Ultrahigh-performance TiNi shape memory alloy by 4D printing

Materials Science & Engineering A 763 (2019) 138166

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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea

Ultrahigh-performance TiNi shape memory alloy by 4D printing a

a,*

a

a

b

H.Z. Lu , C. Yang , X. Luo , H.W. Ma , B. Song , Y.Y. Li a b c

a,b

, L.C. Zhang

T

c,a

National Engineering Research Center of Near-net-shape Forming for Metallic Materials, South China University of Technology, Guangzhou, 510640, China State Key Laboratory of Materials Processing and Die & Mould Technology, Huazhong University of Science and Technology, Wuhan, 430074, China School of Engineering, Edith Cowan University, 270 Joondalup Drive, Joondalup, Perth, WA, 6027, Australia

A R T I C LE I N FO

A B S T R A C T

Keywords: Selective laser melting Shape memory alloys Mechanical properties Shape memory effect

For additively manufactured components, it's widely accepted to have high enough energy input to facilitate nearly full density and low enough energy input to avoid cracking tendency. In this work, ultrahigh-performance Ti50.6Ni49.4 (at.%) shape memory alloy (SMA) was manufactured by selective laser melting (SLM) under high enough energy inputs (155–292 J/mm3). The microstructure, phase transformation behaviors, mechanical and shape memory properties of the SLM-manufactured SMA were investigated by various characterization methods of X-ray diffraction, scanning and transmission electron microscopies, differential scanning calorimetry, room temperature and stress-controlled cyclic tensile tests, etc. Results show that the martensite content and the austenite and martensitic transformation temperatures decrease with the decrease of laser energy input (the increase of laser scanning speed). Interestingly, the SLM-manufactured SMA exhibits ultrahigh tensile strength of 776 MPa and elongation of 7.2% under room-temperature tensile condition. In addition, stress-controlled cyclic tensile tests under 400 MPa indicate that the SLM-manufactured SMA has ultrahigh shape memory effect of 98.7% recovery ratio and 4.99% recoverable strain after ten times loading-unloading cycle. The ultrahigh mechanical and shape memory properties are associated to the combined effects of dispersedly distributed nanosized Ti2Ni precipitates, ultrafine grains and profuse dislocations in the SLM-manufactured SMA. This work substantiates, for the first time, high enough energy input in SLM can be applied to manufacture ultrahighperformance TiNi SMAs.

1. Introduction With the significant advancements in additive manufacturing or 3D printing, the terminology of 4D printing is introduced to witness technological evolution [1]. In essence, 4D printing is 3D printing of intelligent materials. Until now, 4D printing is applicable usually to polymers [2] but rarely to TiNi shape memory alloys (SMAs) [3]. On the one hand, the ratio of Ti and Ni elements in TiNi SMAs, even a slight variation of 0.1at.%, can significantly affect their microstructures and thus mechanical, physical, and functional properties [4]. Specifically, Ni4Ti3 and Ti2Ni are inclined to precipitate in Ni-rich and Ti-rich TiNi SMAs, respectively. On the other hand, additive manufacturing provides great flexibility in obtaining controlled microstructures by nonequilibrium rapid solidification resulted from complicated combination of processing parameters [5]. Undoubtedly, the two aspects increase the difficulty levels in achieving desired properties by tailoring processing parameters and thus microstructures. Meanwhile, there exists a pair of contradiction in additive manufacturing [6], i.e., high enough energy input to completely melt powder particles and thus to obtain

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components with nearly full density, and low enough energy input to minimize temperature gradient and thus residual stress between molten pools and adjacent powder particles in order to avoid cracking tendency. Particularly, such a contradiction is prominent for TiNi SMAs due to low abilities of thermal expansion and conduction of their constituent phases [7], such as B2 TiNi, B19’ TiNi, R phases (TiNi), and so on. To balance this contradiction, a number of attempts have been made recently to manufacture TiNi SMAs by selective laser melting (SLM) [1,8–21]. These efforts mainly focus on Ni-rich TiNi SMAs and confirm exactly that energy inputs applied to TiNi SMAs are relatively low, generally in the range of 30–140 J/mm3 [8,12,14–17,19–21]. Higher energy inputs (> 200 J/mm3) significantly increases impurities pick-up [10]. In addition, it is not difficult to find that mechanical properties were tested mostly under compressive conditions [8,10,14,17,19,20], probably owing to the usual coexistence of micro-pores and microcracks in SLMed parts [22,23]. Among the preferential properties, especially, tensile properties and shape memory effect (SME) of SLMed TiNi SMAs are little investigated. To date, the best tensile properties

Corresponding author. E-mail address: [email protected] (C. Yang).

https://doi.org/10.1016/j.msea.2019.138166 Received 1 May 2019; Received in revised form 11 July 2019; Accepted 15 July 2019 Available online 16 July 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.

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reported are ultimate tensile strength of 606 MPa and elongation of 6.8% for Ni-rich Ni50.1Ti49.9 (chemical composition in at.%, the same hereafter) SMA [12,21]. Correspondingly, the employed SLM parameters in Refs 12 and 21 are laser powers (P) of 250 and 60 W, scanning speed (v) of 1250 and 480 mm/s, powder layer thickness (h) of 120 and 110 μm, hatch spacing (t) of 30 and 25 μm, and energy input (E) of 55.5 and 45.5 J/mm3, respectively. Meanwhile, shape memory property reported so far is 3.54% recoverable strain and 75.2% recovery ratio after 4 cycles at 300 MPa [12]. These scenarios raise an interesting question: can TiNi SMAs be manufactured by 4D printing via SLM under high enough energy input in order to achieve ultrahigh mechanical and shape memory properties? Motivated by this question, we report ultrahigh-performance Ti-rich Ti50.6Ni49.4 SMA produced by high SLM energy inputs of 155–292 J/ mm3. The SLM-produced parts exhibit ultrahigh ultimate tensile strength of 776 MPa and elongation of 7.2%, along with ultrahigh SME of 98.7% recovery ratio and 4.99% recoverable strain after ten times loading-unloading cycle. To our best knowledge, this is the first time that such ultrahigh properties are successfully achieved by SLM in TiNi SMAs. 2. Materials and experimental Fig. 2. The scanning strategy used in the current work.

The as-received TiNi powders were prepared by electrode induction-melting gas atomization technique (AMC Powders Co. Ltd., Beijing). They had a spherical or near spherical morphology with limited small satellites particles (Fig. 1a), together with the particle size distribution of D10 = 21 μm, D50 = 37 μm, and D90 = 61 μm (Fig. 1b) by HORIBA laser scattering particle size distribution analyzer. Their chemical composition were determined to be Ti-rich Ti50.6Ni49.4 (at.%) by inductively coupled plasma-atomic emission spectroscopy, accompanied by minor unavoidable oxygen content below 650 ppm by TC600 Nitrogen/Oxygen Determinator (LECO Co., US). 4D printing was implemented via SLM in a Concept Laser M2 Cusing system in high purity argon atmosphere. The oxygen content of the printed parts was lower than 760 ppm. To achieve tailored microstructures and desired properties, two aspects must be considered simultaneously. First, high enough energy input E (= P/(v⋅h⋅t)) was employed by varying laser scanning speeds (v) of 80, 105, 120, and 150 mm/s, respectively, accompanied by fixing laser power (P) of 70 W, powder layer thickness (h) of 30 μm, and hatch spacing (t) of 100 μm, respectively. Correspondingly, the laser energy inputs were determined to be 292, 222, 194, and 155 J/mm3, respectively. Second, the scan path of laser was alternating ± 45° to the x-axis between consecutive layers (Fig. 2). The as-printed samples had the dimensions of 8 × 8 × 8 and 45 × 10 × 10 mm aligned to x, y and z directions, respectively. The dog-bone shaped samples (ASTM E8/E8M) with a gauge section of 3 mm in width, 2 mm in thickness, and 14 mm in length were cut out parallel to the x-y plane of the as-printed samples to perform roomtemperature tensile and SME tests at a strain rate of 5 × 10−4 in a

universal testing machine (Instron 8862) equipped by an electronic extensometer. Particularly, stress-controlled cyclic tensile experiment was adopted for SME test. In every cycle, the samples were loaded to a very high stress of 400 MPa at room temperature (100 MPa more than of that utilized in Ref. [12]), then unloaded to 20 MPa; the loaded and unloaded strain rate is 5 × 10−4. After ten times cycle, the samples were heated to a temperature of 30 °C higher than Af (austenite transformation finish temperature) by 4 °C/min and held for 5 min to ensure temperature uniformity and enough recovery. Specifically, the recoverable strain after heating (εrec) and recovery ratio (η = εrec/ εtot × 100%, εtot is total irrecoverable strain before heating) were adopted to assess SME of the as-printed samples. To ensure data reproducibility, the reported tensile and SME properties were the average of three individual samples manufactured by the same condition of laser energy input. The 8 × 8 × 8 mm cubic samples was cut from the as-printed samples, and ground for every surfaces with SiC sandpapers with descending roughness of 360, 600, 800, 1200, and 1500 mesh, respectively. The relative density of the as-printed samples was measured to be above 99.5% from the average of five tests by Archimedes method based on the theoretical density (6.45 g/cm3). Phase component was examined by X-ray diffraction (XRD, D/MAX-2500/PC) with Cu Kα radiation. A differential scanning calorimetry (DSC, Netzsch 204F1) was used to investigate phase transformation temperatures at a heating/cooling rate of 10 °C/min in an argon atmosphere based on ASTM F2004-05. The ground cubic samples were polished with

Fig. 1. SEM image (a) and particle size distribution (b) of the as-received Ti-rich Ti50.6Ni49.4 powders. 2

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Fig. 3. (a) The XRD patterns of the as-printed Ti50.6Ni49.4 SMA by various laser energy inputs; (b) The change in the phase volume fraction of martensite and austenite with laser energy inputs.

descending size of 2.5, 1, and 0.5 mm diamond suspension, respectively, subsequently etched for 30 s with a mixture of 50 vol% H2O +40 vol% HNO3+10 vol% HF solution, and finally utilized as the microstructure characterization by the scanning electron microscopy (SEM, Philips XL-30 FEG). Especially, transmission electron microscopy (TEM) was applied to the microstructures analysis. TEM specimens were prepared by cutting from the ground cubic samples, ground further into a thickness of 80 μm using the abrasive paper, punched out φ3 mm disks, twin-jet electro-polished using a Struers TenuPol-5 system inside an mixture solution of 4 vol% perchloric acid and 96 vol% ethanol at −30 °C, and finally ion-beam thinned by a Gatan 691 system. TEM observation was performed on field emission gun high resolution TEM (JEM-2100) at an acceleration voltage of 200 kV. The mean grain size, D ‾ , was determined using individual average length determination in Image Pro-Plus software by counting > 200 grains inside bright-field TEM micrographs. 3. Results The XRD results in Fig. 3 illustrate that all samples printed contain assuredly two phase components of martensite (B19’) and austenite (B2). Based on the integrated areas [24,25] of the diffraction peaks from respective phases in Fig. 3a, the volume fraction of B19’ and B2 were calculated and presented in Fig. 3b. This demonstrates that the phase fraction of B19’ martensite decreases from 99%, 95%, and 90%–78% with the decrease of laser energy inputs from 292, 222, and 194 to 155 J/mm3, respectively. In addition, diffraction peaks from Ti2Ni precipitates may overlap those from B19’. Based on Fig. 4, the increasing laser energy inputs results in higher austenite transformation and martensitic transformation peak temperatures (Ap and Mp) and greater endothermic enthalpies of the two phase transformations, B19′→B2 during heating and subsequent B2→ B19’ during cooling, respectively. As shown in Fig. 5, it can be observed that the SEM microstructures of melting tracks are perpendicular to each other between consecutive layers along the X–Y cross-sections. Magnified observation indicates that melting tracks in the 292 J/mm3-printed sample consist of predominate ultrafine cellular grains. With the decreased laser energy inputs, however, the 222 J/mm3-printed sample is composed of a vast majority of fine dendritic grains and part of ultrafine cellular grains adjacent to molten pools edges. Generally, the formation of dendritic grains is attributed to lower cooling rate of molten pools under the lower laser energy input of 222 J/mm3. Reasonably, the formation of ultrafine cellular grains in the 222 J/mm3-printed sample was attributed to higher cooling rate adjacent to molten pools edges. Notice that the microstructures reported herein are totally different from those of SLMed TiNi SMAs reported so far [8,17], large area of square grains, part small grains or pulled S-shaped grains. Reasonably, this is due to rapid solidification resulted from high enough energy inputs in this

Fig. 4. The DSC scans of the as-printed TiNi SMA by various laser energy inputs.

work. In addition, there are no visible micro-pores and micro-cracks in the as-printed samples, which agree with their nearly full density exceeding 99.5%. Fig. 6 shows the room-temperature tensile engineering stress-strain curves of the as-printed TiNi SMA by various laser energy inputs. Apparently, all samples exhibit excellent combination of high tensile 3

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Fig. 5. The SEM microstructures of the as-printed Ti50.6Ni49.4 SMA: (a) and (b) 292 J/mm3, (c) and (d) 222 J/mm.3.

elastic deformation (Ⅰ), martensite variant reorientation or detwinning (Ⅱ), elastic deformation of reoriented martensite (Ⅲ), and plastic deformation of reoriented martensite (Ⅳ). Noted that, the critical stress of martensite variant reorientation and detwinning, σc, was illustrated in Fig. 6 by frequently-used tangent method [26]. The σc decreases from 147 to 112 MPa for the decreased laser energy inputs from 292 to 155 J/mm3, respectively. Fig. 7 illustrates the stress-controlled cyclic tensile curves and resultant SME during heating the as-printed TiNi SMA. All samples display significant εrec after heating to Af + 30 °C. This is attributed to that the phase transformation of B19′→B2 during heating makes plastic strain in B19’ phase disappear gradually [27–29]. Especially, the asprinted SMA has ultrahigh SME in terms of η of 98.7% and εrec of 4.99% (Table 1 and Fig. 7). Clearly, the values of the εmax and εrec increase with the decrease of laser energy inputs (Table 1). This is consistent with the content change of B2 austenite in the as-printed samples (Fig. 3). The austenite would like to transform into martensite during loading and thus increase the deformation ability. These SME values reported herein are far greater than the corresponding ones (the η of 75.2% and εrec of 3.54% in fourth cycle at 300 MPa [12]) for SLMed Nirich Ni50.1Ti49.9 SMA. Therefore, it is the first time to report such ultrahigh SME for TiNi SMAs by SLM. Reasonably, inevitable irrecoverable strain after heating (εirrec) is attributed to dislocations generation and retained stabilized martensite after cyclic tensile deformation [8,13]. Accordingly, these results consolidate that the terminology 4D printing can provide a possibility to produce shape-complicated TiNi SMAs components with outstanding SME by additive manufacturing. Figs. 8 and 9 present TEM images and corresponding selected area electron diffraction (SAED) patterns of the 292 J/mm3 and 222 J/mm3printed TiNi SMAs. First, cellular grain structure in the 222 J/mm3printed sample (Fig. 9 a) exhibits more regular equiaxed morphology relative to counterparts in the 292 J/mm3-printed one (Fig. 8 a). Such cellular equiaxed structure in the 222 J/mm3-printed sample can be classified into two ultrafine phase regions, i.e., R (Fig. 9 b) and B2

Fig. 6. Room-temperature tensile engineering stress-strain curves of the asprinted TiNi SMA by various laser energy inputs. σc: the critical stress of martensite variant reorientation and detwinning; σUTS: ultimate tensile strength; δ: elongation.

strength and large elongation. Especially, the 222 J/mm3-printed sample has ultrahigh ultimate tensile strength (σUTS) of 776 MPa as well as large elongation (δ) of 7.2%. Surprisingly, these values are 170 MPa and 0.4% higher than the corresponding ones for SLMed Ni-rich Ni50.1Ti49.9 (606 MPa and 6.8% [12]). To our best knowledge, it is the first time to report such ultrahigh tensile properties for TiNi SMAs by SLM. As shown in Fig. 3, the as-printed TiNi SMA consists of predominant B19’ martensite, which would like to undergo several sequential transformation stages during tensile test (Fig. 6) of physically 4

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Fig. 7. The stress-controlled cyclic tensile curves and resultant shape memory effect during heating the as-printed TiNi SMA. εrec: reversible strain after heating; εa: accumulated strain during cyclic tests.

Fig. 10 shows grain size distribution histograms of cellular grains in 292 J/mm3 and 222 J/mm3-printed samples based on TEM observations. From the figure, it is found that the average grain size of the cellular grains in the 292 and 222 J/mm3-printed samples is 770 nm and 587 nm, respectively. The smaller grain size in the 222 J/mm3printed sample is associated to the laser scanning speed as discussed below.

(Fig. 9 c) regions. The formation of R phase is associated to high density dislocations along grain boundaries and interiors of ultrafine grains [4]. In contrast, the cellular structure in the 292 J/mm3-printed sample consists of ultrafine B19’ nanotwin martensites (Fig. 8 a and b), as well as abundant capsule-like nano-sized Ti2Ni precipitates with 30 nm in diameter and 20–100 nm in length (Fig. 8 c). Second, fine dendritic structures in the two samples are composed of interdendritic B19’ nanotwin martensites matrix (Fig. 8 d and 9 d) and B2 austenite branch (Fig. 8 f and 9 f). Noted that, remarkable differences are observed in B2 austenite branches: (1) Nano-sized Ti2Ni precipitates (5–10 nm) are dispersedly distributed in the 222 J/mm3-printed sample (Fig. 9 f), while greater-sized Ti2Ni precipitates (30–150 nm) together with a great deal of dislocations co-exist inside the 292 J/mm3-printed one (Fig. 8 f); (2) Finer nanotwins and their interactions with nano-sized Ti2Ni precipitates are likely to form in the 292 J/mm3-printed one (Fig. 8 e). The precipitation of nano-sized Ti2Ni in our case is similar to that in Ti49.1Ni50.9 [1] and Ti55.7Ni44.3 [11] SMAs by SLM.

4. Discussions 4.1. Microstructure formation mechanism and its effect on shape memory property in the SLM-manufactured SMA As presented above, the as-printed samples are composed of B19’ TiNi martensite and B2 TiNi austenite according to XRD analysis (Fig. 3). Furthermore, TEM observations indicate that the as-printed samples contain minor Ti2Ni precipitates (Figs. 8 and 9), and R phase

Table 1 Summary of the SME for the as-printed TiNi SMA. εmax: maximal strain; εtot: total irrecoverable strain before heating; εrec: recoverable strain after heating; η: recovery ratio after heating. Sample (J/mm3)

292 222 194 155

1st

10th

εmax (%)

εtot (%)

4.30 5.09 5.16 6.06

3.42 3.77 4.33 5.18

± ± ± ±

0.23 0.27 0.39 0.38

± ± ± ±

0.21 0.29 0.33 0.40

εmax (%)

εtot (%)

4.65 5.12 5.29 6.30

3.89 3.93 4.58 5.43

5

± ± ± ±

0.21 0.27 0.34 0.31

± ± ± ±

εrec (%) 0.20 0.15 0.27 0.34

3.57 3.88 4.47 4.99

± ± ± ±

η (%) 0.13 0.10 0.24 0.26

91.8 98.7 97.6 91.9

± ± ± ±

1.3 1.2 0.5 1.0

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Fig. 8. TEM images and corresponding SAED patterns of the 292 J/mm3-printed TiNi SMA: Bright-field images of (a) cellular ultrafine martensites with (b) nanotwins marked by red rectangle in (a), (c) darkfield image of nano-sized Ti2Ni precipitates along grain boundaries of ultrafine martensites marked by yellow rectangle in (a); (d) Bright-field images of dendritic structure with interdendritic B19’ marked by green rectangle (e) and dendritic branch marked by blue rectangle (f), respectively. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)

(TiNi) specifically in the 222 J/mm3-printed sample (Fig. 9). Intrinsically, the formation of these phases in present work is owing to non-equilibrium rapid solidification during SLM process. According to the TiNi binary phase diagram in Fig. 11a [4,30], Ti-rich alloy systems like the alloy composition Ti50.6Ni49.4 in present work should tend to form Ti2Ni in TiNi matrix under equilibrium solidification condition. This can be confirmed by uniformly distributed Ti2Ni inside TiNi matrix for equilibrium solidified Ti50.5Ni49.5 SMA [31]. In contrast, the heating/cooling rate inside molten pools can reach as high as 103–108 K/s during SLM process [6]. Such high cooling rate results in non-equilibrium solidification in SLMed TiNi SMAs. This would cause non-uniform distribution of minor Ti2Ni phase, as seen in TEM observations (Figs. 8 and 9). Noticely, the ultrafine cellular grains and fine dendritic grains (Figs. 5, 8 and 9) reported herein are totally different from those of SLMed TiNi SMAs reported so far. Meanwhile, the 222 J/mm3-printed sample has predominant dendritic grains, finer cellular grains, and specifically formed R phase (Figs. 8 and 9), compared with those of the 292 J/mm3-printed one. The microstructure differences in phase morphology and size and content are related to non-equilibrium rapid solidification during SLM under high enough energy inputs in this work. Generally, during SLM of metallic powders, laser beams scan across powder particles, melting powder particles completely under enough energy inputs to form molten pools. Temperature gradient between the

interior and edge of molten pools and the adjacent powder particles leads to non-equilibrium rapid solidification of molten pools. The higher cooling rate adjacent to molten pools edges relative to the interior of molten pools increases nucleation rates of grains adjacent to molten pools edges [32,33]. This results in the formation of the ultrafine cellular grains adjacent to molten pools edges and the fine dendrite grains inside the interior of molten pools in the 222 J/mm3-printed sample (Fig. 5 c and d). Academically, SLM processing parameters, such as the laser power and scanning speed and resultant the laser energy input, can significantly affect morphology of molten pools and thus microstructures of SLM-manufactured components. Specifically, at a given laser power, the smaller the laser scanning speed, or the greater the laser energy input, the deeper and wider the molten pools. Otherwise, at a given laser scanning speed, the smaller the laser power, or the smaller the laser energy input, the shallower and narrower the molten pools. Herein, the laser powder used is 70 W, lower than the corresponding ones (100–250 W) in Refs. [8,17]; the laser scanning speeds used are 80–150 mm/s, lower than the corresponding ones (175–1500 mm/s) in Refs. [8,17]; the laser energy inputs used are 155–292 J/mm3, greater than the corresponding ones (27.8–158.7 J/ mm3) in Refs. [8,17], respectively. Therefore, these SLM processing parameters, especially high enough energy inputs determined by other SLM parameters, can simultaneously affect the width and depth of the molten pools, bringing about the complicated influences and thus Fig. 9. TEM images and corresponding SAED patterns of the 222 J/mm3-printed TiNi SMA: (a) Bright-field images of cellular grain structure with R phase marked by green rectangle (b) and B2 phase marked by blue rectangle (c), respectively; Bright (d and e) and dark (f)-field images of dendritic structure, (e) interdendritic B19’ marked by red rectangle in (d), and (f) dendritic branch B2 marked by yellow rectangle in (d). (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)

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Fig. 10. Grain size distribution histograms of cellular grains in 292 J/mm3 (a) and 222 J/mm3-printed (b) TiNi SMA.

Fig. 11. (a) The Ti–Ni binary phase diagram; (b) Schematic illustration of phase transformations between austenite and martensite upon loading and unloading, and heating and cooling.

pools in the 222 J/mm3-printed sample have the smaller size (Fig. 10) compared to those in the 292 J/mm3-printed one. The B2 austenite and B19’ martensite are of CsCl type ordered and monoclinic structures, respectively. Essentially, shape memory properties and superelasticity of SMAs originate from phase transformations between B2 ↔ B19’ during loading and unloading, and cooling and heating. During tensile loading with the increased stress, B19’ twinned martensite first deforms elastically and transforms into detwinned martensite (Fig. 11b), followed by elastic deformation of detwinned martensite. The detwinning process of martensite contributes to the two elastic deformation stages in Fig. 6. During tensile unloading, the elastic deformation recovers and the detwinned martensite remains. During heating of the cyclic tensile samples, B19’ detwinned martensite transforms into B2 austenite (Fig. 11b), which results in the recovery of formed tensile strain. The 222 J/mm3-printed sample has dispersedly distributed Ti2Ni precipitates together with higher dislocation density (Figs. 8 and 9). This can significantly reduce the formation of dislocation and slip during the cyclic tensile process, which lead to the higher recoverable strain (εrec) and recovery ratio (η) after heating (Table 1).

totally different SEM microstructures in Fig. 5 from those of SLMed TiNi SMAs in Refs. [8,17]. Regarding the microstructure difference in the 292 and 222 J/mm3printed samples, it can be explained by the difference in the laser energy inputs induced by the laser scanning speed. According to the aforementioned formula E (= P/(v⋅h⋅t)), the higher energy input (E) of 292 J/mm3 corresponds to the smaller scanning speed (v) of 80 mm/s under the fixed laser power (P) of 70 W. Theoretically, the cooling rate of the molten pools is decided by the two aspects, i.e., the laser energy input and the laser scanning speed. The greater the laser energy input, the higher the temperature of the molten pools, and thus the higher the cooling rate of the molten pools. Meanwhile, the slower the laser scanning speed, the lower the cooling rate of the molten pools. Herein, it seems that the laser energy input exert main influence on the microstructure. In other words, the higher cooling rate of the molten pools in the 292 J/mm3-printed sample forms the stronger non-equilibrium solidification condition and consequently cause much more uneven distribution of Ti2Ni precipitates relative to the 222 J/mm3-printed one (Figs. 8 and 9). Meanwhile, the deeper and wider molten pools at the greater laser energy input of 292 J/mm3 leads to the homogeneous nucleation and growth of grains in the molten pools [33,34] and the more homogeneously predominate ultrafine cellular grains in the asprinted SMA (Figs. 5 and 8). On the other hand, the faster laser scanning speed (105 mm/s) would lead to the shallower and narrower the molten pools in the 222 J/mm3-printed sample. This may give rise to the bigger temperature gradient between the interior and edge of the molten pools and the higher cooling rate of the edge of the molten pools. Hence, the cellular grains adjacent to the edge of the molten

4.2. The strengthening mechanism in the SLM-manufactured SMA On the basis of the aforementioned microstructural analyses, the utrahigh tensile properties and SME of the 222 J/mm3-printed sample can be rationalized as combination of various strengthening mechanisms. Firstly, no visible defects, such as microcracks and porosities form in the as-printed samples (Fig. 5), which is the foundation of ultrahigh tensile properties and SME. Secondly, grain refinement strengthening 7

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[35,36] is active in whole tensile process (Fig. 6) because of the mixture of fine dendritic structure and ultrafine cellular structure. Specifically, part of ultrafine cellular grains were equiaxed ultrafine-grained R phases with dislocation networks (Fig. 9 b), lots of aggregative dislocations at grain boundaries could hinder formation and motion of dislocations, which improve the strength and ductility [37–39]. The difficulty in the formation and motion of dislocations during tensile process is effective to enhance SME. Thirdly, the grain-interior nanosized precipitates in dendritic structures (Fig. 9 f) would enhance the strength and ductility according to Orowan mechanism [35]. During tensile test, dislocations form inside dendritic grains and are pinned by the nano-sized precipitates, which would enhance the strength and ductility. These factors make the 222 J/mm3-printed sample had ultrahigh ultimate tensile strength of 776 MPa and large elongation of 7.2% as well as SME of 98.7% recovery ratio. In contrast, the grain-boundary nano-sized Ti2Ni precipitates (Fig. 8 c), interactions between nanotwins and adjacent precipitates (Fig. 8 e) and the precipitation of nano-sized Ti2Ni together with substantial dislocations (Fig. 8 f) in the 292 J/mm3-printed sample are effective to hinder martensite variant reorientation and detwinning, and thus incline to produce stress concentration, dislocations and microcracks. Certainly, this leads to higher σc (147MPa) and lower elongation (5.8%) (Fig. 6) as well as greater εa (0.47%) (Fig. 7) [40,41].

[2] [3]

[4] [5]

[6] [7] [8]

[9]

[10]

[11]

[12]

5. Conclusions

[13]

In summary, ultrahigh-performance Ti-rich Ti50.6Ni49.4 SMA was successfully produced by 4D printing of SLM under high enough energy inputs. Results show that although the as-printed SMA samples contain the B19’ martensite and B2 austenite, the fraction of B19’ martensite and the austenite and martensitic transformation peak temperatures decrease with the decrease of laser energy inputs from 292 to 155 J/ mm3. The formation of fine dendritic grains and ultrafine cellular grains in the as-printed SMA samples are totally different from those of SLMed TiNi SMAs reported so far. Especially, the as-printed SMA has ultrahigh ultimate tensile strength of 776 MPa and large elongation of 7.2% as well as ultrahigh SME of 98.7% recovery ratio, far superior to the corresponding ones of similar TiNi SMAs by SLM in literature. The different microstructure and the ultrahigh mechanical properties and SME herein can be explained by the non-equilibrium rapid solidification of molten pools resulted from the high enough energy inputs by SLM. The results authenticate that high enough energy input is feasible to achieve nearly full density and avoid cracking tendency in TiNi SMAs.

[14]

[15]

[16]

[17]

[18]

[19]

[20]

[21]

Data availability [22]

The raw/processed data required to reproduce these findings cannot be shared at this time as the data also forms part of an ongoing study. [23]

Acknowledgements

[24]

The current work was supported by the National Special Zone Project of Science and Technology Innovation, the National Natural Science Foundation of China (Nos. 51574128 and 51627805), and the Guangdong Natural Science Foundation for Research Team (No. 2015A030312003).

[25]

[26] [27]

Appendix A. Supplementary data

[28] [29]

Supplementary data to this article can be found online at https:// doi.org/10.1016/j.msea.2019.138166.

[30] [31]

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