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Materials Characterization

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Stress-induced α″ phase in a beta Ti–19Nb–1.5Mo–4Zr–8Sn alloy

Yacen Zhanga, Huiqun Liua,⁎, Danqing Yia, Bin Wanga, Yong Jianga, Tao Hua, Yu Xiaoa, Qi Yangb,Jian Wangb, Qi Gaob

a School of Materials Science and Engineering, Central South University, Changsha, Hunan 410083, PR Chinab BAOTI Group Ltd., Baoji, Shanxi 721014, PR China

A R T I C L E I N F O

Keywords:β titaniumα″ martensite packetSelf-accommodating variantsDeformationdegreeHardness

A B S T R A C T

The morphological evolution of the stress-induced α″ martensite with increases in the stretching and rollingdeformation degrees as well as the related orientation relationship were investigated in a β titanium alloy(Ti–19Nb–1.5Mo–4Zr–8Sn)·For a stretching deformation of 1%, the α″ martensite exhibited a packet mor-phology consisting of parallel lamellar structures. On the other hand, in the 2%-rolled sample, the α″ martensiteexhibited a long lath-like distribution, which was accompanied by a few small pieces. With an increase in thedeformation degree, the α″ martensite gradually exhibited a ladder-like structure consisting of long and short α″laths or ellipsoids. The existence of these self-accommodating martensitic variants decreased the elastic energy.The parallel α″ laths present within a single packet were found to have the same orientation. The interfacebetween the α″ and β phases was coherent, and their lattice constant relationships were identified as

≈ ≈″ ″2 a b cβ α α and aβ≈ aα″. Further, with the increase in the deformation degree, the hardness of the alloyfirst decreased and then increased; this was the case both the stretched as well as the rolled samples.

1. Introduction

Metastable β titanium alloys have been the subject of intensive re-search efforts because of their low elastic modulus, high strength-to-weight ratio, and shape-memory behavior [1–3]. The phases in β tita-nium alloys can be broadly classified into equilibrium and metastablephases [4]. The body-centered cubic (BCC) β phase and the hexagonalclose-packed (HCP) α phase are the equilibrium phases. Metastablephases, such as the HCP α′, orthorhombic α″, and HCP ω phases,usually form during quenching, deformation, and aging [5,6].

The α″ phase, which has an orthorhombic structure, has attractedsignificant attention. The shape-memory effect in β titanium alloys isdue to a reversible transformation between the α″ and β phases [7]. Thetransformation of α″ martensite offers the possibility of ensuring ahigher degree of ductility in β titanium alloys [8–10]. The α″ phaseforms during quenching from the β phase region or during the de-formation of the β phase through a stress-induced martensite trans-formation at room temperature [11,12]. In a previous work [13], wehad reported that the α″ phase is also formed during isothermal aging.The microstructure and lattice constants of the isothermal α″ phasewere discussed in detail. We had also elucidated the orientation re-lationship between the isothermal α″ and β phases.

It was found that stress-induced α″ martensite forms during theroom-temperature deformation of a β titanium alloy with a

molybdenum equivalency (MoE) value of less than or equal to 12 wt%[14]. For example, when the β stability MoE value is 12 in the case of aTi–12Mo alloy, the α″ phase forms, along with a complex combinationof ω and α′ phases and {332}⟨113>mechanical twins throughout theplastic straining process [14]. The crystallography of the α″ phase iswell understood, as it has been the subject of extensive experimentalanalyses. Kim et al. [15] reported that, in the case of a Ti–Nb alloy, thelattice constants and transformation strain of α″ martensite changelinearly with the Nb content. Based on an analysis of transmissionelectron microscopy (TEM) micrographs of tensile samples of a Ti–Moalloy, Zhang [12] inferred that α″ laths form heterogeneously at theboundaries owing to the dissociation of the boundary dislocations. Yaoet al. [16] reported that a stress-induced α″ phase forms at the stressplateau stage during the tensile deformation of a single-crystal Ti–Nballoy. The orientation relationship between the α″ and β phases wasalso elucidated. Xu et al. [17] reported the formation of a stress-inducedα″ phase and a (110)α″ twin substructure during rolling deformation.

To some degree, the morphology of the martensite phase in Ti alloysis similar to that in Fe-based alloys. As such, Banerjee andMukhopadhyay found that they can be depicted in the same manner[18]. In the case of the lath martensite formed in Fe-based alloys duringquenching, the austenite grains are filled with several packets of par-allel lath martensite [19]. In a few titanium alloys, the primary β grainsconsist of multiple martensite packets, which exhibit a lath

https://doi.org/10.1016/j.matchar.2018.04.019Received 27 January 2018; Received in revised form 12 April 2018; Accepted 12 April 2018

⁎ Corresponding author.E-mail address: liuhuiqun@csu.edu.cn (H. Liu).

Materials Characterization 140 (2018) 247–258

Available online 13 April 20181044-5803/ © 2018 Elsevier Inc. All rights reserved.

T

morphology. Williams et al. [20] reported the formation of an HCP lathmartensite phase during the quenching of Ti–Cu alloys, showing thatindividual packets contain groups of parallel martensite laths. Zangvilet al. [21] reported the existence of several packets of twin-related α′martensite laths in a Ti–5.3Cu alloy. Zhang et al. [22] observed aparallel acicular martensitic α″ phase in quenched Ti–3.2Mo andTi–4.5Mo alloys. In the case of stress-induced α″ martensite, the de-formation mode and degree both affect its formation and features, in-cluding its packet morphology and orientation relationship. However,the effect of the deformation mode and degree on the morphology andorientation of the stress-induced α″ martensite has not been studied inthe case of β titanium alloys.

In the present study, we continued our efforts to determine whetherα″ martensite packets and their variants are formed during thestretching and rolling deformation of a low-modulusTi–19Nb–1.5Mo–4Zr–8Sn alloy that we had developed previously. The

MoE of this alloy is 6.82. Efforts were made to characterize the evo-lution of the α″ phase under different degrees of deformation. The or-ientation and lattice constant relationships between the α″ and β phaseswere determined. The effect of the deformation degree on the hardnessof the Ti–19Nb–1.5Mo–4Zr–8Snalloywas also studied.

2. Experimental Methods

An ingot of the Ti–19Nb–1.5Mo–4Zr–8Sn alloy was prepared byvacuum-consumable arc melting using high-purity sponge Ti and Zr(99.7 wt%), a Ti–50Nb master alloy, a Ti–80Sn master alloy, and high-purity Mo. The ingot was homogenized at 1230 K for 2 h in a vacuumand then forged into a quadrate (39×230×250mm3). This was fol-lowed by annealing at 1120 K for 1 h. The alloy was then hot-rolledfrom a thickness of 39 mm to 10mm and subsequently cold-rolled to athickness of 2mm. The cold-rolled sheet was cut into

Fig. 1. Dimensions of tensile samples (unit: mm).

Fig. 2. (a) TEM bright-field image, (b) corresponding SAD pattern along [113]β and (c) TEM dark-field image showing nanosized α″ phase precipitated from βmatrixin ST-WQ sample.

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45mm×145mm×2mm and 90mm×40mm×2mm pieces. Thepieces were wrapped in an antioxidation coating before being subjectedto a solution treatment (ST), which was performed at 1173 K for 1 h in atube furnace in a argon atmosphere, and then water quenching (WQ).

Tensile samples (gage dimensions: 50mm×10mm×2mm), asshown in Fig. 1, were cut from the solution-treated45mm×145mm×2mm sheet along the rolling direction (RD).Subsequently, these samples were deformed using an MTS Landmarksystem. An extensometer was used to control the extent of tensile de-formation. Interrupted tensile tests (unloading at strains of 1%, 2.5%,and 5%) were performed at room temperature using an initial strainrate of 2.5× 10−4. In addition, rolling samples were prepared by so-lution-treated pieces with dimensions of 90mm×40mm×2mm. Therolling machine used allowed for precise control over the reduction rateduring each rolling pass. We also measured the thicknesses of thesamples after the rolling process using a vernier calipers in order toensure the deformation degree was at the desired level (2%, 5%, and8%).

Vickers hardness measurements were performed using a 200HBVS-30 microhardness tester with a load of 3 kg and a holding time of 15 s.Seven positions were tested on each sample, with the average taken asthe value for that sample.

The microstructures were examined using a Tecnai G2 F20 TEMsystem operating at 200 kV and an FEI Titan 80-300 TEM system op-erating at 300 kV. The thin foils used for the imaging process wereprepared using a twin-jet electropolisher.

Unit cells of the β, α′, and α″ phases were created using the softwareprograms Materials Studio (MS) and Adobe Illustrator (AI). First, theframe of the unit cells was obtained using the MS program. Then, the AI

program was used to draw the atoms in the respective positions.

3. Results and Discussion

3.1. Microstructure of As-quenched Alloy

Although only extinction stripes could be observed in the TEMbright-field images (Fig. 2a), a TEM dark-field micrograph (Fig. 2c),obtained from the α″ spot circled in Fig. 2b, showed that a dispersednanosized (< 5 nm) α″ phase was formed by the quenching of the βmatrix from 1173 K. This α″ phase was formed because of the presenceof insufficient amounts of β stabilizing elements, which would havehindered the decomposition of the β phase during quenching. It hasbeen reported that metastable phases such as the α′ and α″ phases canbe formed by the quenching of a high-temperature β field, provided theMoE is< 8wt% [23].

3.2. Stress-Induced α″ Phase under Stretch Deformation

TEM micrographs of the stretched samples subjected to a 1%strainare shown in Fig. 3. Several packets consisting of parallel lamellarstructures with thicknesses of approximately 40–60 nm can be observed(see Fig. 3a; packets Ι and ΙΙ). These packets are quadrilaterals withdifferent sizes, and their boundaries are not very clear. The length ofthe parallel α″ lamellae in packet ΙΙ is more than 1.5 μm, while the α″lamellae in packet Ι are much shorter. By comparing these images withthose in Fig. 2, it can be seen that the morphology of the stress-inducedα″ phase under stretching deformation is different from that of thenanosized (< 5 nm) α″ phase that forms during quenching. The

Fig. 3. Typical microstructure of 1%-strain stretching sample. (a) TEM bright-field image of α″ packets; (b) SAD pattern along [113]β taken from packet I; (c) SADpattern along [011]β taken from packet II; (d) TEM dark-field image of α″ packets obtained from dot ″(101)α in (c).

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selected-area diffraction (SAD) pattern of packet Ι was obtained alongthe[113]β zone axis (Fig. 3b). In addition to the primary reflectionsfrom the β matrix, diffuse scattering corresponding to the α″ phase isalso visible in this SAD pattern. This indicates the transformation of theβ phase into the α″ phase. From this diffraction pattern, it can beconcluded that the α″ and β phases satisfy the following orientationrelationship: ″[113]β//[112]α , (110)β//(111)α″, and ″(121)β//(110)α .Fig. 3c shows the SAD pattern of packet ΙΙ taken along the [011]β axis.Weak spots corresponding to the α″ phase can be seen at position 1/2{211}β. The basic orientation relationship between the α″ and β phaseswas identified as ″[011]β//[020]α and (011)β//(002)α″. Further, only afew bright points are present in the dark-field image (Fig. 3d), obtainedfrom the ″(101)α spot in Fig. 3c, because the volume fraction of the α″phase was not very large.

Fig. 4a shows that, with the increase in the degree of deformation inthe 2.5%-strain sample, a ladder-like feature composed of long α″ lathsas well as parallel short platelets, which were present between the longlaths, was formed. This is similar to the results of a previous studywherein small martensite units were observed between the largermartensite units in Fe–31Ni–0.02C after cooling in liquid nitrogen [24].The long laths and the short platelets meet at an angle of almost 50°. Ascan be seen from Fig. 4b, based on the SAD pattern obtained along[113]β and corresponding to area Ι in Fig. 4a, only one set of α″martensite reflections was observed. To observe the morphology of theα″ martensite phase in the 2.5%-strain alloy, a dark-field image wastaken corresponding to the reflection within the red circle shown inFig. 4b; the result is given in Fig. 4c. By comparing the dark-field image(Fig. 4c) with the bright-field image (Fig. 4a), we could confirm that thelong laths in the bright-field image were α″ martensite. These α″ lathsbecame larger than the α″ lamellae observed in the samples with astrain of 1% (Fig. 3). The thicknesses of these laths were 140–150 nm.

Note that the parallel platelets (Fig. 4a, area Ι) observed in the TEMbright-field image cannot be identified clearly in the dark-field image.This implies that these parallel short platelets are not of the α″ phase.This is because the higher-magnification TEM bright-field image of areaII shown in Fig. 4d confirmed the existence of dislocation, which areapproximately parallel to each other. In addition, they are also nearlyparallel to the platelets located between the lamellar α″ phase (Fig. 4a).The motion of these dislocations can affect the growth of the α″ mar-tensite phase [12].

Fig. 5 shows TEM images and the corresponding SAD pattern of the5%-strain sample. The ladder-shaped morphology becomes more ob-vious as the thickness of the α″ laths increases. The SAD pattern takenalong the [100]β axis and shown in Fig. 5b was obtained from area Ι inFig. 5a. Further, Figs. 5a and c show several parallel ellipsoids of the α″phase lying between as well as nearly perpendicular to the α″ laths. Oncomparing these with the parallel platelets observed in Fig. 4a, it can behypothesized that these parallel platelets are closely related to theparallel α″ phase ellipsoids. In addition, a few dislocation stripes thatwere parallel to the α″ phase ellipsoids appeared between the α″ laths.These parallel dislocation stripes were subsequently transformed intothe α″ phase [12]. The reasons the α″ martensite phase exhibited themorphology observed in Figs. 3, 4, and 5 are discussed below.

The elastic energy [25] is a key factor influencing the growth of themartensite phase. The growth behavior of martensite depends on thelocal elastic strain energy. During the primary stages of martensiteformation, the mutual interference between the martensite layers isvery low, such that they are free to grow through the crystal in alldirections. According to Eshelby's elasticity theory [25] and assumingthe martensite phase to be a flat spherical lath (Fig. 6) [26], the changein the elastic energy caused by a martensitic transformation can bedetermined as follows:

Fig. 4. Typical microstructure of 2.5%-strainstretching sample. (a) TEM bright-field image; (b)SAD pattern along [113]β taken from area Ι; (c)dark-field image taken from reflection inside redcircle in (b); (d) higher-magnification TEM bright-field image of area ΙΙ showing high-density disloca-tions. (For interpretation of the references to colourin this figure legend, the reader is referred to the webversion of this article.)

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=g A(c/r)e (1)

where ge is the elastic energy, A is a material constant, c is the half-thickness of the flat spherical lath, and r is its radius [27].

The martensite phase was observed to grow in the radial direction,

leading to a decrease in the elastic energy. However, growth in thethickness direction can lead to an increase in the elastic energy.Therefore, the martensite phase grew into slender lamellae during theinitial stages of the martensite formation process. The initial martensite

Fig. 5. Typical microstructure of 5%-strainstretching samples. (a) TEM bright-field image oftwo variants of α″ martensite; (b) SAD pattern along[100]β taken from area I; (c) dark-field image of twovariants of α″martensite taken from reflection insidered circle in (b); parallel lath-shaped variant A andparallel ellipsoidal variant B—both exhibit ladder-shaped structure. (For interpretation of the refer-ences to colour in this figure legend, the reader isreferred to the web version of this article.)

Fig. 6. Schematic of flat spherical lath [26].

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plates can even span the entire grain. Subsequently, the space availablefor the growth of the martensite plates decreases, causing the plates tobecome shorter [18]. In the present study, the results shown in Fig. 3were found to be consistent with the assumptions mentioned above.The small deformation degree can be regarded as corresponding to theinitial stage of formation of α″ martensite. In packet ΙΙ, the lengths ofthe parallel α″ lamellae were considerably greater than their thick-nesses. Based on Eq. (1), this martensite morphology is helpful for re-ducing ge. The martensite in packet II formed earlier than that in packetI, so the α″ lamellae in packet ΙΙ grew longer readily. The α″ lamellae inpacket Ι were much shorter because of the blocking resulting from themartensite in packet II [18,27].

With the increase in the martensite volume fraction, the martensitetransformation is affected by the distribution of the elastic stress causedby the interactions with the surrounding martensite phase. At this stage,self-accommodation [18] reasonably explains the formation of α″martensite in Fig. 5. The self-accommodating martensitic transforma-tion involves martensite packets, each of which contains several mar-tensite variants. The elastic energy, ge, in Eq. (1), related to thesevariants, can be partially offset, such that the elastic energy of a groupis very low [18,28]. Schroed [29] reported four self-accommodatingvariants of martensite in single crystals of Cu–Zn alloys. In addition,Sarı [28] studied self-accommodating martensites in Cu–Al–Ni alloys bymeans of SEM and TEM. Triangular self-accommodating martensitevariant morphologies have been observed in Ti–Nb, Ti–Ni, and Zr–Nballoys [30–32]. In the present study, the α″ martensites shown in Fig. 5had characteristics very similar to those corresponding to the self-ac-commodating mechanism, in that they were found to have two variantswith ladder-shaped features (Fig. 5c). On comparing Figs. 3, 4, and 5, it

can be seen that the lath-shaped variant A was long and formed beforevariant B (Fig. 5c). This martensite feature formed in the same parentphase grain to accommodate the local stress field. On the other hand,the α″ laths belonging to variant A were parallel to each other, whilethe α″ ellipsoids belonging to variant B were parallel to each other. Thismorphology is evident from the self-similarity of the assemblies ofmartensite [18] and is caused by the boundaries between the martensitevariants being nearly parallel to a habit plane of the parent phase [28].

3.3. Stress-Induced α″ Phase under Rolling Deformation

Fig. 7 shows the distribution characteristics of the α″ phase in a 2%-strain rolled sample. In contrast to the case for the parallel lamellarmartensite phase observed in Fig. 3a, the TEM bright-field image inFig. 7a suggests that the distribution of the α″ phase was not high, giventhat the degree of rolling deformation was relatively small. The TEMdark-field image in Fig. 7b shows that the α″ phase also exhibits a longlath-like structure, which is accompanied by a few small pieces of theα″ phase. The SAD pattern of the 2%-strain rolled sample, obtainedalong the [100]β zone axis and shown in Fig. 7c, indicated the fol-lowing orientation relationships: [100]β//[100]α″ and (011)β//(020)α″. Furthermore, using the reciprocal relationship obtained fromthe SAD pattern, the interplanar spacing of the α″ (020) and α″ (001)planes was measured. Two of the three lattice constants of the α″ phasewere calculated and found to be b= 0.498 nm and c=0.496 nm. Inaddition to the β matrix spots and the α″ diffraction spot, a few addi-tional reflections were also present, as indicated by the blue arrows inFig. 7c. Their existence may be ascribed to subgrains or overlappinggrains in the TEM sample.

Fig. 7. Typical microstructure of 2%-strain rolling sample. (a) TEM bright-field image; (b) TEM dark-field image; (c) SAD pattern along [100]β taken from areawithin yellow circle in (a). (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. 8 shows the TEM microstructure of the α″ martensite phase andthe corresponding SAD pattern in the 5%-strain rolled sample. Fig. 8ashows several packets of the parallel α″ phase variants. To observe themorphology and orientation relationship of the α″ phase variants, fourpackets were selected for further analysis. Two types of structures re-presenting the α″ phase were observed in this 5%-strain rolled sample.Both packet Ι and packet ΙΙ have crossed lamellar structures (Fig. 8a, c,and e). In contrast to the case for the α″ martensite induced by tensiledeformation, we did not observe this feature in the images in Figs. 3, 4,and 5. This was due to the different stress states resulting from tension

and rolling. Packets III and IV exhibited a few ellipsoids or short α″laths lying between the long lath-shaped structures (see Fig. 8a, f, h, i,and k). The SAD pattern obtained along [133]β (see Fig. 8b) includedthree sets of α″ reflections obtained from packet Ι. The orientation re-lationships between the α″ phase and the β phase was identified as[133]β//[001]α″, ″[133]β//[321]α , ″[133]β//[312]α , and

″(011)β//(111)α . Packets Ι and ΙΙ penetrated each other. The TEM dark-field image shown in Fig. 8e, which corresponds to the α″ spot circled inred in Fig. 8d, shows that the parallel α″ lamellae lie on the (011)β axisand have a mean thickness of 100 nm. Fig. 8h, k shows TEM dark-field

Fig. 8. Typical TEM microstructure of 5%-strain rolling sample, (a) bright-field image showing α″ martensite packets; (b) SAD pattern along [133]β of packet Ishowing α″ phase; (c) magnified bright-field image of packet II; (d) SAD pattern along [133]β of packet II; (e) dark-field image obtained from α″ spot circled in red in(d); (f) magnified bright-field image of packet III; (g) SAD pattern along [011]β of packet III; (h) dark-field image obtained from α″ spot circled in red in (g); (i)magnified bright-field image of packet IV; (j) SAD pattern along [113]β of packet IV; (k) TEM dark-field image obtained from α″ spot circled in red in (j). (Forinterpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

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images obtained from the α″ spots circled in red in Fig. 8g, j, respec-tively. Thus, it can be inferred that the parallel α″ lamellae in a singlepacket have variants with the same orientation. Fig. 8d, g, and j arethree different electron diffraction patterns obtained from the 5%-strain

rolled sample. The lattice parameters of the α″ phase were determinedfrom a statistical analysis of the diffraction pattern along [011]β(Fig. 8g) and were found to be a= 0.346 nm and b=0.498 nm. Inaddition, we observed a few martensite laths emanating from the grainboundary; these exhibited “mirror symmetry” along the grainboundary, as shown in Fig. 9. Zhang et al. [12] found that α″martensitelaths emanating from the low-angle grain boundaries and subgrainboundaries in a Ti–Mo alloy.

Fig. 10 shows TEM images and the corresponding SAD pattern of the8%-strain rolled sample. As can be seen from Fig. 10a, closely packedshort α″ laths are located between and nearly perpendicular to the longlath-shaped martensite phase. In contrast to the α″ phase morphologyshown in Fig. 8a, a greater number of short α″ laths appeared betweenthe long lath martensite phase, as the degree of deformation was in-creased. Since the parent phase was segmented by the martensite slice,a new martensite phase had to nucleate and grow in the parent phaseunder stress [27]. Deng et al. [27] reported that small lateral martensitepieces appear in the parent phase domain when the previous martensitesegments the parent phase into small pieces. Further, they thought thatthe radial-direction growth of the martensite phase was blocked by theother martensite slices while the change in the elastic energy caused bythe lateral growth of the martensite slice was greater than that relatedthe radial growth. Therefore, the extent of the martensite transforma-tion at this stage was small. Sarı [28] reported that each self-accom-modating plate group has parallel-sided plate variants. Figs. 7, 8, and10 show that the overall α″ martensite phase developed because of asequential process under rolling deformation. The initially formed lath-like α″ phase was long and partitioned within the β grain. Later, α″variants became more apparent and penetrated each other. Then short

Fig. 9. TEM bright-field image showing α″ lath martensite along β grainboundary in 5%-strain rolling sample.

Fig. 10. Typical TEM microstructure of 8%-strain rolling samples, (a) bright-field image showing long and short lath-like α″ martensite; (b) SAD pattern along[011]β; (c) bright-field image showing existence of dislocations in α″ lath.

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α″ laths began to appear between the long α″ laths and gradually in-creased in number. This process was very similar to the one that occursbecause of stretching deformation. In addition, Fig. 10c shows how theα″ martensite exhibits a ladder-shaped feature under rolling, similar tothat seen in Fig. 5c. The frequent occurrence of this morphology sug-gests that the formation of variant A promotes the nucleation of variantB (Fig. 5c) in its vicinity in such a manner that the stress field created bythe first aids the nucleation of the second [18]. However, there are stilla few small differences between the two. The short α″ phase locatedbetween the long α″ laths in the rolled samples was lath like, and someof these laths can be seen at one end in Fig. 10a. Meanwhile, the shortα″ phase between the long α″ laths in the stretched samples consists ofellipsoids, as shown in Fig. 5a. Note that net-like dislocations can also

be seen within the long α″ laths in the 8%-rolled sample (see Fig. 10c).These features are very similar to the dislocations observed within thelath martensites in Fe-based alloys [33].

Fig. 11a shows an HRTEM image of the 5%-strain stretched sample;the upper-left part shows the β phase, while the lower-right is the α″phase. The corresponding fast Fourier transform (FFT) patterns along[100]β (//[100]α″) are shown in Fig. 11b. Only one set of the α″ isshown. The orientation relationships between the α″ phase and the βphase were determined to be the following: [100]β//[100]α″,

″(011)β//(001)α , and (011)β//(020)α″. Fig. 11c shows an inverse FFT(IFFT) image of the interface between the α″ phase and the β-matrix, inwhich it is possible to see the coherent features. Fig. 11d shows IFFTimages of the β phase along the [100]β axis, while Fig. 11e shows the

Fig. 11. HRTEM images of 5%-strain stretching alloy, (a) HRTEM image along [100]β axis; (b) corresponding FFT pattern along [100]β//[100]α″; (c) IFFT image ofsquare area I in Fig. 11a; (d) IFFT image of square area II in Fig. 11a; (e) IFFT image of square area III in Fig. 11a.

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α″ phase along the [100]α″ axis. The interplanar spacing of the (011)plane in the β phase is 0.247 nm (Fig. 10d). Further, the lattice spacingsof the (020) and (001) planes of the α″ phase are 0.249 nm and0.496 nm, respectively (Fig. 11e). A few studies have reported the lat-tice parameters of the orthorhombic α″ phase [5,13,16]. Table 1 liststhe lattice parameters of the α″ phase as determined in the presentstudy and those reported previously in the literature. In addition, wefocused on the relationship between the lattice constants of the β andα″ phases. In the SAD pattern along [100]β, shownin Fig. 7c, diffractionspot (011)β overlaps diffraction spot (020)α″, while diffraction spot(011)β overlaps diffraction spot ″(002)α . Therefore, the interplanarspacings of the(011)β, (020)α″, and ″(002)α planes are approximatelyequal. The lattice constants of the phases were determined from theinterplanar spacings, which were measured based on the reciprocalrelationships as determined from the SAD pattern. The relationshipbetween the lattice constants of the β and α″ phases was

≈ ≈″ ″2 a b cβ α α . In the SAD pattern along [011]β in Fig. 3c, diffractionspot (200)β overlaps diffraction spot (200)α″. Similarly, we getaβ≈ aα″. In the present study, two sets of basic orientation relationshipsbetween the β and α″ phases were determined from the TEM images.

They are as follows:[100]β//[100]α″, ″(011)β//(001)α , and(011)β//(020)α″ (obtained

from Fig. 7c) and.″[011]β//[020]α , (200)β//(200)α″, and (011)β//(002)α″ (obtained

from Fig. 3c).Moreover, the structural relationship between the α′ and α″ mar-

tensites and their parent β phase can be described using the commonframework of orthohexagonal martensitic transformations [5], asshown in Fig. 12. There is an equivalent c-centered orthorhombic unitcell in the α′, α″, and β lattices. Considering the results of the analysis ofthe relationship between the lattice constants of the β and α″ phases, itis obvious that the size of the orthorhombic structure in the α″ and βphases is almost similar. This means that the Bain distortion is verysmall.

3.4. Effect of Deformation Degree on Change in Hardness

Fig. 13 shows the change in the Vickers hardness as the strain in-creases. The trend in the change in the hardness is the same in both thestretched (Fig. 13a) and the rolled samples (Fig. 13b). When the strainwas very small, the hardness was low. For the stretched samples, theaverage hardness decreased from 242 HV to 207 HV for a deformationof 1%. For the rolled samples, the average hardness decreased from292 HV to 258 HV when the deformation was 2%. This was due to theformation of the orthorhombic α″ phase and is in good agreement withthe results reported by Wan et al. [34]. It is well known that the mar-tensite phase results in strengthening in steels. However, it also de-creases the hardness in a few β titanium alloys. The Ti martensite phaseis softer and weaker than the original BCC structure [35]. Hao et al.[36] reported that the α″ phase exhibits a lower hardness in the βTi–29Nb–13Ta–4.6Zr alloy. Further, Ohmori et al. [37] reported thatthe hardness of the β phase is higher than that of α″ martensite in as-quenched specimens of the β Ti–9.87V–1.78Fe–3.20Al alloy. As thestrain increased, however, so too did the hardness. This can be ascribedto the following two reasons. One is the formation of several α″

Table 1Lattice constants of α″ in β-Ti alloys.

Alloys Phases Lattice constants (nm)

a b c

Ti-19Nb-1.5Mo-4Zr-8Sn [this work] Martensite α″ 0.346 0.498 0.496Ti-19Nb-1.5Mo-4Zr-8Sn [13] Isothermal α″ 0.352 0.494 0.464Ti-8.0Mo-3.9Nb-2.0V-3.1Al [12] Martensite α″ 0.301 0.491 0.453Ti-24Nb-4Zr-8Sn [16] Martensite α″ 0.312 0.473 0.458Ti-15Nb [15] Martensite α″ 0.306 0.495 0.466Ti-18Nb [15] Martensite α″ 0.309 0.492 0.464Ti-20Nb [15] Martensite α″ 0.313 0.487 0.464Ti-22Nb [15] Martensite α″ 0.315 0.484 0.464Ti-24Nb [15] Martensite α″ 0.318 0.481 0.464

Fig. 12. (a) Two unit cells of bcc β phase. Dashed lines indicate edges of equivalent c-centered orthorhombic unit cell. (b) Unit cell of c-centered orthorhombic α″martensite; (c) hcp lattice of α′ phase. Dashed lines indicate edges of equivalent c-centered orthorhombic unit cell [5].

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variants. Martensite laths having two distinct orientation variants canbe separated by packet boundaries. Banerjee and Mukhopadhyay re-ported that packet boundaries can effectively stop slippage and theformation of deformation twin bands [18]. In our research, we foundthat the characteristics of the martensitic variants become more ob-vious as the strain increases. Fig. 8c shows two distinct orientationvariants, which intersect each other in the 5%-strain rolled sample.Figs. 5 and 10 show that a parallel ellipsoid α″ phase formed with afurther increase in the deformation degree. This means that a greaternumber of packet boundaries contribute to the increase in the hardness.Another reason is the formation of dislocations. Dislocations can beseen in Figs. 4, 5, and 10. Note that, in this study, the hardness values ofthe deformed samples were not higher than those of the undeformedsamples.

4. Conclusions

In this study, we studied the formation of stress-induced α″ mar-tensite during the stretching and rolling deformation of a low-modulusTi–19Nb–1.5Mo–4Zr–8Sn alloy. The findings of the study can be sum-marized as follows:

(1) For small deformations, the α″ martensite phase formed exhibited apacket morphology, with a parallel lamellar structure presentwithin. With the increase in the deformation degree, the martensiticdistribution exhibited ladder-like features, and several parallel el-lipsoids or short lath-like α″ structures were present between thelong lath-like α″ phase.

(2) The morphologies and distributions of the α″ phase induced by thetensile and rolling deformations were not essentially different. Adislocation deformation mechanism was observed in the case ofboth deformation types.

(3) Parallel α″ laths within the same packet had the same orientation.The interface between the α″ and β phases was coherent, and therelationships between their lattice constants were determined to be

≈ ≈″ ″2 a b cβ α α and aβ≈ aα″.(4) The formation of the α″ phase led to a decrease in the hardness in

the case of small deformations. Further, α″ phase packet boundariesand dislocations were responsible for the increase in the hardnesswith increase in the degree of deformation.

Acknowledgments

The authors thank BAOTI Group Co. Ltd. for providing the rawmaterials used in this study as well as the use of their facilities forsample fabrication. This work was supported by the Key Projects in the

National Science and Technology (grant number 2014BAC03B05) andthe National Basic Research Program of China(subcontractnumber2014CB644001-2).

Data availability

The raw data required to reproduce these findings are available todownload from http://dx.doi.org/10.17632/bgr468jcd4.1. Theprocessed data required to reproduce these findings are available todownload from http://dx.doi.org/10.17632/hnj2jmzntw.1.

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