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Formation and toughening of metastable phases in TiZrHfAlNb medium entropy alloys
Author’s Accepted Manuscript Formation and toughening of metastable phases in TiZrHfAlNb medium entropy alloys Lu Wang, Chao Fu, Yidong Wu, Qinjia Wang, Xidong Hui, Yandong Wang www.elsevier.com/locate/msea
PII: DOI: Reference:
S0921-5093(18)31711-8 https://doi.org/10.1016/j.msea.2018.12.033 MSA37305
To appear in: Materials Science & Engineering A Received date: 25 September 2018 Revised date: 6 December 2018 Accepted date: 8 December 2018 Cite this article as: Lu Wang, Chao Fu, Yidong Wu, Qinjia Wang, Xidong Hui and Yandong Wang, Formation and toughening of metastable phases in TiZrHfAlNb medium entropy alloys, Materials Science & Engineering A, https://doi.org/10.1016/j.msea.2018.12.033 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Formation and toughening of metastable phases in TiZrHfAlNb medium entropy alloys Lu Wang, Chao Fu, Yidong Wu, Qinjia Wang, Xidong Hui*, Yandong Wang* State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing 100083, PR China [email protected] (Xidong Hui), [email protected] (Yandong Wang) *
Corresponding author
Abstract In this work, the microstructure, phase transformation, and tensile properties were investigated for Ti55-xZr20Hf15Al10Nbx (x =1, 2, 3…, 9, at. %) medium entropy alloys (MEAs). It has been found that the phase composition of as-cast MEAs changes from hexagonal α′ to orthorhombic α′′, then to metastable β, finally to stable β phase at x = 2, 4 and 9, respectively. The phase transformation of Nb4~Nb8 MEAs evolved in the sequence of β+α′′→ω+α′′+β→α+β′→β during heating process. The phase constitution and microstructure of these MEAs strongly influence the tensile properties of the MEAs. The formation of metastable α′′ and β phase in certain compositional and temperature range results in “double yielding” phenomenon and remarkable work-hardening effect. Extraordinary static toughness represented by the product of strength and elongation has achieved for MEA with 8%Nb. The deformation mechanisms for these MEAs changes with the phase composition. For the alloy with most α′′ phase, the deformation is dominated by de-twinning of martensite variants; for the alloys with most β phase, stress-induced martensitic transformation from metastable β to α′′ phase and the refinement of the newborn martensite by the dislocation has been verified. These integrated effects are considered to endow the MEAs excellent deformation ductility. Keywords: Medium entropy alloy; Metastable; Microstructure; Stress induced martensitic transformation; Deformation mechanism.
1.
Introduction
High entropy alloys (HEAs) are a new class of metallic materials developed according to the concept of the high configuration entropy by alloying equimolar or near-equimolar elements in the alloys. This kind of alloys has received considerable attentions in recent years due to their special properties such as solid solution strengthening induced by lattice distortion, sluggish diffusion, excellent cryogenic fracture toughness and high hardness for some ceramic films, etc [1-5]. According to previous study, configurational entropy in HEAs that stabilizes the solid solution phase, especially at high temperatures. Solid solutions with 1
face centered cubic (FCC), body-centered cubic (BCC) and hexagonal close packed (HCP) type of structures have been prepared in a variety of HEAs. Different structure system shows different application potential. Of the HEAs developed to date, the HEAs with BCC type structure exhibit high strength, hardness, melting temperature and phase stability, and are promising for high temperature structural components. Typical examples of BCC type refractory HEAs are TaNbMoW[6], TaNbHfZrTi[7], TaNbHfZrTiAlx[8], NbCrMo0.5Ta0.5TiZr[9] and TiZr0.5NbAlV[10] HEAs. While the HEAs with FCC type show excellent deformation ability, even at the liquid nitrogen temperature. Despite of the above merits of HEAs, one of the most issues is the strength-ductility trade-off dilemma, i.e., increasing the strength leads to decreasing ductility. Recent works have indicated that this puzzle can be solved by relaxing the restrictions of the composition design and reducing the stability of the composed phases [11]. The “metastability-engineering” strategy resulting from the metastable phase has become one of the useful routes for the exploration of the near-infinite compositional space of HEAs. The stress induced martensitic transformation (SIMT) has been found in many different HEAs systems. For example, ductile dual-phase Fe80-xMnxCo10Cr10 (x=45, 40, 35, 30 at.%)[11] HEAs with (FCC+HCP) structure has been prepared by reducing the thermal stability of high-temperature phase and transformation induced work hardening of martensite phase. The ductility of non-equiatiomic Fe40Mn40Co10Cr10 [12] HEA can be improved by the deformation twinning during room temperature tensile process. A BCC type of Ti35Zr27.5Hf27.5Nb5Ta5[13] HEA, which was designed using the ‘d-electron alloy design’ approach, displays a marked transformation-induced plasticity effect. Huang et al.[14] exploited the irreversible SIMT form BCC to HCP structure in the TiZrNbHfTax during tensile test via tailoring the stability of the constituent phases. Zhang et al. [15]revealed that in NbxHfZrTi (x = 1, 0.6, 0.4 and 0.2) HEAs, the phase composition changed from β to β+ω phases as Nb decreases to 0.4, and deformation-induced β→α″ and β→α′ transformations were found in the Nb0.4 and Nb0.2 HEAs, respectively. To achieve “metastability-engineering” strategy, it is necessary to reduce the stability of the composition phase. However, for the high entropy alloy, one of the effective methods is to cut down the configuration entropy. Recently, the multi-principal alloys have been extended to medium entropy alloys (MEAs) along the composition design strategy of HEAs according to the [3]. In this work, we report a new kind of Ti55-xZr20Hf15Al10Nbx (x=1~9 at. %) alloys designed by relaxing the configuration entropy from -xilnxi1.5 to -xilnxi1.2. Among these five elements of HEAs, Nb is the only strong β stabilizer according to the knowledge of traditional titanium alloys, so it is possible to adjust the phase composition and stability by changing the Nb content. These MEAs are endowed not only excellent plasticity but also gigantic product of strength and plastic strain. Besides the excellent mechanical properties, the MEAs designed in this work are also inexpensive and lightweight compared with those of 2
HEAs containing the same kinds of components. In this article, we focus on the compositional dependence of phase composition, phase stability and tensile properties for these MEAs. The deformation mechanism for the plasticizing and toughening is also uncovered.
2.
Experimental
In this work, Ti55-xZr20Hf15Al10Nbx (x=1~9 at. %, denoted as Nb1~Nb9, respectively, in the following context for convenience) MEAs, which have the configuration entropies (-Rxilnxi) in the range of 1.21R to 1.40R, were investigated. The MEA were prepared by arc melting the metallic elements with the purity above 99.95 wt. % inside the chamber filled with low-pressure high-purity argon atmosphere. To ensure chemical homogeneity, the ingots were flipped over and re-melted at least 5 times. After melted, the melt was drop-cast into the water-cooled cupper mold with the size of 10 mm×10 mm×45 mm. The room temperature tensile properties of the MEAs were tested using CMT4105 universal electronic tensile testing machine at a strain rate of 1×10-3 s-1 on the plate-type specimens with the gauge length of 15 mm, width of 3 mm and thickness of 2 mm. The surface of the specimens was polished down to 2000-grit SiC paper. The phase constitutes for the interested alloys were tested by X-ray diffraction (XRD) with CuKα radiation (Rigaku Kmax 2500 V) in a scanning range of 2θ=20-100 at the rate of 10 s-1. The phase transformation temperatures were measured by STA 449C Jupiter model of differential scanning calorimetry (DSC). Protective argon gas was used for purging and protection, and the flow rate of argon gas was 50 ml/min. A heating rate of 20 K/min was adopted. The microstructure of the MEAs was characterized by a Zeiss Supra55 scanning electron microscope (SEM) and a Tecnai G2 F20 200 kV transmission electron microscope (TEM). SEM specimens were first polished down to 2000-grit SiC paper, and electrochemically polished for the final surface clarification using a methanol, 2-butoxy ethanol and perchloric acid solution with a direct voltage of 25 V for 45~60 s at the temperature about 243 K. After polished, the specimens were etched using a solution of HF: HNO3: H2O=1:2:17 (in volume) at room temperature. TEM specimens were punched to Ф3 circle sheets and then ground to about 60 μm, followed by twin-jet electro-polishing using a mixed solution of glacial acetic acid: perchloric acid : ethanol = 4:8:88 (in volume) at a temperature around 243 K.
3.
Results
3.1 Compositional dependence of microstructure The phase constitution and microstructure of the as-cast MEAs were firstly characterized by X-ray diffraction and scanning electron microscopy methods. As shown in Fig.1, the phase composition transforms from hexagonal to orthorhombic and then to cubic structure as Nb content increases from 1% to 9%. In Nb1 MEA, only single hexagonal α′ phase forms. With the Nb content increasing to 2% and 3%, the (100)α′, (002)α′ and (101)α′ peaks significantly 3
broaden. Besides, the (100)α′ peak of Nb3 MEA shifts to low 2θ angle side, and (101)α′ peak shifts to high 2θ angle side, meaning that new phase has formed in this alloy. This conjecture can be further deduced from the peak at 2θ=71°. As Nb content is increased to 4% (Fig. 1(b)), the (100)α′ peak splits into (100)α′′ and (020)α′′, (101)α′ peak splits into (111)α′′ and (021)α′′, while the (002)α′ peak turns into (002)α′′, in addition, the (110)β appears near (002)α′′. These results mean that α′′ phase with orthorhombic structure become the major phase in the MEA with 4% Nb. In Nb5 and Nb6, α′′ phase is much enough so that the (111)α′′ and (021)α′′ peaks can be detected. For Nb7 to Nb9, only single β phase can be observed, without any peaks of α′ or α′′ phase. In all the interested alloys, the XRD patterns only show the existence of α′, α′′ and β phases, and no ω phase is detected, which exists in many metastable Ti alloys. Similar results have been reported in Ti-Zr-Nb-Al[16], Ti-Ta-Al[17], and Ti-V-Al[18] alloys, indicating that the addition of Al suppresses the formation of ω phase. To investigate and validate the microstructural feature with increasing the content of Nb directly, these MEAs were examined by SEM method. As shown in Fig. 2, both the Nb1 and Nb3 MEAs are composed of fine basket-weave and Widmanstatten plates, which commonly exist in near-α titanium alloys [19, 20]. The microstructure of Nb4 alloy contains short basket-weave type of orthorhombic α′′ phase inside the grains of cubic β phase, and no basket-weave lamellar microstructure formed nearby the grain boundary. When the Nb content increases to 5%, the α′′ phase takes the shape of small particles inside the equiaxed grains (in Fig. 2(e)). As the Nb content further increases to 8% (Fig. 2(f)), only equiaxed grains with obvious homogeneous dendritic structure can be observed. TEM investigation was conducted to further reveal the effect of Nb on the microstructure evolution of these MEAs. The refined structures and the selected-area electron diffraction patterns (SADPs) of the Nb1, Nb3, Nb5 and Nb7 MEAs are shown in Fig.3. It is seen that Nb1 MEA consists of different kinds of lathes, from less than 100 nm to 500 nm in width. The SADP confirms that the lath is of hexagonal lattice, i.e., ′ phase, which is in accordance with the XRD result. The Nb3 MEA exhibits two different morphologies. As shown in Fig. 3(b), the large sheet structure with flat and parallel interface on the bottom part is calibrated to be the α′′ phase, while short lamellar structure at the upper part is the α′ phase. In Nb5 MEA, the β and α′′ phase coexist, the α′′ phase takes the shape of plate with about 200 nm in thickness, and has the orientation relationship of [011]β//[001]α′′ with β phase, which has been also found in traditional Ti alloys[21]. In Nb7 MEA, only β phase is observed in the high magnification TEM image, as shown in Fig. 3(d). 3.2
Phase transformation during heating and cooling process
The temperature dependence of phase composition for the interested MEAs was investigated by DSC method in this work. As shown in Fig.4, there are two or three peaks at the heating curves and one peak at the cooling curves (all the peaks marked by sequence 4
numbers). The start and finish temperatures of the phase transformation represented by each peak are listed in Tab.1. In combination with the phase constitute investigated by the XRD and TEM of the MEAs, the Peak1 during the cooling process can be determined to be the transformation from β to α′ and/or α′′. For Nb1 MEA with single α′ phase, as shown in Fig. 4(b), there is only one extreme point on the Peak1, namely the transformation from β to α′. For Nb2 and Nb3 shown in Fig.6(c), the Peak1 displays two extreme points, suggesting that the β phase transforms to α′ and α′′. When cooling these samples from high temperature, a portion of α′ phase with less β-stabilizing element precipitated firstly, then the remnant high temperature β phase containing more β stabilizing elements transfers into orthorhombic α′′ phase. So the extreme point at higher temperature is β→α′, and the one at lower temperature is β→α′′. For Nb4 MEA, the extreme point representing the transformation from β to α′ disappears. As the Nb content further goes up to 7%, the peak during cooling disappears suggesting that no α′′ phase exists with Nb up to 7%. As listed in Tab.1, the Peak1 shifts toward low temperature side with the Nb increasing, meaning that the high temperature β phase becomes more and more stable. The phase transformations reflected by the DSC curves during heating seem to be more complicated than those appeared during the cooling process. For Nb1~Nb3 MEAs, only Peak4 is obvious, which represents the transformations of α′→β (Nb1) and α′/α′′→β (Nb2, Nb3) in the temperature range of 600 ~810 °C. For Nb4~Nb7, three peaks can be obviously found in the range of 130~300 °C, 400~480 °C and 600~810 °C. In order to reveal the structural evolution clearly, we selected Nb6 MEA to have heat treatments at 300 °C and 600 °C for 6h (denoted as Nb6-300 and Nb6-600), respectively. Fig. 5 shows the XRD patterns, DSC curves, TEM and SADP for Nb6-300 and Nb6-600 MEAs. It is seen that the XRD peaks of Nb6-300 are similar to those of as-cast Nb6 MEA. On the DSC curve shown in Fig. 5(b), there is an exothermic peak in the temperature range from 130 °C to 300 °C. The SADP shown in Fig.5(c) indicates that ω phase exists in Nb6-300 sample, so that the phase transformation is β→ω. The precipitation of ω phase is common in traditional titanium alloys such as Ti-Nb[21], Ti-Mo[22], Ti-Nb-Mo-Sn[23] alloys during low temperature annealing. The Nb6-600 MEA is mainly composed of ′ with hcp structure, which can be determined by all the peaks on the XRD pattern except the one located at 2θ=54º. As shown in Fig. 5(d), the bright field TEM image and the corresponding SADP demonstrate that the microstructure is composed of α′ phase and β phase with bcc structure. Based on the above results, the transformations in Nb6 MEA represented by the three peaks during heating can be determined as Peak 2: β→ω; Peak 3: β+α′′+ω→α+β′; Peak 4: α+β′→β. For Nb9 MEA, there are similar phase transformations with those of Nb6. As shown in Fig. 6(a), the peaks of hexagonal ω phase is obvious on the XRD pattern of Nb9 MEA after heated at 300 °C for 6h. It is found that the Nb content has no effect on the transformation temperature, but elevates the volume fraction of the ω phase. Besides, only Peak2 and Peak4 5
appeared on the heating process, which indicates the β→ω and ω→β transformation, respectively. That is to say, the β phase is too stable to decompose into α′ phase at medium temperature during DSC test in this alloy, only thermal ω phase separated out at low temperature. The increase of Nb content seems to have no obvious influence on the start and finish transformation temperature of Peak1, but makes Peak4 move towards low temperature side. In conclusion, the Nb content promote the stability of β phase. 3.3
Phase composition determined tensile properties
From the above section, it has been found that the phase constitution and microstructure of these MEAs are strongly influenced by the addition of Nb. How about the effect of phase composition on the tensile properties of these MEAs? In this respect, typical room-temperature true stress-strain tensile curves of these TiZrHfAlNb alloys are plotted in Fig. 7(a). The yield strength, tensile strength and fracture elongation as a function of Nb content are displayed in Fig. 7(b). From the tensile curves, it can be seen that with the Nb element content lower than 3%, the MEAs exhibit the yield strength above 1200 MPa, the tensile strengths of ~1500 MPa, and the elongations in the range from 3% to 4.5%. However, the yielding strength of Nb4 abruptly decreased to 227 MPa from about 1200 MPa for Nb3 because Nb content is increased by 1%, as shown in Fig. 4(b). The alloys show “double yielding” phenomenon and remarkable work-hardening effect with the Nb content in the range of 4%~8%. For these MEAs, the first yielding strength increases from 227 MP to 735 MPa, the highest tensile strength and increases reaches 1211 MPa for Nb4 and Nb7 MEA, respectively. It is noteworthy that for Nb8 alloy, after yielding at 735 MPa, the plastic deformation continues until 27%, and the “double yielding” phenomenon becomes slightly during this period. Therefore, the MEAs developed in this work exhibit excellent work hardening effect and gigantic product of strength and elongation. Nb4 MEA shows an increase of about 1000 MPa from first yielding stress to fracture strength. And the product of strength and elongation of Nb8 MEA reaches above 30 GPa%, which is comparable to the level of TRIP steel. When the Nb content is increased up to 9%, the “double yielding” together with the work-hardening effect disappears. The SEM images of the fracture surfaces for these TiZrHfAlNb MEAs are shown in Fig. 8. Corresponding to the tensile curves shown in Fig. 7(a), the fractures of Nb1 and Nb3 MEAs with relatively low elongation exhibit pronounced intergranular and transgranular mixed fracture mode. The Nb4~Nb8 MEAs have obvious transgranular fracture characteristics with dimples in different sizes. Among them, the dimples of Nb8 alloy are apparently fewer in amount but larger in size than those of other MEAs, indicating that the deformability of this MEA is the best. The above results of tensile experiments illustrate that the variation of phase composition and microstructure indeed result in mutation of mechanical properties. With the increase of 6
Nb content, the MEAs become more plastic and toughening although the strength decreases. Moreover, the deformation behavior changes from single yielding mode to double yielding mode. The intrinsic mechanisms for these phenomena are discussed in the following sections.
4.
Discussion
4.1
Cooperative effect of the components on the formation of metastable phase In traditional titanium alloys, Nb is a β stabilizing element, and used to promote the
formation of β phase. The α′′ and phase existing in the dual phase region are considered to be metastable[24]. In these TiZrHfAlNb MEAs, it has also been found that Nb is a stronger β-stabilizer to adjust the phase composition. Although the Nb content is insufficient to ensure the formation of β or α′′ phase for Nb1 MEA, so much alloying elements in the α′ phase finally results in high solid solution strengthening effect. With Nb content increased to 2%, the α′ phase is supersaturated, the metastable α′′precipitated out. When the content of Nb increased up to 4%, the metastable α′′ phase becomes dominant. In this MEA a little amount of β phase appears. The volume fraction of β phase goes up with the further increasing of Nb. When Nb content is going up to 9%, the β phase become stable. These results are similar to those for Ti-Nb and Ti-Zr-Nb alloys. For Ti-Nb alloys, the critical contents of Nb for phase transition from hcp α′ phase to orthorhombic α′′ phase and then to bcc β phase are 14% and 27%, respectively[25]. For Ti-30Zr-xNb alloys, however, the corresponding critical values decreased to 7% and 13% [26]. It has been reported that the addition of Zr is able to decrease the martensitic transformation start temperature (Ms) in low Zr-containing β-type alloys, such as Ti-Cr-Zr[27], Ti-Ta-Zr[28], Ti-Nb-Zr[29], Ti-Nb-Ta-Zr[30] and Ti-Nb-Sn-Zr[31], so Zr is considered as a β-stabilizing element in these alloys. And the β-stabilizing effect of Zr appears to be influenced largely by other alloying elements in these alloys, that is to say, the β-stabilizing effect of Zr increases with the increasing content of the β-stabilizing elements Nb, vice versa, the addition of Zr promotes the β stabilizing effect of Nb. The reason may be attributed to that the bonding energy of Ti-Zr-Nb (4.97 eV atom-1) alloy is higher than those of binary Ti-Zr (4.80 eV atom-1) and Ti-Nb (4.80 eV atom-1) alloys[26]. As a member in the same family in the period table of elements, Hf may have the same effect. Therefore, the β-stabilizing effects of Zr, Hf and Nb are promoted by each other. Besides, as reported in Ti-20Zr-10Nb-xAl (x=1, 2, 3, 4) alloys [16], the reverse martensitic transformation start (As) and finish (Af) temperatures decrease with increasing Al content at a rate of approximately 10K per Al atom, and a similar suppression effect of the transformation temperature has also been reported in Ti-24Nb-Al[32] alloy. In these TiZrHfAlNb MEAs, Al could also be considered as a β-stabilizing element. The cooperative addition of Zr, Hf, and Al in these Ti-rich MEAs lowers the critical content of Nb for the formation of metastable α′′ and stable β phase. Therefore, it is seen that through appropriate controlling of Nb addition, the metastable microstructure that contains the 7
composite of α′′ and β phase can be achieved in these MEAs. 4.2
Plasticizing mechanism by stress induced martensitic transformation
It has been confirmed that Nb1~Nb3 MEAs have high strength with low rupture elongation, which is due to the intrinsic properties of the highly alloyed hexagonal phase and supersaturated solid solution strengthening, as shown in Fig.7. The elongation decreases in this compositional range because dislocation movement is hindered at the interface of α′′ and α′. As the Nb content goes up to 4%, “double-yielding” effect appears and holds until the Nb content reaches up to 8%. For the traditional shape memory Ti-alloys such as Ti-Nb-(Zr, Sn, Mo, Al) [21, 23, 29, 33], the stress platform after yielding event corresponds to the reorientation of the twinned martensite variants or SIMT from β to αʺ phase. When the alloy with insufficient β stabilizers cooling down from high temperature, the transformation of β phase into αʺ leads to self-accommodating microstructure in order to minimize the deformation of ex-β grain during transformation. As a consequence, in a given ex-β grain, variants of αʺ martensite are auto-organized by groups in twinning relationship. When a stress is applied, the first activated deformation mechanism is the growth of variants that are favorable oriented to the detriment of others, the self-accommodating microstructure is thus reoriented. However, the metastable β phase retains as the alloy with sufficient β stabilizers cools down. The material may deform by SIMT from metastable β to αʺ phase under applied stress. As for the current MEAs, the deformation mechanism changes with the phase composition and phase stability. For Nb4 MEA, microstructures at different strains were observed by TEM observations. Fig.9 shows the martensitic microstructure which is composed of two different morphologies in Nb4 MEA prior to deformation. One is the martensite laths with the width of several microns, which organize themselves in parallel with inner martensite in them, as shown in Fig.9 (a) and (b). The corresponding SADPs in Fig.9 (c) shows that the inner martensite has orientation relationship [011]αʺ1//[110]αʺ2 with the matrix, which means that they do not share a twinning plane. On the other hand, the {111} type I twinning with [110]αʺ2M// [-110]αʺ2T relationship are discovered between the adjacent martensite laths. The other morphology of martensite phase is in triangle shape, as shown in Fig.9 (d), which are the (011) compound twinning martensites with [100]α″M//[-100]α″T. Fig.10 shows the microstructure of Nb4 MEA subjected to a tensile strain of 3%. The inner martensite was consumed by the matrix compared with torsion. New {111} type I of twins with orientation relationship [011]αʺM// [0-1-1]αʺT are observed. Meanwhile, the (011) compound twinning disappears. Fig.11 is the microstructure after fracture. It can be seen that the second complex twinning microstructure has disappeared, and all the plates organized themselves in parallel, which means that Nb4 was deformed by reorientation of the twinned α″ martensite. 8
To uncover the deformation mechanism of MEAs alloys with more metastable β phase, unloading experiments at 2% and 7% pre-strain of Nb5 MEAs were conducted. From the relative intensities of peaks of the XRD pattern as shown in Fig.12, the content of α′′ increases with the plastic strain. The optical micrographs of fractured Nb5 and Nb8 MEAs as shown in Fig.13. It is seen that the martensitic transformation in Nb5 MEA is hindered by the original α′′. In contrast, the induced martensitic α′′ plates in Nb8 MEA almost grow in the same direction to intersect the whole grain at low strain, and then plates in another direction appeared at high strain. These results illustrate that the SIMT from metastable β to α′′ takes place during loading. The SIMT process of a mixture of β and α″ phases in Ti-33Nb-4Sn[34] alloy was investigated by the in situ synchrotron X-ray diffraction. At the initial deformation stage, the pre-existing α′′ and β phase concurrently experienced elastic deformation. After the first yielding, the SIMT was activated and then it continued during the whole deformation. Especially, the SIMT processed intensely during the stress plateau. Meanwhile, the reorientation of pre-existing α′′ variants occurred at the end of the stress platform. After the stress plateau, the re-oriented and SIMT α′′ variants deformed elastically together with the elastic deformation of residual β phase, leading to a long work hardening stage. Beyond the second yielding period, the complex mixture phase deformed plastically. The same situation happened in the Nb5 and Nb6 MEAs, the critical stress for SIMT at the first stage increases with Nb content due to the enhancement of the β phase stability. Meanwhile, the elongation increases with the Nb content. Specifically, Nb8 shows a long-term work hardening stage until 27% with slight double yielding, indicating that the suppression of original α′′ martensite by Nb content ensures the alloy better deformation ability. Fig. 14 shows the microstructures of Nb8 MEA with 2% and 10% of pre-strain deformation, respectively. After 2% deformation, the SIMT has happened to generate a small amount of martensite in about 100 nm width (Fig. 14(a)). With the deformation increasing to 10%, more martensite plates have appeared. In the β phase between two martensite plates, a large number of dislocations piled up on the interface. It could be seen that some dislocations has cut into the new-born martensite, which can be observed by the dark field TEM image (Fig. 14(b)), and the martensite is divided into several sections. The SIMT results in the intensive strain-hardening effect by dynamic strain-stress partitioning between the β and α′′ phase, and the dislocation movement intensifies the partitioning effect inside the α′′ phase. The partitioning promotes the plastic deformation inside of grains, which can effectively suppress early cracking. Besides, large amount of interfaces are introduced due to the partitioning effect, which then restrict the propagation and movement of dislocation lines. The suppression of cracks and dynamics reduction of mean free path for the dislocations give rise to the combination of the initial work hardening and ductility. Therefore, it is seen that the superior combination of strength and deformation capacity has been achieved by the synergy 9
of SIMT and dislocation movement.
5.
Conclusion
In this research, the microstructure, phase stability and tensile properties of a series of Ti55- Ti55-xZr20Hf15Al10Nbx (x=1~9, at. %, named Nbx) medium entropy alloys were investigated. Based on these results, the conclusion can be drawn as follows: 1. With the increase of β stabilizing element Nb, the phase composition of as-cast MEAs changes from hexagonal α′ to orthorhombic α′′, then to metastable β, finally to stable β phase at 2%, 4% and 9%, respectively. 2. The composition has strong effect on the phase evolution during heating and cooling process. Nb1 MEA exhibits a reversible phase transformation between α′ and β phase. While for Nb2 and Nb3 MEAs with α′′ and α′ phase, the reversible transformation turn into α′′/α′ ↔β. For Nb4~Nb8 MEAs, the phase evolution during heating process is in the sequence of β+α′′→ω+α′′+β→α+β′→β, meaning that the formation of metastable α′′ and β phase in certain compositional and temperature range. For Nb9 MEAs, there is only one ω→β transformation during heating process. 3. The phase constitution and microstructure of these MEAs show intense impact on their tensile properties. For Nb1~Nb3, the MEAs exhibit high yield strength and fracture strengths, and certain elongations. As the Nb content increased from 4% to 8%, the alloys show “double yielding” phenomenon and remarkable work-hardening effect. Nb8 MEA have excellent the product of strength and elongation above 30 GPa%. 4. The deformation mechanisms for these MEAs changes with the phase composition. The
high strength for Nb1~Nb3 alloys attributed to the dislocation slip in hexagonal and orthorhombic structures. In Nb4 MEAs with most α′′ phase, the deformation dominated by the de-twinning and reorientation of martensite variants. In Nb5~Nb8 MEAs with most metastable β phase, the SIMT from metastable β to α′′ phase under external force has been confirmed. Besides refinement of the newborn martensite by the dislocation has also been verified. The integrated effects of SIMT and dislocation movement endow the MEAs excellent deformation capacity.
Acknowledgments This work was financially supported by The National Key Basic Research Program (2016YFB0701402), National Natural Science Foundation of China (Nos. 51571016 and 51531001). Key Laboratory of Research on Hydraulic and Hydro-Power Equipment Surface 10
Engineering Technology of Zhejiang Province (2017SLKL003).
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1114-1117. [18] Z.Y. Yang, X.H. Zheng, W. Cai, Martensitic transformation and shape memory effect of Ti–V–Al lightweight high-temperature shape memory alloys, Scripta Materialia 99 (2015) 97-100. [19] X. Wang, L. Wang, Q.J. Wang, Y.D. Wu, J.J. Si, X.D. Hui, Enhanced mechanical properties and structure stability induced by Si in Ti–8.5Al–1.5Mo alloys, Materials Science and Engineering: A 676 (2016) 304-311. [20] X.J. Jiang, X.Y. Wang, Z.H. Feng, C.Q. Xia, C.L. Tan, S.X. Liang, X.Y. Zhang, M.Z. Ma, R.P. Liu, Effect of rolling temperature on microstructure and mechanical properties of a TiZrAl alloy, Materials Science and Engineering: A 635 (2015) 36-42. [21] H.Y. Kim, Y. Ikehara, J.I. Kim, H. Hosoda, S. Miyazaki, Martensitic transformation, shape memory effect and superelasticity of Ti–Nb binary alloys, Acta Materialia 54(9) (2006) 2419-2429. [22] F. Sun, J.Y. Zhang, M. Marteleur, T. Gloriant, P. Vermaut, D. Laillé, P. Castany, C. Curfs, P.J. Jacques, F. Prima, Investigation of early stage deformation mechanisms in a metastable β titanium alloy showing combined twinning-induced plasticity and transformation-induced plasticity effects, Acta Materialia 61(17) (2013) 6406-6417. [23] D.C. Zhang, S. Yang, M. Wei, Y.F. Mao, C.G. Tan, J.G. Lin, Effect of Sn addition on the microstructure and superelasticity in Ti-Nb-Mo-Sn alloys, Journal of the mechanical behavior of biomedical materials 13 (2012) 156-65. [24] D. Banerjee, J.C. Williams, Perspectives on Titanium Science and Technology, Acta Materialia 61(3) (2013) 844-879. [25] M. Abdel-Hady, K. Hinoshita, M. Morinaga, General approach to phase stability and elastic properties of β-type Ti-alloys using electronic parameters, Scripta Materialia 55(5) (2006) 477-480. [26] M. Abdel-Hady, H. Fuwa, K. Hinoshita, H. Kimura, Y. Shinzato, M. Morinaga, Phase stability change with Zr content in β-type Ti–Nb alloys, Scripta Materialia 57(11) (2007) 1000-1003. [27] S. Ishiyama, S. Hanada, O. Izumi, Effect of Zr, Sn and Al additions on deformation mode and beta phase stability of metastable beta Ti alloys, ISIJ International 31(8) (1991) 807-813. [28] M. Ikeda, S.-y. Komatsu, Y. Nakamura, Effects of Sn and Zr additions on phase constitution and aging behavior of Ti-50 mass% Ta alloys quenched from β single phase region, Materials Transactions 45(4) (2004) 1106-1112. [29] J.I. Kim, H.Y. Kim, T. Inamura, H. Hosoda, S. Miyazaki, Shape memory characteristics of Ti– 22Nb–(2–8)Zr(at.%) biomedical alloys, Materials Science and Engineering: A 403(1-2) (2005) 334-339. [30] X. Tang, T. Ahmed, H. Rack, Phase transformations in Ti-Nb-Ta and Ti-Nb-Ta-Zr alloys, Journal of Materials Science 35(7) (2000) 1805-1811. [31] Y.L. Hao, S.J. Li, S.Y. Sun, R. Yang, Effect of Zr and Sn on Young's modulus and superelasticity of Ti–Nb-based alloys, Materials Science and Engineering: A 441(1-2) (2006) 112-118. [32] M.U. Farooq, F.A. Khalid, H. Zaigham, I.H. Abidi, Superelastic behaviour of Ti–Nb–Al ternary shape memory alloys for biomedical applications, Materials Letters 121 (2014) 58-61. [33] X.J. Jiang, Y.K. Zhou, Z.H. Feng, C.Q. Xia, C.L. Tan, S.X. Liang, X.Y. Zhang, M.Z. Ma, R.P. Liu, Influence of Zr content on β-phase stability in α-type Ti–Al alloys, Materials Science and Engineering: A 639 (2015) 407-411. [34] S. Guo, Y. Shang, J. Zhang, Q. Meng, X. Cheng, X. Zhao, In situ synchrotron X-ray diffraction study of deformation behaviour of a metastable β-type Ti-33Nb-4Sn alloy, Materials Science and Engineering: A 692 (2017) 81-89. 12
(a)
(b)
Fig.1 XRD patterns of as-cast TiZrHfAlNb MEAs, indicating a phase transion αα′′β with the increase of the content of Nb.
13
(a)
(b)
(c)
(d)
(e)
(f)
Fig.2 SEM micrographs of as-cast TiZrHfAlNb MEAs. (a) Nb1; (b) Nb3; (c)and (d) Nb4, (e) Nb5, and (f) Nb8.
14
(a)
(b)
(c)
(d)
Fig.3 TEM bright field images and corresponding SADPs of TiZrHfAlNb MEAs. (a) Nb1, (b) Nb3, (c) Nb5, and (d) Nb7.
15
(a)
(b)
(c)
Fig. 4 DSC curves of as-cast TiZrHfAlNb MEAs. (a) DSC curves of nine HEAs indicating the heating and cooling process with the peaks marked by series numbers, (b) and (c) the curves of Nb1 and Nb3, respectively.
16
(a)
(c)
(a)
(b)
(b)
(d)
Fig.5 Phase transition of Nb6 MEA during heating treatment. (a) XRD patterns heated at 300 °C and 600 °C for 6h; (b) DSC curves with all the peaks marked with corresponding phase transition; (c) and (d) TEM bright field image and corresponding SADPs treated at 300 °C and 600 °C, respectively.
17
(a)
(b)
Fig. 6 (a) XRD patterns of Nb9 MEA at as-cast and treated state at 300 °C for 6h, respectively; (b) DSC curves of Nb9 MEA with the peaks marked to indicate the phase transitions.
18
(a)
(b)
Fig.7 Room temperature tensile properties of as-cast TiZrHfAlNb MEAs. (a) The true stress-strain curves; (b) Dependence of tensile strength, yield strength and elongation as a function of Nb content.
19
(b)
(a)
(a)
(c)
(d)
(e)
(f)
Fig. 8 Fracture morphologies of as-cast TiZrHfAlNb MEAs. (a) Nb1; (b) Nb3; (c) Nb4; (d) Nb5; (e) Nb6; (f) Nb8.
20
(a)
(b)
(c)
(d)
Fig.9 The martensitic microstructure in Nb4 MEAs prior to deformation. (a) and (b) the martensite variants with laths structure; (c) the SADPs of the corresponding area in (b); (d) martensite variants with triangle shape and the corresponding SADPs .
21
(a)
(b)
(c)
(d)
Fig.10 the microstructure of Nb4 MEAs subjected to a tensile strain of 3%. (b) and (d) are the corresponding SADPs of the areas marked by red circles in (a) and (c), respectively.
22
(a)
(b)
(c)
(d)
Fig.11 The microstructure of Nb4 MEAs after fracture. (c) and (d) are the corresponding SADPs of the marked area in (b).
23
Fig.12 Correlation between microstrucrue evolution and stress-strain curve of Nb5 MEAs.
24
(a)
(d)
Fig.13 Optical micrographs of Nb5 (a) and Nb8 (b) MEAs after fracture.
25
(a)
(b)
Fig. 14 TEM bright (a) and dark (b) field images of as-cast Nb8 MEAs at ε=0.1.
26
Table 1. The start and finish temperatures for the transformation presented by the peaks on the DSC curves for the Nbx MEAs.
Transformation temperature during heating /°C Peak 2
Peak 3
Transformation temperature during cooling / °C
Peak 4
Peak 1
Nb1
774
846
760
689
Nb2
750
821
733
676
676
633
Nb3
745
806
685
645
645
604
Nb4
704
785
646
575
Nb5
123
283
410
451
647
764
582
483
Nb6
125
290
422
458
633
737
505
405
Nb7
126
293
427
475
621
724
476
375
Nb8
133
303
593
714
Nb9
137
306
591
689
27
PII: DOI: Reference:
S0921-5093(18)31711-8 https://doi.org/10.1016/j.msea.2018.12.033 MSA37305
To appear in: Materials Science & Engineering A Received date: 25 September 2018 Revised date: 6 December 2018 Accepted date: 8 December 2018 Cite this article as: Lu Wang, Chao Fu, Yidong Wu, Qinjia Wang, Xidong Hui and Yandong Wang, Formation and toughening of metastable phases in TiZrHfAlNb medium entropy alloys, Materials Science & Engineering A, https://doi.org/10.1016/j.msea.2018.12.033 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Formation and toughening of metastable phases in TiZrHfAlNb medium entropy alloys Lu Wang, Chao Fu, Yidong Wu, Qinjia Wang, Xidong Hui*, Yandong Wang* State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing 100083, PR China [email protected] (Xidong Hui), [email protected] (Yandong Wang) *
Corresponding author
Abstract In this work, the microstructure, phase transformation, and tensile properties were investigated for Ti55-xZr20Hf15Al10Nbx (x =1, 2, 3…, 9, at. %) medium entropy alloys (MEAs). It has been found that the phase composition of as-cast MEAs changes from hexagonal α′ to orthorhombic α′′, then to metastable β, finally to stable β phase at x = 2, 4 and 9, respectively. The phase transformation of Nb4~Nb8 MEAs evolved in the sequence of β+α′′→ω+α′′+β→α+β′→β during heating process. The phase constitution and microstructure of these MEAs strongly influence the tensile properties of the MEAs. The formation of metastable α′′ and β phase in certain compositional and temperature range results in “double yielding” phenomenon and remarkable work-hardening effect. Extraordinary static toughness represented by the product of strength and elongation has achieved for MEA with 8%Nb. The deformation mechanisms for these MEAs changes with the phase composition. For the alloy with most α′′ phase, the deformation is dominated by de-twinning of martensite variants; for the alloys with most β phase, stress-induced martensitic transformation from metastable β to α′′ phase and the refinement of the newborn martensite by the dislocation has been verified. These integrated effects are considered to endow the MEAs excellent deformation ductility. Keywords: Medium entropy alloy; Metastable; Microstructure; Stress induced martensitic transformation; Deformation mechanism.
1.
Introduction
High entropy alloys (HEAs) are a new class of metallic materials developed according to the concept of the high configuration entropy by alloying equimolar or near-equimolar elements in the alloys. This kind of alloys has received considerable attentions in recent years due to their special properties such as solid solution strengthening induced by lattice distortion, sluggish diffusion, excellent cryogenic fracture toughness and high hardness for some ceramic films, etc [1-5]. According to previous study, configurational entropy in HEAs that stabilizes the solid solution phase, especially at high temperatures. Solid solutions with 1
face centered cubic (FCC), body-centered cubic (BCC) and hexagonal close packed (HCP) type of structures have been prepared in a variety of HEAs. Different structure system shows different application potential. Of the HEAs developed to date, the HEAs with BCC type structure exhibit high strength, hardness, melting temperature and phase stability, and are promising for high temperature structural components. Typical examples of BCC type refractory HEAs are TaNbMoW[6], TaNbHfZrTi[7], TaNbHfZrTiAlx[8], NbCrMo0.5Ta0.5TiZr[9] and TiZr0.5NbAlV[10] HEAs. While the HEAs with FCC type show excellent deformation ability, even at the liquid nitrogen temperature. Despite of the above merits of HEAs, one of the most issues is the strength-ductility trade-off dilemma, i.e., increasing the strength leads to decreasing ductility. Recent works have indicated that this puzzle can be solved by relaxing the restrictions of the composition design and reducing the stability of the composed phases [11]. The “metastability-engineering” strategy resulting from the metastable phase has become one of the useful routes for the exploration of the near-infinite compositional space of HEAs. The stress induced martensitic transformation (SIMT) has been found in many different HEAs systems. For example, ductile dual-phase Fe80-xMnxCo10Cr10 (x=45, 40, 35, 30 at.%)[11] HEAs with (FCC+HCP) structure has been prepared by reducing the thermal stability of high-temperature phase and transformation induced work hardening of martensite phase. The ductility of non-equiatiomic Fe40Mn40Co10Cr10 [12] HEA can be improved by the deformation twinning during room temperature tensile process. A BCC type of Ti35Zr27.5Hf27.5Nb5Ta5[13] HEA, which was designed using the ‘d-electron alloy design’ approach, displays a marked transformation-induced plasticity effect. Huang et al.[14] exploited the irreversible SIMT form BCC to HCP structure in the TiZrNbHfTax during tensile test via tailoring the stability of the constituent phases. Zhang et al. [15]revealed that in NbxHfZrTi (x = 1, 0.6, 0.4 and 0.2) HEAs, the phase composition changed from β to β+ω phases as Nb decreases to 0.4, and deformation-induced β→α″ and β→α′ transformations were found in the Nb0.4 and Nb0.2 HEAs, respectively. To achieve “metastability-engineering” strategy, it is necessary to reduce the stability of the composition phase. However, for the high entropy alloy, one of the effective methods is to cut down the configuration entropy. Recently, the multi-principal alloys have been extended to medium entropy alloys (MEAs) along the composition design strategy of HEAs according to the [3]. In this work, we report a new kind of Ti55-xZr20Hf15Al10Nbx (x=1~9 at. %) alloys designed by relaxing the configuration entropy from -xilnxi1.5 to -xilnxi1.2. Among these five elements of HEAs, Nb is the only strong β stabilizer according to the knowledge of traditional titanium alloys, so it is possible to adjust the phase composition and stability by changing the Nb content. These MEAs are endowed not only excellent plasticity but also gigantic product of strength and plastic strain. Besides the excellent mechanical properties, the MEAs designed in this work are also inexpensive and lightweight compared with those of 2
HEAs containing the same kinds of components. In this article, we focus on the compositional dependence of phase composition, phase stability and tensile properties for these MEAs. The deformation mechanism for the plasticizing and toughening is also uncovered.
2.
Experimental
In this work, Ti55-xZr20Hf15Al10Nbx (x=1~9 at. %, denoted as Nb1~Nb9, respectively, in the following context for convenience) MEAs, which have the configuration entropies (-Rxilnxi) in the range of 1.21R to 1.40R, were investigated. The MEA were prepared by arc melting the metallic elements with the purity above 99.95 wt. % inside the chamber filled with low-pressure high-purity argon atmosphere. To ensure chemical homogeneity, the ingots were flipped over and re-melted at least 5 times. After melted, the melt was drop-cast into the water-cooled cupper mold with the size of 10 mm×10 mm×45 mm. The room temperature tensile properties of the MEAs were tested using CMT4105 universal electronic tensile testing machine at a strain rate of 1×10-3 s-1 on the plate-type specimens with the gauge length of 15 mm, width of 3 mm and thickness of 2 mm. The surface of the specimens was polished down to 2000-grit SiC paper. The phase constitutes for the interested alloys were tested by X-ray diffraction (XRD) with CuKα radiation (Rigaku Kmax 2500 V) in a scanning range of 2θ=20-100 at the rate of 10 s-1. The phase transformation temperatures were measured by STA 449C Jupiter model of differential scanning calorimetry (DSC). Protective argon gas was used for purging and protection, and the flow rate of argon gas was 50 ml/min. A heating rate of 20 K/min was adopted. The microstructure of the MEAs was characterized by a Zeiss Supra55 scanning electron microscope (SEM) and a Tecnai G2 F20 200 kV transmission electron microscope (TEM). SEM specimens were first polished down to 2000-grit SiC paper, and electrochemically polished for the final surface clarification using a methanol, 2-butoxy ethanol and perchloric acid solution with a direct voltage of 25 V for 45~60 s at the temperature about 243 K. After polished, the specimens were etched using a solution of HF: HNO3: H2O=1:2:17 (in volume) at room temperature. TEM specimens were punched to Ф3 circle sheets and then ground to about 60 μm, followed by twin-jet electro-polishing using a mixed solution of glacial acetic acid: perchloric acid : ethanol = 4:8:88 (in volume) at a temperature around 243 K.
3.
Results
3.1 Compositional dependence of microstructure The phase constitution and microstructure of the as-cast MEAs were firstly characterized by X-ray diffraction and scanning electron microscopy methods. As shown in Fig.1, the phase composition transforms from hexagonal to orthorhombic and then to cubic structure as Nb content increases from 1% to 9%. In Nb1 MEA, only single hexagonal α′ phase forms. With the Nb content increasing to 2% and 3%, the (100)α′, (002)α′ and (101)α′ peaks significantly 3
broaden. Besides, the (100)α′ peak of Nb3 MEA shifts to low 2θ angle side, and (101)α′ peak shifts to high 2θ angle side, meaning that new phase has formed in this alloy. This conjecture can be further deduced from the peak at 2θ=71°. As Nb content is increased to 4% (Fig. 1(b)), the (100)α′ peak splits into (100)α′′ and (020)α′′, (101)α′ peak splits into (111)α′′ and (021)α′′, while the (002)α′ peak turns into (002)α′′, in addition, the (110)β appears near (002)α′′. These results mean that α′′ phase with orthorhombic structure become the major phase in the MEA with 4% Nb. In Nb5 and Nb6, α′′ phase is much enough so that the (111)α′′ and (021)α′′ peaks can be detected. For Nb7 to Nb9, only single β phase can be observed, without any peaks of α′ or α′′ phase. In all the interested alloys, the XRD patterns only show the existence of α′, α′′ and β phases, and no ω phase is detected, which exists in many metastable Ti alloys. Similar results have been reported in Ti-Zr-Nb-Al[16], Ti-Ta-Al[17], and Ti-V-Al[18] alloys, indicating that the addition of Al suppresses the formation of ω phase. To investigate and validate the microstructural feature with increasing the content of Nb directly, these MEAs were examined by SEM method. As shown in Fig. 2, both the Nb1 and Nb3 MEAs are composed of fine basket-weave and Widmanstatten plates, which commonly exist in near-α titanium alloys [19, 20]. The microstructure of Nb4 alloy contains short basket-weave type of orthorhombic α′′ phase inside the grains of cubic β phase, and no basket-weave lamellar microstructure formed nearby the grain boundary. When the Nb content increases to 5%, the α′′ phase takes the shape of small particles inside the equiaxed grains (in Fig. 2(e)). As the Nb content further increases to 8% (Fig. 2(f)), only equiaxed grains with obvious homogeneous dendritic structure can be observed. TEM investigation was conducted to further reveal the effect of Nb on the microstructure evolution of these MEAs. The refined structures and the selected-area electron diffraction patterns (SADPs) of the Nb1, Nb3, Nb5 and Nb7 MEAs are shown in Fig.3. It is seen that Nb1 MEA consists of different kinds of lathes, from less than 100 nm to 500 nm in width. The SADP confirms that the lath is of hexagonal lattice, i.e., ′ phase, which is in accordance with the XRD result. The Nb3 MEA exhibits two different morphologies. As shown in Fig. 3(b), the large sheet structure with flat and parallel interface on the bottom part is calibrated to be the α′′ phase, while short lamellar structure at the upper part is the α′ phase. In Nb5 MEA, the β and α′′ phase coexist, the α′′ phase takes the shape of plate with about 200 nm in thickness, and has the orientation relationship of [011]β//[001]α′′ with β phase, which has been also found in traditional Ti alloys[21]. In Nb7 MEA, only β phase is observed in the high magnification TEM image, as shown in Fig. 3(d). 3.2
Phase transformation during heating and cooling process
The temperature dependence of phase composition for the interested MEAs was investigated by DSC method in this work. As shown in Fig.4, there are two or three peaks at the heating curves and one peak at the cooling curves (all the peaks marked by sequence 4
numbers). The start and finish temperatures of the phase transformation represented by each peak are listed in Tab.1. In combination with the phase constitute investigated by the XRD and TEM of the MEAs, the Peak1 during the cooling process can be determined to be the transformation from β to α′ and/or α′′. For Nb1 MEA with single α′ phase, as shown in Fig. 4(b), there is only one extreme point on the Peak1, namely the transformation from β to α′. For Nb2 and Nb3 shown in Fig.6(c), the Peak1 displays two extreme points, suggesting that the β phase transforms to α′ and α′′. When cooling these samples from high temperature, a portion of α′ phase with less β-stabilizing element precipitated firstly, then the remnant high temperature β phase containing more β stabilizing elements transfers into orthorhombic α′′ phase. So the extreme point at higher temperature is β→α′, and the one at lower temperature is β→α′′. For Nb4 MEA, the extreme point representing the transformation from β to α′ disappears. As the Nb content further goes up to 7%, the peak during cooling disappears suggesting that no α′′ phase exists with Nb up to 7%. As listed in Tab.1, the Peak1 shifts toward low temperature side with the Nb increasing, meaning that the high temperature β phase becomes more and more stable. The phase transformations reflected by the DSC curves during heating seem to be more complicated than those appeared during the cooling process. For Nb1~Nb3 MEAs, only Peak4 is obvious, which represents the transformations of α′→β (Nb1) and α′/α′′→β (Nb2, Nb3) in the temperature range of 600 ~810 °C. For Nb4~Nb7, three peaks can be obviously found in the range of 130~300 °C, 400~480 °C and 600~810 °C. In order to reveal the structural evolution clearly, we selected Nb6 MEA to have heat treatments at 300 °C and 600 °C for 6h (denoted as Nb6-300 and Nb6-600), respectively. Fig. 5 shows the XRD patterns, DSC curves, TEM and SADP for Nb6-300 and Nb6-600 MEAs. It is seen that the XRD peaks of Nb6-300 are similar to those of as-cast Nb6 MEA. On the DSC curve shown in Fig. 5(b), there is an exothermic peak in the temperature range from 130 °C to 300 °C. The SADP shown in Fig.5(c) indicates that ω phase exists in Nb6-300 sample, so that the phase transformation is β→ω. The precipitation of ω phase is common in traditional titanium alloys such as Ti-Nb[21], Ti-Mo[22], Ti-Nb-Mo-Sn[23] alloys during low temperature annealing. The Nb6-600 MEA is mainly composed of ′ with hcp structure, which can be determined by all the peaks on the XRD pattern except the one located at 2θ=54º. As shown in Fig. 5(d), the bright field TEM image and the corresponding SADP demonstrate that the microstructure is composed of α′ phase and β phase with bcc structure. Based on the above results, the transformations in Nb6 MEA represented by the three peaks during heating can be determined as Peak 2: β→ω; Peak 3: β+α′′+ω→α+β′; Peak 4: α+β′→β. For Nb9 MEA, there are similar phase transformations with those of Nb6. As shown in Fig. 6(a), the peaks of hexagonal ω phase is obvious on the XRD pattern of Nb9 MEA after heated at 300 °C for 6h. It is found that the Nb content has no effect on the transformation temperature, but elevates the volume fraction of the ω phase. Besides, only Peak2 and Peak4 5
appeared on the heating process, which indicates the β→ω and ω→β transformation, respectively. That is to say, the β phase is too stable to decompose into α′ phase at medium temperature during DSC test in this alloy, only thermal ω phase separated out at low temperature. The increase of Nb content seems to have no obvious influence on the start and finish transformation temperature of Peak1, but makes Peak4 move towards low temperature side. In conclusion, the Nb content promote the stability of β phase. 3.3
Phase composition determined tensile properties
From the above section, it has been found that the phase constitution and microstructure of these MEAs are strongly influenced by the addition of Nb. How about the effect of phase composition on the tensile properties of these MEAs? In this respect, typical room-temperature true stress-strain tensile curves of these TiZrHfAlNb alloys are plotted in Fig. 7(a). The yield strength, tensile strength and fracture elongation as a function of Nb content are displayed in Fig. 7(b). From the tensile curves, it can be seen that with the Nb element content lower than 3%, the MEAs exhibit the yield strength above 1200 MPa, the tensile strengths of ~1500 MPa, and the elongations in the range from 3% to 4.5%. However, the yielding strength of Nb4 abruptly decreased to 227 MPa from about 1200 MPa for Nb3 because Nb content is increased by 1%, as shown in Fig. 4(b). The alloys show “double yielding” phenomenon and remarkable work-hardening effect with the Nb content in the range of 4%~8%. For these MEAs, the first yielding strength increases from 227 MP to 735 MPa, the highest tensile strength and increases reaches 1211 MPa for Nb4 and Nb7 MEA, respectively. It is noteworthy that for Nb8 alloy, after yielding at 735 MPa, the plastic deformation continues until 27%, and the “double yielding” phenomenon becomes slightly during this period. Therefore, the MEAs developed in this work exhibit excellent work hardening effect and gigantic product of strength and elongation. Nb4 MEA shows an increase of about 1000 MPa from first yielding stress to fracture strength. And the product of strength and elongation of Nb8 MEA reaches above 30 GPa%, which is comparable to the level of TRIP steel. When the Nb content is increased up to 9%, the “double yielding” together with the work-hardening effect disappears. The SEM images of the fracture surfaces for these TiZrHfAlNb MEAs are shown in Fig. 8. Corresponding to the tensile curves shown in Fig. 7(a), the fractures of Nb1 and Nb3 MEAs with relatively low elongation exhibit pronounced intergranular and transgranular mixed fracture mode. The Nb4~Nb8 MEAs have obvious transgranular fracture characteristics with dimples in different sizes. Among them, the dimples of Nb8 alloy are apparently fewer in amount but larger in size than those of other MEAs, indicating that the deformability of this MEA is the best. The above results of tensile experiments illustrate that the variation of phase composition and microstructure indeed result in mutation of mechanical properties. With the increase of 6
Nb content, the MEAs become more plastic and toughening although the strength decreases. Moreover, the deformation behavior changes from single yielding mode to double yielding mode. The intrinsic mechanisms for these phenomena are discussed in the following sections.
4.
Discussion
4.1
Cooperative effect of the components on the formation of metastable phase In traditional titanium alloys, Nb is a β stabilizing element, and used to promote the
formation of β phase. The α′′ and phase existing in the dual phase region are considered to be metastable[24]. In these TiZrHfAlNb MEAs, it has also been found that Nb is a stronger β-stabilizer to adjust the phase composition. Although the Nb content is insufficient to ensure the formation of β or α′′ phase for Nb1 MEA, so much alloying elements in the α′ phase finally results in high solid solution strengthening effect. With Nb content increased to 2%, the α′ phase is supersaturated, the metastable α′′precipitated out. When the content of Nb increased up to 4%, the metastable α′′ phase becomes dominant. In this MEA a little amount of β phase appears. The volume fraction of β phase goes up with the further increasing of Nb. When Nb content is going up to 9%, the β phase become stable. These results are similar to those for Ti-Nb and Ti-Zr-Nb alloys. For Ti-Nb alloys, the critical contents of Nb for phase transition from hcp α′ phase to orthorhombic α′′ phase and then to bcc β phase are 14% and 27%, respectively[25]. For Ti-30Zr-xNb alloys, however, the corresponding critical values decreased to 7% and 13% [26]. It has been reported that the addition of Zr is able to decrease the martensitic transformation start temperature (Ms) in low Zr-containing β-type alloys, such as Ti-Cr-Zr[27], Ti-Ta-Zr[28], Ti-Nb-Zr[29], Ti-Nb-Ta-Zr[30] and Ti-Nb-Sn-Zr[31], so Zr is considered as a β-stabilizing element in these alloys. And the β-stabilizing effect of Zr appears to be influenced largely by other alloying elements in these alloys, that is to say, the β-stabilizing effect of Zr increases with the increasing content of the β-stabilizing elements Nb, vice versa, the addition of Zr promotes the β stabilizing effect of Nb. The reason may be attributed to that the bonding energy of Ti-Zr-Nb (4.97 eV atom-1) alloy is higher than those of binary Ti-Zr (4.80 eV atom-1) and Ti-Nb (4.80 eV atom-1) alloys[26]. As a member in the same family in the period table of elements, Hf may have the same effect. Therefore, the β-stabilizing effects of Zr, Hf and Nb are promoted by each other. Besides, as reported in Ti-20Zr-10Nb-xAl (x=1, 2, 3, 4) alloys [16], the reverse martensitic transformation start (As) and finish (Af) temperatures decrease with increasing Al content at a rate of approximately 10K per Al atom, and a similar suppression effect of the transformation temperature has also been reported in Ti-24Nb-Al[32] alloy. In these TiZrHfAlNb MEAs, Al could also be considered as a β-stabilizing element. The cooperative addition of Zr, Hf, and Al in these Ti-rich MEAs lowers the critical content of Nb for the formation of metastable α′′ and stable β phase. Therefore, it is seen that through appropriate controlling of Nb addition, the metastable microstructure that contains the 7
composite of α′′ and β phase can be achieved in these MEAs. 4.2
Plasticizing mechanism by stress induced martensitic transformation
It has been confirmed that Nb1~Nb3 MEAs have high strength with low rupture elongation, which is due to the intrinsic properties of the highly alloyed hexagonal phase and supersaturated solid solution strengthening, as shown in Fig.7. The elongation decreases in this compositional range because dislocation movement is hindered at the interface of α′′ and α′. As the Nb content goes up to 4%, “double-yielding” effect appears and holds until the Nb content reaches up to 8%. For the traditional shape memory Ti-alloys such as Ti-Nb-(Zr, Sn, Mo, Al) [21, 23, 29, 33], the stress platform after yielding event corresponds to the reorientation of the twinned martensite variants or SIMT from β to αʺ phase. When the alloy with insufficient β stabilizers cooling down from high temperature, the transformation of β phase into αʺ leads to self-accommodating microstructure in order to minimize the deformation of ex-β grain during transformation. As a consequence, in a given ex-β grain, variants of αʺ martensite are auto-organized by groups in twinning relationship. When a stress is applied, the first activated deformation mechanism is the growth of variants that are favorable oriented to the detriment of others, the self-accommodating microstructure is thus reoriented. However, the metastable β phase retains as the alloy with sufficient β stabilizers cools down. The material may deform by SIMT from metastable β to αʺ phase under applied stress. As for the current MEAs, the deformation mechanism changes with the phase composition and phase stability. For Nb4 MEA, microstructures at different strains were observed by TEM observations. Fig.9 shows the martensitic microstructure which is composed of two different morphologies in Nb4 MEA prior to deformation. One is the martensite laths with the width of several microns, which organize themselves in parallel with inner martensite in them, as shown in Fig.9 (a) and (b). The corresponding SADPs in Fig.9 (c) shows that the inner martensite has orientation relationship [011]αʺ1//[110]αʺ2 with the matrix, which means that they do not share a twinning plane. On the other hand, the {111} type I twinning with [110]αʺ2M// [-110]αʺ2T relationship are discovered between the adjacent martensite laths. The other morphology of martensite phase is in triangle shape, as shown in Fig.9 (d), which are the (011) compound twinning martensites with [100]α″M//[-100]α″T. Fig.10 shows the microstructure of Nb4 MEA subjected to a tensile strain of 3%. The inner martensite was consumed by the matrix compared with torsion. New {111} type I of twins with orientation relationship [011]αʺM// [0-1-1]αʺT are observed. Meanwhile, the (011) compound twinning disappears. Fig.11 is the microstructure after fracture. It can be seen that the second complex twinning microstructure has disappeared, and all the plates organized themselves in parallel, which means that Nb4 was deformed by reorientation of the twinned α″ martensite. 8
To uncover the deformation mechanism of MEAs alloys with more metastable β phase, unloading experiments at 2% and 7% pre-strain of Nb5 MEAs were conducted. From the relative intensities of peaks of the XRD pattern as shown in Fig.12, the content of α′′ increases with the plastic strain. The optical micrographs of fractured Nb5 and Nb8 MEAs as shown in Fig.13. It is seen that the martensitic transformation in Nb5 MEA is hindered by the original α′′. In contrast, the induced martensitic α′′ plates in Nb8 MEA almost grow in the same direction to intersect the whole grain at low strain, and then plates in another direction appeared at high strain. These results illustrate that the SIMT from metastable β to α′′ takes place during loading. The SIMT process of a mixture of β and α″ phases in Ti-33Nb-4Sn[34] alloy was investigated by the in situ synchrotron X-ray diffraction. At the initial deformation stage, the pre-existing α′′ and β phase concurrently experienced elastic deformation. After the first yielding, the SIMT was activated and then it continued during the whole deformation. Especially, the SIMT processed intensely during the stress plateau. Meanwhile, the reorientation of pre-existing α′′ variants occurred at the end of the stress platform. After the stress plateau, the re-oriented and SIMT α′′ variants deformed elastically together with the elastic deformation of residual β phase, leading to a long work hardening stage. Beyond the second yielding period, the complex mixture phase deformed plastically. The same situation happened in the Nb5 and Nb6 MEAs, the critical stress for SIMT at the first stage increases with Nb content due to the enhancement of the β phase stability. Meanwhile, the elongation increases with the Nb content. Specifically, Nb8 shows a long-term work hardening stage until 27% with slight double yielding, indicating that the suppression of original α′′ martensite by Nb content ensures the alloy better deformation ability. Fig. 14 shows the microstructures of Nb8 MEA with 2% and 10% of pre-strain deformation, respectively. After 2% deformation, the SIMT has happened to generate a small amount of martensite in about 100 nm width (Fig. 14(a)). With the deformation increasing to 10%, more martensite plates have appeared. In the β phase between two martensite plates, a large number of dislocations piled up on the interface. It could be seen that some dislocations has cut into the new-born martensite, which can be observed by the dark field TEM image (Fig. 14(b)), and the martensite is divided into several sections. The SIMT results in the intensive strain-hardening effect by dynamic strain-stress partitioning between the β and α′′ phase, and the dislocation movement intensifies the partitioning effect inside the α′′ phase. The partitioning promotes the plastic deformation inside of grains, which can effectively suppress early cracking. Besides, large amount of interfaces are introduced due to the partitioning effect, which then restrict the propagation and movement of dislocation lines. The suppression of cracks and dynamics reduction of mean free path for the dislocations give rise to the combination of the initial work hardening and ductility. Therefore, it is seen that the superior combination of strength and deformation capacity has been achieved by the synergy 9
of SIMT and dislocation movement.
5.
Conclusion
In this research, the microstructure, phase stability and tensile properties of a series of Ti55- Ti55-xZr20Hf15Al10Nbx (x=1~9, at. %, named Nbx) medium entropy alloys were investigated. Based on these results, the conclusion can be drawn as follows: 1. With the increase of β stabilizing element Nb, the phase composition of as-cast MEAs changes from hexagonal α′ to orthorhombic α′′, then to metastable β, finally to stable β phase at 2%, 4% and 9%, respectively. 2. The composition has strong effect on the phase evolution during heating and cooling process. Nb1 MEA exhibits a reversible phase transformation between α′ and β phase. While for Nb2 and Nb3 MEAs with α′′ and α′ phase, the reversible transformation turn into α′′/α′ ↔β. For Nb4~Nb8 MEAs, the phase evolution during heating process is in the sequence of β+α′′→ω+α′′+β→α+β′→β, meaning that the formation of metastable α′′ and β phase in certain compositional and temperature range. For Nb9 MEAs, there is only one ω→β transformation during heating process. 3. The phase constitution and microstructure of these MEAs show intense impact on their tensile properties. For Nb1~Nb3, the MEAs exhibit high yield strength and fracture strengths, and certain elongations. As the Nb content increased from 4% to 8%, the alloys show “double yielding” phenomenon and remarkable work-hardening effect. Nb8 MEA have excellent the product of strength and elongation above 30 GPa%. 4. The deformation mechanisms for these MEAs changes with the phase composition. The
high strength for Nb1~Nb3 alloys attributed to the dislocation slip in hexagonal and orthorhombic structures. In Nb4 MEAs with most α′′ phase, the deformation dominated by the de-twinning and reorientation of martensite variants. In Nb5~Nb8 MEAs with most metastable β phase, the SIMT from metastable β to α′′ phase under external force has been confirmed. Besides refinement of the newborn martensite by the dislocation has also been verified. The integrated effects of SIMT and dislocation movement endow the MEAs excellent deformation capacity.
Acknowledgments This work was financially supported by The National Key Basic Research Program (2016YFB0701402), National Natural Science Foundation of China (Nos. 51571016 and 51531001). Key Laboratory of Research on Hydraulic and Hydro-Power Equipment Surface 10
Engineering Technology of Zhejiang Province (2017SLKL003).
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(a)
(b)
Fig.1 XRD patterns of as-cast TiZrHfAlNb MEAs, indicating a phase transion αα′′β with the increase of the content of Nb.
13
(a)
(b)
(c)
(d)
(e)
(f)
Fig.2 SEM micrographs of as-cast TiZrHfAlNb MEAs. (a) Nb1; (b) Nb3; (c)and (d) Nb4, (e) Nb5, and (f) Nb8.
14
(a)
(b)
(c)
(d)
Fig.3 TEM bright field images and corresponding SADPs of TiZrHfAlNb MEAs. (a) Nb1, (b) Nb3, (c) Nb5, and (d) Nb7.
15
(a)
(b)
(c)
Fig. 4 DSC curves of as-cast TiZrHfAlNb MEAs. (a) DSC curves of nine HEAs indicating the heating and cooling process with the peaks marked by series numbers, (b) and (c) the curves of Nb1 and Nb3, respectively.
16
(a)
(c)
(a)
(b)
(b)
(d)
Fig.5 Phase transition of Nb6 MEA during heating treatment. (a) XRD patterns heated at 300 °C and 600 °C for 6h; (b) DSC curves with all the peaks marked with corresponding phase transition; (c) and (d) TEM bright field image and corresponding SADPs treated at 300 °C and 600 °C, respectively.
17
(a)
(b)
Fig. 6 (a) XRD patterns of Nb9 MEA at as-cast and treated state at 300 °C for 6h, respectively; (b) DSC curves of Nb9 MEA with the peaks marked to indicate the phase transitions.
18
(a)
(b)
Fig.7 Room temperature tensile properties of as-cast TiZrHfAlNb MEAs. (a) The true stress-strain curves; (b) Dependence of tensile strength, yield strength and elongation as a function of Nb content.
19
(b)
(a)
(a)
(c)
(d)
(e)
(f)
Fig. 8 Fracture morphologies of as-cast TiZrHfAlNb MEAs. (a) Nb1; (b) Nb3; (c) Nb4; (d) Nb5; (e) Nb6; (f) Nb8.
20
(a)
(b)
(c)
(d)
Fig.9 The martensitic microstructure in Nb4 MEAs prior to deformation. (a) and (b) the martensite variants with laths structure; (c) the SADPs of the corresponding area in (b); (d) martensite variants with triangle shape and the corresponding SADPs .
21
(a)
(b)
(c)
(d)
Fig.10 the microstructure of Nb4 MEAs subjected to a tensile strain of 3%. (b) and (d) are the corresponding SADPs of the areas marked by red circles in (a) and (c), respectively.
22
(a)
(b)
(c)
(d)
Fig.11 The microstructure of Nb4 MEAs after fracture. (c) and (d) are the corresponding SADPs of the marked area in (b).
23
Fig.12 Correlation between microstrucrue evolution and stress-strain curve of Nb5 MEAs.
24
(a)
(d)
Fig.13 Optical micrographs of Nb5 (a) and Nb8 (b) MEAs after fracture.
25
(a)
(b)
Fig. 14 TEM bright (a) and dark (b) field images of as-cast Nb8 MEAs at ε=0.1.
26
Table 1. The start and finish temperatures for the transformation presented by the peaks on the DSC curves for the Nbx MEAs.
Transformation temperature during heating /°C Peak 2
Peak 3
Transformation temperature during cooling / °C
Peak 4
Peak 1
Nb1
774
846
760
689
Nb2
750
821
733
676
676
633
Nb3
745
806
685
645
645
604
Nb4
704
785
646
575
Nb5
123
283
410
451
647
764
582
483
Nb6
125
290
422
458
633
737
505
405
Nb7
126
293
427
475
621
724
476
375
Nb8
133
303
593
714
Nb9
137
306
591
689
27