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Scripta Materialia 78–79 (2014) 53–56

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In situ synchrotron investigation of the deformation behaviorof nanolamellar Ti5Si3/TiNi composite

Daqiang Jiang,a,⇑ Shijie Hao,a Junsong Zhang,a Yinong Liu,b Yang Renc and Lishan Cuia,⇑aDepartment of Materials Science and Engineering, China University of Petroleum-Beijing, Changping, Beijing 102249, China

bSchool of Mechanical and Chemical Engineering, The University of Western Australia, Crawley, WA 6009, AustraliacX-ray Science Division, Argonne National Laboratory, Argonne, IL 60439, USA

Received 1 December 2013; revised 23 January 2014; accepted 27 January 2014Available online 6 February 2014

An in situ nanolamellar Ti5Si3/TiNi composite is prepared by arc melting based on the design principle of load sharing between ahard component and a phase transforming matrix. The composite showed a compressive strength of �2.5 GPa and a fracture strainof �35%. In situ synchrotron X-ray diffraction analysis revealed stage-wise load transfer between the two components and theachievement of 2.1% lattice elastic strain of the brittle ceramic compound Ti5Si3, demonstrating the effectiveness of strain matchingdesign.� 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Composites; Shape memory alloy; NiTi; Synchrotron; Strain matching

A recent research article published in Scienceproposed a new composite design strategy based on theconcept of strain matching between a phase-transform-ing alloy matrix and high-strength nanomaterials, anddemonstrated remarkable mechanical properties [1]. Inthis work, the remarkable mechanical properties wereachieved by matching the magnitude of the uniform lat-tice shear distortion of the martensitic phase transforma-tion in an NiTi shape memory alloy (SMA) and the giantelastic strain limit of high strength Nb nanowires. Thisstrain matching mechanism allows effective load transferbetween the two components, leading to the realizationof extremely large elastic strains in the Nb nanowires(>6%) and the achievement of remarkable mechanicalproperties of the composite [1].

This discovery has broken a long-standing barrier inmetals design [2] and opens new opportunities for high-performance composite materials utilizing SMAs [3]. Infact, much effort has been made in the past to developSMA-based composite materials, such as SMA-fiber-reinforced metal [4–6], polymer [7–9] and cement [10]

http://dx.doi.org/10.1016/j.scriptamat.2014.01.0341359-6462/� 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights

⇑Corresponding authors. Tel.: +86 10 89733975; fax: +86 1089731959; E-mail addresses: dq80jiang@126.com (D. Jiang); lscui@cup.edu.cn (L. Cui)

matrix composites, and TiC/SiC-particle-reinforcedTiNi matrix composites [11–16]. In most cases, suchcomposites are designed to harness and improve the no-vel properties of the SMAs. In this new composite designstrategy, the function of the SMA matrix is different: it isto support the nanomaterials to harness their intrinsicqualities to achieve extraordinary mechanical properties.In this design, two challenges remain: the fabrication ofSMA matrix nanocomposites in bulk forms and under-standing the microscopic mechanisms of strain couplingand load sharing between the constituents.

It is known that many nanomaterials exhibit extraor-dinary mechanical properties, e.g. strengths of the orderof several gigapascals and elastic strains in the range of4–7% [17–22]. Owing to this, it has long been our desireto develop bulk composites using these materials toachieve extraordinary mechanical properties. However,it is difficult to fabricate bulk nanocomposites in indus-trial processes. One possible method is via eutectic solid-ification, which is known to produce microstructureswith nanoscale constituents. The second issue is thedeformation mechanisms of such composites at themicroscopic scale. Shape memory alloys exhibit a num-ber of novel thermal and mechanical properties relatedto their thermoelastic martensitic transformations.These include the ability to recover large non-elastic

reserved.

Figure 1. Microstructure characteristics of the as-cast Ti53Ni41Si6alloy. (a) XRD pattern. (b) SEM backscattered electron micrograph.The inset is a high-magnification SEM image. (c) TEM bright-fieldmicrograph of the eutectic structure.

Figure 2. Stress–strain curve of the Ti5Si3/TiNi composite at roomtemperature.

54 D. Jiang et al. / Scripta Materialia 78–79 (2014) 53–56

(pseudoelastic) deformation spontaneously uponunloading [23], the ability to delay such recovery uponheating at a different location and time (the shape mem-ory effect) [24], and the ability to undergo large non-elas-tic deformation without massive dislocation movementor strain hardening (pseudoplasticity via martensite var-iant detwinning) [25,26]. These ask questions about loadtransfer and sharing at the microscopic scale within acomposite structure, where the SMA component maydeform with practically zero stiffness within a finitestrain range [27,28] whilst the other component deformselastically with continuous stiffening. Such a situation isunprecedented in metal matrix composites using con-ventional elastic–plastic metals, thus raising new ques-tions to be answered.

In this study an in situ nanolamellar Ti5Si3/TiNicomposite was fabricated via eutectic solidification andthe microscopic deformation mechanisms between thebrittle ceramic Ti5Si3 phase and the pseudoplasticB190-TiNi matrix was investigated in situ during com-pression by means of synchrotron diffraction.

The SMA alloy composite system selected for studywas Ti–Ni–Si. This ternary system presents a pseudo-eu-tectic reaction between the near-equiatomic TiNi and theTi5Si3 ceramic compound [29]. In this study, an alloy in-got of �100 g with a nominal composition of Ti53Ni41Si6(at.%) was produced from high-purity elemental metalsof the constituents (purity > 99.8 wt.%) by arc melting.The phase composition was characterized by means ofX-ray diffraction (XRD; Bruker AXS D8) using Cu Ka

radiation. An FEI Quanta 200F scanning electronmicroscope (SEM) and a Tecnai F20 transmission elec-tron microscope (TEM) were used for microstructurecharacterization. Cylindrical specimens 5 mm in diame-ter and 10 mm in length were spark cut from the ingotfor compression testing. Compression test was con-ducted using a WDW-200 material test system at a strainrate of 1 � 10�3 s�1 at room temperature. In situ syn-chrotron high-energy X-ray diffraction (HEXRD) mea-surement during compression test was performed at the11-ID-C beamline of the Advanced Photon Source at Ar-gonne National Laboratory. A high-energy X-ray with abeam size of 0.4 mm � 0.4 mm and a wavelength of0.010798 nm was used to obtain two-dimensional (2-D)diffraction patterns in the transmission geometry.

Figure 1 shows the microstructure characterization ofthe as-cast Ti53Ni41Si6 alloy. The XRD spectrum shownin Figure 1(a) reveals the sharp diffraction peaks of theB190-TiNi and Ti5Si3 phases. The SEM micrograph inFigure 1(b) shows that the alloy has a typical hypoeutec-tic microstructure, consisting of primary TiNi dendrites(bright) and an ultrafine eutectic lamellar structure(gray). The inset in Figure 1(b) shows the TiNi (bright)and Ti5Si3 (dark) eutectic lamellar structures at highermagnification. Figure 1(c) is a TEM bright-field micro-graph of the eutectic structure. The lamellar thicknessesare about 120 nm for TiNi and 50 nm for Ti5Si3. Thevolume fraction of the pre-eutectic TiNi is determined,by means of image analysis, to be �70% and that ofTi5Si3 in the eutectic is determined to be 18% (or 60%in the eutectic).

Figure 2 shows a compression engineering stress–strain curve of the Ti5Si3/TiNi composite at room

temperature, with the stress and strain given as positivevalues. The inset in Figure 2 is the corresponding truestress–strain curve. The engineering stress–strain curvemay be seen in three stages: below 7%, from 7% to13%, and above 13% macroscopic strain. The first stageis characterized by a very low yielding at 120 MPa and avery low modulus of 7 GPa. The second stage has a stiff-ened modulus of 15 GPa prior to the second yielding at1400 MPa. The ultimate strength reached was 2.54 GPaand the maximum plastic strain achieved was 35.6%.

Figure 3 shows in situ synchrotron HEXRD mea-surements during compression of the composite. Fig-ure 3(a)–(c) show the Debye–Scherrer diffraction ringpatterns at different compressive strains. The compres-sion axis is along the vertical direction of the patterns.The patterns shown are the central part of the full pat-tern, revealing the Ti5Si3 (100) ring (the inner ring)and the B190-TiNi (001) ring (the outer ring). In Fig-ure 3(a) the diffraction intensities of B190-TiNi (001)planes in the undeformed sample show a discontinuousbut nevertheless uniform distribution due to the contri-bution of self-accommodating martensite variants. Withincreasing the strain, the diffraction intensities of theB190-TiNi (00 1) planes gradually decreased in the axialdirection and increased in the transverse direction,apparently due to the stress-induced reorientation ofthe self-accommodating martensite variants. This pro-cess can be clearly revealed using the scanned diffractionspectra in the axial direction (±5�) and the transverse

Figure 3. XRD patterns of the Ti5Si3/TiNi composite at differentcompressive strains: (a) 0%, (b) 15%, (c) 33%. (d) Evolution ofdiffraction patterns of B190-TiNi (001) in the directions perpendicularto (left) and parallel with (right) the loading direction.

D. Jiang et al. / Scripta Materialia 78–79 (2014) 53–56 55

direction (±5�), as shown in Figure 3(d). The spectra areobtained for different strains from 0% (the spectrum atthe bottom) to 35% (top). The intensity of the B190-TiNi(001) diffraction in the axial direction (left) can be seento have decreased gradually whilst that in the transversedirection (right) increased during compression, demon-strating the martensite variant reorientation process toalign the B190-TiNi (001) normal axis towards the trans-verse direction upon compression. This is consistentwith the expectation based on crystallographic analysisof the B190 martensite [30].

Figure 4 shows lattice strain measurement duringcompression, in terms of d-spacing dilatation. Fig-ure 4(a) shows the evolutions of the B190-TiNi (00 1)and Ti5Si3 (100) diffraction peaks in the axial directionduring compression. Upon initial loading, both theB190-TiNi (001) and Ti5Si3 (100) peaks shift to lowerd-spacing values, demonstrating the elastic deformationin both phases of the specimen. At higher strain levels,

Figure 4. Microscopic mechanical behavior and internal fracturing ofthe composite. (a) Two sections of diffraction patterns for B190-TiNi(111) and Ti5Si3 (100) perpendicular to the loading direction duringcompression. (b) Evolution of the d-spacing strain for Ti5Si3 (100) andB190-TiNi (111) perpendicular to the loading direction with theapplied macroscopic strain. (c) TEM bright-field image revealinginternal fractures of Ti5Si3 lamellae plates.

both peaks are broadened and their intensities are in-creased. The peak broadening implies increased inho-mogeneity of the internal strain fields, apparentlyrelated to increased defect (dislocations) density. Thegradual increase in diffraction intensity implies grainrotation towards the transverse directions of the nor-mals of the B190-TiNi (001) and Ti5Si3 (100) planes.

The lattice strains of Ti5Si3 (100) and B190-TiNi(111) in the axial direction are calculated from the dif-fraction peak positions. The evolutions of these strainswith respect to the macroscopic strain are shown in Fig-ure 4(b) (in positive values for compressive strain).Upon loading in stage I, up to 7%, the lattice strain ofTi5Si3 increased continuously while the lattice strain ofB190 remained low and practically unchanged. This indi-cates that stress on the B190 martensite during this stageof deformation was almost constant, and the increasingexternal load was mainly borne by the Ti5Si3 phase,implying gradual load transfer from TiNi to Ti5Si3. Thisstage is attributed to the stress-induced martensite vari-ant reorientation deformation of the NiTi matrix andelastic deformation of Ti5Si3 in the eutectic.

In stage II, from 7% to 13%, both the lattice strains ofTi5Si3 and B190 martensite increased. The increase in thelattice strain of B190 martensite with deformation indi-cates the completion of the variant reorientation processand thus the commencement of elastic and plastic defor-mation of the oriented martensite. Whilst the Ti5Si3 lat-tice strain continued to increase with macroscopicdeformation at the same rate, the additional load bear-ing by NiTi, as indicated by its lattice strain increase,implies a more rapid increase in external load in thisstage, as evident in Figure 2. This implies increased loadtransfer to B190-NiTi.

In stage III, from about 13%, the lattice strains ofTi5Si3 (10 0) and B190-TiNi (111) continued to increase,but at a much reduced rate relative to the macroscopicstrain. This is obviously due to the commencement ofmassive plastic deformation of the reoriented martensiteand is also attributed to the microscopic internal frac-turing of the thin and brittle Ti5Si3 lamellae, as evidentin Figure 4(c). The further moderate increase in the truestress on the composite (inset in Fig. 2) is clearly causedby the strain hardening of the B190 martensite. It shouldbe note that the maximum lattice strain reached in theTi5Si3 (100) of the composite is 2.1%. This is twice thatobserved in free-standing Ti5Si3 [31].

This study investigated the deformation behavior ofan in situ Ti5Si3/TiNi composite, in particular the loadtransfer phenomenon between a more compliant det-weening matrix and a hard enforcement phase. Themicrostructure of the composite was nanolamellarTi5Si3/TiNi eutectic in a pre-eutectic B190-TiNi matrix.The composite shows a three stage deformation behav-ior. The first stage is associated with predominantlystress-induced martensite variant reorientation of theB190-NiTi matrix and mild elastic deformation of Ti5Si3,with a very low apparent elastic modulus. The secondstage is associated with concurrent elastic deformationof Ti5Si3, and the elastic and plastic deformation ofthe oriented B190martensite. The third stage is associ-ated with the plastic deformation and strain hardeningof B190-TiNi, and the elastic deformation and fracturing

56 D. Jiang et al. / Scripta Materialia 78–79 (2014) 53–56

of Ti5Si3. The maximum lattice strain of the nanolamel-lar Ti5Si3 was 2.1%. The composite showed an ultimatestrength of 2.54 GPa and a plastic strain of 35.6%.

This work was supported by the National Nat-ural Science Foundation of China (Grant Nos.51001119, 51231008), the National 973 programs ofChina (2012CB619403), Beijing Higher EducationYoung Elite Teacher Project (YETP0686), the Key Pro-ject of Chinese Ministry of Education (313055) and theAustralian Research Council (Grant No.DP140103805). The use of the Advanced Photon Sourcewas supported by the US Department of Energy, Officeof Science and Office of Basic Energy Science, underContract No. DE-AC02-06CH11357.

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