Hydrogen permeability, thermal stability and hydrogen embrittlement of Ni–Nb–Zr and Ni–Nb–Ta–Zr amorphous alloy membranes

Journal of Membrane Science 378 (2011) 42–50

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Hydrogen permeability, thermal stability and hydrogen embrittlement of Ni–Nb–Zr and Ni–Nb–Ta–Zr amorphous alloy membranes Stephen N. Paglieri a,∗ , Narendra K. Pal b , Michael D. Dolan c , Sang-Mun Kim b , Wen-Ming Chien b , Joshua Lamb b , Dhanesh Chandra b , Kevin M. Hubbard d , David P. Moore d a

TDA Research, Inc., Wheat Ridge, CO 80033-1916, USA University of Nevada, Reno, NV 89557-0208, USA c Commonwealth Scientific and Industrial Research Organisation, Division of Energy Technology, Kenmore, QLD 4069, Australia d Los Alamos National Laboratory, Los Alamos, NM 87545, USA b

a r t i c l e

i n f o

Article history: Received 17 December 2010 Received in revised form 22 April 2011 Accepted 24 April 2011 Available online 30 April 2011 Keywords: Hydrogen permeability Amorphous alloy membrane Thermal stability Hydrogen embrittlement Metallic interdiffusion

a b s t r a c t Amorphous alloys are a promising alternative to Pd alloy membranes for hydrogen separation because of their lower cost and comparable hydrogen permeability. A series of amorphous alloy membranes consisting of Ni60 Nb20 Zr20 (at%), (Ni0.6 Nb0.4 )100−x Zrx and (Ni0.6 Nb0.3 Ta0.1 )100−x Zrx (where x = 0, 10, 20 or 30) were prepared by melt spinning and then coating the foil surfaces with a thin (500 nm) layer of Pd using physical vapor deposition (PVD). A (Ni0.6 Nb0.4 )70 Zr30 membrane exhibited the highest hydrogen permeability (1.4 × 10−8 mol m−1 s−1 Pa−0.5 ) of any of the materials, measured in pure hydrogen at 450 ◦ C. Membrane permeability increased with Zr content, but membranes higher in Zr were more susceptible to brittle failure and were more thermally unstable. Decreases in hydrogen permeability were almost always observed during long-term permeability tests at 400 and 450 ◦ C. The addition of Ta slightly increased the thermal stability, but moderately lowered the hydrogen permeability. An AES depth profile of the membrane surface showed that metallic interdiffusion had taken place between the Pd coating and the bulk membrane, which probably accounts for the reduction in hydrogen permeability over time at 400–450 ◦ C. © 2011 Elsevier B.V. All rights reserved.

1. Introduction Hydrogen is a less polluting fuel for transportation and electrical power generation, but key challenges remain for its widespread production, distribution and storage. Presently, wholesale hydrogen production for the oil refining, chemical and petrochemical industries is derived at point-of-use from fossil fuels such as natural gas [1]. In the future, gasification of coal and biomass has the potential to be another large-scale source of hydrogen. Membrane reactors that combine reaction and separation can produce hydrogen more efficiently by steam reforming and water–gas shift [2–7]. Hydrogen separating membranes can also help capture carbon dioxide for sequestration by concentrating it in the retentate. Metal membranes made from Pd alloys are presently used in commercial processes to make ultrapure hydrogen for use in compound semiconductor manufacturing and hydrogen isotope recovery in nuclear facilities [8–10]. Pd alloys have high hydrogen permeabilities and are catalytic for hydrogen dissociation, which

∗ Corresponding author. Tel.: +1 303 940 2335. E-mail addresses: [email protected] (S.N. Paglieri), [email protected] (M.D. Dolan), [email protected] (D. Chandra), [email protected] (K.M. Hubbard). 0376-7388/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.memsci.2011.04.049

enables hydrogen to pass easily through the metal. However, the precious metal Pd is expensive because of its scarcity and high demand for industrial use in catalysts (particularly autocatalysts), electronics and consumer uses such as dentistry and jewelry [11]. Pd also incurs a high environmental cost in terms of the energy consumed in the processes of mining and refining it. Furthermore, the supply of Pd is somewhat limited; globally there are only a few primary deposits of the Pt-group metals from which Pd is obtained. Although alternatives to Pd such as Ni are being investigated as surface catalysts for hydrogen dissociation on amorphous alloy membranes [12], only a thin (usually <1 ␮m thick) layer of Pd surface catalyst is required to accelerate the rate of hydrogen entry and exit from the amorphous membrane. The reactivity of the constituents of the amorphous alloy (such as Nb) towards oxygen results in the formation of surface oxides that impede hydrogen flux [13–15], even in 99.999% pure hydrogen. Therefore, the use of Pd is not eliminated on amorphous membranes, but it is reduced significantly, at least compared to foil or tubular Pd alloy membranes, which are usually at least 10 ␮m thick and often much thicker. Both amorphous or non-crystalline alloys and crystalline, non-Pd alloys with body-centered-cubic (BCC) structures are receiving increased attention for their potential to replace Pd alloys as hydrogen-separating membranes because of their promising

S.N. Paglieri et al. / Journal of Membrane Science 378 (2011) 42–50

attributes such as lower intrinsic cost, hydrogen permeabilities that are similar to Pd alloys and resistance to hydrogen embrittlement [15–21]. Prior work in the field of amorphous membrane materials has focused primarily on ternary and quaternary amorphous alloys based on Ni–Nb–Zr, because these materials combine high hydrogen permeability with mechanical stability in hydrogen [17,18,22–24]. Another practical advantage of amorphous metal membranes is the potential to scale-up their production by using the melt spinning technique to fabricate continuous rolls of thin foil. For example, 100-mm wide sheets of (Ni0.6 Nb0.4 )45 Zr50 Co5 amorphous alloy have been made [25]. Long-term membrane performance at typical operating conditions of 400 ◦ C in hydrogen is related to properties such as the thermal and mechanical stability. Thermal analysis studies have shown that amorphous alloys can crystallize over time at moderate temperatures (>300 ◦ C) and often below their crystallization temperatures (TX ) determined by differential scanning calorimetry (DSC) [26–28]. Crystallization impedes the use of amorphous metal alloys as membranes because it has a negative effect on hydrogen permeability, strength and ductility in hydrogen. Factors that determine the thermal stability include the ratio of Ni to the alloying metals, including their atomic radii, melting points and bond valences [18,29]. Crystallization temperature and the stability of the surface catalyst layer are interrelated and also dictate the thermal stability of the membrane, especially the change in hydrogen permeability over time [25,30,31]. Also, alloys that absorb high concentrations of hydrogen will have higher hydrogen permeability but lesser mechanical stability because of lattice expansion and the resultant stress [20]. The effects of adding other elements such as Hf, Ta, Co, Cu, Mo, Ti, and Al to Ni–Nb, Ni–Zr and Ni–Nb–Zr amorphous metal alloys has been studied in efforts to improve their properties [12,25,28,31–40]. For example, the addition of Co is known to reduce hydrogen embrittlement [12,23,39]. The refractory metal Ta has one of the highest melting points of all metals [41] and it is also very permeable to hydrogen [42–44]. Incorporating this analog of Nb into the Ni–Nb–Zr amorphous alloy should increase its thermal stability without drastically reducing the hydrogen permeability [12]. In fact, this has been verified in thermal analysis studies involving DSC [28,31] and permeability tests [32,45]. Still, the impact of Ta substitution on the hydrogen permeability, thermal stability and resistance to hydrogen embrittlement at typical operating temperatures for metal membranes is not well characterized. Therefore, in an effort to improve the properties of amorphous alloy membranes, we have systematically studied the effects of Ta addition on the performance of a series of Ni–Nb–Zr amorphous alloys with various Zr contents.

2. Methods 2.1. Membrane fabrication and characterization Amorphous alloy membranes were fabricated using planar flow casting [31]. Homogenized crystalline alloy ingots were prepared using argon arc melting, then induction-melted in a boron nitride nozzle under argon and ejected onto a rapidly rotating Cu–Be alloy wheel. The rapidly cooled molten alloy formed continuous, 30-mm wide amorphous alloy ribbons that varied between 50 and 90 ␮m in thickness. The amorphous nature of the free-solidified side of each foil was confirmed with X-ray diffraction (XRD, Cu K␣). Ribbons were coated with 500 nm of Pd on each side (with a 25-nm thick Cr layer to promote adhesion) using RF sputtering (150 W) under argon (10−6 MPa). Auger Electron Spectroscopy was conducted in ultra-high vacuum on untested and tested Pd coated membranes. Sputtering with Ar+ ions was performed to obtain compositional

43

Fig. 1. Schematic of planar membrane testing module. The membrane discs are 25 mm in diameter.

depth profiles. The atomic compositions are approximate because the elemental sensitivity factors used to calculate them were not calibrated using standards for all of the elements or combinations of elements. Different areas of the membrane were analyzed and similar results were obtained. 2.2. Thermal stability The thermal stability of the alloys was examined using a Setaram LabSys DSC system. Non-isothermal measurements were performed under flowing N2 with a heating rate of 20 ◦ C/min. The system temperature was calibrated using high-purity phase transformation and melting point standards. The crystallization temperature (Tx ) is defined as the intersection between the signal baseline and a tangent to the inflection point on the first crystallization exotherm. 2.3. Membrane permeability A membrane disc (25 mm diameter) was cut from the ribbon and sealed in a 316-stainless steel pressure test module using a graphite gasket as shown in Fig. 1. The active membrane area was 2.2 cm2 . Mechanical support was provided by a 0.5-␮m grade porous stainless steel disc (Mott Corp.), with a thin disc of porous alumina paper (Zircar Refractory Composites, Inc.) placed between the porous metal support and the membrane to prevent interdiffusion. All tubing and fittings were made of 316-stainless steel. The test module was centered in a tube furnace and connected to a gas manifold. Hydrogen or nitrogen gases (>99.9% pure or better) were metered in using mass flow controllers. The membrane and module was leak checked by pressurizing with nitrogen at room temperature. To prevent hydrogen embrittlement, both sides of the membrane were purged with nitrogen while heating to either 400 or 450 ◦ C before hydrogen flow was initiated. Even with this precaution, hydrogen induced embrittlement often caused membrane failure soon after the membrane was exposed to hydrogen. Therefore, some membranes were tested at reduced hydrogen partial pressure at higher temperature (450 ◦ C) to avoid cracking. Other studies have used even higher temperatures (500 ◦ C) or smaller diameter samples to avoid this problem [20,28,40,46]. Hydrogen was fed to the upstream (retentate) side of the module and permeated through the membrane. Excess hydrogen was fed to the membrane to avoid concentration polarization [47]. The retentate pressure was controlled by an automatic pressure control valve located downstream of the membrane module. Most membranes

20

25

30

35

were tested at a pressure differential across the membrane of 627 kPa. The feed pressure was 710 kPa and the permeate pressure was local atmospheric pressure (83 kPa) unless otherwise noted. The hydrogen feed pressure stabilized after a few minutes and the flow rate through the membrane was measured using a digital flowmeter. The overall system was controlled by Labview (National Instruments) and in case of overheating, over-pressurization, or a flammable gas leak, the software automatically followed a procedure to shut off the hydrogen and purge the membrane with nitrogen. 2.4. Membrane embrittlement resistance As a qualitative measure of the tendency to embrittle, membranes that survived permeability testing were cooled at 1 ◦ C/min in 203 kPa hydrogen upstream (unless otherwise noted) and 83 kPa hydrogen downstream to determine the temperatures at which they cracked. 3. Results and discussion 3.1. Hydrogen permeability Table 1 contains a compilation of the permeability and embrittlement data for all membranes tested in this study. The highest hydrogen permeability, 1.4 × 10−8 mol m−1 −1 Pa−0.5 , was observed for the membrane that contained the greatest zirconium content; (Ni0.6 Nb0.4 )70 Zr30 . For comparison, Table 1 also displays permeability data at 400–450 ◦ C for the commonly used Pd75 Ag25 alloy (at%) (1.8–2.0 × 10−8 mol m−1 s−1 −0.5 ) [9,48], and an amorphous Ni40 Nb20 Ta5 Zr30 Co5 alloy membrane (1.0 × 10−8 mol m−1 −1 −0.5 ) [12]. The highest hydrogen permeability measured for each alloy is plotted against Zr content in Fig. 2. Permeability increased with Zr content. At 20 or 30 at% Zr, the permeabilities of the alloys that contain Ta were slightly less compared to the (Ni0.6 Nb0.4 )70(80) Zr30(20) alloys tested at the same temperature. The permeability of a Ni60 Nb20 Zr20 membrane was lower than that of the other membranes containing 20 at% Zr, which was probably because it contained less Nb and more Ni, which has a much lower hydrogen permeability than Nb [49]. At lower Zr content (10 at%) there was little variation observed between the performance of the materials. The Ni60 Nb30 Ta10 alloy had the lowest permeability of all membrane compositions tested. 3.2. Long-term hydrogen permeability testing The change in hydrogen permeability over time for membranes with 20 or 30 at% Zr content tested at 400 or 450 ◦ C is shown in Fig. 3.

Table 1 Hydrogen permeabilities and failure conditions for a series of (Ni0.6 Nb0.4 )100−x Zrx and (Ni0.6 Nb0.3 Ta0.1 )100−x Zrx (where x = 0, 10, 20 or 30) amorphous alloy membranes.

Fig. 2. Relationship between Zr content and the hydrogen permeability of amorphous alloys with or without Ta at 400 and 450 ◦ C.

Ni60 Nb30 Ta10 Ni40 Nb20 Ta5 Zr30 Co5 [12] Pd75 Ag25 [9]

15

Ni60 Nb20 Zr20 (Ni0.6 Nb0.4 )90 Zr10 (Ni0.6 Nb0.3 Ta0.1 )90 Zr10

10

Concentration of Zr (at%)

(Ni0.6 Nb0.3 Ta0.1 )80 Zr20

5

(Ni0.6 Nb0.4 )80 Zr20

Ni60Nb30Ta10

0

(Ni0.6Nb0.3Ta0.1)80Zr20 (Ni0.6Nb0.4)90Zr10 Ni60Nb20Zr20

0.065 0.065 0.063 0.078 0.055 0.055 0.049 0.088 0.086 0.085 0.049 0.03 0.198

(Ni0.6Nb0.3Ta0.1)90Zr10

2.1E-09

(Ni0.6 Nb0.3 Ta0.1 )70 Zr30

4.1E-09

450 368 400 345 450 400 450 400 450 344 300 – 81 64 – – –

(Ni0.6Nb0.4)80Zr20

6.1E-09

0 at 50 h – – 5.6 × 10−9 2.7 × 10−10 – – – 0 at 46 h 6.3 × 10−10 1.1 × 10−9 1.2 × 10−9 1.3 × 10−9 at 86 h 2.6 × 10−10 – – –

8.1E-09

1.4 × 10−8 – – 9.8 × 10−9 1.2 × 10−8 8.3 × 10−9 – – 3.8 × 10−9 1.6 × 10−9 1.5 × 10−9 1.3 × 10−9 1.6 × 10−9 2.6 × 10−10 1 × 10−8 1.8 × 10−8 2.0 × 10−8

450°C

450 – – 400 450 400 450 400 450 400 400 400 450 400 400 400 450

(Ni0.6Nb0.3Ta0.1)70Zr30

400°C

1.0E-08

0.054 0.055

450°C, Ta alloy

1.2E-08

1.0E-10

(Ni0.6Nb0.4)70Zr30

400°C, Ta alloy

1.4E-08

(Ni0.6 Nb0.4 )70 Zr30

1.6E-08

517 kPa 119 kPa 32 kPa 18 kPa Broke in H2 after air and N2 purges 260 kPa 308 kPa 119 kPa Dehydrided in N2 203 kPa 119 kPa – 119 kPa 119 kPa – – –

S.N. Paglieri et al. / Journal of Membrane Science 378 (2011) 42–50

Membrane composition (at%) Thickness (mm) Test temperature (◦ C) Highest permeability (mol m−1 s−1 Pa−0.5 ) Permeability at 100 h (mol m−1 s−1 Pa−0.5 ) Failure temperature (◦ C) Failure pressure, P (kPa)

Permeability (mol m-1 s-1 Pa-0.5)

44

1.4E-08

(Ni0.6Nb0.4)70Zr30; 450°C

1.2E-08

(Ni0.6Nb0.3Ta0.1)70Zr30; 450°C

1.0E-08

(Ni0.6Nb0.3Ta0.1)70Zr30; 400°C

Permeability (mol m-1 s-1 Pa-0.5)

Permeability (mol m-1 s-1 Pa-0.5)

S.N. Paglieri et al. / Journal of Membrane Science 378 (2011) 42–50

(Ni0.6Nb0.3Ta0.1)80Zr20; 450°C

8.1E-09 6.1E-09 4.1E-09 2.1E-09 1.0E-10

2.0E-09

1.5E-09

1.0E-09

10

20

30

40

50

(Ni60Nb40)90Zr10; 400°C (Ni0.6Nb0.3Ta0.1)90Zr10; 450°C (Ni0.6Nb0.3Ta0.1)90Zr10; 400°C Ni60Nb30Ta10; 400°C

5.1E-10

1.0E-11

0

45

60

20

0

40

Time (hours)

3.3. Thermal properties Ideally, DSC studies on amorphous alloy membranes should be carried out under hydrogen. For example, Yamaura and Inoue found that the presence of hydrogen stabilized the amorphous phase and increased TX by about 20 ◦ C [12]. However, thermal analysis of non-hydrogenated alloys is still valuable for revealing

Permeability (mol m-1 s-1 Pa -0.5 )

A couple of trends are apparent in the data. First, increasing the Zr content resulted in increased permeability, but decreased thermal stability as has been previously reported [25,31]. Second, Ta lowered the hydrogen permeability but increased the thermal stability, although the increase in thermal stability was not pronounced. At 450 ◦ C, the permeability of the (Ni0.6 Nb0.4 )70 Zr30 membrane was briefly the highest measured in this study, but it decreased rapidly; no flux was detectable after 49 h of testing. A (Ni0.6 Nb0.3 Ta0.1 )70 Zr30 membrane started with a slightly lower permeability at 450 ◦ C, but its performance degraded more slowly because of the stabilizing effect of Ta. The exact reason for the inflection in the permeability curve is not completely understood, but may be related to a combination of the amorphous alloy crystallization kinetics and changes in the properties of the catalytic surface as interdiffusion takes place between the Pd coating and the amorphous alloy. The (Ni0.6 Nb0.4 )70 Zr30 membrane material often embrittled at 400 ◦ C, which made it challenging to test. However, a (Ni0.6 Nb0.3 Ta0.1 )70 Zr30 membrane was successfully tested at 400 ◦ C and had a permeability that was less than that measured at 450 ◦ C for the same material. Whereas a (Ni0.6 Nb0.3 Ta0.1 )70 Zr30 membrane had a slightly lower permeability, it experienced stable performance over a much longer time at 400 ◦ C. A (Ni0.6 Nb0.3 Ta0.1 )80 Zr20 membrane tested at 450 ◦ C had less than half of the initial flux of the (Ni0.6 Nb0.3 Ta0.1 )70 Zr30 membrane and its permeability was reduced to below detectable levels in about 40 h. A Ni60 Nb20 Zr20 membrane began with low, but stable permeability; it took 300 h for its permeability to become too low to measure (Fig. 5). In contrast to the alloys with higher Zr content, the (Ni0.6 Nb0.4 )90 Zr10 and (Ni0.6 Nb0.3 Ta0.1 )90 Zr10 membranes had very stable hydrogen permeability over the course of a 100 h test (Fig. 4). There was little discernable difference in performance between the alloys with or without Ta substitution, other than the permeability of the (Ni0.6 Nb0.3 Ta0.1 )90 Zr10 membrane was higher at 450 ◦ C than at 400 ◦ C. The most stable membrane, Ni60 Nb30 Ta10 , contained no Zr but also had the lowest permeability (2.6 × 10−10 mol m−1 s−1 Pa−0.5 ) by almost a factor of 10 (Fig. 4). The reason that the permeability increased slightly with time for this membrane is not presently understood.

100

120

Fig. 4. Evolution of hydrogen permeability for 10 at% Zr amorphous alloys at 400 and 450 ◦ C with 709 kPa pure hydrogen feed pressure (83 kPa permeate).

1.6E-09 1.4E-09 1.2E-09 1.0E-09 8.0E-10 6.0E-10 4.0E-10 2.0E-10 1.0E-12

0

50

100

15 0

200

25 0

300

35 0

Time (hours) Fig. 5. Evolution of hydrogen permeability for a Ni60 Nb20 Zr20 amorphous alloy at 400 ◦ C with 709 kPa pure hydrogen feed pressure (83 kPa permeate).

composition-stability trends and the trends observed in the data for an inert atmosphere can be used to infer thermal stability trends under hydrogen [31]. The results of the non-isothermal DSC study are shown in Fig. 6. Crystallization temperature (TX ) decreased with increasing Zr content in the alloy. Adding Zr clearly reduces TX , as has been previously demonstrated [28,31]. Comparing the alloys that contained no Zr, an increase in TX of ∼90 ◦ C was observed for the Ni0.6 Nb0.4 alloy with added Ta. For alloys that contained Zr, Ta substitution for Nb raised TX slightly and by approximately the same small amount (10–22 ◦ C) at each Zr composition. Comparing the thermal properties (Fig. 6) and hydrogen permeabilities measured for each alloy (Fig. 2), there appears to be an inverse trend where higher TX equates to lower permeability. In other words, permeability is proportional to Zr content. Further-

800 750 (Ni0.6Nb0.3Ta0.1)100-xZrx

700

TX (°C)

Fig. 3. Evolution of permeability for 20 and 30 at% Zr alloys at 400 and 450 ◦ C with 709 kPa pure hydrogen feed pressure (83 kPa permeate). The (Ni0.6 Nb0.4 )70 Zr30 and (Ni0.6 Nb0.3 Ta0.1 )80 Zr20 membranes were tested at a feed pressure of 184 kPa (83 kPa permeate).

60 80 Time (hours)

650 600

(Ni0.6Nb0.4)100-xZrx

550 500

0

5

10

15 mol% Zr

20

25

30

Fig. 6. Crystallization temperatures (TX ) of Ni–Nb–Zr and Ni–Nb–Zr–Ta amorphous alloys determined using non-isothermal differential scanning calorimetry (DSC). Corresponding TX values refer to the onset of crystallization defined as the intersection between the signal baseline and a tangent to the inflection point on the first crystallization exotherm.

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S.N. Paglieri et al. / Journal of Membrane Science 378 (2011) 42–50

Fig. 7. XRD results for Ni60 Nb20 Zr20 amorphous alloy foil that was either (a) uncoated or (b) Pd coated and unheated or tested at 400 ◦ C for 300 h, and (c) closeup of the Pd(1 1 1) peak showing the shift and broadening after testing.

more, the alloys with substituted Ta (higher TX ) were moderately less permeable at each specific Zr composition. As may be expected, there also seems to be a correlation between TX and thermal stability during-long term testing, with the higher TX materials exhibiting more stable hydrogen flux during long-term permeation tests (Figs. 3 and 4). The membranes that contained Ta appeared to have slightly better thermal stabilities during long-term testing. 3.4. XRD analysis and AES depth profile The XRD results for the Ni60 Nb20 Zr20 membrane are shown in Fig. 7. XRD patterns of the uncoated (untested) membrane showed one broad peak characteristic of an amorphous material (Fig. 7a). For the Pd coated membrane before and after testing (Fig. 7b), the only peaks observed were for the Pd surface coating (except

for an unidentified peak at 39◦ 2) and no crystallization of the underlying amorphous alloy was detected. However, Jayalakshmi et al. have shown that nano-crystallization took place in hydrogen charged Zr- and Ni-based metallic glasses even though no peak could be detected in the XRD traces [26,50]. At temperatures below the glass transition temperature (Tg ) and TX , short-range diffusion of atoms can cause structural relaxation. This local atomic rearrangement and reduction of the free volume could decrease the hydrogen permeability of the membrane [26]. Therefore, further structural characterization of the bulk membrane material using a method with finer resolution such as TEM is required to determine if nanocrystallization is partly responsible for the decrease in hydrogen permeability with time. After testing at 400 ◦ C for 300 h the Pd peaks were shifted to a slightly higher angle, diminished in height, and broadened (Fig. 7c). The slight decrease in lattice parameter may be attributed to alloying with Ni from the bulk membrane or Cr from the adhesion layer. The peak broadening could also be related to crystallize size or microstrain in the Pd layer. The results of AES depth profiles on untested and tested samples of Pd coated Ni60 Nb20 Zr20 membrane are shown in Fig. 8. For the unheated sample, the profile shows the expected multilayer structure of the Pd coated amorphous alloy membrane (Fig. 8a), including the pure Pd coating, the Cr layer, and the NiNbZr substrate. The Cr layer contained a significant concentration of oxygen, which was likely gettered from the PVD chamber during deposition. The multilayer structure of the untested membrane shows distinct layers without significant intermixing. The depth extent of the interface, Pd–Cr and Cr–NiNbZr, is typical for Auger depth profiles and is a function of sputter depth resolution and Auger analysis depth resolution. The shape of the interface/transition between layers is consistent with distinct interfaces. The total depth or thickness of the Pd and Cr(O) layers cannot be quantitatively stated without further analysis of similar standards. The nominally deposited thicknesses of the layers was 500 nm Pd and 25 nm Cr. Although the Auger depth profiles appear to show the Pd is thinner than 500:25 relative to the Cr, this cannot be directly taken from the depth profiles plots as presented. Different metals will sputter at different rates; depending on the conditions, Pd can have sputter yields two times or more than that of Cr. Depth profiles for the heated sample show that there was significant metallic interdiffusion of the layers (Fig. 8b). The interface is shifted to the right because a slightly lower sputtering rate was used. The multilayer structure of the tested sample shows that the distinct layers of the unheated sample had intermixed significantly and the interfaces are much broader. From the outer surface in towards the substrate we observe several interesting features: (1) Cr(O) at the outer surface, (2) significant Cr within the Pd coating layer as well as some O, (3) Ni, Nb and Zr had diffused into the Pd layer, (4) a residual Cr(O) layer (3) Ni, Nb and Zr that is less distinct than that of the unheated sample, (5) O and Pd had diffused into the NiNbZr substrate. The Cr that had diffused into the Pd layer was present at a fairly constant level throughout the Pd layer. Additionally, Cr appears to have diffused to the outer surface where the initial Auger spectra showed Cr and O and not Pd. Pd appears to have diffused through the Cr interface layer as evidenced by the fact that Pd was observed within the NiNbZr (even though Cr was not). The O that was initially present solely within the Cr interface had diffused into the NiNbZr alloy. Although not shown in the depth profile plots, C was analyzed for, but there is a direct overlap with a Pd Auger peak. However, there did not appear to be any significant C contamination within the membranes. The AES results may explain the decrease in hydrogen permeability observed during testing at 400 ◦ C. Metallic interdiffusion can reduce the hydrogen permeability of Pd coated metallic mem-

S.N. Paglieri et al. / Journal of Membrane Science 378 (2011) 42–50 100

80

Approx. Atomic Concentration (%)

(a)

Pd Cr O NiNbZr

Pd

NiNbZr

60

40

Cr

O

20

0 00:00

00:15

00:30

00:45

01:00

01:15

01:30

01:45

Sputter Time (hr:min)

Pd Cr O NiNbZr

Approx. Atomic Concentration (%)

80

Pd

(b)

NiNbZr

60

40

O

20

Cr

0 00:00

00:15

00:30

00:45

01:00

01:15

01:30

01:45

Sputter Time (hr:min) Fig. 8. AES depth profiles on a Ni60 Nb20 Zr20 amorphous membrane coated with 500 nm of Pd that was (a) not heated or (b) tested at 400 ◦ C for 300 h. The sputtering rate for the heated sample (b) was slightly lower, which is why the interface is shifted to the right compared to the unheated sample.

branes [25,51–55]. Surface segregation of Cr and O onto the surface of the tested membrane would be expected to retard the rate of hydrogen dissociation and diffusion into the bulk Pd. Furthermore, considerable O had diffused into the NiNbZr substrate, which could also lower permeability and contribute to embrittlement. 3.5. Hydrogen embrittlement Hydrogen absorption that causes lattice expansion is a common failure mechanism for metal membranes because the fracture strength decreases with increasing hydrogen concentration [40,56]. Results of our qualitative hydrogen embrittlement tests are shown in Table 1. Cooling the membranes in hydrogen after testing was found to provide an estimate of their resistance to hydrogen embrittlement [57]. Also, membranes often failed dur-

47

ing testing, although crack formation usually occurred when the membrane was first exposed to hydrogen or during pressurization with hydrogen. Dehydriding by purging with inert gas at the end of permeation testing could also crack the membrane. Membrane resistance to hydrogen embrittlement increased with decreasing Zr content because of the high solubility of hydrogen in Zr [58,59]. Ta reduced the hydrogen embrittlement even though preliminary measurements using a Sieverts’ apparatus show that Ta increases the hydrogen solubility in the amorphous alloy [60]. Apparently the addition of Ta decreases the tendency for the alloy to form a brittle hydride phase. For example, the (Ni0.6 Nb0.3 Ta0.1 )90 Zr10 and Ni60 Nb30 Ta10 alloys with ≤10 at% Zr that contained Ta did not crack under hydrogen until they were below 100 ◦ C (see Table 1). Ideally, a membrane material would withstand thermal cycling in hydrogen to 25 ◦ C without failing, like Pd75 Ag25 . The (Ni0.6 Nb0.4 )80 Zr20 and (Ni0.6 Nb0.4 )70 Zr30 membranes often embrittled and failed upon initial exposure to hydrogen at 400 ◦ C (see Table 1). For example, at 400 ◦ C a (Ni0.6 Nb0.4 )80 Zr20 alloy membrane broke under a hydrogen pressure differential of 260 kPa (the permeability data in Table 1 was measured just prior to membrane failure), although a (Ni0.6 Nb0.3 Ta0.1 )80 Zr20 alloy membrane broke while attempting to test it at 400 ◦ C and 119 kPa. To reduce the hydrogen concentration in the material and enable the collection of permeability data for the (Ni0.6 Nb0.3 Ta0.1 )80 Zr20 and (Ni0.6 Nb0.4 )70 Zr30 membrane materials, the test temperature was increased to 450 ◦ C and the pressure differential across the membrane was reduced to 101 kPa. The Ta substituted alloys embrittled significantly less, which made it possible to obtain permeability data for the (Ni0.6 Nb0.3 Ta0.1 )70 Zr30 alloy at 400 and 450 ◦ C and a pressure differential across the membrane of 627 kPa. At the end of the permeability test, instead of reducing the temperature at 1 ◦ C/min the hydrogen pressure was slowly increased to assess the membrane’s durability; it cracked at a pressure differential of 517 kPa. A (Ni0.6 Nb0.3 Ta0.1 )80 Zr20 membrane was also successfully tested at 450 ◦ C and a pressure differential across the membrane of 101 kPa. Testing at 450 ◦ C would also be expected to accelerate any degradation mechanisms such as crystallization and interdiffusion between the Pd coating and the membrane, which enabled better comparison of the performance of the various membrane materials under these harsher conditions. Pictures of tested membranes illustrate some interesting features, such as the effects of hydrogen absorption, embrittlement and heat treatment on membrane structure (Fig. 9). Fig. 9a shows the retentate side of a (Ni0.6 Nb0.4 )70 Zr30 membrane that was heated to 400 ◦ C but cracked almost immediately when exposed to hydrogen. Note the horseshoe shaped crack that initiated near the edge of the membrane, just inside sealing surface. Fig. 9b shows the permeate side of a (Ni0.6 Nb0.4 )80 Zr20 membrane that cracked upon exposure to 308 kPa hydrogen at 450 ◦ C. Amorphous alloys are known to undergo devitrifying phase transformations over time at elevated temperatures [28,59,61]. Therefore, crystallization at 400 ◦ C could result in an increased tendency for decohesion upon hydrogen absorption, but here it seems to primarily be the expansion induced by hydrogen absorption that resulted in crack formation, as evident in the photograph in Fig. 9a. Fig. 9c and d shows the retentate and permeate sides of a (Ni0.6 Nb0.3 Ta0.1 )70 Zr30 membrane after testing at 450 ◦ C for 60 h (data shown in Table 1 and Figs. 2 and 3). The membrane cracked when it was exposed to a 3 min air purge, probably because of the stress caused by dehydriding. Partial crystallization of the amorphous structure may have also resulted in decreased ductility that increased the tendency for cracking, but a more detailed structural

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Fig. 9. Pictures of membranes tested at 400 or 450 ◦ C. (a) Retentate side of a (Ni0.6 Nb0.4 )70 Zr30 membrane that cracked when exposed to hydrogen at 400 ◦ C. Note the hairline crack. (b) Permeate side of a (Ni0.6 Nb0.4 )80 Zr20 membrane that cracked upon exposure to 308 kPa hydrogen at 450 ◦ C. Note: the graphite gasket is underneath the membrane, stuck to the retentate side. (c and d) Retentate and permeate sides of a (Ni0.6 Nb0.3 Ta0.1 )70 Zr30 membrane showing the effects of hydrogen absorption. The membrane was tested for ∼60 h at 450 ◦ C and then it cracked during a 3 min air purge. (e and f) Retentate and permeate sides of a (Ni0.6 Nb0.4 )80 Zr20 membrane heated to 450 ◦ C under flowing nitrogen and then cooled, but not exposed to hydrogen. The blue–green color is believed to be PdOx . (For interpretation of the references to color in this figure legend, the reader is referred to the web version of the article.)

analysis (such as TEM) would be required to determine if that was a factor here. Membrane cracking was often observed to initiate around the inner edge of the graphite gasket where the membrane was constrained. For example, Fig. 9b is the permeate side of a (Ni0.6 Nb0.4 )80 Zr20 membrane that cracked around the perimeter upon exposure to 308 kPa hydrogen at 450 ◦ C. The (Ni0.6 Nb0.4 )80 Zr20 membrane expanded slightly, but not as much as the (Ni0.6 Nb0.4 )70 Zr30 and (Ni0.6 Nb0.3 Ta0.1 )70 Zr30 membranes shown in Fig. 9a, c and d. These two membranes with higher Zr content bowed out towards the retentate side when exposed to hydrogen. This happened not only because the membrane was prevented from expanding away from the high pressure feed towards the permeate side by the porous stainless steel frit support, but the higher hydrogen pressure on the retenate side of the membrane resulted in a higher hydrogen concentration in the membrane causing it to bow out in that direction [62].

A (Ni0.6 Nb0.4 )80 Zr20 membrane was heated to 450 ◦ C, but was determined to have formed a pinhole during heating as determined by nitrogen pressure tests before and after heat-up. After cooling in nitrogen and removal from the test module, the retentate surface of the membrane appeared blue–green, most likely because of the formation of Pd-oxide (PdOx ) on the surface (Fig. 9e). However, XRD was unable to detect the presence of PdOx , probably because the layer was too thin. The permeate side also had a blue–green blotch (Fig. 9f), where it looked like the oxygen had bled through a small defect in the membrane such as a pinhole. The oxygen was likely gettered from the industrial grade nitrogen used as a purge gas on the retentate side of the membrane during heat up and cooling. The permeate side of the membrane was purged by a different nitrogen cylinder. Other tested membranes did not exhibit this coloration because the PdOx that may have formed would have been reduced by hydrogen during testing, since PdOx is not a stable oxide and is easily reduced [63,64].

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4. Summary and conclusions The performance of a series of melt spun amorphous alloy membranes consisting of Ni60 Nb20 Zr2 0 (at%), (Ni0.6 Nb0.4 )100−x Zrx or (Ni0.6 Nb0.3 Ta0.1 ) 100−x Zrx (where x = 0, 10, 20 or 30), was evaluated. Here is an overview of the results. 4.1. Hydrogen permeability Permeability increased with increasing Zr content and the initial permeabilities of the (Ni0.6 Nb0.4 )70 Zr30 and (Ni0.6 Nb0.3 Ta0.1 )70 Zr30 membranes approached that of Pd75 Ag25 (1.0–1.4 vs. 1.8–2.0 × 10−8 mol m−1 s−1 Pa−0.5 at 400–450 ◦ C). The partial substitution of Nb with Ta resulted in slightly lower permeability. Alloys with ≤10 at% Zr content had very low permeabilities; <2 × 10−9 mol m−1 s−1 Pa−0.5 .

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through judicious alloy development to increase the thermal and structural stability of the membrane and coating while maintaining high hydrogen permeability. Acknowledgements This material is based upon work supported by the Department of Energy National Energy Technology Laboratory under Award Number DE-FE0000998. Amorphous alloys were prepared by Michael Kellam at the Commonwealth Scientific and Industrial Research Organisation. The authors wish to express their appreciation to Sarah J. DeVoss, Rita Dubovik, Vladimir Y. Belits, Tyler B. Gleditsch and Kerry A. Libberton for assistance with experimental setup, XRD, data collection and photography. References

4.2. Thermal stability The rapid decrease in hydrogen flux with time at 400 or 450 ◦ C for alloys containing higher Zr contents was attributed to metallic interdiffusion between the Pd surface coating and the membrane alloy and also possibly to changes in bulk membrane structure that were undetectable by XRD, which is consistent with the results of previous studies [25,26,28,65,66]. Ta seemed to modestly increase the thermal stability of the amorphous Ni–Nb–Zr alloy by decreasing the rate of permeation decay during long term tests with only a slight impact on hydrogen permeability. Alloys containing ≤10 at% Zr displayed very stable, but low, hydrogen permeabilities throughout the 100 h tests. Crystallization temperatures (TX ) measured using nonisothermal DSC displayed lower thermal stability at higher Zr contents. Therefore, higher TX meant lower hydrogen permeability for these alloys. The addition of Ta resulted in a modest increase in TX for the alloys that contained Zr. 4.3. AES depth profile analysis An AES depth profile of a Ni60 Nb20 Zr20 membrane that was tested for 300 h showed Cr diffusion into and through the Pd coating to the surface, mostly O on the surface, and Pd and O diffusion into the NiNbZr. The change in membrane composition near the surface due to metallic interdiffusion was most likely the primary cause for the reduction in hydrogen permeability of the membranes over time at 400–450 ◦ C. 4.4. Hydrogen embrittlement Mechanical stability or resistance to brittle fracture decreased with increasing Zr content, although substituting Ta for Nb clearly decreased the membranes’ susceptibility to hydrogen embrittlement. For example, (Ni0.6 Nb0.4 )70 Zr30 often embrittled too extensively to conduct permeability testing at 400 ◦ C. Incorporation of Ta enabled the testing of (Ni0.6 Nb0.3 Ta0.1 )70 Zr30 membranes at 400 ◦ C at relatively high hydrogen pressure, but this material was still unsound after long term hydrogen permeability testing. Pictures of the tested membranes showed that the materials expanded significantly in the presence of hydrogen, which contributed to crack formation. Even though some of the amorphous alloy membranes we tested exhibited relatively high hydrogen permeabilities that were comparable to Pd alloys, these particular compositions of amorphous alloys may not be suitable for long term hydrogen separation at high temperatures because they are unstable with respect to hydrogen flux over time and they are also prone to embrittlement. However, it may be possible to create a functional membrane material

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