate at nanoprecipitates (such as Ni3Ti, Ni3Al, Ni3Mo, etc.), i.e., stage І; and then the nucleus will grow up during heating and reach stage ІI; further input of heat will lead to the further growing up and new nucleations, i.e., stage ІII; after that, the nucleate crystals will merge together through continuous nucleation and growth, so large sized particles (see in stage ІV) come into being. Precipitates may also encounter with slight growth in the whole stages. The particles have a spherical shape in order to minimize surface energy. A large region of amorphous structures can still be clearly observed in Fig. 9a after aging heat treatment. It is mainly because that the crystallization temperature (Tx) of

Fe-based amorphous usually reaches up to 600 °C [51], so age-treated at 490 °C might have caused only a minor quantity of crystallization of amorphous structures, even though the nanoprecipitates decreased the degree of undercooling needed for nucleation. 3.4. Phase transformation XRD patterns of as-fabricated and heat-treated SLM specimens are shown in Fig. 11a. Volumetric percentages of martensite α and austenite γ phases of SLM specimens are summarized in Table 2 by applying Rietveld Refinement analysis. The as-fabricated specimen consists of a majority of martensite (α) phases and a small minority of austenite (γ) phases. The amount of γ phase in age-treated specimen increased in comparison to that of as-fabricated specimen. Complete martensite phase is obtained after solution treatment, and a few volume fraction austenite phase (γ111) emerge again after integrated solution and age treatments. As shown in Fig. 11b, the phase transformations can be explained with the help of the Fe-Ni binary metastable diagram, which plotted the transformations between austenite and martensite upon heating and cooling [52]. The metastable diagram indicates that martensite starts to transform from austenite at the point of the Ms temperature and even a very low cooling rate produces a fully martensitic structures. The solution treatment in fully austenitizing region can

Fig. 11. XRD analysis the as-fabricated and heat-treated SLM specimens (a) and (b) metastable phase diagram of the Fe-Ni system (Ref. [52]).

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Table 2 Volumetric percentages of martensite α and austenite γ phases for SLM produced specimens. Specimens

α (%)

γ (%)

As-fabricated Age-treated Solution treated Solution-age treated

89.7 89.0 100 93.5

10.3 11.0 0 6.5

adequately dissolve the alloying elements to the matrix and obtain supersaturated homogeneous α-Fe solid solutions. So solution treatment is applied for making full use of alloying elements and producing completely uniform martensitic matrix after cooling down. Besides, it is reported that, reverted austenite (γ′) will be inevitably present during the age treatment [23,32], so the increased content of austenite after age treatment was ascribed to the inevitable reversion of martensite to a more stable reverted γ′ phase. Besides, age treatment results primarily from the precipitation of intermetallic compounds in maraging steels, as aforementioned in Section 3.3, precipitations take place preferentially on dislocations and within the lath martensite to produce a fine uniform particle distribution. The major hardener is Ti, Mo and A1 elements. The η-Ni3Ti is the dominate precipitate and is quickly formed due to the most rapid interaction between Ni and Ti during aging. The Ni3Al precipitates will form through substitution of the remaining Ti in the matrix by Al. The Ni3Mo will initially form because of its better lattice fitting with the martensitic matrix, and further aging can result in the in-situ transformation of metastable Ni3Mo to the equilibrium Fe2Mo (hexagonal C14-type) phase [52]. The cobalt is not involved in the age-hardening reaction, but aims to decrease solubility of Mo in the martensitic matrix and thus increase the amount of Ni3Mo precipitate formed during age hardening. The aging time in our work is limited (6 h), so the main precipitates are η-Ni3Ti, Ni3Al and Ni3Mo phases, Fe2Mo precipitate is not discovered. Consequently, the phase transformations during heat treatments can be reasonably summarized as follows: αþγ

490 °C aging


840 °C solution

α þ γ þ γ0 þ Ni3 ðTi; Al; MoÞ precipitates α

490 °C aging

α þ γ0 þ Ni3 ðTi; Al; MoÞ precipitates

3.5. Mechanical properties analysis The hardness features for horizontal and vertical cross sections of the as-fabricated and heat-treated SLM specimens are displayed in Fig. 12, and the representative engineering stress-strain curves of the as-

Fig. 13. Representative engineering stress-strain curves of the as-fabricated and heattreated grade 300 maraging steels produced by SLM.

fabricated and heat-treated SLM specimens under uniaxial tension are shown in Fig. 13. The measured mechanical properties are summarized in Table 3. The hardness of the as-fabricated specimen is within range of 35–36 HRC, by contrast, the hardness is significantly improved and reached 53–55 HRC after age-hardening treatment. Meanwhile, age treatment leads to an improvement of tensile performances. The ultimate tensile strength (UTS) increases from 1165 MPa to 2014 MPa, which increases by about 73%. However, the break elongation (El) reduces from 12.44% to 3.28% after aging, which indicates that the desired strength significantly increased at the cost of aggravating troublesome brittleness upon aging. As revealed in Table 3, the hardness and tensile strengths (UTS and yield strength YS) of the as-fabricated and agingtreated SLM specimens absolutely meet the wrought standard requirements, while only the elongation of aging-treated specimen is inferior to the standard value. Metallic materials combining high-strength with novel toughness are always of such desirable. Therefore, an attempt, solution treatment before aging, has been made to improve strength and ductility. Elated to find that, although the hardness and strength slightly decrease, the ductility of solution-treated specimen (El = 14.4%) increases in comparison to as-fabricated specimen (El = 12.44%). As revealed by the XRD pattern in Fig. 11, maraging steels exhibit a fully martensitic structure after solution treatment. Due to the high nickel content and the virtual absence of carbon, martensite in maraging steels is no longer a supersaturated solid solution of carbon in α-Fe; instead, it is an iron nickel martensite, i.e., a solid solution of nickel (as well as some Co and Mo) in α-Fe. Hence, these alloys are soft and deformable with a measured hardness of about 30 HRC. For improved strength and ductility, the integrated solution and subsequent age treatment were adapted. Notably, the hardness (52–54 HRC) and strength (UTS = 1943 MPa, YS = 1882 MPa), as well as elongation (5.6%) of solutionaged specimens exactly located in the standard ranges. Fractographies of as-fabricated and heat-treated specimens are revealed in Fig. 14. The as-fabricated specimen broke after substantial plastic deformation and a great quantity of big and deep dimples formed in the fracture morphologies (Fig. 14a), which is a typical trans granular Table 3 Comparison of the mechanical properties of grade 300 maraging steels fabricated by SLM and conventional wrought processing methods.

Fig. 12. Hardness of the as-fabricated and heat-treated grade 300 maraging steels produced by SLM.



YS (MPa)

El (%)


SLM as-fabricated SLM aged SLM solution SLM solution-aged Wrought [21] Wrought aged [52,53]

1165 ± 7 2014 ± 9 1025 ± 5 1943 ± 8 1000–1170 1930–2050

915 ± 7 1967 ± 11 962 ± 6 1882 ± 14 760–895 1862–2000

12.44 ± 0.14 3.28 ± 0.05 14.40 ± 0.35 5.60 ± 0.08 6–15 5–7

35–36 53–55 28–29 52–54 35 52


C. Tan et al. / Materials and Design 134 (2017) 23–34

Fig. 14. Fracture morphologies of (a) as-fabricated, (b) aged and (c) solution-aged grade 300 maraging steels produced by SLM.

ductile fracture. During the tensile process, the specimen undergoes a large plastic deformation and the micro-cavities arise at the precipitates or imperfections. Meanwhile, large stresses, caused by the cavities, will lead to more micro-cavities in return. Eventually, the cavities conjoin together and the fast growing tears cause the fracture. As for the agehardened specimen, shown in Fig. 14b, the cleavage fractures are obviously present, and the dimples are shallow and limited, so the plastic deformation is insufficient, and its main fracture mode is a brittle fracture. Besides, as exhibited in the Fig. 14c, the solution-age treated specimen with presence of smaller dimples integrates ductile and brittle fractures. Therefore, the integrated solution-aging treatment for this SLM produced maraging steels, is a more preferable treatment method for obtaining novel mechanical performances. Reasons accounting for this SLM produced ultra-high strength and novel ductility grade 300 maraging steel can be explained as follow. Two strengthening mechanisms can account for the improvement of the hardness and strength upon age hardening, i.e., precipitation strengthening and strain strengthening. The relatively soft (30 HRC) and ductile low carbon lath martensite in maraging steel normally contains a high dislocation density in the order of 1011–12/cm2 [23]. During aging treatment, as revealed in Fig. 9, an uniform and dense distribution of Ni3(Ti, Al, Mo) intermetallic nanoprecipitates formed in the martensite matrix, these precipitates serve to strengthen the martensitic matrix. The strengthening can be explained by the following Orowan relationship:   Gb λ−d Þ In 2πKðλ−dÞ  2b 1 1 1 ¼ þ1 K 2 1−v

8 > > < σA ¼ σ0 þ > > :

where σA is the yield strength of age hardened specimen, σ0 is the yield strength of martensite matrix, which equals to the yield strength of solution-treated specimens (i.e. 962 MPa). G is the shear modulus of the matrix, b is the Burgers vector, λ is the interspace of nanoprecipitates, v is Poisson's ratio of the matrix and K is related to v. Fig. 9a exhibited that the dispersive distributed needle-shaped nanoprecipitates are about 8 nm in diameter and 30 nm in length, it can be simplified as a sphere of equivalent volume of diameter d = 14 nm. The λ is observed as about 25 nm in average. Taking G = 71 GPa, v = 0.3 and Burgers vector b = 0.249 nm [23,24], the yield strength of 2142 MPa for age hardened specimen was theoretically calculated. The calculation agrees reasonably with the measured value (1967 MPa). Therefore, the precipitated-phase strengthening is one of the main hardening mechanism. Besides, the precipitate particles are semi-coherent with the martensite matrix, which will cause little distortion and increase the amount of dislocations in the matrix (as observed in Fig. 9a). So the distortion and dislocations in the matrix will lead to strain strengthening [54]. The strain strengthening can be characterized by Ludwik equation σ = σo + Kεn, where σo is the yield strength, K is the strength coefficient and n is the strain hardening exponent. The value of strain hardening exponent is influenced by the microstructural characteristics during

aging period, such as dislocation density, the size and volume fraction of precipitates, and domain boundaries, which act as a barrier to dislocation movement [26]. With regard to novel ductility of the grade 300 maraging steels, it can be explained as follow. Firstly, the barely carbon is kept below 0.03 wt% in order to avoid the formation of brittle phases, such as TiC, which will adversely affect the ductility and toughness. The carbon is no longer added as a harder of martensites. Secondly, the amount of pinned dislocations was decreased due to low carbon and nitrogen concentration, so that great amounts of glissile dislocations can relax stress through regional plastic deformation and prevent a larger stress concentration at the stress concentration locations. Thirdly, martensite distortion during age hardening triggered high density dislocations, which is conducive to the formation of dispersive precipitates. The precipitates increase the motion resistance of dislocations and hinder the long range movement of dislocations; however, the short range movement of dislocations is still feasible. As a consequence, the maraging steels are able to combine ultra-high strength with novel ductility. 4. Conclusions High-performance grade 300 maraging steel were successfully fabricated by selective laser melting. Heat treatments including aging and integrated solution-aging treatments were crucial to improve the mechanical performances. The primary conclusions are summarized as follows: (1) High formation qualities including a low surface roughness (4.16–4.79 μm) and a high relative density (99.98%) were obtained with the optimized process parameters. A plenty of nanocrystals could be clearly observed on the fresh surface of as-fabricated specimens owing to the high solidification rate. (2) The submicron cellular and acicular microstructures could be formed in response to the instant melting and rapid solidification during SLM. The microstructural evolutions were mainly determined by degree of undercooling, interfacial energy and the ratio of temperature gradient to solidification rate (G/R). (3) Nanoprecipitates triggered by intrinsic heat treatment (IHT) and amorphous phases generated through extremely fast solidification were clearly observed in as-fabricated specimens. HRTEM images clearly exhibited massive sized needle-shaped (about 6–10 nm in diameter and 15–45 nm in length) nanoprecipitates Ni3X (X = Ti, Al, Mo) and 50–60 nm sized spherical core-shell structural nanoparticles embedded in amorphous matrix of age hardened specimens. These precipitates accounted for precipitation strengthening and followed the Orowan mechanism. The austenite reversion and probable phase transformations were revealed by XRD and TEM. (4) The integrated solution-aging treatment was a more preferable heat treatment method for this kind of SLM fabricated maraging steels. The hardness, tensile strengths as well as ductility of asfabricated and solution-aged specimens reached standard wrought level.

C. Tan et al. / Materials and Design 134 (2017) 23–34

The results demonstrated that the high-performance grade 300 maraging steel which was comparable to the standard wrought levels could be produced by SLM additive manufacturing and subsequent solution-aging heat treatments. The findings in this work are useful for industrial guidance and extensive applications of the SLM produced maraging steels. Besides, the results also indicate the possibility of producing full amorphous metals and in-situ age hardening of maraging steels without need of further heat treatments by laser based additive manufacturing.

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