Design of Nb-based multi-phase alloy membranes for high hydrogen permeability and suppressed hydrogen embrittlement

Journal of Membrane Science xxx (xxxx) xxx

Contents lists available at ScienceDirect

Journal of Membrane Science journal homepage: http://www.elsevier.com/locate/memsci

Design of Nb-based multi-phase alloy membranes for high hydrogen permeability and suppressed hydrogen embrittlement Erhu Yan a, d, *, R.N. Min a, P. Zhao b, **, R.D.K. Misra c, P.R. Huang a, Y.J. Zou a, H.L. Chu a, H. Z. Zhang a, F. Xu a, L.X. Sun a, *** a

Guangxi Key Laboratory of Information Materials, Guilin University of Electronic Technology, Guilin, 541004, PR China College of Materials Science and Engineering, Qingdao University of Science & Technology, Qingdao, 266044, China Department of Metallurgical, Materials and Biomedical Engineering, University of Texas, El Paso 500 W. University Avenue, El Paso, TX, 79968-0520, USA d Department of Energy, Materials and Telecommunications, INRS-EMT, Quebec, J3X 1S2, Canada b c

A R T I C L E I N F O

A B S T R A C T

Keywords: Nb-HfCo alloys Phase diagram Hydrogen permeability DBTC

Nb-Hf-Co alloys, consisting of a multi-phase microstructure with primary BCC-(Nb, Hf) and eutectic {BCC-(Nb, Hf) þ HfCo} phases, are a potential alloys for hydrogen separation and purification. However, the compositional window, comprising of the above phases, has not yet been established because of lack of important information concerning the phase diagrams. To address this need, this study first calculated the phase diagram and related phase equilibria information of this class of alloys, by using the CALPHAD approach, and established a new compositional window suitable for hydrogen permeation. The alloys in this window exhibit very high hydrogen permeability. In particular, Nb45Hf27.5Co27.5 exhibits the highest permeability of 5.32 � 10 8 mol H2 m 1s 1Pa 0.5 at 673 K, which is the best among all the Nb-based multi-phase alloys. Secondly, the boundary for ductile to brittle transition-hydrogen concentration (DBTC) of the Nb-HfCo multi-phase alloy was determined for the first time and found to occur at around 0.72–0.82 H/M, which reveals the fundamental reasons for brittle failure occurring during hydrogen permeation. From the point of DBTC, a new design concept for Nb-based hydrogen permeable membranes is proposed and the factors affecting their brittle failure are analysed. Thirdly, to verify the feasibility of the above concept, Nb45Hf27.5-xTixCo27.5 (x ¼ 0 … 10) alloys were designed and comprehensively studied. It is demonstrated that Nb45Hf20Ti7.5Co27.5 alloy exhibits excellent hydrogen permeability with enhanced hydrogen embrittlement under appropriate permeation conditions. The significance of the study is that the best composition range for hydrogen-permeable Nb-HfCo alloys is explored, and some novel hydrogen-permeable alloys can be easily developed by analysing their absorption pressure-composition isotherms in combination with the newly explored DBTC.

1. Introduction Nowadays, hydrogen-selective alloy membranes are promising ma­ terials for hydrogen purification technology because of their good strength and nearly infinite selectivity [1–4]. However, the currently used palladium-based alloys [5,6] are expensive, which severely limits their large-scale applications. Considering the low-cost and higher hydrogen permeability, group 5B metals such as vanadium (V), niobium (Nb), and tantalum (Ta) have attracted significant attention [1,4,7–10]. Nevertheless, in pure form, these metals suffer severely from hydrogen

embrittlement (HE) due to their high hydrogen solubility [8–12]. The solution to the HE problem is to alloy with other metals, such as Ti, Ni, Co, Zr, Hf, W, Mo, Cu or a combination of them [1,2,8,11–16]. Corre­ spondingly, some new hydrogen permeable alloys, such as Nb-Ti-Ni [17, 18], Nb-Ti-Co [19], Nb-Ti-Fe [20], Nb-Zr-Ni [21] and Nb-Hf-Ni [12], have been recently developed. Based on the studies carried out on different alloys, it is noted that the studied alloys are characterized by a multi-phase microstructure, i. e., primary BCC-Nb and eutectic (BCC–Nb þ TiNi/TiCo/HfNi/ZrNi/ TiFe). The former mainly controls the hydrogen permeability, whereas

* Corresponding author. Guangxi Key Laboratory of Information Materials, Guilin University of Electronic Technology, Guilin, 541004, PR China. ** Corresponding author. *** Corresponding author. E-mail addresses: [email protected] (E. Yan), [email protected] (P. Zhao), [email protected] (L.X. Sun). https://doi.org/10.1016/j.memsci.2019.117531 Received 1 August 2019; Received in revised form 23 September 2019; Accepted 4 October 2019 Available online 7 October 2019 0376-7388/© 2019 Elsevier B.V. All rights reserved.

Please cite this article as: Erhu Yan, Journal of Membrane Science, https://doi.org/10.1016/j.memsci.2019.117531

E. Yan et al.

Journal of Membrane Science xxx (xxxx) xxx

Fig. 1. Solidification and microsegregation of Nb-Hf-Co alloys. (a) Schematic illustration of the melting apparatus. Inset: the appearance of as-prepared Nb30Hf35Co35 sample; (b) Schematic illustration of the dendritic growth in the three-dimensional space; (c) Solute redistribution at the S/L interface (i.e. the po­ sition Q in (b)) during the dendritic growth; (d) Five basic solidified morphology considered in the present work. (i)–(v) represent spherical, cylindrical, plate-like, inward-cylinder and inward-spherical dendrite, respectively, which correspond to the location shown in Fig. 1(b).

the latter imparts embrittlement resistance [17–21]. Thus, the micro­ structure can be tailored by modifying the chemistry and enhance the properties. However, there is a continued challenge of compositional window with multi-phase structures in different alloys. Another problem is that the ternary phase diagram of these alloy systems is incomplete [1, 13,22,23]. Studies are required to explore the desired compositions. Hence, it is necessary to determine phase diagrams and related phase equilibria, for each alloy system. In a previous study on Nb-Hf-Co system [24,25] it was shown that alloys with >30 at.% Nb, containing similar multi-phase structure, exhibited excellent permeability with enhanced HE. The high permeability of these alloys was attributed to their high hydrogen solubility, which renders them as an alternative to Pd-based alloys. However, the compositional window in this alloy series, appro­ priate for hydrogen permeation, has not been studied. It is still unknown whether the alloys exhibit the desired properties (e.g. higher perme­ ability), and requires further study. In addition, it is widely accepted that the mechanical properties are intrinsically related to the concentration of absorbed hydrogen in the alloys [1,16,20,21,26–28]. Meanwhile, degradation of the resistance to HE, due to hydrogen absorption, has been widely reported in the above 5B group metals and their alloys [29–32]. To address this aspect, the DBTC for Nb-based single-phase alloys (e.g. Nb–5W, Nb–5Ru and Nb–5W-(5–10)Mo etc.) was investigated by Yukawa et al. [33,34]. They observed that DBTC occurred at a concentration level of ~0.23 H/M (atomic hydrogen to metal ratio), which is the upper limit for dissolved hydrogen in these alloys. If this value is exceeded, HE will inevitably occur during hydrogen permeation. A similar case was observed in V-based single-phase alloys (e.g. V–W, V–Mo and V–Co [35,36]) and the upper critical concentration limit was ~0.2 H/M. However, numerous studies have shown absence of cracking or brittle fracture, due to HE occurs in Nb-based multi-phase alloys such as Nb40Ti30Ni30 [17,18,37], (Nb35W5)Ti30Ni30 [38], Nb35Mo5Ti30Ni30 [39] and Nb40Hf30Ni30 [12,

40,41] etc., even at dissolved hydrogen concentration of as high as 0.6. A question arises if there is a fundamental upper limit of concentration, in these multi-phase alloy membranes, to prevent them from brittle failure. For this reason, Dolan et al. [42] investigated the DBTC of V–Ti–Ni multi-phase alloy for the first time and found that its value was in the range of 0.35–0.45 H/M, which is significantly higher than that in single-phase ones (~0.2 H/M). Similar to the above analysis of DBTC in V-based alloys, there may be a big difference in the DBTC values be­ tween Nb-based single-phase and multi-phase alloys. A number of studies [33,34,43] have been made on Nb-based single-phase alloys in terms of DBTC and improvement in resistance to HE. However, the DBTC of the multi-phase alloys remains unclear, and the relevant brittle fracture problem, closely related with DBTC, has not been addressed. Considering the above discussion, we first studied the phase diagram of the Nb-Hf-Co alloy system using CALPHAD, and the solidification path, microstructure, and hydrogen permeation characteristics, of different alloys were studied. This contributes to establishing a new compositional window suitable for hydrogen permeation, and to develop novel hydrogen-selective alloy membranes with excellent desired properties. Secondly, the DBTC boundary of the Nb-Hf-Co multiphase alloys was determined from a series of the hydrogen permeation/ absorption tests, as well as the in-situ small punch (SP) test, and a new design concept for Nb-based multi-phase hydrogen permeable mem­ branes is proposed, from the point of view of the newly established DBTC. Lastly, using this concept, Nb45Hf27.5-xTixCo27.5 (x ¼ 7.5, 10) hydrogen-permeable alloys, which possess good resistance to HE and excellent hydrogen permeability, were designed and successfully developed. In brief, the present study introduces a new strategy for designing and developing Nb-based alloy membranes with improved HE resistance and durability, while minimising the associated permeability penalty.

2

E. Yan et al.

Journal of Membrane Science xxx (xxxx) xxx

Table 1 The values of the geometric weight coefficient vector, γ(fS) ¼ γ(γ1, γ2, γ3, γ4, γ5), used in the present calculations of solidification paths.

2. Experimental and numerical procedures

model, reported by us, was adopted in the calculation process (see Ref. [44]). This model considers the dynamic changes in dendrite dimension/morphology (which is assumed to be spherical, cylindrical, plate-like, inward-cylinder, and inward-spherical dendrite, see Fig. 1 (d)), and also includes the SBD effects, solidification rate, etc., which covers all possible influential factors affecting the solidification process. For the solidified morphology in Fig. 1(d), their corresponding values of the geometric weight coefficient vector, γ(fS) ¼ γ(γ 1, γ 2, γ3, γ4, γ5), are shown in Table 1. In addition, four different solidification patterns may occur in the present alloy system, according to the analysis of the Thermo-Calc calculation results (described later). They are: (a) primary phase solid­ ification, (b) binary eutectic reaction, (c) ternary eutectic reaction, and (d) ternary quasi-peritectic reaction. All possible solidification reactions were considered in the construction of the iterative calculation program. In each calculation iterative step, the number, variety, and names of the phases are obtained through the Thermo-Calc software, and then compared with the results of the previous step. Correspondingly, the liquid composition and solidification temperature of each alloy were calculated and saved. The solid volume fraction, fS, is taken as the control variable in the primary phase or binary eutectic solidification, while the temperature T was the control variable in the ternary eutectic/ quasi-peritectic solidification reaction.

2.1. Theoretical method and model The Nb-Hf-Co phase diagram was investigated by using the CAL­ PHAD method. Each solution phase is treated as the substitutional so­ lution model with Redliche-Kister formula. During the calculation process, the Gibbs free energy of all possible phases was calculated at a given temperature, pressure, and composition. The solidification paths of various alloys were calculated using the software Thermo-Calc, coupled with the micro-segregation model. Thermodynamic data (e.g. liquid/solid surface temperature (TL/TS), phase compositions (CL/CS), partition coefficient (kL/kS)) were obtained by FORTRAN calculation program coupled with Thermo-Calc TQ6-interface. Generally, the leading phase (marked as point P in Fig. 1 (a)) tends to nucleate and grow in the mushy region. During its growth (see Fig. 1(b)), solute redistribution occurs at the S/L interface, Fig. 1(c), which can be approximately described by the classical Scheil-Gulliver equation. But the results calculated by this model produce a large error because of assumptions such as no diffusion in the solid phase, and infinitely fast diffusion in the liquid phase, at different temperatures. Similar problem exists on using other models such as the Lever-rule equation, and the Clyne-Kurz equation, which ignore the dendrite morphology, solidifi­ cation rate, or the extent of solid-back diffusion (SBD). All these pa­ rameters can influence the solidification behaviour such as micro-scale solute redistribution, and solidification path, especially during the later stages of the solidification process. So, the unified micro-segregation 3

E. Yan et al.

Journal of Membrane Science xxx (xxxx) xxx

Table 2 Invariant reactions with liquid phases of Nb-Hf-Co system in Nb-rich region by CALPHAD method. Number

U1 E1 E2

Invariant reactions

L þ HfCo ↔ (Nb,Co) þ Nb6Co7 L ↔ (Nb,Hf) þ Nb6Co7 þ (Nb,Co) L ↔ (Nb, Hf) þ HfCo þ Hf2Co

Reaction type

Temperature (� C)

Compositions of the liquid phases (at.%) x (Nb)

x (Hf)

x (Co)

II

1432.3

38.5

22.4

39.1

I

1065.8

52.5

5.9

41.6

I

1321.1

8.9

57.8

33.3

673 K. Permeability at each temperature was calculated by Eq. (1), � � Φ P0:5 P0:5 D⋅K P0:5 P0:5 u d u d ¼ (1) J¼ L L

Fig. 2. The Nb-Hf-Co ternary alloy compositions used in this study is plotted on the calculated liquidus projections using the CALPHAD approach. These alloys are divided into four sub-regions marked by solid triangles (denominated by I), solid squares (II), solid circles (III) and solid diamonds (IV) according to their characteristics in microstructures, solidification paths and hydrogen transport behaviours (described later). The blue number, in brackets above the solid circles, is the hydrogen permeability value (in 10 8 mol H2 m-1s-1Pa-1/2 unit) for each alloy at 673 K. Similarly, the blue NG in brackets indicates that the hydrogen permeability of alloys cannot be measured due to HE. (For inter­ pretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)

Where L represents the thickness of the membrane, D and K represent the hydrogen diffusivity and solubility, respectively, and their product is defined as the hydrogen permeability, i.e. Φ ¼ D � K. The hydrogen absorption of several alloys was measured via the Sieverts’ method, described previously [45]. Before testing, Pd-coated samples of each alloy were activated by four absorbing-desorbing cy­ cles between 673 K and 300 K, at 2 MPa. Subsequently, their absorption isotherms were measured at 523–673 K, up to about 0.7 MPa, and the corresponding pressure-composition-temperature (PCT) curves, with varying hydrogen pressure, were recorded.

2.2. Experimental procedures 2.2.1. Membrane preparation and characterisation Nb-Hf-Co alloys (~35 g) were prepared using a vacuum selfconsuming arc melting furnace, under argon atmosphere, and were denoted as 1#, 2# … 50#, respectively (see Fig. 2). The Nb45Hf27.5xTixCo27.5 (x ¼ 0 … 10) alloys were prepared using identical melting method. The purity of the raw materials was greater than 99.95 mass%. Each sample was turned over and re-melted several times until homo­ geneity was obtained. The disc-type samples (Φ16 mm � 0.7 mm thick) were cut from the arc-melted ingots using electrical discharge wire machining, and ground and polished using 1 μm alumina powders. Subsequently, a thin Pd coating, ~200 nm thick, was deposited on the surface of these samples, by using a magnetron sputtering apparatus for H2 dissociation and reassociation (H2 ↔ 2e- þ 2Hþ). The pre-permeation alloy microstructure was observed using scan­ ning electron microscopy (FEI Quanta 200) in back-scattered electron mode (BSE). Surfaces were rinsed using alcohol prior to analysis. Quantitative composition analysis of the phases was performed by the energy dispersive X-ray spectrometer (EDX). XRD patterns of the sam­ ples were recorded using a Cu Kα radiation source (λ ¼ 1.5418 Å) operating at 40 kV and 40 mA over the 2-theta range 20–90� , with a step size of 0.05� .

2.2.3. In-situ SP test The mechanical performance of several alloys were measured using the in-situ SP test method. Prior to the measurements, the samples (8 mm � 8 mm � 0.7 mm) were sectioned using electrical discharge machining, and then polished mechanically and a Pd coating layer was deposited, using the experimental procedure described above. The sample was placed in the apparatus and then it was evacuated and heated to 523 K. Subsequently, hydrogen gas was introduced into the apparatus and kept at 0.35 MPa for 30 min. Next, the specimen was punched by a steel ball (Φ3 mm) with a constant loading rate of 0.3 mm/ min, and the corresponding load-deflection curves were recorded. Detailed information of the test is documented in Ref. [46]. 3. Results and discussion 3.1. Calculation of phase diagram and solidification paths Fig. 2 shows the calculated liquidus projection of the ternary Nb-HfCo system. The primary BCC-(Nb, Hf) solidification region is surrounded by four different binary eutectic transformations, i.e., L→BCC ðNb; HfÞ þ BCC ¼ ðNb; COÞ, L→BCC ðNb; HfÞþNb6 Co7 , L→BCC ðNb; HfÞ þ HfCo, and L→BCC ðNb; HfÞ þ Hf 2 Co, in the order of increasing Hf content. These transformations include three ternary invariant trans­ formations, marked as U1, E1 and E2, respectively, and details of each reaction are summarised in Table 2. The L→BCC ðNb; HfÞ þ GfCo eutectic transformation has a saddle point “S” around 1475 � C and with decreasing temperature, it reaches two invariant transformations, i.e., U1 and E2 in lower and higher Hf regions, respectively. The brittle intermetallic compounds, such as Nb6Co7 and Hf2Co [24,25], can be formed after these two reactions, which make the alloy membranes prone to brittle fracture prior to usage. Therefore, the precipitation of these two brittle phases should be avoi­ ded during the solidification process. In addition, the experimental al­ loys in the BCC-(Nb, Hf) phase region are classified into four groups

2.2.2. Permeation and solubility testing Hydrogen permeation across the membranes was measured using constant pressure method. Disc-type samples were first sealed with two annealed copper gaskets and then placed into the permeation cell. Subsequently, the system was evacuated to 8 � 10-3 Pa, leak tested and purged with air three times with pure H2. Once the vacuum was ob­ tained, the sample was heated to 523 K. During the test, the upstream (Pu) sides of the membrane was supplied with a high purity H2 at 0.3–0.5 MPa pressure. Meanwhile, the pressure of downstream (Pd) sides was stabilized at 0.1 MPa. The steady-state H2 flux (J) was then recorded by a calibrated mass flow meter. After this procedure, the subsequent permeation experiments were conducted at 573 K, 623 K and 4

E. Yan et al.

Journal of Membrane Science xxx (xxxx) xxx

Fig. 3. Calculated solidification paths for Nb-Hf-Co alloys in the BCC-(Nb, Hf) phase region. (a) Solidification paths of typical alloys 1# (Nb90Hf5Co5), 18# (Nb55Hf20Co25), 28# (Nb45Hf27.5Co27.5) and 38# (Nb45Hf35Co20) in four regions. (b)–(e) The temperature vs. solid fraction (T-fs) for alloys from the regions I, II, III and IV, respectively.

(labelled as I, II, III and IV) depending on their characteristics in solid­ ification paths, crystal structures, and hydrogen permeation behaviours (discussed further below). The solidification paths and T-fs curves for the typical alloys in each region are shown in Fig. 3. The alloys (e.g. 1# or 2#) in region I, only undergo one transformation, L→BCC ðNb; HfÞ, see Fig. 3(a) and (b). In addition to the above transformation, two additional reactions, i.e., L→ BCC ðNb; HfÞþNb6 Co7 , and L→BCC ðNb; HfÞþNb6 Co7 þ BCC ðNb; CoÞ, occur for the alloys in region II, except for alloys 8# (Nb45Ti15Co40) and 15# (Nb40Ti20Co40), which is attributed to their particular location in the eutectic valley, as plotted in Fig. 3(a). For the alloys (e.g. 28# and 29#) in region III, the solidification begins with the formation of primary BCC ðNb; HfÞphase. Then the liquid composition

changes along the eutectic line and the temperature decreases, and the solidification is terminated on this line at about ~1320 � C, with the formation of a binary eutectic fBCC ðNb; HfÞ þHfCog microstructure, as seen in Fig. 3(a) and (d). Among these alloys, Nb30Hf35Co35 alloy (#31), as an exception, only experienced the latter transformation, which is consistent with previous results [24]. Besides, the volume fraction of primary BCC ðNb; HfÞ increases with the increase of Nb content (31#→30#→29#→28#→24#). Accordingly, a declining trend of the volume fraction of binary eutectic phase is observed. In contrast with region III, alloys (e.g. 38# and 39#) in region IV undergo the additional ternary eutectic reaction, L→BCC-(Nb, Hf) þ HfCo þ Hf2Co (see Fig. 3(e)) as well as the above two similar reactions of the alloys in region III. 5

E. Yan et al.

Journal of Membrane Science xxx (xxxx) xxx

Fig. 4. XRD diffractograms of the as-cast Nb-Hf-Co alloys. (a)–(d) represent the typical alloys in regions I, II, III and IV (as shown in Fig. 2), respectively, and the insets in (a) are the SEM micrograph and EDS result of alloy 1# (Nb90Hf5Co5).

Based on these simulated results, it may be noted that small changes in the alloy composition in the BCC ðNb; HfÞ phase region, may in­ troduces changes in the solidification path with a corresponding impact on the final solidification microstructure. Besides, the alloys in region III are most promising for the fabrication of the hydrogen permeable membranes. The ability to understand the phase equilibrium informa­ tion, and discern the solidification paths saves time and expense asso­ ciated with the experimental research, but also provides the basis for the optimisation design of Nb-based multi-phase alloys.

in Fig. 5(a)). The formation of the above three phases can be attributed to the ternary eutectic reaction at E1, as predicted by the calculated results in Fig. 3 (a) and (c). Alloys 15# and 8# are all entirely composed of eutectic {BCC-(Nb, Hf)þ Nb6Co7}, which can be attributed to their presence in the eutectic valley and their solidification terminates in the ternary eutectic E1, see Fig. 3(a). Similar analysis, by using XRD and SEM, was also performed for other alloys in the BCC-(Nb, Hf) phase region, as discussed above. As clearly illustrated in Fig. 4(c) and (d), the samples in region III consist only of BCC ðNb; HfÞ and Bf HfCo phases, whereas the additional Hf2Co Bragg peak appeared in the XRD patterns of the alloys in region IV. These results were further confirmed by SEM analysis, as shown in Fig. 5 (d)–(i). For alloys, 24# and 28# in region III, a multi-phase microstructure is observed, see Fig. 5 (d) and (e). Nevertheless, alloy 31# consists entirely of eutectic {BCC-(Nb, Hf) þ HfCo}, which shows a fine morphology with a branched tree-like shape, as illustrated in Fig. 5 (f). A similar hypoeutectic microstructure was seen for specimens (e.g. 42# and 38#) in region IV (see Fig. 5 (g)–(i)). Additionally, the ternary eutectic, consisting of BCC ðNb; HfÞ,HfCo and Hf 2 Co, solidifies as a divorced eutectic (marked by the blue square), which is distributed around the solidified binary eutectic (marked by the red ellipse). Based on this analyses, the sequence of solidification of the alloys in this region is clear: initially BCC(Nb, Hf) phase forms and grows in under-cooled alloy melts. Solidification proceeds further along the eutectic valley, with the formation of a binary eutectic, and finally terminates at eutectic point E2 (see Fig. 3(a)). In contrast, the samples in region III only un­ dergo the former two reactions and terminate in the eutectic valley. Besides, the volume fraction of binary eutectic fBCC ðNb; HfÞ þHfCog

3.2. Microstructure and hydrogen permeability 3.2.1. Microstructure To further confirm the above results, XRD/EDS analysis and micro­ structural observations were carried out for the typical alloys in each region of Fig. 2, as illustrated in Figs. 4 and 5. In Fig. 4 (a), onlyt a single BCC ðNb; HfÞ phase is identified and the composition of alloy 1# is Nb4.61 at.%Hf-2.81 at.%Co (see the inset in Fig. 4 (a)), indicating that the segregation of Hf and Co occurs during its solidification. For samples (e. g. 12#, 18#, and 15#) in region II, BCC-(Nb, Hf), BCC-(Nb, Co) and HCP-Nb6Co7 phases are identified, as shown in Fig. 4(b). Additionally, another unidentified phase (labelled ‘■‘), appears to exist, presumably due to the solute distribution during the non-equilibrium process of solidification. Further observation of the microstructure, in Fig. 5(a)–(c), it is clear that primary BCC ðNb; HfÞ dendrites, and multi-phase eutectic (red circle), appears in alloys 12# and 18#, surrounded by a small amount of ternary eutectic structure consisting of BCC-(Nb, Hf), BCC-(Nb, Co) and HCP-Nb6Co7 phases (marked as ‘1’, ‘2’ and ‘3’, respectively, see the inset 6

E. Yan et al.

Journal of Membrane Science xxx (xxxx) xxx

Fig. 5. BSE images of typical as-cast samples in regions II, III and IV. (a)–(c) represent typical alloys 12# (Nb65Hf15Co20), 18# (Nb55Hf20Co25) and 15# (Nb40Hf20Co40) in region II; (d)–(f) represent typical alloys 24# (Nb50Hf25Co25), 28# (Nb45Hf27.5Co27.5) and 31#(Nb30Hf35Co35) in region III; (g)–(i) represent typical alloys 42# (Nb40Hf40Co20), 38#(Nb45Hf35Co20) and 34# (Nb50Hf30Co20) in region IV. Inset in (a): the magnified ternary eutectic morphology corresponding to local region (the blue square).

increases with decreasing Nb content (e.g. 24#→28#→31#), and it reaches the maximum of ~100 vol% for alloy 31#, which corresponds with the calculated results in Fig. 3. Theoretical/experimental results both reveal that alloys in region III contain the desired, dual-phase, microstructural characteristics that are expected to be the most promising candidates for hydrogen purification. However, if HE fracture occurs in region III, alloys in this region would be unacceptable. To test this hypothesis, further investigation on hydrogen permeation behaviour is required. In this paper, a first attempt to explore this “compositional window” by measuring their hydrogen permeability and susceptibility to HE was conducted, and good results were achieved, which is discussed below.

[31]. In Fig. 6 (a) a linear relationship between J � L and ΔP0.5, for typical alloy 29#, is found at each temperature, and a similar linear relationship appears in other alloys. These results indicate that the hydrogen trans­ port reaction through these alloy membranes follows Eq. (1). Moreover, the Φ values increase with increasing temperature and Nb content (31#→30#→29#→28#), and this trend can be found in other alloy series such as Nb–Ti–Ni/Co and Nb–Zr–Ni/Co [17,18,31,47], as shown in Fig. 6 (b). Besides, the permeabilities of the present alloys are significantly higher than pure Pd, especially at 673 K. Notably, the Nb45Hf27.5Co27.5 alloy (28#) exhibited a hydrogen permeability of 5.32 � 10 8 mol H2 m 1s 1Pa 0.5 at 673 K, which is superior to that (5.21 � 10 8 mol H2 m 1s 1Pa 0.5) of the Nb40Hf32Co28 alloy [48], previously reported to be the best in this alloy series. Nevertheless, this alloy was prone to brittle failure at 573 K and below, probably because of excessive hydrogen absorption. In contrast, alloys 29#-31# can permeate continuously at each temperature, indicating that these alloys possess an improvement in the resistance to HE, which could be attributed to more eutectic phases in their microstructure. In addition, Nb-Hf-Co alloys possess a pronounced higher hydrogen

3.2.2. Hydrogen permeability In Fig. 6, the hydrogen transport properties, between 523 K and 673 K, of alloys from region III are presented. When the test temperature is higher than 673 K, the degradation in permeability is rapid, due to interdiffusion between the catalyst Pd layer and the substrates. This problem can be avoided or substantially mitigated by placing an interlayer (as a diffusion barrier) between the Pd coating and the substrate 7

E. Yan et al.

Journal of Membrane Science xxx (xxxx) xxx

Fig. 6. Hydrogen permeation characteristics of the as-cast alloys from region III. (a) The relationship between (J � L) and ΔP0.5 for alloy 29# (Nb40Hf30Co30); (b) Temperature dependence of the hydrogen permeability Φ in the form of an Arrhenius plot for alloys 28#-31# and pure Pd; (c) Comparison of hydrogen permeability of representative multi-phase NbxM(100-x)/2N(100-x)/2 alloys (M and N represent Hf, Zr or Ti and Co or Ni, respectively) and pure Pd.

permeability as compared to other Nb-based multi-phase alloys, as seen in Fig. 6(c) and Table 3. Typically, the Φ value of Nb45Hf27.5Co27.5 alloy at 673 K is about 1.2–3.3 times higher than that for Nb–Zr(Hf)Ni, Nb–Ti (Ni)Co alloys and pure Pd [12,19,40,41,49,50,52]. Even for eutectic Nb30Hf35Co35 alloy (31#) without any primary BCC ðNb; HfÞ phase, its permeability (3.32 � 10 8 mol H2 m 1s 1Pa 0.5) at 673 K is obvi­ ously higher than that of other hypoeutectic Nb40Ti30Co30 (3.08 � 10 8 mol H2 m 1s 1Pa 0.5) [47,50] and/or Nb50Ti25Ni25 (2.81 � 10 8 mol H2 m 1s 1Pa 0.5) [17,18] alloys. The high perme­ ability of Nb30Hf35Co35 originates from the contributions of more BCC-(Nb, Hf) in the eutectic, as previously reported [24,25]. This may be an advantage of Nb-Hf-Co alloy over other alloy series, in terms of the hydrogen permeability. Aside from primary BCC ðNb; HfÞ, the eutectic has a great contribution to the hydrogen permeation. Hence, this allows the attainment of a combination of high permeability and strong resis­ tance to HE by properly adjusting the fraction of the primary BCC ðNb; HfÞ and eutectic phases. In contrast, for other alloys such as Nb-Ti-Co Nb-Ti-Ni and Nb-Hf-Ni [12,17,47,50], HE resistance improvement leads to a high fraction of eutectic in the microstructure (and vice versa), which is accompanied by a decrease of permeability. It is difficult to balance these two properties which conflict with each other. From this point of view, the Nb-HfCo alloys are a good candidate for hydrogen separation membranes.

3.3. The compositional window suitable for hydrogen permeation For ease of comparison, the hydrogen permeability, Φ, for all alloys in region III are summarised in Fig. 2. For each alloy membrane, at least three repeated permeation tests were made at each temperature, in order to assure the reliability and consistency of results. The alloys, 21#, 22#, 24#-27# and 32 #, in region III, cracked when they were first exposed to hydrogen or during pressurisation with hydrogen, probably because of the stress caused by hydriding [33]. This means that these alloys cannot be considered as possible candidates for H2 separation membranes. Therefore, region III in Fig. 2 is not the expected compo­ sitional window, and the above alloys should be excluded. Correlating the above simulated and experimental results (Figs. 2–6), a promising compositional window, i.e. region III0 in Fig. 2, suitable for hydrogen permeation has been successfully identified. Alloys in this region, comprising of multi-phase microstructure, show superior hydrogen separation performance, and would be potential candidates for hydrogen separation membranes. On the contrary, the alloys in re­ gion I, which consist of a single-phase BCC-(Nb, Hf) structure, experi­ ence rapid cracking that limits their utilisation as H2 separation membranes. A similar problem of HE is found for the alloys in regions II and IV, due to the formation of some brittle Co6Nb7 and Ti2Co phases. Thus, only region III’ is a suitable and promising compositional window, and alloys in this region are of potential interest for applications in H2 8

E. Yan et al.

Journal of Membrane Science xxx (xxxx) xxx

of hydrogen across the membrane is more suitable for the description of the hydrogen transport progress, as proposed by Dolan et al. [55] and Yukawa et al. [33,56]. In addition, the hydrogen concentration increases with decreasing temperature, which can be attributed the fact that the absorption pro­ cess of hydrogen into these alloys is an endothermic reaction and a fall in the temperature is helpful for the absorption [26,38,57,58]. An increase in Nb content (29#→28#→24#) further increase the hydrogen con­ centration under the same testing conditions. For instance, at 0.5 MPa and at 523 K, the hydrogen absorption, C0.5, is about 0.77 H/M for Nb40Hf30Co30 (29#), whereas that is increased to 0.93 H/M for Nb50Hf25Co25 (24#). The observed solubility trends for these alloys followed those seen for the hydrogen permeability (see Fig. 6(b) and Table 3). According to Eq. (1), the higher the hydrogen solubility, the hydrogen flow permeating the membrane greatly increases. However, excessive hydrogen solubility will lead to the degradation of the HE resistance, for instance, Nb50Hf25Co25 (24#), Fig. 7(c). So, one of the most challenging problems for the alloy membrane design is how to control of solubility up to a point where HE is inhibited, while also increasing the permeability (or the H2 flux).

Table 3 Summary of hydrogen permeability (Φ), solubility (K) and diffusivity (D) for NbHf-Co alloys. Φ values of other Nb-based multi-phase alloys and pure Pd are also shown for comparison. Alloy

Temperature [K]

Φ[10-8 mol H2 m-1 s-1 Pa-0.5]

K[moH2 m3 Pa-0.5]

D[109 -1 m s 1 ]

Nb50Hf25Co25

673 623 573 523 673 623 573 523 673 623 573 523 673

NG NG NG NG 5.32 3.46 NG NG 4.96 3.23 2.08 1.12 3.32 4.64 4.31

32.42 27.51 23.87 23.87 31.87 26.73 23.54 18.73 31.75 26.61 22.85 17.51 24.41 – –

– – – – 1.66 1.29 – – 1.56 1.21 0.91 0.65 1.36 – –

3.99

–

–

3.08

–

–

3.47

–

–

2.81

–

–

–

–

Nb45Hf27.5Co27.5

Nb40Hf30Co30

Nb30Hf35Co35 Nb40Zr30Ni30 [49] Nb40Hf30Ni30 [12, 40,41] Nb60Ti20Co20 [19, 50] Nb40Ti30Co30 [19, 50] Nb56Ti22Ni22 [17, 52] Nb50Ti25Ni25 [17, 52] Pd [5,32]

1.61 -8

-1

-1

-0.5

Experimental error: Φ, �0.03 (10 mol H2 m s Pa Pa-0.5); K, �0.02 (10-9 m-1 s-1).

3.4.2. DBTC of Nb-based alloys To address this issue, the PCT curves, shown in Fig. 7, are summar­ ised together and redrawn in Fig. 8(a) in which three regions, marked by ‘I’, ‘II’ and ‘III’, are divided based on their susceptibility to HE. The upper and lower limits of the H/M value for region I are 0.72 and 0.82, respectively, which are closely related with the hydrogen pressures. The area, i.e. region II, to the left hand side of region I, lies in a lower hydrogen concentration (<0.72 H/M), and the alloys (e.g. Nb40Hf30Co30), with their PCT curves located in this region, exhibit an excellent resistance to HE, making them good candidates for hydrogen permeation. However, as for region III, the hydrogen concentration in the alloys is up to 0.93 H/M (exceeding the range of region I, 0.72–0.82 H/M), and as a consequence, the corresponding specimens (e. g. Nb50Hf25Co25) embrittle rapidly due to their higher hydrogen solu­ bility than region I. These results indicate that the embrittlement can be inhibited by controlling the concentration below the critical value, 0.72 H/M, and above or below the concentration of region I is the most important criteria to referee to determine whether the HE occurs or not for these membranes during hydrogen permeation. Thus, it is inferred that region I is the BDTC for Nb-HfCo multi-phase alloys, and its values are determined to be 0.72–0.82 H/M. When the dissolved hydrogen in the alloys exceeds this critical value, especially 0.82 H/M, HE will inevitably occur under the given testing conditions. To further confirm the above results, the mechanical properties of the typical alloys (29#, 28# and 24#) were further investigated by the in-situ SP tests at 673 K and 0.35 MPa, and their load-deflection (L-D) curves are shown in the inset of Fig. 8(a). The maximum deflection decreases with increasing hydrogen concentrations, whereas a declining trend of the maximum failure load is observed. In the case of test con­ ditions D and E, for alloys 29# and 28#, the values of maximum load and deflection are relatively high, indicating that ductile fracture occurs for these alloys. This can be attributed to the relatively low hydrogen concentrations (0.65 H/M and 0.7 H/M), which are below the DBTC of Nb-HfCo alloys. However, when the H/M value is increased to 0.81, i.e., test condition F, the characteristic brittle fracture, due to the HE, takes place for the alloy 24#. These results indicate that the boundary for the ductile-to-brittle transition exists in the region of the concentration be­ tween 0.7 and 0.81 at 0.35 MPa, which correspond with the results in Figs. 6 and 8(a). These results are particularly important, because the resistance to HE (or durability) of these alloy membranes can be improved by properly adjusting the hydrogen content below this critical value, i.e. the DBTC in Fig. 8(a). In addition, the above DBTC (i.e. region I in Fig. 8 (a)) was also plotted in the Nb–H binary phase diagram in order to further analyse the mechanism of HE in Nb-based alloys, as shown in Fig. 8 (b). The DBTC

); D, �0.02 (moH2 m-3

separation and purification. For example, Nb30Hf35Co35 alloy (31#), consists of a fully lamellar eutectic structure and has a relatively low melting point (see Fig. 3), which is suitable for the development of new kinds of alloy membranes by special processing methods such as direc­ tional solidification [51], rolling [39,52] and/or a melt-spinning tech­ nique [1,32,53]. If the alloy membranes can be successfully prepared by the above techniques, their properties such as hydrogen permeability, or flux, would be further enhanced. Work in this area is currently ongoing. 3.4. Hydrogen solubility and DBTC 3.4.1. Hydrogen solubility It is known that the absorbed hydrogen concentration in Nb-based alloys has a significant effect on mechanical strength, especially resis­ tance to HE [29,32,54]. This embrittlement problem is exacerbated when the membranes are cooled to room temperature [42]. To analyse the causes of HE for Nb-Hf-Co alloys, the hydrogen sorption properties of typical samples (24#, 28# and 29#) were investigated, as shown in Fig. 7. Clearly, Sieverts’ law, i.e. C ¼ K⋅P0.5, is only valid at very low hydrogen pressures (<0.01 MPa), see Fig. 7 (a)–(c). On the contrary, it does not hold when the pressure exceeds 0.01 MPa. Despite this, there is almost a linear relationship between C and P0.5 in the pressure range of 0.1–0.5 MPa for each measured temperature, which can be written as follows: C ¼ K⋅P0.5 þ α

(2)

where α is a constant, and its values are illustrated in Fig. 7 (d)–(f). Substituting Eq. (2) into Fick’s first law of diffusion equation, i.e. J ¼ D (Cu–Cd)/L, the same form of expression with Eq. (1), i.e. J ¼ D⋅K 0.5 (P0.5 u -Pd )/L, was obtained. This means that hydrogen permeability can be defined in Eq. (1), although the pressure-concentration relationship does not obey Sieverts’ law. Nevertheless, the above pseudo-Sieverts’ equation (C ¼ K⋅P0.5 þ α) is not applicable when hydrogen pressure is below 0.1 MPa or above 0.5 MPa. In these cases, the chemical potential 9

­

E. Yan et al.

Journal of Membrane Science xxx (xxxx) xxx

Fig. 7. Pressure-composition isotherms of Nb-Hf-Co alloys in the form of a Sieverts’ plot (i.e. C vs. P0.5). (a) (d) Nb40Hf30Co30 (29#); (b) (e) Nb45Hf27.5Co27.5 (28#); (c) (f) Nb50Hf25Co25 (24#). The right ones are the corresponding magnified parts of PCT curves and the linear fittings with pressure ranging from 0.1 to 0.5 MPa.

reported Nb-based multi-phase alloys such as Nb40Ti30Ni30, Nb35Mo5 Ti25Co30 and Nb35W5Ti25Co30, are about 0.35–0.6 H/M, which is lower than that of region I (0.72–0.82 H/M). This might explain why they did not undergo brittle fracture due to HE, during their hydrogen perme­ ation process, even if their hydrogen concentration exceeded the DBTC of the Nb-based single-phase alloys (i.e., region I’, 0.23–0.33 H/M), as mentioned above. In a word, these results break through the conven­ tional limits that the DBTC for Nb-based alloys merely occurs at a hydrogen concentration level of ~0.23 H/M, and hence provides an important clue to the design of Nb-based permeable membranes against the HE.

region (or value) in pure Nb, and its single-phase alloys proposed by Gahr et al. and Yukawa et al. [33–36,46,59,60], are also included for comparison. In the case of pure Nb, the DBTC below 573 K is at ~0.3 H/M, while its critical value gradually increases at higher tem­ peratures, especially at T > 623 K. Compared with that of pure Nb, the DBTC region of the Nb-based single-phase alloys (marked as I0 ) was found to be shifted to the lower hydrogen content region, i.e. H/M � 0.23–0.33, whereas it shifts greatly towards the right side (H/M � 0.72–0.82, marked as I), for Nb-based multi-phase alloys. These results reveal that multi-phase alloys have more advantages, because they show a better tolerance for the high hydrogen concentration (or pressure) than their corresponding single-phase counterparts, thus achieving the targeted flux value without a mechanical penalty. Besides, the dissolved hydrogen concentrations (see the circle in Fig. 8(b)) of the 10

E. Yan et al.

Journal of Membrane Science xxx (xxxx) xxx

Fig. 8. The DBTC (region I) for Nb-Hf-Co multi-phase alloys. (a) The isotherms are divided into three sub-regions marked as I (DBTC), II and III according to their susceptibility to HE; (b) The boundary for the ductile-to-brittle transition for pure Nb, Nb-based single-phase and multi-phase alloys plotted on the Nb–H binary phase diagram. The number, in brackets on the right side of the alloy, is the pressure (in MPa unit) at the upstream side of the membranes during the hydrogen permeation testing. Inset in (a): L-D curves for the alloys 29#, 28# and 24# measured at 673K in hydrogen atmosphere.

3.5. Concept for the design of Nb-based alloy membranes

curve, DBTC position and the HE susceptibility of the membranes, a new design concept for Nb-based multi-phase alloy membranes is proposed, as shown in Fig. 9. Three typical PCT curves at a certain temperature, T, are present in the picture, which are marked as (i), (ii) and (iii) from right to left, respectively. The feed-side and permeate-side’s pressures of the mem­ brane are fixed to be Pu and Pd. Then, for the alloy membrane having the PCT curve (i), the hydrogen concentration difference across the mem­ brane is high and a large K (i.e., K1) can be obtained, which is beneficial to increase the H2 flux according the equation in the bottom right corner of Fig. 9. Nevertheless, the hydrogen content (e.g. at hydrogen pressure

As mentioned above, the DBTC, CDBTC , of Nb-based alloys can be expressed as follows: CDBTC ¼ 1

μi

(3)

Where μi is 0.67–0.77 H/M for a single-phase alloy, whereas it is 0.18–0.28 H/M for a multi-phase alloy. In each case, the absorbed hydrogen content in the membrane should be kept below the critical value, otherwise, they will face the risk of HE. Based on the work of Zhu et al. and Suzuki et al. [38,56,61], considering the shape of the PCT 11

E. Yan et al.

Journal of Membrane Science xxx (xxxx) xxx

Fig. 9. Schematic illustration of the concept of Nb-based multi-phase alloy design based on the PCT curve and the DBTC in Fig. 8.

Pu) greatly exceeds the region of the DBTC, thus resulting in severe HE fracture for the membrane. Two alternative methods can be adopted, one being to reduce the pressure of the upstream (Pu→Pu’, i.e. from position ① to position ②) and the other to move the PCT curve toward the upper left orientation, for instance, by alloying. Then, for the alloy membrane having the PCT curve (ii), the maximum hydrogen absorp­ tion capacity (position ③) is below the lower limit value of DBTC (i.e. 0.72) at the same hydrogen pressure, Pu. Meanwhile, the pressure dif­ ference (Pu–Pd) across the membrane becomes larger than that (Pu’-Pd) for the PCT curve (i). In such a case, it is expected to increase the H2 flux and improve the capability of the membrane against HE. However, if the PCT curve is shifted excessively to the left-hand region, like PCT curve (iii), the hydrogen concentration difference (Cu’-Cd’) and solubility (K3) become small, which reduces the H2 flux. Thus, the membrane which has the PCT curve (ii) would be the ideal case. From this point of view, it is necessary to control the hydriding property of the Nb-based alloy by composition design or proper alloying with other metals. It has been confirmed that adding small amounts of a transition metal (e.g. Fe, Ti, Ni, Co, Zr, Mo, Ru or Rh [2,13,17–21,32, 37–41,43,50,57,62]) can attenuate the hydrogen uptake of group 5B metals and alter the shape of the PCT curve (or the stability of hydrides), thereby suppressing HE. For example, Dolan et al. [26] have successfully developed a multiphase V85Ni10Ti5 alloy by partial substitution of V with Ni and Ti, which shows a high hydrogen permeability of 9.3 � 10 8 mol H2 m 1s 1Pa 0.5 at 673 K. Nevertheless, the type and/or quantities of selected elements should be precisely controlled so as to avoid the sharp decrease of the dissolved hydrogen concentration in the alloy (such as PCT curve (iii) in Fig. 9), which will reduce the perme­ ability or the steady-state H2 flux. In brief, the new design concept for hydrogen permeable membranes provides a practical method to improve the resistance to embrittlement of Nb-based alloy membranes while minimising the associated permeability penalty.

different hydrogen transport behaviours at temperatures between 523 K (brittle fracture due to HE) and 623 K (good resistance to HE), see Figs. 6 and 8. Besides, there are two reasons for the substitution of Hf by Ti. First, the element Ti, in the same IVA group as Hf, has a similar physi­ cochemical performance, which is likely to form the multi-phase struc­ ture required for H2 permeation after partial substitution. Second, Ti shows a lower hydrogen affinity than Hf [13,45,63], and there is a possibility that appropriate substitution of Hf by Ti can form new hydrogen permeable alloys with excellent comprehensive performances. Fig. 10(a) shows the SEM images of Nb45Hf27.5-xTixCo27.5 alloys. For x ¼ 0, i.e. Nb45Hf27.5Co27.5 (28#) alloy, its SEM micrographs can be seen in Fig. 5(e). Obviously, all the alloys exhibit a typical multi-phase microstructural feature, but with more eutectic {BCC-Nbþ(Hf, Ti)Co} phases around the BCC-Nb phase boundary. XRD patterns in Fig. 10(b) further confirmed that each alloy contained these phases. Besides, the substitution of Hf by Ti induces a lower hydrogen concentration, and the corresponding PCT curves are shifted toward the upper left region, as shown in Fig. 10(c). It should be noted that the hydrogen content at 0.3 MPa (i.e., position ①) is ~0.75 H/M for Nb45Hf20Ti2.5Co27.5 (x ¼ 2.5) alloy, which is much higher than the critical value (0.72 H/M) of the DBTC in Fig. 8. On the contrary, in the case of Nb45Hf27.5xTixCo27.5 (x ¼ 5, 7.5 and 10) alloys, their hydrogen concentration (e.g. position ③) are close to or less than the boundary of the DBTC at identical pressure. Thus, it is inappropriate for taking Nb45Hf20Ti2.5Co27.5 (x ¼ 2.5) alloy as the hydrogen permeable mem­ brane because its adsorption isotherm has a similar shape with the PCT curve (i) in Fig. 9. That is, this alloy is likely to fracture due to its higher hydrogen solubility. In contrast, the degradation of mechanical prop­ erties caused by HE may be avoided, or substantially mitigated, for other Nb45Hf27.5-xTixCo27.5 alloys. To test this hypothesis, the hydrogen permeation test for all the membranes was performed, as shown in Fig. 10(d). All the samples displayed qualitatively similar behaviour at the beginning of the mea­ surements, i.e., the H2 flux rose rapidly when they were exposed during pressurisation with hydrogen. However, Nb45Hf25Ti2.5Co27.5 shows an abrupt increase in H2 flux within a few minutes, indicating that failure of this membrane occurs by cracking because of HE. This can be attributed to the fact that its hydrogen concentration (~0.75 H/M) was above the

3.6. Design of novel Nb-based hydrogen permeable alloys Following this concept, Nb45Hf27.5-xTixCo27.5 (x ¼ 0 … 10) alloys were designed and developed, as an example, in this study. This alloy series was chosen because the base Nb45Hf27.5Co27.5 alloy exhibits 12

E. Yan et al.

Journal of Membrane Science xxx (xxxx) xxx

Fig. 10. Design of Nb45Hf27.5-xTixCo27.5 (x ¼ 0 … 10) alloy membranes with high hydrogen flux and strong resistance to hydrogen embrittlement following the concept proposed in Fig. 9. (a) BSE micrographs; (b) XRD patterns; (c) PCT curves at 573 K, and (d) Hydrogen permeation flux with varying time.

curve (iii) in Fig. 9, thus further reducing the H2 flux. For instance, the hydrogen flux, J, of Nb45Hf20Ti7.5Co27.5 is ~3.4 (2.1) times higher than that for Nb30Mo10Ti30Ni30 (pure Pd) under the same test conditions, see Fig. 10(d) and Table 4. From this point of view, the Nb45Hf20Ti7.5Co27.5 alloy designed by the concept shown in Fig. 9, is a much better candidate for hydrogen-permeable materials, because it exhibits much improved resistance to embrittlement and a larger H2 flux. It is noteworthy that the DBTC is not the only consideration when designing, or selecting, an alloy as a hydrogen permeation membrane. The intrinsic properties of the alloys such as strength, stiffness, fracture toughness, heat expansion coefficient and corrosion-resistance, are also important [16,31,57,64]. Besides, phase stability at 473 K, and below, is a key consideration, since the formation of some brittle metal hydrides (e.g. β-hydride in pure V and Nb metals [42,57,59,65]) can lead to potentially catastrophic failure. In particular, the hydride phase transi­ tion (such as α→β in pure Pd, Nb or V [58,66–68]) should be avoided in practice through the implementation of proper operating procedures. For instance, the permeation cell with membrane should be in vacuum before heating to the desired working temperature in order to avoid the αþβ miscibility gap. Nonetheless, the control of solubility below DBTC through proper compositional substitutions is the practical tool for Nb-based alloy membrane design. By using this strategy, it is possible to develop novel hydrogen permeable alloys to meet the 2015 U. S. DOE performance targets of ~100 m3 h-1 m-2 H2 flux and ~5 years durability [26,57].

Table 4 Hydrogen transport properties of Nb45Hf27.5-xTixCo27.5 (x ¼ 0 … 10) alloy membranes. The results of Nb30Mo10Ti30Ni30 alloy are shown for comparison. Alloy

Φ[10-8 mol H2 m-1 s-1 Pa-0.5]

K[mol H2 m-3 Pa-0.5]

D[109 m-1 s1 ]

J[10-3 mol H2 m-2 s-1]

Nb45Hf27.5Co27.5 Nb45Hf27.5Ti2.5Co27.5 Nb45Hf27.5Ti5Co27.5 Nb45Hf27.5Ti7.5Co27.5 Nb45Hf27.5Ti10Co27.5 Nb30Mo10Ti30Ni30 [2, 39]

NG NG 2.11 1.93 1.84 0.62

23.54 22.13 21.18 20.57 19.85 15.12

– – 0.95 0.95 0.93 0.41

NG NG 7.66 7.44 7.11 2.39

maximum of DBTC, i.e., 0.72 H/M. When the concentration in the membrane is below 0.72 H/M, Nb45Hf22.5Ti5Co27.5 (x ¼ 5) alloy is sub­ ject to HE failure after hydrogen permeation for 27min, implying that its resistance against HE is, to some extent, improved. This is another proof that the DBTC of the Nb-based multi-phase alloy membrane is 0.72–0.82 H/M. A further improvement of HE resistance is clearly observed for other Nb45Hf27.5-xTixCo27.5 (x ¼ 7.5 and 10) alloys with lower hydrogen concentrations, and they did not fail during hydrogen permeation for 150 min. Despite this, the dissolved hydrogen content should not be too low (e.g. Nb30Mo10Ti30Ni30 [2,39] in Fig. 10(c)), otherwise, the PCT curve is shifted excessively to the left region like PCT 13

E. Yan et al.

Journal of Membrane Science xxx (xxxx) xxx

4. Conclusions

[8] G. Song, M.D. Dolan, M.E. Kellam, S. Zambelli, V-Ni-Ti multi-phase alloy membranes for hydrogen purification, J. Alloy. Comp. 509 (2011) 9322–9328. [9] Y.S. Jo, C.H. Lee, S.Y. Kong, K.Y. Lee, C.W. Yoon, S.W. Nam, J. Han, Characterization of a Pd/Ta composite membrane and its application to a large scale high-purity hydrogen separation from mixed gas, Separ. Purif. Technol. 200 (2018) 221–229. [10] E.H. Yan, H.R. Huang, R.N. Min, F. Xua, P. Zhao, Y.J. Zou, H.L. Chu, H.Z. Zhang, L. X. Sun, Design and characterizations of novel Nb-ZrCo hydrogen permeationalloys for hydrogen separation applications, Mater. Chem. Phys. 212 (2018) 282–291. [11] K. Hashi, K. Ishikawa, T. Matsuda, K. Aoki, Microstructure and hydrogen permeability in Nb-Ti-Co multiphase alloys, J. Alloy. Comp. 425 (2006) 284–290. [12] F. Shi, X.P. Song, Effect of niobium on the microstructure, hydrogen embrittlement, and hydrogen permeability of NbxHf(1-x)/2Ni(1-x)/2 ternary alloys, Int. J. Hydrogen Energy 35 (2010) 10620–10623. [13] J.W. Phair, R. Donelson, Developments and design of novel (non-palladium-based) metal membranes for hydrogen separation, Ind. Eng. Chem. Res. 45 (2006) 5657–5674. [14] Y. Shimpo, S. Yamaura, M. Nishida, H. Kimura, A. Inoue, Development of meltspun Ni-Nb-Zr-Co amorphous alloy for high-performance hydrogen separating membrane, J. Membr. Sci. 286 (2006) 170–173. [15] M.D. Dolan, S. Hara, N.C. Dave, K. Haraya, M. Ishitsuka, A.Y. Ilyushechkin, K. Kita, K.G. McLennan, L.D. Morpeth, M. Mukaida, Thermal stability, glass-forming ability and hydrogen permeability of amorphous Ni64Zr36-xMx, (M¼Ti, Nb, Mo, Hf, Ta or W) membranes, Separ. Purif. Technol. 65 (2009) 298–304. [16] M.D. Dolan, M.E. Kellam, K.G. McLennan, D. Liang, G. Song, Hydrogen transport properties of several vanadium-based binary alloys, Int. J. Hydrogen Energy 38 (2013) 9794–9799. [17] K. Hashi, K. Ishikawa, T. Matsuda, K. Aoki, Hydrogen permeation characteristics of multi-phase Ni-Ti-Nb alloys, J. Alloy. Comp. 368 (2004) 215–220. [18] W.M. Luo, K. Ishikawa, K. Aoki, Hydrogen permeability in Nb-Ti-Ni alloys containing much primary (Nb, Ti) phase, Mater. Trans. 46 (2005) 2253–2259. [19] K. Hashi, K. Ishikawa, T. Matsuda, K. Aoki, Microstructure and hydrogen permeability in Nb-Ti-Co multiphase alloys, J. Alloy. Comp. 425 (2006) 284–290. [20] K. Ishikawa, S. Watanabe, K. Aoki, Microstructure and hydrogen permeability in Nb-TiFe alloys, J. Alloy. Comp. 566 (2013) 68–72. [21] K. Ishikawa, T. Takano, T. Matsuda, K. Aoki, High hydrogen permeability in the Nb-Zr-Ni eutectic alloy containing the primary body-centered-cubic (Nb, Zr) phase, Appl. Phys. Lett. 87 (2005), 081906. [22] S.X. Huang, X.D. Zhang, Y. Jiang, Y.R. Jiang, C. Mao, D. Wu, L.G. Zhang, L.B. Liu, Experimental investigation of Ti-Nb-Co ternary system at 1000 � C, Mater. Des. 115 (2017) 170–178. [23] H.H. Xu, Y. Du, Z.H. Yuan, S.T. Li, J.C. Schuster, Y.H. He, Phase equilibria of the Fe-Nb-Ti system at 900 � C, J. Alloys. Comp. 396 (2005) 151–155. [24] E.H. Yan, X.Z. Li, D.M. Liu, M. Rettenmayr, Y.Q. Su, J.J. Guo, Nb-HfCo alloys with pronounced high hydrogen permeability: a new family of metallic hydrogen permeation membranes, Int. J. Hydrogen Energy 39 (2014) 8385–8389. [25] E.H. Yan, L.X. Sun, F. Xu, Y.J. Zou, H.L. Chu, H.Z. Zhang, Y.X. Sun, Changes in microstructures and hydrogen permeability of Nb30Hf35Co35 eutectic alloy membranes by annealing, Int. J. Hydrogen Energy 41 (2016) 1401–1407. [26] M.D. Dolan, G. Song, D. Liang, M.E. Kellam, D. Chandra, J.H. Lam, Hydrogen transport through V85Ni10M5 alloy membranes, J. Membr. Sci. 373 (2011) 14–19. [27] S. Tosti, A. Santucci, A. Pietropaolo, S. Brutti, O. Palumbo, F. Trequattrini, A. Paolone, Hydrogen sorption properties of V85Ni15, Int. J. Hydrogen Energy 43 (2018) 2817–2822. [28] M.L. Martin, M. Dadfarnia, A. Nagao, S.A. Wang, P. Sofronis, Enumeration of the hydrogen-enhanced localized plasticity mechanism for hydrogen embrittlement in structural materials, Acta Mater. 165 (2019) 734–750. [29] F. Stefano, S. Barison, S. Boldrini, A. Ferrario, M. Romano, F. Montagner, E. Miorin, M. Fabrizio, L. Armelao, Hydrogen separation by thin vanadium-based multilayered membranes, Int. J. Hydrogen Energy 43 (2018) 3235–3243. [30] M.D. Dolan, D.M. Viano, M.J. Langley, K.E. Lamb, Tubular vanadium membranes for hydrogen purification, J. Membr. Sci. 549 (2018) 306–311. [31] E.H. Yan, R.N. Min, H.R. Huang, P. Zhao, P.R. Huang, Y.J. Zou, H.L. Chu, S.H. Sun, L.X. Sun, Multiphase Nb-TiCo alloys: the significant impact of surface corrosion on the structural stability and hydrogen permeation behaviour, Int. J. Hydrogen Energy 44 (2019) 16684–16697. [32] S.N. Paglieri, N.K. Pal, M.D. Dolan, S.M. Kim, W.M. Chien, J. Lamb, D. Chandra, K. M. Hubbard, D.P. Moore, Hydrogen permeability, thermal stability and hydrogen embrittlement of Ni-Nb-Zr and Ni-Nb-Ta-Zr amorphous alloy membranes, J. Membr. Sci. 378 (2011) 42–50. [33] A. Suzuki, H. Yukawa, T. Nambu, Y. Matsumoto, Y. Murata, Analysis of hydrogen mobility in Nb-based alloy membranes in view of new description of hydrogen permeability based on hydrogen chemical potential, J. Alloy. Comp. 645 (2015) 107–111. [34] Y. Awakura, T. Nambu, Y. Matsumoto, H. Yukawa, Hydrogen solubility and permeability of Nb-W-Mo alloy membrane, J. Alloy. Comp. 509S (2011) S877–S880. [35] A. Suzuki, H. Yukawa, S. Ijiri, T. Nambu, Y. Matsumoto, Y. Murata, Alloying effects on hydrogen solubility and hydrogen permeability for V-based alloy membranes, Mater. Trans. 56 (2015) 1688–1692. [36] H. Yukawa, T. Nambu, Y. Matsumoto, V-W alloy membranes for hydrogen purification, J. Alloy. Comp. 509S (2011) S881–S884. [37] K. Ishikawa, K. Yonehara, Effects of tungsten addition on hydrogen absorption and permeation properties of Nb40Ti30Ni30 alloy, J. Alloy. Comp. 749 (2018) 634–639. [38] K.J. Zhu, X.Z. Li, Z.F. Zhu, R.C. Chen, Y.Q. Su, J.J. Gou, M. Rettenmayr, D.M. Liu, Analysis of W/Mo alloying on hydrogen permeation performance of dual phase Nb-

In this study, the ternary phase diagram, solidification behaviour and hydrogen transport properties of Nb-Hf-Co alloys were studied. It is shown that three ternary invariant transformations, and five different solidification paths, exist in the BCC-(Nb, Hf) phase region, where a new compositional window suitable for hydrogen permeation was explored and established for the first time. The alloys in this window have been proven to exhibit very high hydrogen permeability. Furthermore, the boundary for the DBTC of Nb-based multi-phase alloys was determined and demonstrated to be at around 0.72–0.82 H/M, which is different from that (~0.23–0.23 H/M) in the single-phase alloy. Based on this, the alloy membranes with the hydrogen concentration close to or below, 0.72 H/M exhibit a high hydrogen permeability without incurring the HE resistance penalty. Finally, a concept to develop Nb-based multiphase hydrogen permeable membranes is proposed on the basis of DBTC, coupled with their PCT curves. Following this concept, a novel alloy series, i.e. Nb45Hf27.5-xTixCo27.5 (x ¼ 7.5 and 10), was successfully developed, and exhibited an optimum combination of high H2 perme­ ation flux and strong resistance to HE under appropriate permeation conditions. Data Availability Statement The raw/processed data required to reproduce these findings cannot be shared at this time as the data also forms part of an ongoing study. Declaration of competing interest The authors declared that they have no conflicts of interest to this work. No conflict of interest exits in the submission of this manuscript, and manuscript is approved by all authors for publication. We declare that we do not have any commercial or associative interest that repre­ sents a conflict of interest in connection with the work submitted. Acknowledgements This research was founded by National Natural Science Foundation of China (51701048, 51761009, 51671062 and 51801041), National Key Research and Development Program (2018YFB1502105), the Innovation Project of GUET Graduate Education (2019YCXS109) and Guangxi Scientific Technology Team (2017AD23029). Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi. org/10.1016/j.memsci.2019.117531. References [1] E.H. Yan, H.R. Huang, S.H. Sun, Y.J. Zhou, H.L. Chu, L.X. Sun, Development of NbTi-Co alloy for high-performance hydrogen separating membrane, J. Membr. Sci. 565 (2018) 411–424. [2] X.Z. Li, X. Liang, D.M. Liu, R.R. Chen, F.F. Huang, R. Wang, M. Rettenmayr, Y. Q. Su, J.J. Guo, H.Z. Fu, Design of (Nb, Mo)40Ti30Ni30 alloy membranes for combined enhancement of hydrogen permeability and embrittlement resistance, Sci. Rep. 7 (2017) 209. [3] Y.S. Lee, J.H. Shim, J.Y. Suh, A finite outlet volume correction to the time lag method: the case of hydrogen permeation through V-alloy and Pd membranes, J. Membr. Sci. 585 (2019) 253–259. [4] W. Wang, K. Ishikawa, K. Aoki, Microstructural change-induced lowering of hydrogen permeability in eutectic Nb-TiNi alloy, J. Membr. Sci. 351 (2010) 65–68. [5] H.Y. Jia, P. Wu, G.F. Z, S.C. Eduardo, A. Serrano, G.R. Castro, H.Y. Xu, C.L. Sun, A. Goldbach, High-temperature stability of Pd alloy membranes containing Cu and Au, J. Membr. Sci. 544 (2017) 151–160. [6] S. Pati, R.A. Jat, N.S. Anand, D.J. Derose, K.N. Karn, S.K. Mukerjee, S.C. Parida, PdAg-Cu dense metallic membrane for hydrogen isotope purification and recovery at low pressures, J. Membr. Sci. 522 (2017) 151–158. [7] I.S. Sipatov, N.I. Sidorov, E.A. Pastukhov, I.E. Gabis, V.A. Piven, A.A. Esin, S. V. Pryanichnikov, A.A. Vostryakov, Hydrogen permeability and structure of vanadium alloy membranes, Pet. Chem. 57 (2017) 483–488.

14

E. Yan et al.

[39] [40] [41] [42] [43] [44]

[45]

[46]

[47] [48]

[49] [50] [51] [52]

Journal of Membrane Science xxx (xxxx) xxx

Ti-Ni alloys based on hydrogen chemical potentials, J. Membr. Sci. 584 (2019) 290–299. D.M. Liu, X.Z. Li, H.Y. Geng, R.C. Chen, M. Rettenmayr, Y.Q. Su, H. Li, J.J. Gou, H. Z. Fu, Development of Nb35Mo5Ti30Ni30 alloy membrane for hydrogen separation applications, J. Membr. Sci. 553 (2018) 171–179. F. Shi, Microstructure and hydrogen permeability of Nb40Hf30Ni30 ternary alloy, Int. J. Hydrogen Energy 35 (2010) 10556–10559. F. Shi, H.T. Xiao, Hydrogen storage properties of Nb-Hf-Ni ternary alloys, Int. J. Hydrogen Energy 38 (2013) 2318–2324. M.D. Dolan, M.A. Kochanek, C.N. Munnings, K.G. McLennan, D.M. Viano, Hydride phase equilibria in V-Ti-Ni alloy membranes, J. Alloy. Comp. 622 (2015) 276–281. F. Liu, Z.J. Xu, Z.M. Wang, M.Y. Dong, J.Q. Deng, Q.R. Yao, Y. Ma, J.X. Zhang, N. Wang, Z.H. Guo, Structures and mechanical properties of Nb-Mo-Co(Ru) solid solutions for hydrogen permeation, J. Alloy. Comp. 756 (2018) 26–32. E.H. Yan, X.Z. Li, M. Rettenmayr, D.M. Liu, Y.Q. Su, J.J. Gou, D.M. Xu, H.Z. Fu, Design of hydrogen permeable Nb-Ni-Ti alloys by correlating the microstructures, solidification paths and hydrogen permeability, Int. J. Hydrogen Energy 39 (2014) 3505–3516. X.Z. Li, D.M. Liu, R.R. Chen, E.H. Yan, X. Liang, M. Rettenmayr, Y.Q. Su, J.J. Gou, H.Z. Fu, Changes in microstructure, ductility and hydrogen permeability of Nb-(Ti, Hf)Ni alloy membranes by the substitution of Ti by Hf, J. Membr. Sci. 484 (2015) 47–56. T. Nambu, K. Shimizu, Y. Matsumoto, R. Rong, N. Watanabe, H. Yukawa, M. Morinaga, I. Yasuda, Enhanced hydrogen embrittlement of Pd-coated niobium metal membrane detected by in situ small punch test under hydrogen permeation, J. Alloy. Comp. 446–447 (2007) 588–592. Y. Saeki, Y. Yamada, K. Ishikawa, Relationship between hydrogen permeation and microstructure in Nb-TiCo two-phase alloys, J. Alloy. Comp. 645 (2015) s32–s35. E.H. Yan, L.X. Sun, F. Xu, D.M. Xu, S.J. Qiu, C.L. Xiang, H.Z. Zhang, Y.X. Sun, Changes in microstructure, solidification path and hydrogen permeability of NbHf-Co alloy by adjusting Hf/Co ratio, Int. J. Hydrogen Energy 41 (2016) 1391–1400. T. Takano, K. Ishikawa, T. Matsuda, K. Aoki, Hydrogen permeation of eutectic NbZr-Ni alloy membranes containing primary phases, Mater. Trans. 45 (2004) 3360–3362. W.M. Luo, K. Ishikawa, K. Aoki, Highly hydrogen permeable Nb-Ti-Co hypereutectic alloys containing much primary bcc-(Nb, Ti) phase, Int. J. Hydrogen Energy 37 (2012) 12793–12797. K. Kishida, Y. Yamaguchi, K. Tanaka, H. Inui, S. Tokui, K. Ishikawa, K. Aoki, Microstructures and hydrogen permeability of directionally solidified Nb-Ni-Ti alloys with the Nb-NiTi eutectic microstructure, Intermetallics 16 (2008) 88–95. K. Ishikawa, S. Tokui, K. Aoki, Hydrogen permeation in anisotropic Nb-TiNi twophase alloys formed by forging and rolling, Int. J. Hydrogen Energy 42 (2017) 11411–11421.

[53] S.I. Yamaura, A. Inoue, Effect of surface coating element on hydrogen permeability of melt-spun Ni40Nb20Ta5Zr30Co5 amorphous alloy, J. Membr. Sci. 349 (2010) 138–144. [54] K.H. Kim, H.C. Park, J. Lee, E. Cho, S.M. Lee, Vanadium alloy membranes for high hydrogen permeability and suppressed hydrogen embrittlement, Scr. Mater. 68 (2013) 905–908. [55] M.D. Dolan, K.G. McLennan, J.D. Way, Diffusion of atomic hydrogen through V-Ni alloy membranes under nondilute conditions, J. Phys. Chem. C 116 (2012) 1512–1518. [56] H. Yukawa, T. Nambu, Y. Matsumoto, N. Watanabe, G.X. Zhang, M. Morinaga, Alloy design of Nb-based hydrogen permeable membrane with strong resistance to hydrogen embrittlement, Mater. Trans. 49 (2008) 2202–2207. [57] M.D. Dolan, Non-Pd BCC alloy membranes for industrial hydrogen separation, J. Membr. Sci. 362 (2010) 12–18. [58] N.A. Al-Mufachi, N.V. Rees, R. Steinberger-Wilkens, Hydrogen selective membranes: a review of palladium-based dense metal membranes, Renew. Sustain. Energy Rev. 47 (2015) 540–551. [59] S. Gahr, H.K. Birnbaum, Hydrogen embrittlement of niobium-III. High temperature behavior, Acta Metall. 26 (1978) 1781–1788. [60] S. Gahr, M.L. Grossbeck, H.K. Birnbaum, Hydrogen embrittlement of Nb IMacroscopic behavior at low temperatures, Acta Metall. 25 (1977) 125–134. [61] A. Suzuki, H. Yukawa, T. Nambu, Y. Matsumoto, Y. Murata, Alloy design of Vbased hydrogen permeable membrane under given temperature and pressure condition, Int. J. Hydrogen Energy 42 (2017) 22325–22329. [62] K. Komiya, Y. Shinzato, H. Yukawa, M. Morinaga, I. Yasuda, Measurement of hydrogen permeability of pure Nb and its alloys by electrochemical method, J. Alloy. Comp. 404–406 (2005) 257–260. [63] M. Tada, Phase diagram of Hf-H alloy, J. Mater. Sci. 25 (1990) 2934–2936. [64] P.Y. Li, Z. Wang, Z.H. Qiao, Y.N. Liu, X.C. Cao, W. Li, J.X. Wang, S.C. Wang, Recent developments in membranes for efficient hydrogen purification, J. Mater. Sci. 495 (2015) 130–168. [65] N.A.A. Rusman, M. Dahari, A review on the current progress of metal hydrides material for solid-state hydrogen storage applications, Int. J. Hydrogen Energy 41 (2016) 12108–12126. [66] L.S. McLeod, F.L. Degertekin, A.G. Fedoro, Non-ideal absorption effects on hydrogen permeation through palladium-silver alloy membranes, J. Mater. Sci. 339 (2009) 109–114. [67] S. Yun, S.T. Oyama, Correlations in palladium membranes for hydrogen separation: a review, J. Membr. Sci. 375 (2011) 28–45. [68] M.V. Goltsova, Reverse hydride transformations in the Pd-H system, Int. J. Hydrogen Energy 31 (2006) 223–229.

15