Ti–Ni shape memory alloy composite with the larger recoverable strain

Ti–Ni shape memory alloy composite with the larger recoverable strain

Composites Communications 23 (2021) 100583 Contents lists available at ScienceDirect Composites Communications journal homepage: www.elsevier.com/lo...

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Composites Communications 23 (2021) 100583

Contents lists available at ScienceDirect

Composites Communications journal homepage: www.elsevier.com/locate/coco

The higher compressive strength (TiB+La2O3)/Ti–Ni shape memory alloy composite with the larger recoverable strain Xiaoyang Yi a, b, Guijuan Shen a, Xianglong Meng a, *, Haizhen Wang b, **, Zhiyong Gao a, Wei Cai a, Liancheng Zhao a a b

School of Materials Science and Engineering, Harbin Institute of Technology, Harbin, 150001, China College of Nuclear Equipment and Nuclear Engineering, Yantai University, Yantai, 264005, China

A R T I C L E I N F O

A B S T R A C T

Keywords: Ti-Ni shape Memory alloy composite Microstructure Thermal analysis Mechanical properties

In the present study, in-situ TiB whiskers and La2O3 particles showing a network structure were designed and introduced into Ti–Ni composites. With LaB6 content increasing, the amount of in-situ TiB whiskers and La2O3 particles also increases. In addition to the in-situ reinforcements, Ti2Ni second phase also precipitates. The B19’⇌B2 martensitic transformation can be detected in all Ti–Ni composites, irrespective of LaB6 contents. It should be noted that R-phase transformation also appears for Ti–Ni composite added 0.5wt%LaB6. The martensitic transformation temperatures firstly decrease and then increase with the increased LaB6 content, which can be attributed to the changing of matrix composition. As LaB6 content increases from 0wt% to 0.97wt %, the evolution of network structure is as follows: quasi-continuous network structure → continuous network structure → quasi-continuous network structure. In proportion, the mechanical properties of Ti–Ni composites firstly decrease and then increase. The superior mechanical properties with the highest fracture strain of 35.9% and the largest fracture stress of 2476 MPa can be obtained in Ti–Ni composite with adding 0.97wt%LaB6, which is due to the larger capacity of bearing the loading for reinforcements showing a quasi-continuous network structure. Meanwhile, at the pre-strain of 7%, the maximum strain recovery ratio of Ti–Ni composite is 88.34%.

1. Introduction Binary Ti–Ni shape memory alloys possess the unique functional performances including shape memory effect (SME) and superelasticity (SE). The superior functional properties origin from a reversible thermoelastic martensitic transformation between a higher temperature austenite phase (B2) and a lower temperature martensite phase (B19’) [1,2]. Apart from the SME and SE, the higher strength is necessary to widen the industrial applicability of Ti–Ni alloys [3]. The effective measures of improving the mechanical properties mainly consist of aging treatment, thermo-mechanical treatment, alloying and introduc­ tion of reinforcement etc. [4–11]. In contrast, the power metallurgy is the most effective ways to introduce reinforcements into the metal matrix composite. To date, the ex-situ reinforcements in Ti–Ni com­ posites mainly consist of nano-scale Al2O3 particle, SiO2 and TiC etc., while the in-situ reinforcements in Ti–Ni composites are primarily composed of TiC, TiO2 and Ni2Si and so on [12–16]. Moreover, the Ti–Ni

based composites reinforced by in-situ reinforcement phase are featured with more superior mechanical and functional performances due to the perfect interface between the in-situ reinforcement and matrix, compared with the Ti–Ni composites reinforced by ex-situ re­ inforcements. For instance, the critical stress for stress induced martensitic transformation and two-way shape memory effect strain increase with the increasing of TiC content, while the shape memory effect strain deteriorates due to TiC addition [9]. The CNT/Ti–Ni com­ posite not only shows the larger elastic modulus, higher yield strength and fracture strength, but also exhibits the superior SME and SE [17]. Under the condition of 8% pre-strain, the maximum SME strain recovery ratio is 92.5% and largest SE strain recovery ratio of 86.3% can be ob­ tained in 1 vol%CNT/Ti–Ni composite [17]. Nevertheless, Ti–Ni com­ posites prepared from elemental Ti and Ni as well as the reinforcement are characterized by the lower density, lower degree of purity and poor mechanical properties [10,11]. The above drawbacks can be avoided to some extent, when pre-alloyed Ti–Ni powders are chosen as the original

* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (X. Meng), [email protected] (H. Wang). https://doi.org/10.1016/j.coco.2020.100583 Received 8 July 2020; Received in revised form 15 November 2020; Accepted 26 November 2020 Available online 1 December 2020 2452-2139/© 2020 Elsevier Ltd. All rights reserved.

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materials. However, the reinforcement in recent investigation about Ti–Ni composites is focus on the ex-situ phase [13]. Recent studies also reveal that the outstanding mechanical properties can be achieved by changing the distribution state of reinforcements. Compared with the metal based composites enhanced by the uniformly distributed re­ inforcements, the metal based composites with the reinforcements showing the network structure can provide the greater strength owing to the ability of bearing a larger loading of reinforcement with a network structure [18]. Thus, in the present study, it is expected to introduce the in-situ TiB and La2O3 reinforcements into the Ti–Ni shape memory alloy to improve the strength and shape recovery characteristics. In addition, the distri­ bution state of in-situ reinforcement is adjusted and the content of in-situ reinforcements on the martensitic transformation behavior and me­ chanical properties as well as the strain recovery characteristics of Ti–Ni composite are investigated. 2. Experimental procedures 2.1. Theoretical basis In order to introduce the in-situ TiB and La2O3 reinforcements in the Ti–Ni composites, the available raw materials of LaB6 ceramic particles were adopted. The change of Gibbs free energy (ΔG) and enthalpy change (ΔH) as a function of temperature for in-situ reaction (12Ti+2LaB6+3[O]→12TiB+La2O3) during the hot pressed sintering in our reported studies reveal that introducing in-situ TiB and La2O3 phase is reasonable in thermodynamically by adopting LaB6 ceramic particles as original materials [19]: 2.2. Preparation of Sample Fig. 2. The microstructural feature of Ti–Ni composite with the different LaB6 contents (a): 0wt%; (b): 0.3wt%; (c): 0.5wt%; (d): 0.7wt%; (e):0.97wt%; (f): 1.5wt%.

In the present study, the gas-atomized Ti50Ni50 alloy powders with the particle size ranging from 50 μm to 100 μm was offered by AVIC Maite powder metallurgy technology (Beijing) Co., Ltd. It can be seen from Fig. 1(a) that the gas-atomized Ti–Ni alloy powders were ideally smooth and spherical. Nevertheless, some tiny satellite balls attached at the surface of larger alloy powders also can be observed. The LaB6 ceramic particles with a purity of 99.9% were purchased from Beijing Haoke Technology Co., Ltd, as shown in Fig. 1(b). To form in-situ TiB and La2O3 reinforcements with the network structure, the lower energy ball milling was employed. The experimental procedures of the prepa­ ration of samples and sampling mode were illustrated in Fig. 2. In addition, the details of ball milling and hot-pressed sintering has been reported in our previous literature [20].

2.3. Characterization of Samples The identification of phase constituents in Ti–Ni composites was performed on an X-Ray diffraction (XRD) equipped with Cu Kα radiation at room temperature. The microstructural features of Ti–Ni composites were characterized by the field-emission scanning electron microscope (FE-SEM) and transmission electron microscope (TEM). The samples for the XRD analysis and SEM observation were in polished condition. The TEM foils were electro-polishing using a twin-jet method, which

Fig. 1. The illustration showing the preparation and characterization of Ti–Ni composites. 2

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executed at around − 30◦ C. The effect of LaB6 content on the martensitic transformation behaviors of Ti–Ni composites was investigated by adopting the differential scanning calorimetry (DSC) with a heating and cooling rate of 20 ◦ C/min. The compressive test was measured by an Instron-5569 machine with a crosshead speed of 0.2 mm/min at room temperature. Strain recovery characteristics were obtained at room temperature by comparing the samples length before deformation, after various applied strains and heating above the finish temperature of reverse martensitic temperature. The detailed testing procedures have been stated in our published paper [21]. 3. Results and discussions 3.1. Microstructural features Fig. 2 shows the backscattered electron (BSE) images of Ti–Ni com­ posites added the different contents of LaB6 ceramic particles. Obvi­ ously, the microstructure of sintered Ti–Ni alloy without LaB6 addition is characterized by the light white matrix and grey second phases, as shown in Fig. 2(a). The grey second phase distributes regularly, mainly appearing at the surface of original Ti–Ni alloy powders and forming the network structure. Combined with the results in our previous literatures [22], it can be speculated that the grey second phase is Ti2Ni phase. The LaB6 addition results in the appearance of black whisker reinforcement and white reinforcement with a spherical shape in Ti–Ni composites. Moreover, smaller amount of LaB6 ceramic particles result in the increasing of Ti2Ni second phase. Lots of Ti2Ni second phases connected closely, forming the continuous network structure. The matrix is divided into multiple isolated units. With the increased LaB6 content, the amount of black and white reinforcement continuously increases. However, the number of grey Ti2Ni second phase gradually decreases owing to LaB6 addition, as illustrated in Fig. 2(b)~(f). The reduction of grey Ti2Ni second phase is related to the consumption of titanium in matrix owing to the occurrence of in-situ reaction. Similarly, the excessive LaB6 ceramic particle results in the network structure evolving from quasi-continuous to continuous network structure. In order to identify the in-situ reinforcements in Ti–Ni composites, the HADDF images and STEM-EDS map scanning micrographs are per­ formed and the results are shown in Fig. 3(a)~(f). It can be observed that there are larger volume of second phases, which are rich in Ti and poor in Ni and can be considered as Ti2Ni second phase. Besides, the smaller particles containing La and O element as well as the tiny whiskers with Ti element and B element also can be detected. Meanwhile, the selected area electron diffraction (SAED) patterns taken from the different re­ inforcements are shown in Fig. 3(g)~(i). The results further confirms that the in-situ reinforcements are TiB and La2O3 reinforcements.

Fig. 3. (a): HADDF-STEM image and (b~f): corresponding STEM-EDS map scanning results as well as the (g) TEM image and (h~i) SAED pattern of Ti–Ni composite with 0.97wt%LaB6 ceramic particle.

of Ti in the matrix and further causing the decrease of martensitic transformation temperatures. Thus, when the content of LaB6 ceramic particle increases from 0.3wt% to 0.5 at.%, the martensitic trans­ formation temperatures of Ti–Ni composite decrease. The subsequent increase of martensitic transformation temperatures of Ti–Ni composites with the further increased LaB6 content can be attributed to the intro­ duction of La element. It has been revealed that the martensitic trans­ formation temperatures of Ti–Ni alloy can be increased by 62◦ C per 1 at. %La addition [24]. Besides, the mismatch internal stress can be intro­ duced owing to the different coefficient of thermal expansion between the in-situ reinforcements and matrix [25]. The amount of in-situ rein­ forcement rises with the increased LaB6 content. In turn, the mismatch internal stress become gradually larger and further elevate the martensitic transformation temperatures [26]. This may be another factor in raising the martensitic transformation temperatures of Ti–Ni composites with the excessive content of LaB6 ceramic particles. To verify the transformation in Ti–Ni composite with 0.5wt%LaB6, the partial DSC curves are carried out and the results are shown in Fig. 4 (c). The results confirm that the extra exothermic peak during cooling in Ti–Ni composite containing 0.5wt%LaB6 can be ascribed to R-phase transformation. This conclusion is judged from the narrower trans­ formation temperature hysteresis. The transformation temperature hysteresis of R-phase is smaller than 5◦ C [27]. Meanwhile, the R-phase transformation is featured with the higher stability of thermal cycling [28]. As shown in Fig. 4(d), the Ti–Ni composite added 0.5wt%LaB6 shows the perfect thermal cycling stability. With undergoing 15 thermal cycles, the transformation temperatures of R-phase transformation almost keep constant, while the peak temperature of forward martens­ itic transformation (Mp) and peak temperature of reverse martensitic

3.2. Martensitic transformation behaviors Fig. 4(a) shows the DSC curves of Ti–Ni composites with the various contents of LaB6 ceramic particles. Merely one endothermic peak and exothermic peak corresponding to B19’⇌ B2 transformation are observed in heating and cooling process in all Ti–Ni composites, apart from Ti–Ni composite with 0.5wt%LaB6. In addition. With the LaB6 content increasing, the martensitic transformation temperatures firstly decrease and then increases, as illustrated in Fig. 4(b). As the LaB6 content is increased from 0wt% to 0.5wt%, both start temperature of reverse martensitic transformation (As) and start temperature of for­ ward martensitic transformation (Ms) decrease from 71.5◦ C and 71.3◦ C to 60.0◦ C and 40.0◦ C, respectively. For comparison, the martensitic transformation temperatures of sintered Ti–Ni alloy are shown in Fig. S1. Upon the content of LaB6 ceramic particles is more than 0.5wt%, both As and Ms increase continuously to 81.5◦ C and 73.0◦ C. It is well known that the martensitic transformation temperatures of binary Ti–Ni shape memory alloys drop with the decreasing of Ti/Ni ratio [23]. In the present Ti–Ni composites, the LaB6 addition can lead to the consumption 3

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Fig. 4. The martensitic transformation behavior of Ti–Ni composite (a): DSC curves of Ti–Ni composite with various LaB6 contents; (b): Dependence of martensitic transformation temperatures on the content of LaB6 ceramic particle; (c): the partial DSC curves of Ti–Ni composite added with 0.5wt%LaB6; (d): effect of thermal cycling number on the martensitic transformation temperatures of Ti–Ni composite added with 0.5wt%LaB6.

transformation (Ap) decrease by 4.4◦ C and 3.7◦ C. However, the mech­ anism for the presence of R-phase transformation in the Ti–Ni compos­ ites is unclear and the possible reasons are processing.

higher than that of Ti–Ni composites enhanced by the in-situ rein­ forcement with the homogeneous distribution, as shown in Fig. 5(c) [12, 13,17]. Compared with the uniformly distributed in-situ reinforcements, the reinforcements showing the network structure can increase the volume fraction of reinforcement in local zones under the same volume fraction of reinforcement condition, leading to the larger stress con­ centration factor and the improved load-bearing capacity of the re­ inforcements, which can be confirmed the results of finite element method (FEM) in Fig. S2 [18]. However, the further increased LaB6 content would make the mechanical properties of Ti–Ni composite worsen. As well, the excessive LaB6 addition results in the formation of continuous network structure, which promotes the occurrence of stress concentration around the continuous network structure and deteriora­ tion of mechanical properties. Meanwhile, no obvious stress plateau in the compressive stress-strain curves is observed for Ti–Ni composites. Based on the existing literatures, the absence of stress plateau in the present Ti–Ni composites can be attributed to the stress transfer from Ti–Ni matrix to in-situ reinforcements with a quasi-continuous network structure [26]. The superposition of elastic deformation of in-situ rein­ forcement and reorientation/de-twining of martensite variants causes the disappearance of the stress plateau.

3.3. Mechanical properties Fig. 5(a) displays the compressive stress-strain curves of Ti–Ni composites with various LaB6 contents at room temperature. It can be seen that the Ti–Ni composites can be divided into three stages: <5% (I), 5%~13% (II), >13% (III) of applied strain, regardless of the LaB6 con­ tents. By comparison, 0.5wt%LaB6 addition decreases both the fracture strain and stress of Ti–Ni composite. The deterioration of mechanical properties for Ti–Ni composites with minor LaB6 ceramic particle can be ascribed to the formation of continuous network structure constructed by numbers of Ti2Ni second phases. The Ti–Ni composite is divided into many isolated matrix units by the continuous network structure, which is not favor to the transfer of stress and coordination of strain between the adjacent matrix units. With the LaB6 content increasing, the volume fraction of in-situ reinforcements in the quasi-continuous network structure zone also raises. Thus, the mechanical properties then are enhanced owing to the higher strength and hardness of in-situ TiB and La2O3 reinforcements, as shown in Fig. 5(b). The maximum fracture strain of 35.9% and largest strength of 2476 MPa can be obtained in Ti–Ni composite by adjusting the content (0.97wt%) of LaB6 ceramic particle, which is comparable to the strength of reported Ti–Ni based composites reinforced by nanowires [29–31]. Moreover, the strength of Ti–Ni composite with a quasi-continuous network structure is evidently

3.4. Strain recovery characteristics The strain recovery characteristics of Ti–Ni composites are listed in Table 1. The minor LaB6 addition causes the slight reduction of shape recovery ratio owing to the formation of continuous network structure. 4

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Fig. 5. (a): The mechanical properties of Ti–Ni composite with different LaB6 contents; (b) The evolution of mechanical properties and microstructure with the increased LaB6 ceramic particle: (c): Comparison of mechanical properties for Ti–Ni composites.

deformation is relatively less. Besides, the microstructure of Ti–Ni composites added 0.7wt% and 0.97wt%LaB6 is featured with the quasicontinuous network structure. The stress and strain can effectively transfer and coordinate between the adjacent matrix units, which can avoid the stress concentration and favor to the strain recovery characteristics.

Table 1 The strain recovery characteristics of Ti–Ni composites with various LaB6 con­ tents under the 7% pre-strain condition. LaB6 content (wt.%)

εE (%)

εSME(%)

εR(%)

Shape recovery ratio(%)

0 0.3 0.5 0.7 0.97 1.5

3.93 4.55 4.14 4.14 4.14 4.56

1.64 1.02 1.84 2.04 2.04 1.42

5.58 5.57 5.98 6.18 6.18 5.98

79.67 79.59 85.42 88.34 88.34 85.48

4. Conclusions In the present study, the in-situ TiB phase and La2O3 reinforcements with a network structure were synthesized successfully in Ti–Ni com­ posite through lower energy milling and hot pressed sintering. The amount of in-situ TiB and La2O3 reinforcements steadily increases with the increased LaB6 content. The quasi-continuous network structure can be constructed by controlling the moderate LaB6 ceramic particle. The minor or excessive LaB6 addition led to the formation of the continuous network structure mainly constituted by massive of Ti2Ni phase or insitu reinforcements. The martensitic transformation temperatures of Ti–Ni composites firstly decreased owing to the consumption of Ti element and then increased due to the introduction of La element and internal stress fields. Minor LaB6 addition deteriorate the mechanical properties and strain recovery characteristics as a result of the constraint of continuous network structure. The superior mechanical performances showing the maximum fraction stress and strain of 2476 MPa and 35.9% can be obtained in Ti–Ni composites added with 0.97wt%LaB6. In the meanwhile, as the LaB6 content was adjusted as 0.7wt% and 0.97wt%, Ti–Ni composites can show the largest recoverable strain of 6.18% with the pre-strain of 7%.

When the LaB6 content is further increased, the amount of in-situ rein­ forcement forming the quasi-continuous network structure gradually increases, which can share the much external loading and avoid the plastic deformation of matrix. However, the continuous network struc­ ture formed by massive of in-situ reinforcements can facilitate to the stress accumulation at the vicinity of continuous network structure, and making it easier to introduce the dislocation. Hence, the strain recovery characteristics firstly increase and then decrease, as the content of LaB6 ceramic particle increases from 0.3wt% to 1.5wt%. The maximum recoverable strain of 6.18% containing 4.14% elastic recovery strain and 2.04% shape memory effect strain can be obtained in Ti–Ni composites by tailoring the LaB6 content of 0.7wt% and 0.97wt%. The achievement of excellent strain recovery characteristics can be concluded as follows: In-situ TiB and La2O3 reinforcements are major in Ti–Ni composites with adding 0.7wt% and 0.97wt%LaB6. Compared with the network struc­ ture formed by Ti2Ni second phase, the network structure constructed by in-situ reinforcements can bear more loads owing to higher elastic modulus and strength of in-situ reinforcements. Under the equal external stress condition, the stress on the matrix is small and the plastic 5

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CRediT authorship contribution statement

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Xiaoyang Yi: Writing - original draft, preparation; design and perform the corresponding testing. Xianglong Meng: preparation of the specimens of characterization, Funding acquisition. Haizhen Wang: Writing - original draft, preparation; design and perform the corre­ sponding testing. Zhiyong Gao: Writing - review & editing, and dis­ cussions.. Wei Cai: Writing - review & editing, and discussions.. Liancheng Zhao: Writing - review & editing, and discussions.. 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 the present paper. Acknowledgements This work was supported by the National Natural Science Foundation of China (Grant Nos. 51871080, 51931004, 51571073 and 11875046). Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi. org/10.1016/j.coco.2020.100583. References [1] W.M. Huang, Z. Ding, C.C. Wang, J. Wei, Y. Zhao, H. Purnawali, Shape memory materials, Mater. Today 13 (2010) 54–61. [2] K. Otsuka, X. Ren, Physical metallurgy of Ti-Ni-based shape memory alloys, Prog. Mater. Sci. 50 (2005) 511–678. [3] S.A. Shabalovskaya, Surface, corrosion and biocompatibility aspects of nitinol as an implant material, Bio Med. Mater. Eng. 12 (2002) 69–109. [4] X.Y. Yi, K.S. Sun, W.H. Gao, H.Z. Wang, B. Sun, W. Yao, X.L. Meng, Z.Y. Gao, W. Cai, The precipitation behaviors, martensite transformation and superelasticity in the super-high stress assisted aged Ni-rich Ti-Ni alloy, Intermetallics 104 (2019) 8–15. [5] X.Y. Yi, W.H. Gao, H.Z. Wang, W. Yao, X.L. Meng, Z.Y. Gao, W. Cai, L.C. Zhao, Dependence of aging parameters on precipitation behavior, martensitic transformation and mechanical properties in the stress assisted aged Ni-Ti alloy, Mater. Sci. Eng., A 736 (2018) 354–363. [6] V. Demers, V. Brailovski, S.D. Prokoshkin, K.E. Inakekyan, Thermomechanical fatigue of nanostructured Ti–Ni shape memory alloys, Mater. Sci. Eng., A 513–514 (2009) 185–196. [7] P.C. Jiang, Y.F. Zheng, Y.X. Tong, F. Chen, B. Tian, L. Li, D.V. Gunderov, R. Z. Valiev, Transformation hysteresis and shape memory effect of an ultrafinegrained TiNiNb shape memory alloy, Intermetallics 54 (2014) 133–135. [8] Yi X Y, Meng X L, Cai W, Zhao L C. Larger Strain Recovery Characteristics of Ti-93. [9] K. Johansen, H. Voggenreiter, G. Eggeler, On the effect of TiC particles on the tensile properties and on the intrinsic two way effect of NiTi shape memory alloys produced by powder metallurgy, Mater. Sci. Eng., A 273–275 (1999) 410–414. [10] X. Feng, J.H. Sui, W. Cai, Processing of multi-walled carbon nanotube-reinforced TiNi composites by hot pressed sintering, J. Compos. Mater. 45 (2011) 1553–1557. [11] F.C. Yen, K.S. Hwang, Shape memory characteristics and mechanical properties of high-density powder metal TiNi with post-sintering heat treatment, Mater. Sci. Eng., A 528 (15) (2011) 5296–5305.

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