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High strength and ductility AlCrFeNiV high entropy alloy with hierarchically heterogeneous microstructure prepared by selective laser melting
Journal of Alloys and Compounds 813 (2020) 152196
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High strength and ductility AlCrFeNiV high entropy alloy with hierarchically heterogeneous microstructure prepared by selective laser melting Haili Yao a, Zhen Tan a, *, Dingyong He a, b, Zhenlu Zhou a, Zheng Zhou a, Yunfei Xue c, Li Cui a, Lijia Chen a, Guohong Wang a, Ying Yang d a
College of Materials Science and Engineering, Beijing University of Technology, Beijing, 100124, PR China Beijing Engineering Research Center of Eco-materials and LCA, Beijing, 100124, PR China School of Materials Science and Engineering, Beijing Institute of Technology, Beijing, 100086, PR China d Beijing Center for Physical and Chemical Analysis, Beijing, 100089, PR China b c
a r t i c l e i n f o
a b s t r a c t
Article history: Received 27 June 2019 Received in revised form 3 September 2019 Accepted 6 September 2019 Available online 10 September 2019
Almost fully dense AlCrFeNiV HEA consisting of face-centred-cubic (FCC) solid solution and L12 nano phase was prepared from gas-atomized alloy powder with optimized selective laser melting (SLM) processing parameters. Microstructure characterization reveals the presence of hierarchical structures including columnar grains, sub-grains, L12 nano phase and dislocations in SLMed HEA. Unique columnar grains ranging from several tens of microns up to 200 mm grow along the direction of the temperature gradient. High cooling speed and non-equilibrium solidification during SLM process induced the formation of sub-grains in every columnar grain, accompanied with the heterogeneous distribution of dislocations and L12 nano phase. The SLMed AlCrFeNiV HEA exhibited an outstanding combination of high strength (ultimate tensile strength ~1057.47 MPa) and excellent ductility (plastic strain ~30.3%). The characteristic hierarchically heterogeneous structure contributes to the increase of strength without losing ductility. The sub-grains contribute significantly to the enhanced strength through dislocation hardening. The excellent ductility is correlated with the progressive work-hardening mechanism regulated by the heterogeneous distribution dislocation and L12 nano phase within sub-grains. © 2019 Elsevier B.V. All rights reserved.
Keywords: Selective laser melting (SLM) High entropy alloy Hierarchically heterogeneous microstructure High strength-ductility L12 nano phase
1. Introduction As a revolutionary alloy, high entropy alloys (HEAs) contain five or more principle alloying elements in near-equimolar ratios to stabilize solid solution phases by maximizing configurational entropy [1]. Because of the various multiple elements composition and corresponding specific structure, several HEAs exhibit excellent mechanical properties that conventional alloys are difficult to achieve [2e4]. For example, CrMnFeCoNi HEA with single-phase face-centred cubic (FCC) solid solution shows outstanding combinations of strength (1280 MPa) and fracture toughness (exceed 300 MPa m1/2) at cryogenic temperatures down to 77 K [5].
* Corresponding author. Department: College of Materials Science and Engineering, Beijing University of Technology, 100 Pingleyuan Street, Chaoyang District, Beijing, 100124, PR China. E-mail address: [email protected] (Z. Tan). https://doi.org/10.1016/j.jallcom.2019.152196 0925-8388/© 2019 Elsevier B.V. All rights reserved.
Metastable Fe50Mn30Co10Cr10 high-entropy dual-phase alloy can simultaneously achieve greatly improved strength and ductility [6]. Fabrication processes especially the thermomechanical (TM) treatment also play a significant role on the achievement of optimal mechanical performance of HEAs [5,6]. However, the treatments including cold forging, cross rolling and recrystallization annealing are particular time-consuming work to HEAs [7e10]. Selective laser melting (SLM), as a typical technology of additive manufacturing (AM), can three dimensionally (3D) fabricate components with intricated shape and refined resolution [11e13]. Now, SLM has been employed on the fabricating HEAs including AlCoCrFeNi [14e16], CoCrFeNiMn [17,18], and AlCrCuFeNi [19]. These SLMed HEAs showed improved strength or ductility compared with the casting HEAs. Grain refinement caused by extremely high cooling rate was regarded as a significant reason for the improvement of mechanical performance of the SLMed HEAs, which was usually obtained through traditional TM treatments [8,10,20]. Thus, SLM could fabricate HEA with good mechanical
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performance without TM treatment. Besides the grain refinement, a hierarchically heterogeneous microstructure could also be found in SLMed specimens because of the forming mode of SLM including melt pools connection and layer-layer deposition [21,22]. Recent researches also demonstrated that dislocation network or subgrains also contributed to the enhancement of strength and ductility [23,24]. Accordingly, the improvement of strength and ductility of SLMed HEAs may be the comprehensive effect of grain refinement and hierarchically heterogeneous microstructure. AlCrFeNiV HEA was employed in the present research, this alloy could achieve ultrahigh tensile strength (1.9 GPa) while retaining good ductility (9%) due to the spinodal orderedisorder nanostructure obtained through TM treatment [25]. The specific high cooling rate and non-equilibrium solidification during SLM may induce the characteristic microstructure of HEA, and leading to the significant difference of mechanical performance. Consequently, microstructure and mechanical property of the SLMed AlCrFeNiV HEA were investigated. The similarities and differences of SLMed HEA, as-cast HEA and HEA after TM treatment were also compared to estimate the availability of SLM for fabricating high performance HEAs.
sample with relative density of 99.88% was fabricated with a laser power of 140 W, scanning speed of 900 mm/s and scanning pitch of 0.05 mm. Tensile test samples were prepared with the optimum parameters and the size diagram is shown in Fig. 2. 2.3. Characterization of microstructure and mechanical properties The relative density of each final part was estimated by the Image method. Based on the ImageJ software, the relative density of SLMed samples was measured by processing optical microscope (OM) images of side view, and the average of eight individual measurements was taken as the relative density. The X-ray diffraction (XRD) patterns of raw powder, samples built by SLM and casting were collected on a Rigaku Ultima IV X-ray diffraction. The microstructures of the surface and cross section were characterized by SEM, energy-dispersive X-ray spectroscopy (EDS, achiS-4800), electron backscattered diffraction (EBSD, FEIQUANTA FEG 650) and transmission electron microscope (TEM, FEI talos-F200X). The mechanical properties were tested by INSTRON 5985 electronic universal material testing machine.
2. Materials and experiment methods
3. Results
2.1. Materials
3.1. Microstructure
The raw Al0.5Cr0.9FeNi2.5V0.2 high entropy alloy (HEA) powder prepared by gas-atomized was employed in the present research. The microstructure of powder particles was observed by achiS-4800 scanning electron microscope (SEM). Fig. 1a shows the SEM images of the raw HEA powders, and most of the powders exhibited spherical shape. Dendrites could be found in the cross section of powder (Fig. 1b). The size distribution of the powders was mainly in the range of 20 mme80 mm through Tod Laser Particle Sizer. The oxygen content of powder measured by Bruker G8 Galileo analyzer was 0.014%.
Fig. 3 shows the XRD patterns of raw powder, as-cast HEA and SLMed HEA. FCC phase could be found on all the XRD patterns, which indicated the similar phase composition. Fig. 4 shows the SEM micrographs of SLMed HEA. Pores were significantly restrained both on the top view and the side view along the building direction, while various structure morphologies were observed. As shown in Fig. 4a, the melted scan tracks are visible as ovals with irregular distribution. Each oval on the top view contains a number of cellular grains. On the side view, Fig. 4c shows arcshaped features (width and depth are ~ 100 mm and ~60 mm, respectively), and the distance between these arcs ranges ~ 50 mm, which suggests that the laser beam melted the material to a depth of more than one layer of added powder (~30 mm). The arcs were occupied by the columnar grains (grain boundaries are marked by white dotted lines) and they were oriented normal in the solidification direction. The width and length of the columnar grain were mainly ~15 mm and 75e200 mm, respectively. The columnar grains grow throughout more than one melted scan track along the building direction. Therefore, SLMed HEA sample was mainly composed by columnar grains. SEM images in Fig. 4d presents details of the columnar grains of the samples. Columnar grains possessed very fine cellular-strip sub-grains, which is mainly attributed to rapid cooling of the molten pool. The EBSD analysis depicted in Fig. 5 provides insights into the
2.2. Experimental equipment and process The process was conducted by a SLM machine (EOS M100). The laser beam has maximum laser power of 200 W and the diameter of the focused laser spot size is ~50 mm. Cubic samples with dimensions of 8 8 8 mm were fabricated to acquire the optimal parameters. The laser scanning path of each layer was rotated 67 by the path of the previous layer. The pre-heating temperature of substrate was set to 353 K. The process was carried out under argon atmosphere with the concentrations of O2 controlled 0.1%. During SLM process, the relative density of sample was affected by the combined effects of the laser power (P, W), the scanning speed (v, mm/s), the hatching space (h, mm) and the layer thickness (d, mm). Orthogonal experiments were designed to explore the effect of processing parameters on the relative density of samples. The
Fig. 1. SEM images of the AlCrFeNiV HEA powders: (a) Surface morphology; (b) Cross section morphology.
Fig. 2. Configuration of SLMed AlCrFeNiV HEA for quasi-static tensile test.
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Fig. 3. XRD patterns of the raw powders, as-cast, and SLMed AlCrFeNiV HEA.
Fig. 4. SEM images of the SLMed AlCrFeNiV HEA: (a) Top view 200; (b) Top view 1000; (c) Side view 200; (d) Side view 1000.
microstructural features and texture development in the SLMed HEA. Fig. 5a shows the inverse pole figure (IPF) map on the side view (along the building direction). Grains under EBSD exhibited a ripple pattern instead of a traditional faceted morphology, while molten pool tracks could not be observed in IPF map. According to the size of molten pool and grain distribution in SEM images (Fig. 4d), the white dotted line outlines the boundary of the molten pool. Most of the elongated grains through multiple melt pools were parallel to the building direction. The red color grains presented preferred orientations of <001> parallel to the building direction. Obviously, the strong orientation of the FCC phase was <100> along the building direction [22,23,26]. The color and morphology variation within different grains suggests slight variations of crystallographic orientation among the constituent subgrains. Fig. 5b represents the grain boundaries map with various angles. SLMed HEA contained a large fraction of low-angle grain boundaries (LAGBs, 2 e10 ). It was found that a significant fraction
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(68.22%) of low angle grain boundaries (LAGBs) with misorientation angles lower than 10 could be identified. The sub-grains which appeared in each columnar grain contributed to the high amount of LAGBs. High angle grain boundaries (HAGBs) only accounted for few proportions, as shown in Fig. 5d. Orientation gradients or local misorientations across grains were also observed, as evidenced in the kernel average misorientation (KAM) map shown in Fig. 5c. A comparison of Fig. 5b and d indicates the direct correlation between orientation deviation and LAGBs. The estimated average grain size (Fig. 5e) based on the HAGBs is ~69.32 mm, with a large standard deviation, leading to a grain area distribution that spans several orders of magnitude. The present EBSD data revealed the unconventional structure of SLMed HEA including ripple-like shapes grain, broad grain-size distribution, and large fraction of LAGBs. Fig. 6 shows further investigation about the sub-grains in SLMed HEA. Two different sub-grains morphologies with cellular (diameters ~ 500 nm) and strip (widths 30 nme60 nm) shapes were distributed in different grains, as shown in Fig. 6b and c (corresponding to the selected area I and area II in Fig. 6a, respectively). Due to complex temperature field formed by the layer-layer processing feature and the laser scanning path of SLM, there is no conspicuous regularities of distribution of strip and cell sub-grains in a single molten pool. TEM image of the HEA in Fig. 6d indicated that the sub-grains boundaries are consisted of dislocation walls. And the distribution of dislocations inner sub-grains also exhibit notable heterogeneity (Fig. 6d). Some sub-grains exhibited almost dislocationfree, while some of them contained a high density of dislocations, as shown in Fig. 6e and f (corresponding to the area III and area IV in Fig. 6d). And the dislocations always concentrated at sub-grain boundaries. Fig. 7b shows the selected area electron diffraction (SAED) pattern corresponding to circular area of Fig. 7a, which indicates that SLMed HEA contain two phases, FCC phase and L12 nano phase (namely Ni3Al). HRTEM image and corresponding fast fourier transformation (FFT) pattern of the selected square region in Fig. 7c also demonstrated the existence of L12 nano phase, and the L12 nano phase could be seen in the FFT pattern of region (marked by the white arrow). Coherent L12 nano phase with the length about 5e10 nm can be observed in Fig. 7c. The lattice constants of FCC and L12 phase were very similar (0.3597 nm for FCC and 0.3575 nm for L12), and the interplanar spacings of FCC and L12 nano phase were approximately 0.207 nm, which was consistent with the results shown in XRD (~0.207 nm). 3.2. Mechanical properties Fig. 8 shows the true stress-strain curves of the SLMed HEA, ascast HEA and HEA after TM treatment under tensile test at room temperature. The yield stress and the ultimate tensile strength of SLMed HEA (651.36 MPa, 1057.21 MPa) were superior to those of casting samples (489.63 MPa, 873.09 MPa). Although the strength of SLMed tensile samples was lower than that of the HEA after TM treatment [25], the ductility of SLM samples improved remarkably. Fig. 9 shows SEM images of the tensile fracture morphology of SLMed samples after tensile tests, respectively. The size of ductile dimples was much smaller than those obtained in as-cast HEA [27e29]. A great number of dimples with small size indicated more surface energy produced during fracture process, which reflected higher strength. In addition, the size of dimples depended to a great extent on the size of sub-grains [30]. As shown in Fig. 9b (a magnification of the rectangular area in Fig. 9a), the average dimple size was approximate 500 nm which was almost the same as that of sub-grains (Fig. 6c).
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Fig. 5. EBSD analysis of SLMed AlCrFeNiV HEA: (a) IPF map; (b) Image quality (IQ) map with LAGBs and HAGBs superimposed; (c) Kernel average misorientation (KAM) map; (d) Misorientation angle plot corresponding to (b); (e) Grain size distribution.
4. Discussion 4.1. Formation of hierarchically heterogeneous microstructure Microstructure characterization reveals the heterogeneity of hierarchical structures including columnar grains, sub-grains, L12 nano phase and dislocations in SLMed AlCrFeNiV. Columnar grains, as the typical grain morphology of SLM, usually formed through epitaxial growth due to the re-melting of the previous depositing layer and the significant temperature gradient between the depositing layers [21,31]. Columnar grains in the present study exhibited preferential growth aligned closely with the maximum heat flow direction at the solid-liquid interface. A schematic diagram of re-melting of adjacent layers during SLM process was shown in Fig. 10. Columnar grains showed various growth direction and size within a single molten pool compared with other metal materials. Because of the lower thermal conductivity of HEA, the heat from the central of molten pool could not be transferred to the surrounding of heat source in time, which increases the temperature gradient and promotes the epitaxial growth of grains [32e34]. Compared with metal materials with high thermal conductivity, such as Cu alloy, the columnar grains of HEA samples have a longer span of molten pool. Sub-grains have been observed in a range of metals or alloys prepared by SLM [17,24,35], however, there is no consensus about the formation mechanism of sub-grains. SLM is normally related to a very rapid melting and solidification process which is far from equilibrium condition, and the microstructure is strongly affected
by the high temperature gradient, high cooling rate and large undercooling depending on the metal property, specimen geometry, and laser parameters [36e38]. The formation of the sub-grains in the present HEA was related to the Bernard convection. Bernard convection occurred under the driving of surface tension when the temperature gradient of the liquid layer exceeded the critical temperature gradient [39,40]. Larger vertical temperature gradient induced the vertical melt convection and cellular sub-grains formed. With laser scanning, horizontal temperature gradient in the center of the laser was dominant, and the horizontal melt convection promoted the formation of strip sub-grains. In general, deformation tends to induce the formation of dislocations, however, the present SLMed HEA exhibited a relative high content of dislocation without deformation. This could be attributed to the rapid heating and cooling rate in local micro-region during SLM. The rapid heating and cooling rate not only increased the vacancy concentration [24], but also generated large thermal residual stresses in HEA. The vacancy concentration could be regarded as the source of dislocation and residual stress could produce dislocations, and the dislocation multiplication resulted in the high-density dislocations distribution in HEA. In addition, heterogeneous distribution of dislocations could also be found in sub-grains. Different cooling rate of sub-grains led to the variety of vacancy concentration and residual stress level, inducing the different quantity of dislocations. Besides, dislocations multiplication and interaction occurred easily in smaller sub-grains for the smaller mean free path of dislocation and hindering effect of subgrain boundaries.
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Fig. 8. True stress-strain curves of the as-cast, SLMed AlCrFeNiV HEA and corresponding result in Ref.25 after quasi-static tensile test.
Fig. 6. (a) SEM images of the SLMed AlCrFeNiV HEA; (bec) Magnification micrograph of area I and area II in (a); (d) Bright-field TEM images of the cellular sub-grain; (eef) Magnification micrograph of area III and area IV in (d). Fig. 9. (a) SEM images of the fracture surfaces of SLMed AlCrFeNiV HEA after quasistatic tensile test; (b) Magnification micrograph of the selected square region in (a).
reasons for the formation of L12 nano phase. Although the spinodal decomposition was extremely restrained in SLMed HEA due to the high cooling rate, supersaturated FCC phase, sever lattice distortion and residual stress existed in the SLMed HEA. All of the factors mentioned above provide the driving force for the formation of L12 nano phase. The subsequent laser scanning, which could be regarded as localized heat treatment supplied the thermal dynamic conditions.
4.2. Improvements of strength and ductility
Fig. 7. (a) Bright field TEM images of the SLMed AlCrFeNiV HEA; (b) Selected area electron diffraction (SAED) pattern of the circle region in (a); (ced) HRTEM image and corresponding FFT pattern of the selected square region in (c).
The most remarkable finding was the formation of L12 nano phase in SLMed AlCrFeNiV HEA, and the L12 nano phase was usually obtained in cast AlCrFeNiV HEA through TM treatment [25]. And in the HEA through TM treatment, high Ni/Al ratio, spinodal decomposition and the ordering of the disorder FCC phase were the
Fig. 11 exhibits comparison of the yield stress and elongation of SLMed HEA with different element systems including FeCoCrNi [37], CoCrFeMnNi [17,38,41], AlFeCoCrNi [16]. The present SLMed AlCrFeNiV HEA exhibited higher yielding strength and plasticity. All the structure in the macro-, micro-, and nano-level has contributed to the overall tensile properties of the present HEA. Fig. 12 shows the schematic diagram of the dislocation formation and motion, accompanied by the interaction between dislocations and obstacles (including L12 nano phase, pre-existence dislocations, sub-grain boundaries and grain boundaries). Dislocations tended to appear firstly in the sub-grains with low dislocation density, and dislocations motion would encounter L12 nano phase and pre-existence dislocations inner sub-grains. L12 nano phase provided strong diffuse attractive obstacles to trap moving dislocations, and create antiphase boundaries (APBs) on the slip planes of the ordered L12 nano phase when they are sheared by dislocations. Simultaneously, the highly coherent and low misfit FCCL12 interfaces could minimize the elastic strain accumulation resulting
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Fig. 10. Schematic diagram of grain growth during the re-melting of adjacent forming layers in SLMed AlCrFeNiV HEA.
Fig. 11. Yield strength and elongation of SLMed AlCrFeNiV HEA and SLMed FeCoCrNi [37], AlFeCoCrNi [14] and CoCrFeMnNi [17,38] HEAs from reference under quasi-static tensile condition.
from dislocation shear and hence prevented crack initiation at these interfaces. The interaction between the newly formed dislocations and pre-existence dislocations caused more dislocations multiplication, which also contributed to the strength improvement of SLMed HEA. Due to the dislocation network of sub-grain boundary, it acts as stable and soft barriers that hinder dislocation motion for strength, but meanwhile guaranteed a continuous plastic flow by allowing dislocation from transmitting. Thus, sub-grain boundaries possess stronger dislocation storage capacity and dislocations will be more easily trapped and stored in sub-grain boundaries so that dislocation slip is particularly difficult [17]. When the dislocation density in one sub-grain increases to be close to that of other sub-grains, the strength of sub-grains tends to be consistent, thus achieving the synergistic movement of dislocations in the tensile process. The misorientation between sub-grains may also contribute to the stability of the sub-grain boundaries. In addition, dislocations pileup at sub-grain boundaries caused the stress concentration. The sub-micron size dimples on the fracture surface are the results of breaking sub-grains. With the increase of stress, dislocations passed through subgrain boundaries to grain boundaries, which would become new obstacles to dislocation motion. Dislocations pile up near grain boundaries, which resulting in more serious strain hardening. Although L12 nano phase and sub-grain boundary could contribute to the improvement of strength, their strengthen effect was still far behind the grain boundary. The grain size in the SLMed HEA was much larger than that of the HEA after TM treatment (~2.8 mm) [25], which exhibited lower strength. The plastic deformation mainly depends on the dislocation motion, and the effect of hierarchically heterogeneous structure on the dislocation motion plays a significant role in mechanical performance of SLMed HEA. Hierarchical structure from nano-scale (L12 nano phase, pre-existence dislocations) to sub-micron scale (sub-grain boundaries and grain boundaries) hinders the dislocation motion progressively. And heterogeneous distribution of the dislocation density and sub-grains morphology also made the deformation progressively. As a consequence, the SLMed AlCrFeNiV HEA achieves the high strength and ductility.
Fig. 12. Schematic diagram of the interaction between dislocations and obstacles (including L12 nano phase, pre-existence dislocations, sub-grain boundaries and grain boundaries).
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5. Conclusions 1) SLMed AlCrFeNiV HEA consisting of FCC phase and L12 nano phase shows a hierarchical heterogeneous structure including columnar grains, sub-grains, L12 nano phase and dislocations. Cellular or strip sub-grains were distributed in each columnar grain, and dislocations and L12 nano phase were distributed in sub-grains. 2) High cooling speed and non-equilibrium solidification during SLM process caused extremely high temperature gradient and large surface tension gradient, accompanied with the Bernard convection along the vertical or horizontal direction, leading to the formation of cellular or strip sub-grains. Dislocations formed due to the increased the vacancy concentration and large thermal residual stresses produced by rapid heating and cooling rate in SLMed process. Supersaturated FCC phase, sever lattice distortion, residual stress, and subsequent laser scanning provide the driving force for the formation of L12 nano phase. 3) The SLMed AlCrFeNiV HEA exhibited an outstanding combination of high strength (ultimate tensile strength ~1057.47 MPa) and excellent ductility (plastic strain ~30.3%). Hierarchical structure from nano-scale (L12 nano phase, pre-existence dislocations) to sub-micron scale (sub-grain boundaries and grain boundaries) hindered the dislocation motion progressively. And heterogeneous distribution of the dislocation density and subgrains morphology also made the deformation progressively, achieving the balance of high strength and ductility.
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Acknowledgement This work is supported by National Natural Science Foundation of China (Grant No. 51901004 and 51771005), financial supports by National Key R&D Program of China (No. 2017YFB0305800) and Beijing Natural Science Foundation (Grant No. 2194067).
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
High strength and ductility AlCrFeNiV high entropy alloy with hierarchically heterogeneous microstructure prepared by selective laser melting Haili Yao a, Zhen Tan a, *, Dingyong He a, b, Zhenlu Zhou a, Zheng Zhou a, Yunfei Xue c, Li Cui a, Lijia Chen a, Guohong Wang a, Ying Yang d a
College of Materials Science and Engineering, Beijing University of Technology, Beijing, 100124, PR China Beijing Engineering Research Center of Eco-materials and LCA, Beijing, 100124, PR China School of Materials Science and Engineering, Beijing Institute of Technology, Beijing, 100086, PR China d Beijing Center for Physical and Chemical Analysis, Beijing, 100089, PR China b c
a r t i c l e i n f o
a b s t r a c t
Article history: Received 27 June 2019 Received in revised form 3 September 2019 Accepted 6 September 2019 Available online 10 September 2019
Almost fully dense AlCrFeNiV HEA consisting of face-centred-cubic (FCC) solid solution and L12 nano phase was prepared from gas-atomized alloy powder with optimized selective laser melting (SLM) processing parameters. Microstructure characterization reveals the presence of hierarchical structures including columnar grains, sub-grains, L12 nano phase and dislocations in SLMed HEA. Unique columnar grains ranging from several tens of microns up to 200 mm grow along the direction of the temperature gradient. High cooling speed and non-equilibrium solidification during SLM process induced the formation of sub-grains in every columnar grain, accompanied with the heterogeneous distribution of dislocations and L12 nano phase. The SLMed AlCrFeNiV HEA exhibited an outstanding combination of high strength (ultimate tensile strength ~1057.47 MPa) and excellent ductility (plastic strain ~30.3%). The characteristic hierarchically heterogeneous structure contributes to the increase of strength without losing ductility. The sub-grains contribute significantly to the enhanced strength through dislocation hardening. The excellent ductility is correlated with the progressive work-hardening mechanism regulated by the heterogeneous distribution dislocation and L12 nano phase within sub-grains. © 2019 Elsevier B.V. All rights reserved.
Keywords: Selective laser melting (SLM) High entropy alloy Hierarchically heterogeneous microstructure High strength-ductility L12 nano phase
1. Introduction As a revolutionary alloy, high entropy alloys (HEAs) contain five or more principle alloying elements in near-equimolar ratios to stabilize solid solution phases by maximizing configurational entropy [1]. Because of the various multiple elements composition and corresponding specific structure, several HEAs exhibit excellent mechanical properties that conventional alloys are difficult to achieve [2e4]. For example, CrMnFeCoNi HEA with single-phase face-centred cubic (FCC) solid solution shows outstanding combinations of strength (1280 MPa) and fracture toughness (exceed 300 MPa m1/2) at cryogenic temperatures down to 77 K [5].
* Corresponding author. Department: College of Materials Science and Engineering, Beijing University of Technology, 100 Pingleyuan Street, Chaoyang District, Beijing, 100124, PR China. E-mail address: [email protected] (Z. Tan). https://doi.org/10.1016/j.jallcom.2019.152196 0925-8388/© 2019 Elsevier B.V. All rights reserved.
Metastable Fe50Mn30Co10Cr10 high-entropy dual-phase alloy can simultaneously achieve greatly improved strength and ductility [6]. Fabrication processes especially the thermomechanical (TM) treatment also play a significant role on the achievement of optimal mechanical performance of HEAs [5,6]. However, the treatments including cold forging, cross rolling and recrystallization annealing are particular time-consuming work to HEAs [7e10]. Selective laser melting (SLM), as a typical technology of additive manufacturing (AM), can three dimensionally (3D) fabricate components with intricated shape and refined resolution [11e13]. Now, SLM has been employed on the fabricating HEAs including AlCoCrFeNi [14e16], CoCrFeNiMn [17,18], and AlCrCuFeNi [19]. These SLMed HEAs showed improved strength or ductility compared with the casting HEAs. Grain refinement caused by extremely high cooling rate was regarded as a significant reason for the improvement of mechanical performance of the SLMed HEAs, which was usually obtained through traditional TM treatments [8,10,20]. Thus, SLM could fabricate HEA with good mechanical
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performance without TM treatment. Besides the grain refinement, a hierarchically heterogeneous microstructure could also be found in SLMed specimens because of the forming mode of SLM including melt pools connection and layer-layer deposition [21,22]. Recent researches also demonstrated that dislocation network or subgrains also contributed to the enhancement of strength and ductility [23,24]. Accordingly, the improvement of strength and ductility of SLMed HEAs may be the comprehensive effect of grain refinement and hierarchically heterogeneous microstructure. AlCrFeNiV HEA was employed in the present research, this alloy could achieve ultrahigh tensile strength (1.9 GPa) while retaining good ductility (9%) due to the spinodal orderedisorder nanostructure obtained through TM treatment [25]. The specific high cooling rate and non-equilibrium solidification during SLM may induce the characteristic microstructure of HEA, and leading to the significant difference of mechanical performance. Consequently, microstructure and mechanical property of the SLMed AlCrFeNiV HEA were investigated. The similarities and differences of SLMed HEA, as-cast HEA and HEA after TM treatment were also compared to estimate the availability of SLM for fabricating high performance HEAs.
sample with relative density of 99.88% was fabricated with a laser power of 140 W, scanning speed of 900 mm/s and scanning pitch of 0.05 mm. Tensile test samples were prepared with the optimum parameters and the size diagram is shown in Fig. 2. 2.3. Characterization of microstructure and mechanical properties The relative density of each final part was estimated by the Image method. Based on the ImageJ software, the relative density of SLMed samples was measured by processing optical microscope (OM) images of side view, and the average of eight individual measurements was taken as the relative density. The X-ray diffraction (XRD) patterns of raw powder, samples built by SLM and casting were collected on a Rigaku Ultima IV X-ray diffraction. The microstructures of the surface and cross section were characterized by SEM, energy-dispersive X-ray spectroscopy (EDS, achiS-4800), electron backscattered diffraction (EBSD, FEIQUANTA FEG 650) and transmission electron microscope (TEM, FEI talos-F200X). The mechanical properties were tested by INSTRON 5985 electronic universal material testing machine.
2. Materials and experiment methods
3. Results
2.1. Materials
3.1. Microstructure
The raw Al0.5Cr0.9FeNi2.5V0.2 high entropy alloy (HEA) powder prepared by gas-atomized was employed in the present research. The microstructure of powder particles was observed by achiS-4800 scanning electron microscope (SEM). Fig. 1a shows the SEM images of the raw HEA powders, and most of the powders exhibited spherical shape. Dendrites could be found in the cross section of powder (Fig. 1b). The size distribution of the powders was mainly in the range of 20 mme80 mm through Tod Laser Particle Sizer. The oxygen content of powder measured by Bruker G8 Galileo analyzer was 0.014%.
Fig. 3 shows the XRD patterns of raw powder, as-cast HEA and SLMed HEA. FCC phase could be found on all the XRD patterns, which indicated the similar phase composition. Fig. 4 shows the SEM micrographs of SLMed HEA. Pores were significantly restrained both on the top view and the side view along the building direction, while various structure morphologies were observed. As shown in Fig. 4a, the melted scan tracks are visible as ovals with irregular distribution. Each oval on the top view contains a number of cellular grains. On the side view, Fig. 4c shows arcshaped features (width and depth are ~ 100 mm and ~60 mm, respectively), and the distance between these arcs ranges ~ 50 mm, which suggests that the laser beam melted the material to a depth of more than one layer of added powder (~30 mm). The arcs were occupied by the columnar grains (grain boundaries are marked by white dotted lines) and they were oriented normal in the solidification direction. The width and length of the columnar grain were mainly ~15 mm and 75e200 mm, respectively. The columnar grains grow throughout more than one melted scan track along the building direction. Therefore, SLMed HEA sample was mainly composed by columnar grains. SEM images in Fig. 4d presents details of the columnar grains of the samples. Columnar grains possessed very fine cellular-strip sub-grains, which is mainly attributed to rapid cooling of the molten pool. The EBSD analysis depicted in Fig. 5 provides insights into the
2.2. Experimental equipment and process The process was conducted by a SLM machine (EOS M100). The laser beam has maximum laser power of 200 W and the diameter of the focused laser spot size is ~50 mm. Cubic samples with dimensions of 8 8 8 mm were fabricated to acquire the optimal parameters. The laser scanning path of each layer was rotated 67 by the path of the previous layer. The pre-heating temperature of substrate was set to 353 K. The process was carried out under argon atmosphere with the concentrations of O2 controlled 0.1%. During SLM process, the relative density of sample was affected by the combined effects of the laser power (P, W), the scanning speed (v, mm/s), the hatching space (h, mm) and the layer thickness (d, mm). Orthogonal experiments were designed to explore the effect of processing parameters on the relative density of samples. The
Fig. 1. SEM images of the AlCrFeNiV HEA powders: (a) Surface morphology; (b) Cross section morphology.
Fig. 2. Configuration of SLMed AlCrFeNiV HEA for quasi-static tensile test.
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Fig. 3. XRD patterns of the raw powders, as-cast, and SLMed AlCrFeNiV HEA.
Fig. 4. SEM images of the SLMed AlCrFeNiV HEA: (a) Top view 200; (b) Top view 1000; (c) Side view 200; (d) Side view 1000.
microstructural features and texture development in the SLMed HEA. Fig. 5a shows the inverse pole figure (IPF) map on the side view (along the building direction). Grains under EBSD exhibited a ripple pattern instead of a traditional faceted morphology, while molten pool tracks could not be observed in IPF map. According to the size of molten pool and grain distribution in SEM images (Fig. 4d), the white dotted line outlines the boundary of the molten pool. Most of the elongated grains through multiple melt pools were parallel to the building direction. The red color grains presented preferred orientations of <001> parallel to the building direction. Obviously, the strong orientation of the FCC phase was <100> along the building direction [22,23,26]. The color and morphology variation within different grains suggests slight variations of crystallographic orientation among the constituent subgrains. Fig. 5b represents the grain boundaries map with various angles. SLMed HEA contained a large fraction of low-angle grain boundaries (LAGBs, 2 e10 ). It was found that a significant fraction
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(68.22%) of low angle grain boundaries (LAGBs) with misorientation angles lower than 10 could be identified. The sub-grains which appeared in each columnar grain contributed to the high amount of LAGBs. High angle grain boundaries (HAGBs) only accounted for few proportions, as shown in Fig. 5d. Orientation gradients or local misorientations across grains were also observed, as evidenced in the kernel average misorientation (KAM) map shown in Fig. 5c. A comparison of Fig. 5b and d indicates the direct correlation between orientation deviation and LAGBs. The estimated average grain size (Fig. 5e) based on the HAGBs is ~69.32 mm, with a large standard deviation, leading to a grain area distribution that spans several orders of magnitude. The present EBSD data revealed the unconventional structure of SLMed HEA including ripple-like shapes grain, broad grain-size distribution, and large fraction of LAGBs. Fig. 6 shows further investigation about the sub-grains in SLMed HEA. Two different sub-grains morphologies with cellular (diameters ~ 500 nm) and strip (widths 30 nme60 nm) shapes were distributed in different grains, as shown in Fig. 6b and c (corresponding to the selected area I and area II in Fig. 6a, respectively). Due to complex temperature field formed by the layer-layer processing feature and the laser scanning path of SLM, there is no conspicuous regularities of distribution of strip and cell sub-grains in a single molten pool. TEM image of the HEA in Fig. 6d indicated that the sub-grains boundaries are consisted of dislocation walls. And the distribution of dislocations inner sub-grains also exhibit notable heterogeneity (Fig. 6d). Some sub-grains exhibited almost dislocationfree, while some of them contained a high density of dislocations, as shown in Fig. 6e and f (corresponding to the area III and area IV in Fig. 6d). And the dislocations always concentrated at sub-grain boundaries. Fig. 7b shows the selected area electron diffraction (SAED) pattern corresponding to circular area of Fig. 7a, which indicates that SLMed HEA contain two phases, FCC phase and L12 nano phase (namely Ni3Al). HRTEM image and corresponding fast fourier transformation (FFT) pattern of the selected square region in Fig. 7c also demonstrated the existence of L12 nano phase, and the L12 nano phase could be seen in the FFT pattern of region (marked by the white arrow). Coherent L12 nano phase with the length about 5e10 nm can be observed in Fig. 7c. The lattice constants of FCC and L12 phase were very similar (0.3597 nm for FCC and 0.3575 nm for L12), and the interplanar spacings of FCC and L12 nano phase were approximately 0.207 nm, which was consistent with the results shown in XRD (~0.207 nm). 3.2. Mechanical properties Fig. 8 shows the true stress-strain curves of the SLMed HEA, ascast HEA and HEA after TM treatment under tensile test at room temperature. The yield stress and the ultimate tensile strength of SLMed HEA (651.36 MPa, 1057.21 MPa) were superior to those of casting samples (489.63 MPa, 873.09 MPa). Although the strength of SLMed tensile samples was lower than that of the HEA after TM treatment [25], the ductility of SLM samples improved remarkably. Fig. 9 shows SEM images of the tensile fracture morphology of SLMed samples after tensile tests, respectively. The size of ductile dimples was much smaller than those obtained in as-cast HEA [27e29]. A great number of dimples with small size indicated more surface energy produced during fracture process, which reflected higher strength. In addition, the size of dimples depended to a great extent on the size of sub-grains [30]. As shown in Fig. 9b (a magnification of the rectangular area in Fig. 9a), the average dimple size was approximate 500 nm which was almost the same as that of sub-grains (Fig. 6c).
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Fig. 5. EBSD analysis of SLMed AlCrFeNiV HEA: (a) IPF map; (b) Image quality (IQ) map with LAGBs and HAGBs superimposed; (c) Kernel average misorientation (KAM) map; (d) Misorientation angle plot corresponding to (b); (e) Grain size distribution.
4. Discussion 4.1. Formation of hierarchically heterogeneous microstructure Microstructure characterization reveals the heterogeneity of hierarchical structures including columnar grains, sub-grains, L12 nano phase and dislocations in SLMed AlCrFeNiV. Columnar grains, as the typical grain morphology of SLM, usually formed through epitaxial growth due to the re-melting of the previous depositing layer and the significant temperature gradient between the depositing layers [21,31]. Columnar grains in the present study exhibited preferential growth aligned closely with the maximum heat flow direction at the solid-liquid interface. A schematic diagram of re-melting of adjacent layers during SLM process was shown in Fig. 10. Columnar grains showed various growth direction and size within a single molten pool compared with other metal materials. Because of the lower thermal conductivity of HEA, the heat from the central of molten pool could not be transferred to the surrounding of heat source in time, which increases the temperature gradient and promotes the epitaxial growth of grains [32e34]. Compared with metal materials with high thermal conductivity, such as Cu alloy, the columnar grains of HEA samples have a longer span of molten pool. Sub-grains have been observed in a range of metals or alloys prepared by SLM [17,24,35], however, there is no consensus about the formation mechanism of sub-grains. SLM is normally related to a very rapid melting and solidification process which is far from equilibrium condition, and the microstructure is strongly affected
by the high temperature gradient, high cooling rate and large undercooling depending on the metal property, specimen geometry, and laser parameters [36e38]. The formation of the sub-grains in the present HEA was related to the Bernard convection. Bernard convection occurred under the driving of surface tension when the temperature gradient of the liquid layer exceeded the critical temperature gradient [39,40]. Larger vertical temperature gradient induced the vertical melt convection and cellular sub-grains formed. With laser scanning, horizontal temperature gradient in the center of the laser was dominant, and the horizontal melt convection promoted the formation of strip sub-grains. In general, deformation tends to induce the formation of dislocations, however, the present SLMed HEA exhibited a relative high content of dislocation without deformation. This could be attributed to the rapid heating and cooling rate in local micro-region during SLM. The rapid heating and cooling rate not only increased the vacancy concentration [24], but also generated large thermal residual stresses in HEA. The vacancy concentration could be regarded as the source of dislocation and residual stress could produce dislocations, and the dislocation multiplication resulted in the high-density dislocations distribution in HEA. In addition, heterogeneous distribution of dislocations could also be found in sub-grains. Different cooling rate of sub-grains led to the variety of vacancy concentration and residual stress level, inducing the different quantity of dislocations. Besides, dislocations multiplication and interaction occurred easily in smaller sub-grains for the smaller mean free path of dislocation and hindering effect of subgrain boundaries.
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Fig. 8. True stress-strain curves of the as-cast, SLMed AlCrFeNiV HEA and corresponding result in Ref.25 after quasi-static tensile test.
Fig. 6. (a) SEM images of the SLMed AlCrFeNiV HEA; (bec) Magnification micrograph of area I and area II in (a); (d) Bright-field TEM images of the cellular sub-grain; (eef) Magnification micrograph of area III and area IV in (d). Fig. 9. (a) SEM images of the fracture surfaces of SLMed AlCrFeNiV HEA after quasistatic tensile test; (b) Magnification micrograph of the selected square region in (a).
reasons for the formation of L12 nano phase. Although the spinodal decomposition was extremely restrained in SLMed HEA due to the high cooling rate, supersaturated FCC phase, sever lattice distortion and residual stress existed in the SLMed HEA. All of the factors mentioned above provide the driving force for the formation of L12 nano phase. The subsequent laser scanning, which could be regarded as localized heat treatment supplied the thermal dynamic conditions.
4.2. Improvements of strength and ductility
Fig. 7. (a) Bright field TEM images of the SLMed AlCrFeNiV HEA; (b) Selected area electron diffraction (SAED) pattern of the circle region in (a); (ced) HRTEM image and corresponding FFT pattern of the selected square region in (c).
The most remarkable finding was the formation of L12 nano phase in SLMed AlCrFeNiV HEA, and the L12 nano phase was usually obtained in cast AlCrFeNiV HEA through TM treatment [25]. And in the HEA through TM treatment, high Ni/Al ratio, spinodal decomposition and the ordering of the disorder FCC phase were the
Fig. 11 exhibits comparison of the yield stress and elongation of SLMed HEA with different element systems including FeCoCrNi [37], CoCrFeMnNi [17,38,41], AlFeCoCrNi [16]. The present SLMed AlCrFeNiV HEA exhibited higher yielding strength and plasticity. All the structure in the macro-, micro-, and nano-level has contributed to the overall tensile properties of the present HEA. Fig. 12 shows the schematic diagram of the dislocation formation and motion, accompanied by the interaction between dislocations and obstacles (including L12 nano phase, pre-existence dislocations, sub-grain boundaries and grain boundaries). Dislocations tended to appear firstly in the sub-grains with low dislocation density, and dislocations motion would encounter L12 nano phase and pre-existence dislocations inner sub-grains. L12 nano phase provided strong diffuse attractive obstacles to trap moving dislocations, and create antiphase boundaries (APBs) on the slip planes of the ordered L12 nano phase when they are sheared by dislocations. Simultaneously, the highly coherent and low misfit FCCL12 interfaces could minimize the elastic strain accumulation resulting
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Fig. 10. Schematic diagram of grain growth during the re-melting of adjacent forming layers in SLMed AlCrFeNiV HEA.
Fig. 11. Yield strength and elongation of SLMed AlCrFeNiV HEA and SLMed FeCoCrNi [37], AlFeCoCrNi [14] and CoCrFeMnNi [17,38] HEAs from reference under quasi-static tensile condition.
from dislocation shear and hence prevented crack initiation at these interfaces. The interaction between the newly formed dislocations and pre-existence dislocations caused more dislocations multiplication, which also contributed to the strength improvement of SLMed HEA. Due to the dislocation network of sub-grain boundary, it acts as stable and soft barriers that hinder dislocation motion for strength, but meanwhile guaranteed a continuous plastic flow by allowing dislocation from transmitting. Thus, sub-grain boundaries possess stronger dislocation storage capacity and dislocations will be more easily trapped and stored in sub-grain boundaries so that dislocation slip is particularly difficult [17]. When the dislocation density in one sub-grain increases to be close to that of other sub-grains, the strength of sub-grains tends to be consistent, thus achieving the synergistic movement of dislocations in the tensile process. The misorientation between sub-grains may also contribute to the stability of the sub-grain boundaries. In addition, dislocations pileup at sub-grain boundaries caused the stress concentration. The sub-micron size dimples on the fracture surface are the results of breaking sub-grains. With the increase of stress, dislocations passed through subgrain boundaries to grain boundaries, which would become new obstacles to dislocation motion. Dislocations pile up near grain boundaries, which resulting in more serious strain hardening. Although L12 nano phase and sub-grain boundary could contribute to the improvement of strength, their strengthen effect was still far behind the grain boundary. The grain size in the SLMed HEA was much larger than that of the HEA after TM treatment (~2.8 mm) [25], which exhibited lower strength. The plastic deformation mainly depends on the dislocation motion, and the effect of hierarchically heterogeneous structure on the dislocation motion plays a significant role in mechanical performance of SLMed HEA. Hierarchical structure from nano-scale (L12 nano phase, pre-existence dislocations) to sub-micron scale (sub-grain boundaries and grain boundaries) hinders the dislocation motion progressively. And heterogeneous distribution of the dislocation density and sub-grains morphology also made the deformation progressively. As a consequence, the SLMed AlCrFeNiV HEA achieves the high strength and ductility.
Fig. 12. Schematic diagram of the interaction between dislocations and obstacles (including L12 nano phase, pre-existence dislocations, sub-grain boundaries and grain boundaries).
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5. Conclusions 1) SLMed AlCrFeNiV HEA consisting of FCC phase and L12 nano phase shows a hierarchical heterogeneous structure including columnar grains, sub-grains, L12 nano phase and dislocations. Cellular or strip sub-grains were distributed in each columnar grain, and dislocations and L12 nano phase were distributed in sub-grains. 2) High cooling speed and non-equilibrium solidification during SLM process caused extremely high temperature gradient and large surface tension gradient, accompanied with the Bernard convection along the vertical or horizontal direction, leading to the formation of cellular or strip sub-grains. Dislocations formed due to the increased the vacancy concentration and large thermal residual stresses produced by rapid heating and cooling rate in SLMed process. Supersaturated FCC phase, sever lattice distortion, residual stress, and subsequent laser scanning provide the driving force for the formation of L12 nano phase. 3) The SLMed AlCrFeNiV HEA exhibited an outstanding combination of high strength (ultimate tensile strength ~1057.47 MPa) and excellent ductility (plastic strain ~30.3%). Hierarchical structure from nano-scale (L12 nano phase, pre-existence dislocations) to sub-micron scale (sub-grain boundaries and grain boundaries) hindered the dislocation motion progressively. And heterogeneous distribution of the dislocation density and subgrains morphology also made the deformation progressively, achieving the balance of high strength and ductility.
[11]
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Acknowledgement This work is supported by National Natural Science Foundation of China (Grant No. 51901004 and 51771005), financial supports by National Key R&D Program of China (No. 2017YFB0305800) and Beijing Natural Science Foundation (Grant No. 2194067).
[22]
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