Thermal Stability of the HfNbTiVZr High-Entropy Alloy - Inorganic

Article Cite This: Inorg. Chem. XXXX, XXX, XXX−XXX

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Thermal Stability of the HfNbTiVZr High-Entropy Alloy Victor Pacheco,*,† Greta Lindwall,‡ Dennis Karlsson,† Johan Cedervall,† Stefan Fritze,† Gustav Ek,† Pedro Berastegui,† Martin Sahlberg,† and Ulf Jansson† †

Department of ChemistryÅngström Laboratory, Uppsala University, Box 523, Uppsala SE-75120, Sweden Department of Materials Science and Engineering, KTH Royal Institute of Technology, Stockholm SE-10044, Sweden

‡

Inorg. Chem. Downloaded from pubs.acs.org by UNIV OF SOUTH DAKOTA on 12/14/18. For personal use only.

S Supporting Information *

ABSTRACT: The multicomponent alloy HfNbTiVZr has been described as a single-phase high-entropy alloy (HEA) in the literature, although some authors have reported that additional phases can form during annealing. The thermal stability of this alloy has therefore been investigated with a combination of experimental annealing studies and thermodynamic calculations using the CALPHAD approach. The thermodynamic calculations show that a single-phase HEA is stable above about 830 °C. At lower temperatures, the most stable state is a phase mixture of bcc, hcp, and a cubic C15 Laves phase. Annealing experiments followed by quenching confirm the results from thermodynamic calculations with the exception of the Laves phase structure, which was identified as a hexagonal C14 type instead of the cubic C15 type. Limitations of the applied CALPHAD thermodynamic description of the system could be an explanation for this discrepancy. As-synthesized HfNbTiVZr alloys prepared by arc-melting form a single-phase bcc HEA at room temperature. In situ annealing studies of this alloy show that additional phases start to form above 600 °C. This indicates that the observed HEA is metastable at room temperature and stabilized by a slow kinetics during cooling. X-ray diffraction analyses using different cooling rates and annealing times show that the phase transformations in this HEA are slow and that completely different phase compositions can be obtained depending on the annealing procedure. In addition, it has been shown that the sample preparation method (mortar grinding, heat treatment, etc.) has a significant influence on the collected diffraction patterns and therefore on the phase identification and analysis. The novel idea of designing alloys with five or more principal elements has significantly increased the design possibilities. In order to accelerate the discovery of new HEAs with targeted properties, the phase formation mechanisms and the factors governing phase stability need to be fully understood, since simple trial and error experiments cannot be used to investigate the whole HEA compositional range effectively.13 Although a variety of phase formation rules, based on atomic size differences (δ), electronegativity (χ), thermodynamic parameters (ΔHmix and Ω), and valence electron concentration (VEC), have been proposed,1 this topic is still not fully understood. Many HEA systems have been found to solidify as singlephase solid solutions, for example, CoCrFeMnNi (the so-called Cantor alloy)14 and CoCrCuFeNi, which form ccp solid solutions, while alloys based on refractory transition metals

1. INTRODUCTION Traditional alloying methods in metallic materials are based on the incorporation of additional elements to a system dominated by one or two main components, e.g., steel and nickel-based superalloys. Recently, a new alloying approach has been demonstrated, where five or more elements are combined in equimolar or near equimolar compositions. These alloys often exhibit solid solutions of simple ccp (cubic close packed, sometimes called FCC or face centered cubic in the literature) or bcc (body centered cubic) structures. The solid solution is assumed to be stabilized by high configurational entropy, and these are therefore named high-entropy alloys (HEAs).1,2 The research on HEAs has been extensive because of their excellent mechanical properties at high temperatures3,4 as well as cryogenic temperatures,5 high specific strength,6 high crack resistance,7 and good corrosion resistance.8−10 Recently, an excellent hydrogen storage capacity has also been reported for the refractory alloy HfNbTiVZr.11,12 © XXXX American Chemical Society

Received: October 17, 2018

A

DOI: 10.1021/acs.inorgchem.8b02957 Inorg. Chem. XXXX, XXX, XXX−XXX

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Inorganic Chemistry such as HfNbZrTaTi15 and MoNbTaVW usually form bcc solid solutions.1,16,17 However, it is known that some HEAs, which solidify as single-phase solid solutions, experience phase transformations when exposed to heat treatments.18,19 An excellent example is the equimolar AlCoCrFeNi, which can form a bcc solid solution but transforms by spinodal decomposition into a disordered Fe−Cr-rich bcc phase and an ordered (B2) Ni−Al-phase.20 Secondary phases have also been reported for the CoCrFeMnNi alloy at temperatures below 800 °C, even though it is considered to be a prototype for a single-phase solid solution HEA.21 It is possible that many alloys that claimed to be single-phase HEAs are only stabilized by kinetics and that they will form secondary phases upon aging or after high-temperature annealing. The formation of such secondary phases can be beneficial or detrimental for the properties of the alloys. Therefore, an evaluation of thermal stability is essential. The refractory HfNbTiVZr alloy has an excellent hydrogen storage capacity and is therefore interesting for future storage applications. This alloy is described as an HEA in the literature, but contradictory results have been reported regarding phase stability. Arc-melted equimolar alloys are clearly single solid solutions.11,12 Furthermore, Feuerbacher et al. have reported a single bcc solid solution after arc-furnace synthesis followed by an annealing at 1500 °C.16 Fazakas et al. have reported one bcc single-phase after synthesis by induction melting followed by suction casting.17 However, this group has identified two structures after heat treatment at 900 °C during 10 min: a bcc solid solution and a cubic C15 Laves phase.17 The formation of the Laves phase in this system is further supported by the predictions from the Laves phase formation criteria proposed by Yurchenko et al.,22 since the alloy is in the area of the electronegativity and atomic size differences in which HEAs with Laves phases are expected. In order to clarify the phase formation mechanism and thermal stability in the HfNbTiVZr refractory system, the alloy has been evaluated at different annealing times and temperatures using X-ray diffraction (XRD), scanning electron microscopy (SEM), and transmission electron microscopy (TEM). The phase stability is discussed and compared with thermodynamic calculations performed using the CALPHAD method. The effects of different sample preparation methods have also been investigated.

water. After this procedure heat treatments were performed at 600, 800, 1000, and 1200 °C for 48 h and 600, 800, and 1000 °C for 1 h. The crystalline structure was evaluated using a Bruker D8 Advance Diffractometer with Cu Kα radiation and a θ−2θ setup. In situ XRD measurements with a wavelength of 0.20711 Å were conducted at the P02.1 beamline at PETRA III in the German electron synchrotron (DESY). These experiments were performed in transmission mode using a PerkinElmer XRD1621 detector and a 5 s exposure time for each diffraction pattern. The experimental setup consisted of sapphire capillaries connected to a vacuum system and to an Ar gas source.23 After the system was evacuated, it was filled with Ar gas up to a slight overpressure of 0.1 bar. The sample was then heated to 900 °C with a heating rate of 15 °C/min using a Kanthal wire wrapped around the sapphire capillary. The temperature was monitored using a thermocouple located inside the sapphire tube. Detailed explanations of this experimental setup can be found in ref 23. The data reduction from 2D images of the diffraction pattern was conducted using the software Fit2D.24 The microstructure was evaluated with a Zeiss Merlin SEM with a secondary electron (SE) detector and an energy-dispersive X-ray spectrometer (EDS). The samples for electron microscopy analysis were prepared by standard metallographic techniques through grinding with SiC paper. For the final polishing, a mixture of SiO2 and H2O2 was used to unveil the microstructure. Transmission electron microscopy studies were carried out using a FEI Titan Themis 200 (scanning) transmission electron microscope, equipped with probe corrector and a SuperX EDS detector. The TEM sample was prepared using the in situ lift-out technique in an FEI Strata DB 235 focused ion beam (FIB). The region of interest was coated with platinum to protect it from ion beam damage during preparation, and a final polishing step using 5 keV ions was used to minimize ion damage to the lamella. The electron diffraction patterns were evaluated using the Crystbox software.25

3. THERMODYNAMIC CALCULATIONS The thermodynamic properties of the HfNbTiVZr system were studied using the CALPHAD approach. CALPHAD is a method to describe phase-based properties of multicomponent systems.26 In the case of CALPHAD thermodynamic modeling, the Gibbs free energy is the property modeled for each phase in the system of interest. Since many phases are solid solutions with varying composition, the compound energy formalism (CEF) is commonly used26−28 and the Gibbs free energy for a phase is modeled as Gm =

srf

Gm − TcnfSm +

phys

Gm + EGm

(1)

where srfGm is the Gibbs energy of an unreacted mixture of the constituents in a phase; physGm is the Gibbs energy contribution given by a physical model, e.g., a model for magnetic transitions; and cnfSm is the configurational entropy term. The nonideal contribution to the Gibbs energy is captured by the EGm term expressed as

2. EXPERIMENTAL SECTION HfNbTiVZr alloys were prepared through electric arc-melting synthesis in a high-purity Ar atmosphere. Raw elements of highpurity Hf (Johnson Matthey, Materials and Technology U.K, 99.6%), Nb (Cerac, 99.8%), Ti (Chempur, 99.995%), V (MRC, 99.95%), and Zr (Cerac, 99.8%) were melted to obtain an equiatomic composition. Each alloy was remelted at least five times to ensure complete homogenization. The chemical composition of the alloys was verified using X-ray fluorescence (XRF), and only slight deviations from the target composition were observed (±3 at. %). Different types of heat treatments were conducted to evaluate the thermal stability. Heat treatments with a duration of 1 h at 600, 800, and 1000 °C were conducted in a high-vacuum system with a base pressure of 8 × 10−8 Torr. The samples were heated and cooled with a rate of 5 °C/min (slow cooling). Additional heat treatments with fast cooling (water quenching) were conducted using a muffle furnace. In this case, the samples were placed in Al2O3 crucibles and sealed in evacuated quartz ampules to prevent oxidation. A steel ampule was used for the 1200 °C experiment. The ampules were heated with 5 °C/min to the holding temperature and finally quenched in cold

E

Gm =

∑ PI (Y )L I i

Ii

1

+

∑ PI (Y )L I 2

I2

2

+ ... (2)

It includes sums over interaction parameters, L, multiplied by the product of the constituent fractions, P(Y), of the first order, second order, etc. To construct a CALPHAD thermodynamic description, the Gibbs energy functions are parametrized using computational and experimental data on thermochemistry and phase equilibria and the parameters are stored in CALPHAD databases. Equilibrium phase fractions and compositions can then be obtained by minimizing the Gibbs energy of the B


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Inorganic Chemistry

lower temperatures, the hexagonal closed packed (hcp) phase as well as a second low-temperature bcc phase (bcc LT) also become stable. In Figure 2a, the equilibrium composition of the C15 phase is shown and a relatively constant composition over its temperature stability range is observed, i.e., at about (Hf, Zr)1V2 with some solubility of Ti on the V lattice sites. The hcp phase contains comparable amounts of Hf, Ti, and Zr according to the calculations but changes with temperature as shown in Figure 2b. The bcc LT phase is rich in Nb with increasing solubility of the other elements with temperature as shown in Figure 2c. As can be seen in Figure 2d, the bcc HT phase has an equimolar composition corresponding to an HEA. The present calculations compare relatively well with the experimental results below. Hence, extrapolation from the binary system descriptions seems to be a good approximation for the current system in order to identify the conditions for which a single-phase bcc region is stable. All five constituents have bcc as their equilibrium structure for large temperature intervals, and the binary phase diagrams for eight of the 10 binary subsystems contain a temperature interval where a single bcc phase is stable over the whole composition range. The exceptions are the V−Hf and V−Zr binary phase diagrams for which the C15 Laves phase becomes stable when the solubility limit of the bcc phase is exceeded, see Figure 3a,b. The ideal configuration entropy’s contribution to the Gibbs free energy of a phase increases with the number of constituents in a solid solution and is an important factor for the stability of HEAs. Although it is important to keep in mind that not all highly concentrated solid solutions are stable since their stability depends on both nonideal contributions to the excess Gibbs energy and competing phases, it could be of interest to study the thermochemical properties of a phase and how the individual contributions to Gibbs free energy vary with composition and temperature. This is possible when using the CALPHAD approach, and in Figure 4, the calculated energy curves for the HfNbTiVZr bcc phase are shown as a function of temperature at the equiatomic composition. The standard element reference (SER)30 is used as the reference state. In addition to the Gibbs free energy curve, both the enthalpy and the entropy parts are shown in the figure. The calculation shows that the entropy contribution, −TSm, is the dominating part of the Gibbs free energy for the bcc phase for this composition. Furthermore, it can be concluded that the excess entropy contribution is small and that the entropy of the phase is primarily given by the ideal configuration entropy. This is demonstrated by comparing the entropy of mixing with the ideal mixing entropy for a five-component system which is −13.38 J/mol K. The entropy of mixing according to calculation using the TCHEA3 description is −13.05 J/mol K at 1273 K (approximately constant with varying temperature), and hence, the excess entropy contribution is less than 0.5 J/mol K. The calculated values should be regarded with some caution since none of the ternary subsystems in the TCHEA3 database are assessed and the excess energy thus is lacking potential nonideal contributions from ternary and higher-order interactions. Nevertheless, the small excess entropy is in accordance with ab initio molecular dynamics simulations for a similar system, equiatomic HfNbTaTiVZr, by Gao et al.31 The results above indicate that the CALPHAD description extrapolated from the binary systems may be an adequate approximation for the bcc phase of this system. However, this

system accounting for all phases’ contributions at the condition of interest. The strength of the CALPHAD approach is the possibility to combine constituent subsystems (unary, binary, ternary, etc.) to extrapolate into multicomponent space, and hence, the method is suitable for commercial materials which normally contain several elements and phases. However, the number of subsystems within a multicomponent system that need to be modeled to establish a complete CALPHAD thermodynamic description increases rapidly with the number of components. Consequently, multicomponent commercial CALPHAD databases are often dedicated to a particular alloy system, e.g., Febased or Ni-based alloy systems. The increased interest in HEA systems has enhanced the demand for CALPHAD databases where all the subsystems within a multicomponent system are evaluated over the whole composition interval and commercial databases for HEAs are starting to emerge. In the current work, the HEA thermodynamic database by ThermoCalc Software, TCHEA3,29 was utilized. As will be discussed further in the Results and Discussion section, this database cannot be regarded as complete in the case of the HfNbTiVZr system since only the binary subsystems have been fully assessed. Hence, calculations for the multicomponent system use extrapolations from the binary descriptions; i.e., no ternary or higher-order interactions are included in the excess energy contribution in eq 1. Nevertheless, the present calculations provide a valuable basis for discussing the thermodynamics of this five-component system and help identifying subsystems for which further CALPHAD modeling is needed.

4. RESULTS AND DISCUSSION 4.1. Thermodynamic Calculations. Figure 1 shows the calculated equilibrium phase fractions as a function of

Figure 1. Calculated equilibrium phase fractions as a function of temperature for the HfNbTiVZr system. The HfNbTiVZr HEA (bcc HT) is stable between 830 and 1490 °C.

temperature for the equiatomic HfNbTiVZr system using the TCHEA3 thermodynamic database.29 According to the calculations, the equilibrium microstructure for this alloy consists of a single bcc phase at high temperatures (bcc HT), from the solidus at 1490 °C down to about 830 °C. At this temperature, the cubic C15 Laves phase becomes stable. At C


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Inorganic Chemistry

Figure 2. Calculated elemental composition as a function of temperature of (a) the C15-Laves phase, (b) the hcp phase, (c) the low-temperature bcc LT phase, and (d) the high-temperature bcc HT phase.

Figure 3. Calculated binary phase diagrams for (a) the V−Hf system and (b) the V−Zr system.

Laves phase, a small composition region exists within which the C14 Laves phase is stable.32 Furthermore, a NbV2 C14 Laves phase has been predicted to be stable by first-principles calculations33 which is, although never observed experimentally in the binary Nb−V system, an indication that the Nb−V containing subsystems need to be further studied both theoretically and experimentally to develop a reliable thermodynamic description of the current HEA system. 4.2. Structure and Composition of As-Synthesized HfNbTiVZr. After synthesis in the arc-furnace, the HfNbTiVZr alloy solidifies into a bcc single-phase solid solution as confirmed by X-ray diffraction (Figure 5). The microstructure

might not be the case for all phases within the system. The experimental characterization of the bulk samples in the present work suggests that the hexagonal C14 Laves phase forms during annealing at lower temperatures. The calculations, on the other hand, predict that the cubic C15 phase is the stable Laves phase in this system. This disagreement could be a consequence of the description being limited to extrapolation of the binaries and thus fail to predict phases that would be stable in a ternary system. For the current case, the incompleteness of the description of the Hf−Nb−V system is likely to be the cause since experimental studies have suggested that, in addition to a broad phase field of the C15 D


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Inorganic Chemistry

in Figure 6a. As can be seen, the alloy consists of a single-phase bcc from room temperature up to about 600 °C, where new phases start to form. In this region, the unit cell parameter of the bcc phase changes from 3.3640(4) to 3.4075(5) Å, most probably due to thermal expansion. Also, a change in the peak width can be seen with the peaks getting sharper as the temperature increases. For example, at room temperature the fwhm for the (110) reflection (4.95°) is 0.1001°; for 600 °C, it decreases to 0.0946°, and at approximately 800 °C it reaches 0.0447°. The new phases formed above 600 °C correspond to a C14 Laves phase and two hexagonal phases. Figure S1 in Supporting Information (SI) shows a detailed fitting with the Rietveld refinement method of the diffractogram recorded at 850 °C. It is clear that the peaks cannot be attributed to a single hcp phase. The two phases have different unit cell parameters; for example, at 780 °C, one phase (Hf−Zr-rich) has unit cell parameters of 3.2319(5) and 5.1541(9) Å for a and c, respectively, while the other hcp phase (Ti-rich) exhibits cell parameters of 3.1356(9) and 5.0044(9) Å for a and c, respectively. This suggests that one phase is Ti-rich and the second is more rich in Hf and Zr. The variation of the phases’ weight fractions as a function of temperature is shown in Figure 6b. It should be noted that the phase transformation kinetics of an alloy is dependent on the heating rate. However, the results in Figure 6 can be still used as a guideline to roughly evaluate the stability range of the as-synthesized bcc phase. The phase transformation of a single solid solution bcc phase to a mixture of bcc, hcp, and C14 Laves phase causes an increase in the cell parameter of the bcc phase from 3.4075(5) Å at approximately 600 °C to 3.4257(3) Å at 780 °C. This can be explained from the calculated compositions of the phases shown in Figure 2 as well as the phase fraction curve in Figure 6b. As can be seen, a significant amount of the C14 Laves phase is formed during annealing. This phase has the composition V2(Hf, Zr, Nb) and leads to a reduced V-content in the remaining bcc phase. Vanadium is the smallest atom, which leads to an increase in the average atomic radius of the remaining bcc phase and therefore an increase of the cell parameter. 4.3. Equilibrium Phases in the HfNbTiVZr System. To experimentally determine the equilibrium phases, as-synthesized HfNbTiVZr samples were heat-treated for 48 h at 600, 800, 1000, and 1200 °C. After annealing, the samples were rapidly quenched to room temperature. Figure 7 shows diffractograms from the annealed and quenched samples. As can be seen, the phase content in the annealed samples is in general agreement with the thermodynamic calculations in

Figure 4. Calculated Gibbs free energy, enthalpy, and entropy contribution for the bcc phase at equiatomic composition as a function of temperature. The SER state is taken as the reference state for the elements.

of the alloy in this condition is also shown in Figure 5. As can be seen from EDS, the alloy exhibits a single-phase structure with coarse grains. Additionally, it is clear that the elements are distributed homogeneously without apparent segregation. The unit cell parameter for the bcc phase can be determined to 3.377(2) Å, which is in good agreement with an expected cell parameter of 3.371 Å (assuming that we have a randomly ordered HEA and that Vegard’s law can be applied). The fact that the alloy exhibits a single-phase bcc solid solution after arc-melting synthesis is not surprising since similar results have been observed by other groups.12,16,17 From the diffractogram and the EDS results in Figure 5, it is easy to draw the incorrect conclusion that HfNbTiVZr is a true HEA at low temperatures. However, the thermodynamic calculations above suggest that a mixture of a bcc, hcp, and Laves C15 phases should be the most stable phases at low temperatures and that the single-phase alloy observed in Figure 5 therefore is metastable and a result of slow kinetics during cooling. Consequently, a phase transformation to a more stable structure is expected at higher temperatures. The thermal stability of the metastable as-synthesized alloy was therefore investigated by in situ annealing experiments at the German electron synchrotron (DESY) in Hamburg, Germany. The outcomes of these measurements are shown

Figure 5. (a) XRD pattern and (b) EDS elemental map of the HfNbTiVZr alloy in the as-synthesized condition. λ = 1.5418 Å. E


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Inorganic Chemistry

Figure 6. (a) Densitometric view of the powder X-ray diffractograms recorded in situ during heating between room temperature (RT) and 850 °C with 15 °C/min and (b) variation of phase content (wt %) calculated using the Rietveld method. λ = 0.207 Å.

and steel ampules). However, in all cases they exhibit very low intensities stemming from low phase fractions which can be regarded to have a negligible influence on the rest of the microstructure. The observed cell parameter of the quenched bcc sample (1200 °C) in Figure 7 is 3.368(2) Å, which is almost identical to the cell parameter of 3.377(2) Å for the as-synthesized sample in Figure 5. Additionally, these values are in good agreement with the value calculated through Vegard’s law, 3.371 Å (see section 4.2). The hcp phase at 600 °C has cell parameters a = 3.202(2) Å and c = 5.092(2) Å. This is close to what is expected from an hcp solid solution phase dominated by Zr and Hf (a = 3.206 Å and c = 5.087 Å for Hf0.5Zr0.534). The thermodynamic calculations in Figure 2 indeed suggest that the hcp phase should be dominated by these elements. The results above show that a multiphase mixture of bcc, hcp, and a Laves phase is formed during the annealing of the as-synthesized alloy. The XRD data shows that the Laves phase formed is hexagonal (C14), while a cubic C15 Laves phase is predicted in the thermodynamic calculations. The calculations also predicted that the Laves phase should have the composition V2(Hf, Zr). To clarify the structure and composition of the Laves phase formed during annealing, a more detailed investigation was done using TEM. According to the XRD data, this sample should contain precipitates of the C14 Laves phase. Figure 8 shows that such precipitates are indeed formed with an anisotropic shape with a width to height ratio of about 2:1. EDS shows that they are enriched in V but contain no or very little Ti. In addition, the EDS maps shows the presence of Hf and some Zr and Nb. The EDS maps suggest that the volume density of Hf is about the same in the matrix and the precipitates, while the Nb and Zr contents are lower in the precipitates than in the surrounding bcc matrix. Consequently, it can be confirmed that the Laves phase has the composition V2(Hf, Zr, Nb). Electron diffraction of the precipitates give diffraction spots that can be attributed to a hexagonal structure, i.e., C14. No coherence can be observed between the precipitates and the bcc matrix. Within the precipitates, a lamellar microstructure is often observed (Figure 8b), which could be attributed to defects such as stacking faults. Previous reports on the transformation from the C15 to C14 phase in the V2Hf system at 1000 °C indicate

Figure 7. X-ray diffractograms of the HfNbTiVZr alloy after 48 h heat treatments at 600, 800, 1000, and 1200 °C followed by a rapid quenching. λ = 1.5418 Å.

Figure 1. Annealing experiments for longer times showed the same results as the 48 h experiments. At 600 °C a phase mixture of bcc, hcp, and a hexagonal C14 Laves phase is observed. As described above, the thermodynamic calculation suggests the presence of a bcc, a hcp, and a cubic C15 Laves phase at this temperature. When the annealing temperature is increased to 800 °C, the hcp phase disappears. A further increase to 1000 and 1200 °C leads to a more or less singlephase bcc phase, in agreement with the thermodynamic predictions in Figure 1. The experimental results confirm that HfNbTiVZr at high temperatures can be considered to be a HEA, but that the formation of a single-phase bcc solid solution at room temperature is due to slow kinetics for the formation of the Laves and hcp phases. It should be noted that Figure 7 shows a few extremely weak peaks between 2θ 30° and 40° for the samples annealed at 1200 °C. These peaks are different depending on the method used for the preparation of the samples for the heat treatment (heat treatments in quartz F


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Inorganic Chemistry

Figure 8. (a) Z-contrast image of the HEA heat-treated at 1000 °C for 1 h with slow cooling, showing precipitates that range to several hundreds of nanometers in size. The insets show the elemental distribution of the different elements in the HEA. (b) A lamellar structure can be seen within the precipitate in the bright field image. (c) The diffraction pattern of the precipitate oriented in the [3̅ 2̅ 1] zone axis.

Figure 9. XRD patterns after 1 h heat treatments at 600, 800, and 1000 °C. (a) Fast cooling (water quenching) and (b) slow cooling (5 °C/min) from the corresponding temperatures. A heating rate of 5 °C/min was used to reach the respective temperature in each case. λ = 1.5418 Å.

relative intensities of the C14 Laves phase compared to the bcc phase are higher after 1 h compared to those after 48 h, suggesting that the amount of Laves phase is reduced with time. This can be understood from the heating procedure. The heating rate for this sample is about 5 °C/min, which is slower than that for the in situ study in Figure 6. Consequently, during the heating procedure a significant amount of Laves phase and hcp phase will be formed from the bcc solid solution. As the heating continues, the relative phase fraction of the hcp and Laves phase (see Figure 1) will decrease in a reversed process, where they are dissolved to form a solid solution HEA with a bcc structure. This process is kinetically limited. The hcp phase is most likely rapidly dissolved while the dissolution process of the Laves phase is slower. After 1 h, a significant amount still remains, while a phase content closer to equilibrium is attained after 48 h. An additional complexity is observed when different cooling rates after annealing are considered. This is illustrated in Figure 9b showing diffractograms from samples annealed for 1 h at

that the Nb content plays a major role in stabilizing the C14 Laves phase, which is consistent with our results.35 4.4. Influence of Annealing Time and Cooling Rates on Phase Composition. The results above clearly show that HfNbTiVZr alloy is a single-phase HEA at high temperatures, but only metastable as a solid solution at temperatures below 830 °C. Starting with an as-synthesized metastable solid solution phase, a time-dependent phase transformation where kinetics will influence the phase content at different annealing times is expected. This is illustrated in Figure 9a showing diffractograms for samples annealed for only 1 h at 600, 800, and 1000 °C. As can be seen, they show clear differences compared with the corresponding diffractograms in Figure 7 acquired after 48 h. After 1 h at 600 °C, the sample consists of mainly bcc and a small amount of Laves phase but with no indication of an hcp phase. The 600 °C diffractogram in Figure 9a looks very similar to the results obtained by Fazakas et al.17 The diffractogram acquired after 1 h at 800 °C also looks very different compared to that for the 48 h sample. The G


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Inorganic Chemistry 600, 800, and 1000 °C followed by a much slower cooling rate (5 °C/min). Starting with the sample annealed at 600 °C for 1 h, a completely different result from a rapidly quenched 600 °C sample can be seen. The slow cooling rate seems to lead to the formation of two bcc phases, in line with the thermodynamic calculations in Figure 1. This is not observed in the quenched sample. Also, the slowly cooled 800 °C sample shows a different phase content compared to that of the quenched sample. In this case a significant amount of the hcp phase is formed, most likely due to an increased time in the multiphase region during cooling, allowing for this phase to form. Additionally, a clear difference can be seen when comparing samples heat-treated for 1 h at 1000 °C; besides the bcc phase, in the slowly cooled sample a hcp and C14 Laves phase can be seen, while the quenched sample shows only C14 besides the bcc phase. It is noted that some additional peaks can be seen in the diffractogram of the samples slowly cooled from 800 and 1000 °C (approximately at 29.3°, 36°, and 36.7° in 2θ); these unidentified peaks most probably correspond to intermediate phases in the quinary system (the system has most likely not reached thermodynamic equilibrium due to the heating procedure). The results in Figure 9 emphasize an important factor in the interpretation of XRD data from HEA systems: The rates of some phase transformations can be very slow in these complex alloy systems, and too short annealing times or too slow cooling rates can give misleading results. This is of course valid for all alloy systems where phase transformations occur. However, HEA systems with five or more elements can have slower kinetics due the complexity of the system where different phases with various structures and compositions are formed. In the HEA literature, many annealing studies are fairly short, and the cooling rate after the heat treatment is not always given. The present results suggest that these are important factors to consider. 4.5. Influence of Sample Preparation on Phase Composition. An additional parameter which can affect the results is the sample preparation. After synthesis, the solid pieces obtained after arc-melting and annealing can be analyzed with XRD as a solid piece. A disadvantage with this procedure is that the texture of the grains can affect the interpretation of the XRD data. This can be avoided by a mortar grinding step where the as-synthesized sample is crushed into powder. Such a step, however, adds extra energy to the system allowing for additional phase transformations, addition of defects, and stresses. During our work with the HfNbTiVZr system, it has been observed that mortar grinding of this alloy significantly affected the results. This is demonstrated in Figure 10 showing the diffractograms from an as-synthesized sample after 1 h annealing at 800 °C followed by a rapid water quenching. If this sample is analyzed as a solid piece, a diffractogram with a bcc and a C14 Laves phase (initial condition, no mortar grinding) is observed. If the sample is ground for 10 min, a quite different diffractogram is observed. In this case, only the two strongest Laves phase peaks at 36.7° and 40° can be seen. In fact, they are the expected peaks from the C15 Laves phase. There are two possible explanations for this behavior: (i) a mechanically induced transformation of the C14 phase to C15 caused by the additional energy added during the mortar grinding process and (ii) reduction in grain size and addition of stresses to the systems caused line broadening and led to broader and less intense peaks (making it difficult to separate C14 from C15).

Figure 10. XRD for different sample conditions. Bottom: XRD pattern after 800 °C for 1 h followed by water quenching. Middle: XRD pattern after pulverization by mortar grinding. Top: XRD pattern after heat treatment for 30 days at 400 °C. λ = 1.5418 Å.

Without additional TEM studies these changes in the diffractograms cannot be explained. Interestingly, however, a subsequent annealing of the ground sample at 400 °C for 30 days regenerates the typical peaks for the C14 Laves phase. This can be due to a stress release leading to sharper peaks and/or a phase transformation to a thermodynamically more stable condition.

5. CONCLUDING REMARKS We have investigated the HfNbTiVZr system with different annealing and cooling procedures combined with thermodynamic calculations using a CALPHAD approach. The most important conclusions are the following: 1. The HfNbTiVZr alloy is a single-phase HEA with a bcc structure above about 830 °C. This can be concluded from the thermodynamic modeling and is confirmed by the 48 h annealing experiments. At lower temperatures the thermodynamically stable phase composition is a mixture of bcc, hcp, and a Laves phase. 2. The CALPHAD description of the HfNbTiVZr system extrapolated from the binary subsystems seem to be an adequate approximation although further work is needed for a more complete picture of the equilibrium structures of this system, in particular when it comes to thermodynamic properties of the Laves phases in the ternary subsystems. 3. The as-synthesized alloy showing only one solid solution phase in Figure 5 is metastable at room temperature and stabilized by slow kinetics of the equilibrium phases. Experimental results from as-synthesized samples showing a solid solution at room temperature are not enough to define a multicomponent alloy system as a single-phase high-entropy system. Annealing experiments and/or thermodynamic modeling are required for such conclusions. It should be noted that many papers in the literature identifying new HEA systems are based on only as-synthesized samples. A semantic question is of course if an alloy which is thermodynamiH


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Inorganic Chemistry

(6) Zou, Y.; Maiti, S.; Steurer, W.; Spolenak, R. Size-Dependent Plasticity in an Nb25Mo25Ta 25W25 Refractory High-Entropy Alloy. Acta Mater. 2014, 65, 85−97. (7) Fritze, S.; Malinovskis, P.; Riekehr, L.; von Fieandt, L.; Lewin, E.; Jansson, U. Hard and Crack Resistant Carbon Supersaturated Refractory Multicomponent Nanostructured Coatings. Sci. Rep. 2018, 8, 14508. (8) Qiu, Y.; Thomas, S.; Gibson, M. A.; Fraser, H. L.; Birbilis, N. Corrosion of High Entropy Alloys. npj Mater. Degrad. 2017, 1, 15. (9) Yeh, J. W.; Chen, S. K.; Lin, S. J.; Gan, J. Y.; Chin, T. S.; Shun, T. T.; Tsau, C. H.; Chang, S. Y. Nanostructured High-Entropy Alloys with Multiple Principal Elements: Novel Alloy Design Concepts and Outcomes. Adv. Eng. Mater. 2004, 6, 299−303. (10) Malinovskis, P.; Fritze, S.; Riekehr, L.; von Fieandt, L.; Cedervall, J.; Rehnlund, D.; Nyholm, L.; Lewin, E.; Jansson, U. Synthesis and Characterization of Multicomponent (CrNbTaTiW)C Films for Increased Hardness and Corrosion Resistance. Mater. Des. 2018, 149, 51−62. (11) Sahlberg, M.; Karlsson, D.; Zlotea, C.; Jansson, U. Superior Hydrogen Storage in High Entropy Alloys. Sci. Rep. 2016, 6, 36770. (12) Karlsson, D.; Ek, G.; Cedervall, J.; Zlotea, C.; Møller, K. T.; Hansen, T. C.; Bednarčík, J.; Paskevicius, M.; Sørby, M. H.; Jensen, T. R.; Jansson, U.; Sahlberg, M. Structure and Hydrogenation Properties of a HfNbTiVZr High-Entropy Alloy. Inorg. Chem. 2018, 57, 2103− 2110. (13) Zhang, Y.; Zhou, Y. J.; Lin, J. P.; Chen, G. L.; Liaw, P. K. SolidSolution Phase Formation Rules for Multi-Component Alloys. Adv. Eng. Mater. 2008, 10, 534−538. (14) Cantor, B.; Chang, I. T. H.; Knight, P.; Vincent, A. J. B. Microstructural Development in Equiatomic Multicomponent Alloys. Mater. Sci. Eng., A 2004, 375−377, 213−218. (15) Couzinié, J. P.; Dirras, G.; Perrière, L.; Chauveau, T.; Leroy, E.; Champion, Y.; Guillot, I. Microstructure of a near-Equimolar Refractory High-Entropy Alloy. Mater. Lett. 2014, 126, 285−287. (16) Feuerbacher, M.; Lienig, T.; Thomas, C. A Single-Phase Bcc High-Entropy Alloy in the Refractory Zr-Nb-Ti-V-Hf System. Scr. Mater. 2018, 152, 40−43. (17) Fazakas, E.; Zadorozhnyy, V.; Varga, L. K.; Inoue, A.; Louzguine-Luzgin, D. V.; Tian, F.; Vitos, L. Experimental and Theoretical Study of Ti20Zr20Hf 20Nb20 × 20 (X = v or Cr) Refractory High-Entropy Alloys. Int. J. Refract. Hard Met. 2014, 47, 131−138. (18) Lee, C.; Song, G.; Gao, M. C.; Feng, R.; Chen, P.; Chen, Y.; An, K.; Guo, W.; Poplawsky, J.; Li, S.; Samaei, A. T.; Chen, W.; Hu, A.; Choo, H.; Liaw, P. K.; Brechtl, J. Lattice Distortion in a Strong and Ductile Refractory High-Entropy Alloy. Acta Mater. 2018, 160, 158. (19) Chen, S. Y.; Tong, Y.; Tseng, K. K.; Yeh, J. W.; Poplawsky, J. D.; Wen, J. G.; Gao, M. C.; Kim, G.; Chen, W.; Ren, Y.; Feng, R.; Li, W. D.; Liaw, P. K. Phase Transformations of HfNbTaTiZr HighEntropy Alloy at Intermediate Temperatures. Scr. Mater. 2019, 158, 50−56. (20) Munitz, A.; Salhov, S.; Hayun, S.; Frage, N. Heat Treatment Impacts the Micro-Structure and Mechanical Properties of AlCoCrFeNi High Entropy Alloy. J. Alloys Compd. 2016, 683, 221− 230. (21) Otto, F.; Dlouhý, A.; Pradeep, K. G.; Kuběnová, M.; Raabe, D.; Eggeler, G.; George, E. P. Decomposition of the Single-Phase HighEntropy Alloy CrMnFeCoNi after Prolonged Anneals at Intermediate Temperatures. Acta Mater. 2016, 112, 40−52. (22) Yurchenko, N.; Stepanov, N.; Salishchev, G. Laves-Phase Formation Criterion for High-Entropy Alloys. Mater. Sci. Technol. 2017, 33, 17−22. (23) Hansen, B. R. S.; Møller, K. T.; Paskevicius, M.; Dippel, A. C.; Walter, P.; Webb, C. J.; Pistidda, C.; Bergemann, N.; Dornheim, M.; Klassen, T.; Jørgensen, J. E.; Jensen, T. R. In Situ X-Ray Diffraction Environments for High-Pressure Reactions. J. Appl. Crystallogr. 2015, 48, 1234−1241. (24) Hammersley, A. P. FIT2D: An Introduction and Overview. Eur. Synchrotron Radiat. Facil. Int. Rep. ESRF97HA02T. 1997, 68, 58.

cally stabilized as a solid solution by configurational entropy at high temperatures can also be called an HEA at metastable conditions at room temperature. 4. The phase composition of the HfNbTiVZr alloy is strongly dependent on the annealing time and the cooling rates after the annealing experiments. Our results suggest that the kinetics can be slow and that correct interpretation of the equilibrium phase content may require extensive annealing studies with well-defined cooling rates. This may seem to be an obvious statement, but it should be noted that few detailed studies of this have been published in the literature and that many studies are published without annealing experiments or with rather short annealing times. Finally, we conclude that the preparation method (e.g., use of mortar grinding) may affect the results giving misleading or erroneous results. In particular, this is important in studies of metastable materials.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.inorgchem.8b02957.

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XRD pattern of HfNbTiVZr at 850 °C refined with the Rietveld method (PDF)

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Victor Pacheco: 0000-0001-8500-1632 Dennis Karlsson: 0000-0002-5511-5986 Johan Cedervall: 0000-0003-0336-2560 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This research is funded by the Swedish Foundation for Strategic Research (SSF) within the Swedish national graduate school in neutron scattering (SwedNess) and by the Swedish Foundation for Strategic Research (SSF), through the project “SSF − Development of processes and Materials in AM”. Parts of this research were performed at the beamline P02.1 at the German electron synchrotron (DESY), Hamburg. The authors wish to thank Dr. Jo-Chi Tseng for her help and support during these experiments.

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REFERENCES

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Thermal Stability of the HfNbTiVZr High-Entropy Alloy - Inorganic

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