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Investigation on the microstructure and mechanical behaviors of a laser formed Nb-Ti-Al alloy
Materials Characterization 162 (2020) 110193
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Investigation on the microstructure and mechanical behaviors of a laser formed Nb-Ti-Al alloy ⁎
Zhiwu Shia, , Jinlai Liub, Hua Weic, Hongyu Zhangb, Xiaofeng Sunb, Qi Zhengb,
T
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a
Department of Materials Sciences and Manufacturing Processes, AECC Commercial Aircraft Engine Co., Ltd., Shanghai 200241, China Superalloys Division, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, China c College of Civil Engineering and Architecture, Zhejiang University, Hangzhou 310058, China b
A R T I C LE I N FO
A B S T R A C T
Keywords: Laser forming Nb-Ti-Al alloy Microstructure Ti(O,C) Double diffraction Mechanical behavior
Nb-Ti-Al alloys have attracted much attention as a potential candidate for high-temperature structural applications beyond the scope of Ni-based superalloys. Previously, Nb-Ti-Al alloys were prepared by arc-melting or hot-pressing. In this study, laser forming of the Nb-23Ti-15Al alloy is explored, and the microstructure and mechanical behaviors are systematically investigated. The results indicate that the nearly defectless Nb-Ti-Al alloy with fine dendrites could be obtained through laser forming. Three phases, β, δ and Ti(O,C), present in the alloy, and the β/δ matrix displays a columnar microstructure. Moreover, the Ti(O,C) phase, a solid-solution of TiO or TiC, appears in the alloy. Ti(O,C) is formed due to the introduction of the O and C from the elemental powder, which has a face-centered cubic (FCC) structure and a moderate lattice parameter between that of TiO and TiC. Two morphologies of Ti(O,C), large cobblestone-like particles and small dispersive particles, are observed in the alloy. Forbidden reflections of δ phase occur due to double diffraction and slight superlattice reflections of β phase arise. Furthermore, refined β/δ phases and dispersed small Ti(O,C) result in higher microhardness and fracture toughness. In brief, our results indicate that laser forming is a potential method for manufacturing Nb-Ti-Al alloys with prominent properties.
1. Introduction
properties requirements in a single part [18,19]. Previous studies witnessed the laser forming of Nb-Si based alloys [14,20–28]. Brice et al. [20] deposited Nb-35Si (at.%) binary alloy using elemental blends of powders. The alloy consists of homogenous Nb5Si3, Nb3Si and metastable NbSi2 phases. Dehoff et al. [21] laser formed Nb-Ti-Cr-Si alloys employing elemental powder blends and obtained significantly refined microstructure. Besides, the influence of process parameters such as scan speed [28], laser power [25–27] on Nb-Si based alloys was also studied. However, Nb-Ti-Al alloys produced through laser forming, as well as the microstructure and mechanical behaviors of the alloys, have not been reported before. In addition, Rozmus et al. [16] found dispersed TiO phase in hotpressed Nb-Ti-Al alloys, but did not characterize TiO in detail. The investigation of Dymek et al. [29] demonstrates that in the processing of Nb based alloys via powder metallurgy, O element tends to be incorporated due to the high chemical activity of the raw powder. Shi et al. [13] found various kinds of twinned Ti(O,C) phases in hot-pressed Nb-23Ti-15Al alloy, including lamellar, coner and three-fold Ti(O,C) twins, and investigated the formation mechanisms of the twins. Until now, whether the oxide or carbide can appear in laser formed Nb-Ti-Al
Nb-based alloys, including Nb-Si alloys, Nb-Ti-Al alloys etc., have been developed remarkably in the last several decades due to their outstanding high-temperature properties [1–6]. With the temperature capability beyond nickel-based superalloys, Nb-Ti-Al alloys are regarded as potential high-temperature structure materials in aero-engines [4–14]. Previously, Nb-Ti-Al alloys were manufactured via arcmelting, casting, forging or hot-pressing. For the arc-melted Nb-Ti-Al buttons of casting or aging state, δ phase appears at β grain boundaries as plates, or assumes a shape of massive Widmanstätten plates precipitating from the grain boundaries or inside β grains [4,8,15]. Rozmus et al. [9,10,16] prepared Nb-Ti-Al alloys via hot-pressing and obtained fine-scale microstructure with blocky β phase distributing in δ matrix. Shi et al. [17] found that the microstructure of the hot-pressed Nb-Ti-Al alloys, blocky or columnar, depends on the hot-pressing temperature. Laser forming, an advanced manufacture technique, can be used to form near-net-shape complex parts with high performance rapidly, without the employment of molds. Moreover, graded compositions in components can also be created, which can help to balance various
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Corresponding authors. E-mail addresses: [email protected] (Z. Shi), [email protected] (Q. Zheng).
https://doi.org/10.1016/j.matchar.2020.110193 Received 12 November 2019; Received in revised form 6 February 2020; Accepted 6 February 2020 Available online 07 February 2020 1044-5803/ © 2020 Published by Elsevier Inc.
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The etched metallographic microstructure is demonstrated in Fig. 1c–f. Apart from the interdendritic Ti(O,C), β and δ phases constitute the dendrite arms. Deduced from the Nb-Ti-Al ternary diagrams [30,31], the initially solidified phase of the Nb-23Ti-15Al alloy from liquid state is β phase. Then δ phase precipitates in β phase during cooling, as well as during the cyclic temperature rising and falling for the cladding of later layers. Similarly, for the Nb-Ti-Cr-Si alloys constructed through laser engineered net shaping by Dehoff et al. [21] and the Nb-26Ti-22Si-6Cr-3Hf-2Al alloy prepared by Dicks et al. using direct laser fabrication [28], a large number of precipitates, such as αNb5Si3, appear in the as-deposited state. Dehoff et al. believe that early deposited layers undergo a significant heat treatment owing to the deposition of other layers. However, although the temperature is high, the heat treatment period is quite short, and in this study the cladding only lasted for about 15 min. Moreover, since Al and Nb atoms diffuse sluggishly, and Ti slows down the growth rate of the δ phase [4], δ phase precipitates much slowly in Nb-Ti-Al alloys during heat treatment. For example, tens or hundreds of hours are needed for the precipitation of a large amount of δ phase for the alloy held at high temperature [4]. Therefore, the reason for the sufficient precipitating of δ phase in the laser formed Nb-based alloys still remains indistinct. The δ phase in the laser formed Nb-23Ti-15Al alloy mainly takes the columnar shape, and thus equiaxed phases showing the transverse section of δ columns are observed in some regions (Fig. 1d), while regular strips revealing the longitudinal section are more frequently found (Fig. 1e). Additionally, microstructure exhibiting the mixed shapes is also often seen, as shown in Fig. 1f. From the longitudinal microstructure, it can be seen that the interface of β and δ phases is flat, and a bunch of δ strips precipitate in a similar direction. The δ phase is thick and continuous while the β phase is thinner and separated by δ strips. It is noteworthy that the duplex structure resembles that of the Nb-23Ti-15Al alloy prepared by hot-pressing at 1440 °C, which is also featured by a duplex structure of β and δ phases, and columnar β and δ phases are separated by each other [17]. Additionally, the microstructure of the Nb-23Ti-15Al alloy prepared by hot-pressing at 1350 °C is remarkably different, which is characterized by blocky morphology [17]. The discrepancy is attributed to different formation mechanisms of δ phase, since δ phase is formed by consuming σ and β phases in the alloy hot pressed at 1350 °C, while δ phase directly precipitates from the β phase in the alloy laser formed or hot pressed at 1440 °C. TEM observations are also in accordance with the SEM results. Fig. 3a shows β and δ strips and Fig. 3b presents the equiaxed phases, the transverse sections of β and δ strips. In addition, TEM tilting experiments demonstrate that β and δ phases have orientation relationships of (021)δ//(110)β, which might be the reason for the formation of the regular columnar phases, since precipitating of δ phase from β phase with preferred orientation relationships can reduce the β/δ interface energy. The orientation relationships of cobblestone-like Ti (O,C) and β or δ phases have also been investigated. Through tilting a large number of Ti(O,C) particles, no specific orientation relationship of Ti(O,C)/δ or Ti(O,C)/β has been found. As aforementioned, the hot-pressed Nb-23Ti-15Al alloy also contains Ti(O,C) phase. In contrast to the aggregated Ti(O,C) phase in the hot-pressed Nb-23Ti-15Al alloy [32], the Ti(O,C) in the laser formed Nb-23Ti-15Al alloy disperses uniformly and no obvious agglomeration is observed. On one hand, fine-scale and homogeneously dispersed Ti (O,C) bring an advantage of dispersion strengthening for the alloy. On the other hand, the dispersed Ti(O,C) results in better fracture toughness than the aggregated Ti(O,C), since the latter might act as the crack initiation in consequence of its brittleness. Moreover, careful inspection reveals that Ti(O,C) particles can be distinctly divided into two groups. Apart from the large cobblestone-like Ti(O,C) with a dimension of several micrometers, the small ones of about several hundreds of nanometers can be observed via TEM, as indicated by arrows in Fig. 3. Fig. 4 illustrates the EPMA results of phases composition distribution in the laser formed Nb-23Ti-15Al alloy. The β/δ matrix is mainly
alloy is not clear, and the influence on the alloy properties has not been reported. This study has explored the laser forming technology of a NbTi-Al alloy, and analyzed the microstructure and mechanical behaviors of the alloy in detail. 2. Experimental The raw materials used in this study are Nb, Ti and Al elemental powders, the purities of which are 99%, 99% and 99.5% (wt%), respectively. The powders were mixed to a composition of Nb-23Ti-15Al (at.%) and then mechanically alloyed in a Shunchi PMQW4L 3-D planetary ball mill filled with argon at room temperature. During mechanical alloying, the milling speed was 250 rpm. Too large or too small particles are not suitable for laser forming due to their poor flowability, therefore, after milling for a proper duration, particles within the range of 100–400 meshes were sieved out for laser forming, and the rest powder continued to be milled until applicable. Afterwards, the Nb-Ti-Al alloy was laser formed using the mechanically alloyed powder via a DL-HL-T5000B CO2 laser with a maximum power of 5 kW. The equipment comprises coaxial powder-feed system and multi-axis numerical control systems. The cladding was carried out in an argon atmosphere to avoid oxidation. During cladding, a power of approximately 2 kW and a laser travel speed of 7 mm/s were used, and travel directions rotated 90° from previous layer. The laser formed samples were 20 mm × 20 mm × 20 mm cubes. To observe the microstructure of the alloy, the samples were cut via wire electrical discharge machining. After grinding and polishing, the samples were etched in a solution of 65 mL H2O + 15 mL H2SO4 + 10 mL HF + 5 mL HNO3 + 5 mL H2O2 for 10 s to 20 s, and then microstructure examination was carried out by a JEOL JSM-5800 scanning electron microscopy (SEM). A Rigaku D/MAX 2500 X-ray diffraction (XRD) with Cu Kα radiation (λ = 1.5406 Å) was used to analyze the phase types qualitatively and lattice parameters of the phases quantitatively. Electron probe microanalysis (EPMA) was conducted via a CAMECA SX-100 to determine compositions of the phases. Vickers microhardness tests were conducted at a load of 1000 g for a dwell time of 15 s. After SEM observation, the crack lengths of the Vickers hardness indentations were measured using micrographs for the calculation of fracture toughness. The foils for transmission electron microscopy (TEM) examination were mechanically thinned to < 50 μm, followed by electropolishing in a twin jet polisher with a solution of 10 vol% H2SO4 in methanol at −30 °C. Afterwards, the electropolished foils were further thinned via ion milling. Microstructure observation and selected area electronic diffraction (SAED) analysis of the phases were performed using a JEOL JEM-2100 TEM with a voltage of 200 kV. 3. Results and discussion 3.1. Microstructure of the Nb-23Ti-15Al alloy Fig. 1 shows the SEM microstructure of the laser formed Nb-23Ti15Al alloy. Combining with the XRD results in Fig. 2, it can be seen that the alloy contains β, δ and Ti(O,C) phases, as marked in Fig. 1d–f. Fig. 1a and b presents the unetched metallographical microstructure, in which the dendritic morphology is observed since the cobblestone-like black Ti(O,C) particles distribute in interdendritic regions and outline the dendrites. The dendrites are quite fine and no directional characteristic is revealed. Moreover, it can also be seen that the low-magnification microstructures are not uniform, and Fig. 1a reveals coarser microstructure in the left and finer microstructure in the right. However, no significant difference was observed in high-magnification examination of the two types of microstructures. In addition, the laser formed Nb-Ti-Al alloys were nearly defectless, except for some small pores with diameters from several micrometers to several tens of micrometers, as demonstrated in Fig. 1a. 2
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Fig. 1. SEM micrographs showing microstructures of the laser formed Nb-23Ti-15Al alloy: (a) (b) polished and unetched sample, (c)–(f) etched sample.
fabricated by the two methods. For the laser formed alloy, the volume percentage of β phase is about one-third lower, while the fractions of δ and Ti(O,C) phases are higher than those of the hot-pressed alloy. In the laser formed alloy, δ phase has a largest percentage in the three phases of about 59%, while Ti(O,C) is of about 10%. 3.2. SAED analysis and Ti(O,C) twins SAED techniques were used to further determine and analyze the crystal structure of the existing phases in the laser formed alloy. Fig. 6 exhibits typical SAED patterns of matrix phases in the Nb-23Ti-15Al alloy, which confirms that the matrix includes β phase with a bodycentered cubic (BCC) structure and δ phase with an A15 structure. Through careful analyses, Fig. 6a and b is determined to be [001] and [011] patterns of β phase, and Fig. 6c and d is determined to be [001] and [111] patterns of δ phase. It is noteworthy that Fig. 6b reveals faint {100} superlattice reflections of β phase, which indicates that weak B2 ordering of β phase takes place in the laser formed Nb-23Ti-15Al alloy. On the contrary, in the hot-pressed Nb-23Ti-15Al alloy, remarkable superlattice reflections and antiphase domains are observed, revealing the appearance of prominent B2 ordering [32]. The Nb-23Ti-15Al alloys produced in the two routines show different degrees of B2 ordering, which can be explained by the difference in cooling rate of the two manufacturing methods. The cooling rate of laser forming process is much higher than that of hot-pressing, and hence more sufficient B2 ordering occurs in the hot-pressed alloy.
Fig. 2. XRD pattern of the laser formed Nb-23Ti-15Al alloy.
rich in Nb and Al, while the dark particles are abundant in Ti, O and C. Fig. 5 shows the phase volume fractions of β, δ and Ti(O,C) in the laser formed and the 1440 °C hot-pressed Nb-23Ti-15Al alloys [17], which are obtained through image analyses using plenty of SEM images. The results demonstrate different phase contents in the alloy 3
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Fig. 3. TEM bright field images showing two kinds of microstructures in the laser formed Nb-23Ti-15Al alloy: (a) a region of β and δ strips, (b) a region of equiaxed β and δ.
It is interesting that Ti(O,C) twins are observed in the laser formed Nb-23Ti-15Al alloy. Fig. 8 shows the morphology and SAED pattern of a Ti(O,C) twin. The Ti(O,C) twin is quite small with a diameter of about 150 nm and situates at the β/δ phase boundary. Instead of a regular lamellar twin, the Ti(O,C) twin belongs to a corner twin, that is, the grain is divided into two twinning parts by a {111} boundary. Ti(O,C) twins are seldom observed in the laser formed alloy, and no other kinds of Ti(O,C) twins, for example, lamellar or three-fold Ti(O,C) twins are observed. However, these Ti(O,C) twins are frequently found in the hotpressed Nb-23Ti-15Al alloy, as reported by Shi et al. [13]. The discrepancy can be reasoned by the fabrication conditions. According to Ref. [13], the Ti(O,C) twins are formed through oriented attachment of Ti(O,C) particles, which consists of the contact, rotation, attachment and Ostwald ripening of Ti(O,C) grains. Since the solidification and cooling is rapid during laser forming, the hindrance for rotation and attachment of Ti(O,C) grains is remarkably severe. Therefore, only a
Typical SAED results for the large cobblestone-like Ti(O,C) particles are shown in Fig. 7. According to the SAED patterns, the phase has a face-centered cubic (FCC) structure and a lattice parameter of 4.28 Å. Similarly, the smaller Ti(O,C) particles described in Section 3.1 are also of the FCC structure and the lattice parameter of the phase is 4.27 Å. Using the XRD results shown in Fig. 2, the lattice parameter of the FCC phase is calculated to be 4.27 Å, which is in accordance with the SAED results. It is notable that the lattice parameter of Ti(O,C) is between that of FCC TiO (4.18 Å, ICDD PDF card no. 65-2900) and TiC (4.32 Å, ICDD PDF card no. 65-8804) since Ti(O,C) is the solid-solution of TiO or TiC, as pointed out in Ref. [32]. The formation of Ti(O,C) phase mainly results from the introduction of O and C in the elemental powders, the contents of O and C are 0.46 wt% and 0.20 wt% respectively in the mixed powders. Nb-Ti-Al alloy produced using elemental powders is difficult to get rid of O and C, which are extraordinary active.
Fig. 4. EPMA compositions mapping of the laser formed Nb-23Ti-15Al alloy (a), showing the distribution of (b) Nb, (c) Al, (d) Ti, (e) C, (f) O. Red color represents high element concentration while blue color corresponds to low concentration. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.) 4
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few Ti(O,C) corner twins are formed and no higher fold twins are observed in the laser formed alloy.
SAED pattern of the δ phase shown in Fig. 6c and the extinction rule. Some visible diffractions spots, for example, {100} and {300}, are forbidden according to the extinction rule. To make it clear, SAED patterns were taken from areas with different thicknesses: ordinary thickness area, thin area (near the edge of a TEM foil) and extremely thin area (at the edge of a TEM foil). The experimental results of the [112] SAED patterns are exhibited in Fig. 9. In the ordinary thickness area (Fig. 9a), the {111} extinction reflections have an intensity approximately equal to that of the allowed reflections. As the thickness reduces, the {111} reflections are relatively much weaker although still appear in thin area (Fig. 9b). In contrast, the extremely thin area (Fig. 9c) shows nearly disappeared {111}. It can be deduced that as the thickness of δ phase decreases, the forbidden reflections tend to conform to the system absence law. In thin enough areas, the reflection only involves kinematics effect. On the contrary, in thicker areas, the interaction of diffraction beam with transmission beam and other diffraction beams cannot be ignored. In electric diffraction, when the diffraction beam is strong, it is rediffracted within the thick crystal, and then double diffraction occurs. Therefore, the appearance of forbidden reflections of δ phase is caused by double diffraction. Although appearance of forbidden reflections of δ phase has been observed in some previous literature, only a few have studied this subject [33].
3.3. Double diffraction of δ phase
3.4. Microhardness and fracture toughness
δ phase has an A15 crystal structure, in which Al occupies 2a site while Nb occupies 6c site. According to Ref. [26], the reflections appear when h + k + l = 2n or h = 2n + 1, k = 4n and l = 4n + 2. It is interesting to note that there is a large discrepancy between the [001]
Vickers microhardness of the alloy was determined to be 745.8 HV, which is much higher than the 1440 °C hot-pressed alloy (620.6 HV [17]) with the same composition. Using crack lengths at the corners of the indentations, fracture toughness (KIC) of the alloy is calculated by
Fig. 5. Phase volume percentages of the laser formed and the 1440 °C hotpressed Nb-23Ti-15Al alloys.
Fig. 6. SAED patterns of β and δ phases in the laser formed Nb-23Ti-15Al alloy. The arrows indicate superlattice reflections. 5
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Fig. 7. Typical SAED patterns of Ti(O,C) phase in the laser formed Nb-23Ti-15Al alloy.
Fig. 8. TEM micrographs showing the morphology (a) and SAED pattern (b) of a Ti(O,C) twin in the laser formed Nb-23Ti-15Al alloy.
microstructure compared with the hot-pressed alloy. Statistics results of SEM images show that the thickness of β/δ columnar pairs in the two alloys is 1.45 μm and 1.75 μm, respectively. Refining of phases acts as a role similar to the refined crystals, giving rise to increased microhardness and fracture toughness. Secondly, as previously mentioned, Ti (O,C) distributes much more homogeneously in the laser formed alloy, which contribute to the promoted fracture toughness. Thirdly, as noted in Section 3.1, the percentages of δ and Ti(O,C) phases are higher in the laser formed alloy, and hence larger microhardness of the alloy can be expected. In addition, according to the microhardness test results and the corresponding mechanism analysis, it is reasonable to deduce that the laser formed Nb-23Ti-15Al alloy might show better other mechanical properties, such as tensile and compression properties, which still require further studies.
[34]: 3 5
1 HV ⎞ E 25 KIC = 0.035al− 2 ⎛ ⎝ Φ ⎠
(1)
where a is the half-diagonal of the indentations, l is the crack length, HV is the Vickers hardness, Φ is the constraint factor and Φ = 3 [34], E = 123 GPa [35]. The calculated fracture toughness is 4.2 MPa·m1/2, which is also higher than that of the hot-pressed alloy (3.7 MPa·m1/2) [17]. Both the laser formed and the 1440 °C hot-pressed Nb-23Ti-15Al alloys exhibit columnar microstructure. However, the laser formed alloy shows much higher microhardness and fracture toughness. The reason can be attributed to the microstructure differences of the alloys. First of all, the laser formed alloy shows thinner columnar 6
Materials Characterization 162 (2020) 110193
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Fig. 9. SAED patterns of δ phase in the laser formed Nb-23Ti-15Al alloy in areas of different foil thicknesses, B = [112]: (a) ordinary thickness area, (b) thin area, (c) extremely thin area. The arrows indicate forbidden reflections.
4. Conclusions [6]
This study explores laser forming of the Nb-23Ti-15Al alloy, and microstructure as well as mechanical behaviors of the alloy have been investigated in-depth. Main conclusions are drawn as following:
[7] [8]
1. Nearly defectless Nb-23Ti-15Al alloy was produced via laser forming, which mainly contains columnar β and δ phases, as well as Ti(O,C) phase. 2. Two kinds of Ti(O,C) microstructures appear in the alloy due to the incorporation of C and O from the elemental powder: large cobblestone-like particles and small dispersive particles. Ti(O,C) has a FCC structure and a lattice parameter between that of TiO and TiC. 3. Double diffractions take place in δ phase, resulting in moderate intensity of the forbidden reflections. Slight ordering of β phase appears and weak superlattice reflections were observed. 4. Due to refined β/δ phases and dispersed small Ti(O,C), the laser formed Nb-23Ti-15Al alloy exhibits high microhardness and fracture toughness.
[9]
[10]
[11]
[12]
[13]
[14]
[15]
Declaration of competing interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.
[16]
[17]
Acknowledgments [18]
This work was supported by grants from the NNSF of China (Nos., 51601192, 51671189). The authors are very grateful to Guang Yang, Chun Shang and Baolei Cui for the laser forming processing.
[19]
[20]
Data availability statement [21]
The data generated during and/or analyzed during the current study are available from the corresponding author on reasonable request.
[22]
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Intermetallics 4 (1996) 13–22. [32] Z. Shi, H. Wei, H. Zhang, D. Wu, T. Jin, X. Sun, Q. Zheng, Investigation of microstructure in hot-pressed Nb-23Ti-15Al alloy, J. Alloys Compd. 636 (2015) 61–66. [33] M. Aindow, L.S. Smith, J. Shyue, M.H. Loretto, H.L. Fraser, On the origins of ‘forbidden’ 100-type spots in electron diffraction patterns from the A15 compounds Nb3Al, Cr3Si and V3Si, Philos. Mag. Lett. 69 (1994) 23–30. [34] A.T. Wieg, Y. Kodera, Z. Wang, C. Dames, J.E. Garay, Thermomechanical properties of rare-earth-doped AlN for laser gain media: the role of grain boundaries and grain size, Acta Mater. 86 (2015) 148–156. [35] L. Murugesh, K.T.V. Rao, R.O. Ritchie, Crack growth in a ductile-phase-toughened Nb/Nb3Al in situ intermetallic composite under monotonic and cyclic loading, Scr. Metall. Mater. 29 (1993) 1107–1112.
[27] I. Goncharov, D.V. Masaylo, A. Orlov, N.G. Razumov, A. Obrosov, The effect of laser power on the microstructure of the Nb-Si based in situ composite, fabricated by laser metal deposition, Key Eng. Mater. Trans. Tech. Publ. 822 (2019) 556–562. [28] R. Dicks, F. Wang, X. Wu, The manufacture of a niobium/niobium-silicide-based alloy using direct laser fabrication, J. Mater. Process. Technol. 209 (2009) 1752–1757. [29] S. Dymek, M. Dollar, K. Leonard, Synthesis and characterization of mechanically alloyed Nb3Al-base alloys, Mater. Sci. Eng. A 239–240 (1997) 507–514. [30] A. Hellwig, M. Palm, G. Inden, Phase equilibria in the Al-Nb-Ti system at high temperatures, Intermetallics 6 (1998) 79–94. [31] G.L. Chen, X.T. Wang, K.Q. Ni, S.M. Hao, J.X. Cao, J.J. Ding, X. Zhang, Investigation on the 1000, 1150 and 1400 °C isothermal section of the Ti-Al-Nb system,
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Investigation on the microstructure and mechanical behaviors of a laser formed Nb-Ti-Al alloy ⁎
Zhiwu Shia, , Jinlai Liub, Hua Weic, Hongyu Zhangb, Xiaofeng Sunb, Qi Zhengb,
T
⁎
a
Department of Materials Sciences and Manufacturing Processes, AECC Commercial Aircraft Engine Co., Ltd., Shanghai 200241, China Superalloys Division, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, China c College of Civil Engineering and Architecture, Zhejiang University, Hangzhou 310058, China b
A R T I C LE I N FO
A B S T R A C T
Keywords: Laser forming Nb-Ti-Al alloy Microstructure Ti(O,C) Double diffraction Mechanical behavior
Nb-Ti-Al alloys have attracted much attention as a potential candidate for high-temperature structural applications beyond the scope of Ni-based superalloys. Previously, Nb-Ti-Al alloys were prepared by arc-melting or hot-pressing. In this study, laser forming of the Nb-23Ti-15Al alloy is explored, and the microstructure and mechanical behaviors are systematically investigated. The results indicate that the nearly defectless Nb-Ti-Al alloy with fine dendrites could be obtained through laser forming. Three phases, β, δ and Ti(O,C), present in the alloy, and the β/δ matrix displays a columnar microstructure. Moreover, the Ti(O,C) phase, a solid-solution of TiO or TiC, appears in the alloy. Ti(O,C) is formed due to the introduction of the O and C from the elemental powder, which has a face-centered cubic (FCC) structure and a moderate lattice parameter between that of TiO and TiC. Two morphologies of Ti(O,C), large cobblestone-like particles and small dispersive particles, are observed in the alloy. Forbidden reflections of δ phase occur due to double diffraction and slight superlattice reflections of β phase arise. Furthermore, refined β/δ phases and dispersed small Ti(O,C) result in higher microhardness and fracture toughness. In brief, our results indicate that laser forming is a potential method for manufacturing Nb-Ti-Al alloys with prominent properties.
1. Introduction
properties requirements in a single part [18,19]. Previous studies witnessed the laser forming of Nb-Si based alloys [14,20–28]. Brice et al. [20] deposited Nb-35Si (at.%) binary alloy using elemental blends of powders. The alloy consists of homogenous Nb5Si3, Nb3Si and metastable NbSi2 phases. Dehoff et al. [21] laser formed Nb-Ti-Cr-Si alloys employing elemental powder blends and obtained significantly refined microstructure. Besides, the influence of process parameters such as scan speed [28], laser power [25–27] on Nb-Si based alloys was also studied. However, Nb-Ti-Al alloys produced through laser forming, as well as the microstructure and mechanical behaviors of the alloys, have not been reported before. In addition, Rozmus et al. [16] found dispersed TiO phase in hotpressed Nb-Ti-Al alloys, but did not characterize TiO in detail. The investigation of Dymek et al. [29] demonstrates that in the processing of Nb based alloys via powder metallurgy, O element tends to be incorporated due to the high chemical activity of the raw powder. Shi et al. [13] found various kinds of twinned Ti(O,C) phases in hot-pressed Nb-23Ti-15Al alloy, including lamellar, coner and three-fold Ti(O,C) twins, and investigated the formation mechanisms of the twins. Until now, whether the oxide or carbide can appear in laser formed Nb-Ti-Al
Nb-based alloys, including Nb-Si alloys, Nb-Ti-Al alloys etc., have been developed remarkably in the last several decades due to their outstanding high-temperature properties [1–6]. With the temperature capability beyond nickel-based superalloys, Nb-Ti-Al alloys are regarded as potential high-temperature structure materials in aero-engines [4–14]. Previously, Nb-Ti-Al alloys were manufactured via arcmelting, casting, forging or hot-pressing. For the arc-melted Nb-Ti-Al buttons of casting or aging state, δ phase appears at β grain boundaries as plates, or assumes a shape of massive Widmanstätten plates precipitating from the grain boundaries or inside β grains [4,8,15]. Rozmus et al. [9,10,16] prepared Nb-Ti-Al alloys via hot-pressing and obtained fine-scale microstructure with blocky β phase distributing in δ matrix. Shi et al. [17] found that the microstructure of the hot-pressed Nb-Ti-Al alloys, blocky or columnar, depends on the hot-pressing temperature. Laser forming, an advanced manufacture technique, can be used to form near-net-shape complex parts with high performance rapidly, without the employment of molds. Moreover, graded compositions in components can also be created, which can help to balance various
⁎
Corresponding authors. E-mail addresses: [email protected] (Z. Shi), [email protected] (Q. Zheng).
https://doi.org/10.1016/j.matchar.2020.110193 Received 12 November 2019; Received in revised form 6 February 2020; Accepted 6 February 2020 Available online 07 February 2020 1044-5803/ © 2020 Published by Elsevier Inc.
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The etched metallographic microstructure is demonstrated in Fig. 1c–f. Apart from the interdendritic Ti(O,C), β and δ phases constitute the dendrite arms. Deduced from the Nb-Ti-Al ternary diagrams [30,31], the initially solidified phase of the Nb-23Ti-15Al alloy from liquid state is β phase. Then δ phase precipitates in β phase during cooling, as well as during the cyclic temperature rising and falling for the cladding of later layers. Similarly, for the Nb-Ti-Cr-Si alloys constructed through laser engineered net shaping by Dehoff et al. [21] and the Nb-26Ti-22Si-6Cr-3Hf-2Al alloy prepared by Dicks et al. using direct laser fabrication [28], a large number of precipitates, such as αNb5Si3, appear in the as-deposited state. Dehoff et al. believe that early deposited layers undergo a significant heat treatment owing to the deposition of other layers. However, although the temperature is high, the heat treatment period is quite short, and in this study the cladding only lasted for about 15 min. Moreover, since Al and Nb atoms diffuse sluggishly, and Ti slows down the growth rate of the δ phase [4], δ phase precipitates much slowly in Nb-Ti-Al alloys during heat treatment. For example, tens or hundreds of hours are needed for the precipitation of a large amount of δ phase for the alloy held at high temperature [4]. Therefore, the reason for the sufficient precipitating of δ phase in the laser formed Nb-based alloys still remains indistinct. The δ phase in the laser formed Nb-23Ti-15Al alloy mainly takes the columnar shape, and thus equiaxed phases showing the transverse section of δ columns are observed in some regions (Fig. 1d), while regular strips revealing the longitudinal section are more frequently found (Fig. 1e). Additionally, microstructure exhibiting the mixed shapes is also often seen, as shown in Fig. 1f. From the longitudinal microstructure, it can be seen that the interface of β and δ phases is flat, and a bunch of δ strips precipitate in a similar direction. The δ phase is thick and continuous while the β phase is thinner and separated by δ strips. It is noteworthy that the duplex structure resembles that of the Nb-23Ti-15Al alloy prepared by hot-pressing at 1440 °C, which is also featured by a duplex structure of β and δ phases, and columnar β and δ phases are separated by each other [17]. Additionally, the microstructure of the Nb-23Ti-15Al alloy prepared by hot-pressing at 1350 °C is remarkably different, which is characterized by blocky morphology [17]. The discrepancy is attributed to different formation mechanisms of δ phase, since δ phase is formed by consuming σ and β phases in the alloy hot pressed at 1350 °C, while δ phase directly precipitates from the β phase in the alloy laser formed or hot pressed at 1440 °C. TEM observations are also in accordance with the SEM results. Fig. 3a shows β and δ strips and Fig. 3b presents the equiaxed phases, the transverse sections of β and δ strips. In addition, TEM tilting experiments demonstrate that β and δ phases have orientation relationships of (021)δ//(110)β, which might be the reason for the formation of the regular columnar phases, since precipitating of δ phase from β phase with preferred orientation relationships can reduce the β/δ interface energy. The orientation relationships of cobblestone-like Ti (O,C) and β or δ phases have also been investigated. Through tilting a large number of Ti(O,C) particles, no specific orientation relationship of Ti(O,C)/δ or Ti(O,C)/β has been found. As aforementioned, the hot-pressed Nb-23Ti-15Al alloy also contains Ti(O,C) phase. In contrast to the aggregated Ti(O,C) phase in the hot-pressed Nb-23Ti-15Al alloy [32], the Ti(O,C) in the laser formed Nb-23Ti-15Al alloy disperses uniformly and no obvious agglomeration is observed. On one hand, fine-scale and homogeneously dispersed Ti (O,C) bring an advantage of dispersion strengthening for the alloy. On the other hand, the dispersed Ti(O,C) results in better fracture toughness than the aggregated Ti(O,C), since the latter might act as the crack initiation in consequence of its brittleness. Moreover, careful inspection reveals that Ti(O,C) particles can be distinctly divided into two groups. Apart from the large cobblestone-like Ti(O,C) with a dimension of several micrometers, the small ones of about several hundreds of nanometers can be observed via TEM, as indicated by arrows in Fig. 3. Fig. 4 illustrates the EPMA results of phases composition distribution in the laser formed Nb-23Ti-15Al alloy. The β/δ matrix is mainly
alloy is not clear, and the influence on the alloy properties has not been reported. This study has explored the laser forming technology of a NbTi-Al alloy, and analyzed the microstructure and mechanical behaviors of the alloy in detail. 2. Experimental The raw materials used in this study are Nb, Ti and Al elemental powders, the purities of which are 99%, 99% and 99.5% (wt%), respectively. The powders were mixed to a composition of Nb-23Ti-15Al (at.%) and then mechanically alloyed in a Shunchi PMQW4L 3-D planetary ball mill filled with argon at room temperature. During mechanical alloying, the milling speed was 250 rpm. Too large or too small particles are not suitable for laser forming due to their poor flowability, therefore, after milling for a proper duration, particles within the range of 100–400 meshes were sieved out for laser forming, and the rest powder continued to be milled until applicable. Afterwards, the Nb-Ti-Al alloy was laser formed using the mechanically alloyed powder via a DL-HL-T5000B CO2 laser with a maximum power of 5 kW. The equipment comprises coaxial powder-feed system and multi-axis numerical control systems. The cladding was carried out in an argon atmosphere to avoid oxidation. During cladding, a power of approximately 2 kW and a laser travel speed of 7 mm/s were used, and travel directions rotated 90° from previous layer. The laser formed samples were 20 mm × 20 mm × 20 mm cubes. To observe the microstructure of the alloy, the samples were cut via wire electrical discharge machining. After grinding and polishing, the samples were etched in a solution of 65 mL H2O + 15 mL H2SO4 + 10 mL HF + 5 mL HNO3 + 5 mL H2O2 for 10 s to 20 s, and then microstructure examination was carried out by a JEOL JSM-5800 scanning electron microscopy (SEM). A Rigaku D/MAX 2500 X-ray diffraction (XRD) with Cu Kα radiation (λ = 1.5406 Å) was used to analyze the phase types qualitatively and lattice parameters of the phases quantitatively. Electron probe microanalysis (EPMA) was conducted via a CAMECA SX-100 to determine compositions of the phases. Vickers microhardness tests were conducted at a load of 1000 g for a dwell time of 15 s. After SEM observation, the crack lengths of the Vickers hardness indentations were measured using micrographs for the calculation of fracture toughness. The foils for transmission electron microscopy (TEM) examination were mechanically thinned to < 50 μm, followed by electropolishing in a twin jet polisher with a solution of 10 vol% H2SO4 in methanol at −30 °C. Afterwards, the electropolished foils were further thinned via ion milling. Microstructure observation and selected area electronic diffraction (SAED) analysis of the phases were performed using a JEOL JEM-2100 TEM with a voltage of 200 kV. 3. Results and discussion 3.1. Microstructure of the Nb-23Ti-15Al alloy Fig. 1 shows the SEM microstructure of the laser formed Nb-23Ti15Al alloy. Combining with the XRD results in Fig. 2, it can be seen that the alloy contains β, δ and Ti(O,C) phases, as marked in Fig. 1d–f. Fig. 1a and b presents the unetched metallographical microstructure, in which the dendritic morphology is observed since the cobblestone-like black Ti(O,C) particles distribute in interdendritic regions and outline the dendrites. The dendrites are quite fine and no directional characteristic is revealed. Moreover, it can also be seen that the low-magnification microstructures are not uniform, and Fig. 1a reveals coarser microstructure in the left and finer microstructure in the right. However, no significant difference was observed in high-magnification examination of the two types of microstructures. In addition, the laser formed Nb-Ti-Al alloys were nearly defectless, except for some small pores with diameters from several micrometers to several tens of micrometers, as demonstrated in Fig. 1a. 2
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Fig. 1. SEM micrographs showing microstructures of the laser formed Nb-23Ti-15Al alloy: (a) (b) polished and unetched sample, (c)–(f) etched sample.
fabricated by the two methods. For the laser formed alloy, the volume percentage of β phase is about one-third lower, while the fractions of δ and Ti(O,C) phases are higher than those of the hot-pressed alloy. In the laser formed alloy, δ phase has a largest percentage in the three phases of about 59%, while Ti(O,C) is of about 10%. 3.2. SAED analysis and Ti(O,C) twins SAED techniques were used to further determine and analyze the crystal structure of the existing phases in the laser formed alloy. Fig. 6 exhibits typical SAED patterns of matrix phases in the Nb-23Ti-15Al alloy, which confirms that the matrix includes β phase with a bodycentered cubic (BCC) structure and δ phase with an A15 structure. Through careful analyses, Fig. 6a and b is determined to be [001] and [011] patterns of β phase, and Fig. 6c and d is determined to be [001] and [111] patterns of δ phase. It is noteworthy that Fig. 6b reveals faint {100} superlattice reflections of β phase, which indicates that weak B2 ordering of β phase takes place in the laser formed Nb-23Ti-15Al alloy. On the contrary, in the hot-pressed Nb-23Ti-15Al alloy, remarkable superlattice reflections and antiphase domains are observed, revealing the appearance of prominent B2 ordering [32]. The Nb-23Ti-15Al alloys produced in the two routines show different degrees of B2 ordering, which can be explained by the difference in cooling rate of the two manufacturing methods. The cooling rate of laser forming process is much higher than that of hot-pressing, and hence more sufficient B2 ordering occurs in the hot-pressed alloy.
Fig. 2. XRD pattern of the laser formed Nb-23Ti-15Al alloy.
rich in Nb and Al, while the dark particles are abundant in Ti, O and C. Fig. 5 shows the phase volume fractions of β, δ and Ti(O,C) in the laser formed and the 1440 °C hot-pressed Nb-23Ti-15Al alloys [17], which are obtained through image analyses using plenty of SEM images. The results demonstrate different phase contents in the alloy 3
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Fig. 3. TEM bright field images showing two kinds of microstructures in the laser formed Nb-23Ti-15Al alloy: (a) a region of β and δ strips, (b) a region of equiaxed β and δ.
It is interesting that Ti(O,C) twins are observed in the laser formed Nb-23Ti-15Al alloy. Fig. 8 shows the morphology and SAED pattern of a Ti(O,C) twin. The Ti(O,C) twin is quite small with a diameter of about 150 nm and situates at the β/δ phase boundary. Instead of a regular lamellar twin, the Ti(O,C) twin belongs to a corner twin, that is, the grain is divided into two twinning parts by a {111} boundary. Ti(O,C) twins are seldom observed in the laser formed alloy, and no other kinds of Ti(O,C) twins, for example, lamellar or three-fold Ti(O,C) twins are observed. However, these Ti(O,C) twins are frequently found in the hotpressed Nb-23Ti-15Al alloy, as reported by Shi et al. [13]. The discrepancy can be reasoned by the fabrication conditions. According to Ref. [13], the Ti(O,C) twins are formed through oriented attachment of Ti(O,C) particles, which consists of the contact, rotation, attachment and Ostwald ripening of Ti(O,C) grains. Since the solidification and cooling is rapid during laser forming, the hindrance for rotation and attachment of Ti(O,C) grains is remarkably severe. Therefore, only a
Typical SAED results for the large cobblestone-like Ti(O,C) particles are shown in Fig. 7. According to the SAED patterns, the phase has a face-centered cubic (FCC) structure and a lattice parameter of 4.28 Å. Similarly, the smaller Ti(O,C) particles described in Section 3.1 are also of the FCC structure and the lattice parameter of the phase is 4.27 Å. Using the XRD results shown in Fig. 2, the lattice parameter of the FCC phase is calculated to be 4.27 Å, which is in accordance with the SAED results. It is notable that the lattice parameter of Ti(O,C) is between that of FCC TiO (4.18 Å, ICDD PDF card no. 65-2900) and TiC (4.32 Å, ICDD PDF card no. 65-8804) since Ti(O,C) is the solid-solution of TiO or TiC, as pointed out in Ref. [32]. The formation of Ti(O,C) phase mainly results from the introduction of O and C in the elemental powders, the contents of O and C are 0.46 wt% and 0.20 wt% respectively in the mixed powders. Nb-Ti-Al alloy produced using elemental powders is difficult to get rid of O and C, which are extraordinary active.
Fig. 4. EPMA compositions mapping of the laser formed Nb-23Ti-15Al alloy (a), showing the distribution of (b) Nb, (c) Al, (d) Ti, (e) C, (f) O. Red color represents high element concentration while blue color corresponds to low concentration. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.) 4
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few Ti(O,C) corner twins are formed and no higher fold twins are observed in the laser formed alloy.
SAED pattern of the δ phase shown in Fig. 6c and the extinction rule. Some visible diffractions spots, for example, {100} and {300}, are forbidden according to the extinction rule. To make it clear, SAED patterns were taken from areas with different thicknesses: ordinary thickness area, thin area (near the edge of a TEM foil) and extremely thin area (at the edge of a TEM foil). The experimental results of the [112] SAED patterns are exhibited in Fig. 9. In the ordinary thickness area (Fig. 9a), the {111} extinction reflections have an intensity approximately equal to that of the allowed reflections. As the thickness reduces, the {111} reflections are relatively much weaker although still appear in thin area (Fig. 9b). In contrast, the extremely thin area (Fig. 9c) shows nearly disappeared {111}. It can be deduced that as the thickness of δ phase decreases, the forbidden reflections tend to conform to the system absence law. In thin enough areas, the reflection only involves kinematics effect. On the contrary, in thicker areas, the interaction of diffraction beam with transmission beam and other diffraction beams cannot be ignored. In electric diffraction, when the diffraction beam is strong, it is rediffracted within the thick crystal, and then double diffraction occurs. Therefore, the appearance of forbidden reflections of δ phase is caused by double diffraction. Although appearance of forbidden reflections of δ phase has been observed in some previous literature, only a few have studied this subject [33].
3.3. Double diffraction of δ phase
3.4. Microhardness and fracture toughness
δ phase has an A15 crystal structure, in which Al occupies 2a site while Nb occupies 6c site. According to Ref. [26], the reflections appear when h + k + l = 2n or h = 2n + 1, k = 4n and l = 4n + 2. It is interesting to note that there is a large discrepancy between the [001]
Vickers microhardness of the alloy was determined to be 745.8 HV, which is much higher than the 1440 °C hot-pressed alloy (620.6 HV [17]) with the same composition. Using crack lengths at the corners of the indentations, fracture toughness (KIC) of the alloy is calculated by
Fig. 5. Phase volume percentages of the laser formed and the 1440 °C hotpressed Nb-23Ti-15Al alloys.
Fig. 6. SAED patterns of β and δ phases in the laser formed Nb-23Ti-15Al alloy. The arrows indicate superlattice reflections. 5
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Fig. 7. Typical SAED patterns of Ti(O,C) phase in the laser formed Nb-23Ti-15Al alloy.
Fig. 8. TEM micrographs showing the morphology (a) and SAED pattern (b) of a Ti(O,C) twin in the laser formed Nb-23Ti-15Al alloy.
microstructure compared with the hot-pressed alloy. Statistics results of SEM images show that the thickness of β/δ columnar pairs in the two alloys is 1.45 μm and 1.75 μm, respectively. Refining of phases acts as a role similar to the refined crystals, giving rise to increased microhardness and fracture toughness. Secondly, as previously mentioned, Ti (O,C) distributes much more homogeneously in the laser formed alloy, which contribute to the promoted fracture toughness. Thirdly, as noted in Section 3.1, the percentages of δ and Ti(O,C) phases are higher in the laser formed alloy, and hence larger microhardness of the alloy can be expected. In addition, according to the microhardness test results and the corresponding mechanism analysis, it is reasonable to deduce that the laser formed Nb-23Ti-15Al alloy might show better other mechanical properties, such as tensile and compression properties, which still require further studies.
[34]: 3 5
1 HV ⎞ E 25 KIC = 0.035al− 2 ⎛ ⎝ Φ ⎠
(1)
where a is the half-diagonal of the indentations, l is the crack length, HV is the Vickers hardness, Φ is the constraint factor and Φ = 3 [34], E = 123 GPa [35]. The calculated fracture toughness is 4.2 MPa·m1/2, which is also higher than that of the hot-pressed alloy (3.7 MPa·m1/2) [17]. Both the laser formed and the 1440 °C hot-pressed Nb-23Ti-15Al alloys exhibit columnar microstructure. However, the laser formed alloy shows much higher microhardness and fracture toughness. The reason can be attributed to the microstructure differences of the alloys. First of all, the laser formed alloy shows thinner columnar 6
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Fig. 9. SAED patterns of δ phase in the laser formed Nb-23Ti-15Al alloy in areas of different foil thicknesses, B = [112]: (a) ordinary thickness area, (b) thin area, (c) extremely thin area. The arrows indicate forbidden reflections.
4. Conclusions [6]
This study explores laser forming of the Nb-23Ti-15Al alloy, and microstructure as well as mechanical behaviors of the alloy have been investigated in-depth. Main conclusions are drawn as following:
[7] [8]
1. Nearly defectless Nb-23Ti-15Al alloy was produced via laser forming, which mainly contains columnar β and δ phases, as well as Ti(O,C) phase. 2. Two kinds of Ti(O,C) microstructures appear in the alloy due to the incorporation of C and O from the elemental powder: large cobblestone-like particles and small dispersive particles. Ti(O,C) has a FCC structure and a lattice parameter between that of TiO and TiC. 3. Double diffractions take place in δ phase, resulting in moderate intensity of the forbidden reflections. Slight ordering of β phase appears and weak superlattice reflections were observed. 4. Due to refined β/δ phases and dispersed small Ti(O,C), the laser formed Nb-23Ti-15Al alloy exhibits high microhardness and fracture toughness.
[9]
[10]
[11]
[12]
[13]
[14]
[15]
Declaration of competing interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.
[16]
[17]
Acknowledgments [18]
This work was supported by grants from the NNSF of China (Nos., 51601192, 51671189). The authors are very grateful to Guang Yang, Chun Shang and Baolei Cui for the laser forming processing.
[19]
[20]
Data availability statement [21]
The data generated during and/or analyzed during the current study are available from the corresponding author on reasonable request.
[22]
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