Article pubs.acs.org/cm
Tunable Nanoporous Metallic Glasses Fabricated by Selective Phase Dissolution and Passivation for Ultrafast Hydrogen Uptake Wei Jiao,† Pan Liu,†,‡ Huaijun Lin,§ Wei Zhou,‡ Zhi Wang,† Takeshi Fujita,† Akihiko Hirata,†,¶ Hai-Wen Li,∥ and Mingwei Chen*,†,‡,⊥,@ †
WPI Advanced Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200030, China § Jiangsu Collaborative Innovation Center for Advanced Inorganic Function Composites, School of Materials Science and Engineering, Nanjing Tech University, Nanjing 210009, China ∥ International Research Center for Hydrogen Energy, Kyushu University, Fukuoka 819-0395, Japan ⊥ CREST, JST, 4-1-8 Honcho Kawaguchi, Saitama 332-0012, Japan @ Department of Materials Science and Engineering, Johns Hopkins University, Baltimore, Maryland 21218, United States ¶ Mathematics for Advanced Materials-OIL, AIST-Tohoku University, Sendai 980-8577, Japan ‡
S Supporting Information *
ABSTRACT: Realizing a large specific area in disordered metallic glasses is of great scientific and technological importance. Here we report a nanoporous multicomponent metallic glass fabricated by the combination of selective phase dissolution and passivation of a spinodally decomposed glassy precursor. The nanoporous metallic glass shows superior hydrogen uptake performance by taking advantage of the large specific surface area of the nanoporous structure and the high diffusivity of hydrogen in metallic glasses. The facile route of selective corrosion and passivation, decoupling the galvanic corrosion and alloy stability, opens a new avenue for functionalizing metallic glasses as a large-surface area and lightweight material for various structural and functional applications.
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INTRODUCTION Arising from a disordered atomic structure, metallic glasses (MGs) have many unique properties in mechanics, chemistry, and physics that cannot be obtained from crystalline alloys.1−4 However, the applications of MGs, such as in catalysis,5,6 sensing,7 gas absorption,8 dye degradation,9 and membrane filters,10 are often limited by the small specific surface area. One of the appropriate solutions for achieving a large specific surface area is the introduction of pores into MGs. As a result, porous MGs have been fabricated by water vapor release, salt dissolution, and low-pressure infiltration of hallow carbon microspheres.11−13 These porous MGs with large pore and skeleton sizes of ∼200 μm show exceptional mechanical properties, particularly in compression ductility.14 However, the large pore sizes cannot help too much in the improvement in the specific surface area, which is critical for many functional applications. By using an anodic alumina oxide as the template, one-dimensional (1D) pores with a size of ∼100 nm have been fabricated in 1 μm thick MG films by physical vapor deposition.15 However, the low porosity and 1D nature of pore channels still cannot solve the issue of the small specific surface area of MGs. By using MGs as the precursors, nanoporous metals and alloys with a large surface area and © 2017 American Chemical Society
high porosity can be achieved by selectively dissolving less stable components through galvanic corrosion based on the disparity of electrode potentials among the component elements.16−20 Nevertheless, the resultant nanoporous metals and alloys become crystalline, and thus, the unique properties of the MGs cannot be preserved.10 By contrast, the selective phase dissolution method can partially preserve the atomic and microscale structure of biphase precursors.21−25 Because of the complex electrochemical dissolution behavior in multiphase MGs, the relationship between the initial phase-separated microstructure and the final structure is not straightforward.20,25 On the basis of the disparity in the transpassive dissolution behavior of various elements, selective dissolution can be achieved through excavation of the crystalline phase with a lower transpassive potential.21,22 Here we report the successful fabrication of three-dimensional (3D) bicontinuous nanoporous MGs by selective corrosion of spinodally decomposed MG precursors and passivation. In addition, the pore size can be intentionally tuned by the underlying phase Received: March 13, 2017 Revised: May 2, 2017 Published: May 3, 2017 4478
DOI: 10.1021/acs.chemmater.7b01038 Chem. Mater. 2017, 29, 4478−4483
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or 29) precursor alloys. The precursor ingots were remelted and injected onto a rotating copper roller in an induction furnace under an argon atmosphere. The cooling rate of melt-spun ribbons was adjusted by tuning the speed of the rotating copper roller from 500 to 4K rpm. The ribbons were cut in small pieces ∼3 mm in width and ∼20 mm in length for the fabrication of nanoporous metallic glasses. Nanoporous Glassy Alloy Fabrication. The ribbons were freely corroded in 0.1 mol L−1 nitric acid for 24 h to obtain the nanoporous product. To accelerate the fabrication process, the ribbons were also selectively etched through galvanic corrosion at a fixed potential of 0.4 V, which is lower than the breakdown potential of passivity (Vbp). Electrochemical corrosion was conducted in an electrochemical workstation (lvium Technology). A three-electrode setup was used for the corrosion, in which a Ag/AgCl electrode is used as the reference electrode, a Pt sheet as the counter electrode, and a precursor ribbon as the working electrode. The selectively etched ribbons were rinsed five times with deionized water and alcohol to remove the remaining chemicals and then dried in a vacuum chamber. Structural Characterization and Chemical Composition Analysis. The structures of the precursor alloys and nanoporous products were determined by X-ray diffraction (XRD) by using a Rigaku Ultima X-ray diffractometer with Cu Kα radiation. Highresolution transmission electron microscopy (HR-TEM), selected area electron diffraction (SAED), high-angle annular dark-field scanning TEM (HAADF-STEM), and energy-dispersive X-ray spectroscopy (EDS) were performed by using a Cs-corrected transmission electron microscope (JEOL JEM-2001F) equipped with energy-dispersive Xray spectroscopy. The morphology of the nanoporous glassy alloys was observed by using a JEOL JIB-4600F scanning electron microscope. DSC experiments were performed with a PerkinElmer 8500 instrument under a purified argon atmosphere at a heating rate of 20 K/min. Hydrogen Uptake Experiment. Hydrogenation experiments for all samples were conducted at 200 °C with an initial pressure of around 2.3 MPa for 1 h, in a Sieverts-type automatic gas reaction controller (AMC Co.). Prior to the hydrogenation experiments, the samples were vacuumed at 200 °C for 2 h under a pressure of ∼2.7 × 10−3 Pa. The loading amounts of the nanoporous metallic glass for hydrogenation measurements were ∼0.1 g. TG experiments were performed in a Rigaku TG/DTA-8120/s instrument under a He flow of 100 mL/min at a heating rate of 5 K/min, and the differential scanning calorimeter (DSC) experiments were conducted on a Rigaku Thermo plus EV02 DSC8231 instrument. For the reference sample of the nanoporous crystalline counterpart, it is obtained through annealing at 600 °C for 2 h. The specific area is the same as that of nanoporous metallic glass. No pore collapse or ligament growth in the nanoporous crystalline counterpart has been observed via scanning electron microscopy (SEM) either.
separation process. The combination of a prominent disordered nature and a large surface area makes nanoporous MGs applicable for gas absorption, as demonstrated by the high hydrogen uptake capacity at ultrahigh rates.
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EXPERIMENTAL METHODS
Fabrication Principle. In conventional selective etching, the formation of nanoporosity is considered to be the result of a spinodal decomposition like reaction at electrode−electrolyte interfaces.26 In this study, spinodally decomposed MGs are utilized as the precursors for nanopore formation by selective phase dissolution and passivation. We designed a pseudobinary glass system of Zr47Cu46Al7 and Y47Cu46Al7, which has a miscibility gap (Figure 1a) originating from
Figure 1. Schematic illustration of the fabrication principle. (a) Miscibility gap in the phase diagram of the pseudobinary system of Zr47Cu46Al7 and Y47Cu46Al7. (b) 3D interconnected network of each component located in the spinodal decomposition region. (c) Distinct chemical corrosion behavior of two components. (d) Porous structure formed via preferential corrosion of the Y47Cu46Al7 phase. The light blue part on the surface of the ligament highlights the formation of the passivation layer. the positive heat of mixing (+35 kJ/mol) between Zr and Y.27 By tuning the contents of two components into the spinodal decomposition region, we found two interpenetrated and bicontinuous glass phases form a 3D network-like structure as represented in Figure 1b.28,29 Two glass components have distinct corrosion behaviors and electrochemical stabilities in a nitric acid solution as shown in Figure 1c. Compared to the Y47Cu46Al7 glassy phase, the Zr47Cu46Al7 glassy component is relatively inert because of surface passivation. The rate of corrosion of the Zr47Cu46Al7 glassy phase is ∼4 orders of magnitude lower than that of the Y47Cu46Al7 phase, characterized by corrosion current density icorr at open circuit potentials. Furthermore, the corrosion current density of the Zr47Cu46Al7 glassy phase further decreases when the applied potentials are >0.4 V greater than passivation potential Epassivation. The stability of the Zr47Cu46Al7 glassy phase can be attributed to the inherent passivity propensity of Zr.30 The passivation layer is expected to slow and even prevent the further corrosion of the Zr47Cu46Al7 phase. By utilizing the difference in the electrochemical stability of two glassy phases, the Y47Cu 46Al7 component can be selectively dissolved while the relatively stable Zr47Cu46Al7 glass is retained as the skeletons of a nanoporous structure (Figure 1d). Precursor Preparation. Master alloys Zr47Cu46Al7 and Y47Cu46Al7 were made by arc-melting the mixtures of pure elements in a Tigettered high-purity argon atmosphere. To ensure the homogeneity of the chemical composition, the ingots were reversed and remelted at least three times. The pseudobinary glass precursors were prepared by mixing Zr47Cu46Al7 and Y47Cu46Al7 in the weight ratios calculated from the relative content of Zr and Y in the Cu46Zr(47−x)Al7Yx (x = 18, 23,
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RESULTS AND DISCUSSION Spinodally Decomposed Glass Precursor. The continuous phase separation in the designed precursor alloys is illustrated by STEM observation and EDS analysis as shown in Figure 2. The contrast variation in the HAADF-STEM image indicates that the precursor is chemically inhomogeneous (Figure 2a). Electron diffraction (SAED), featuring the characteristic diffraction halo, verifies the glassy nature of the precursor (Figure 2b). The chemical heterogeneity, originating from the spinodal decomposition, is further confirmed by the STEM−EDS mapping (Figure 2c−f). Zr and Y are distributed in spatial distinct regions with a nearly identical domain size of ∼20 nm. The spatial distribution of Cu is uniform and does not show obvious partitions in either Zr-rich or Y-rich glass domains. The distribution of Al is nearly the same as that of Zr (Figure 2f), indicating that Al slightly enriches in the Zr-rich glassy phase. The bright regions in the HAADF image correspond to the Zr-rich glassy phase, while the dark region is associated with the Y-rich one. The fluctuation in the 4479
DOI: 10.1021/acs.chemmater.7b01038 Chem. Mater. 2017, 29, 4478−4483
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Figure 2. Characterization of the structure and chemical composition of the precursor. (a) HAADF-STEM image of the precursor. The marked region is used for energy-dispersive X-ray spectroscopy element mapping. (b) Selected area electron diffraction of the precursor. STEM−EDS mapping of (c) Zr, (d) Al, (e) Cu, and (f) Y in the precursor. The scale bar is 50 nm for panel a, 5 nm−1 for panel b, and 20 nm for panels c−f.
Figure 3. Morphology of the porous alloy. (a and b) SEM images of the residual alloy (panel b is highly magnified). The speed of the copper roller is 4K rpm. The pore sizes were tuned by cooling rate and characterized by the speed of the copper roller: (c) 2K, (d) 1K, and (e) 500 rpm. (f) Variation of the pore size and specific surface area (SSA) with the rotation rates of the copper roller. Scale bars are 500 nm for panel a, 100 nm for panel b, and 1 μm for panels c−e.
chemical composition of the precursor is further confirmed by small-angle X-ray scattering (SAXS) (Figure S1). Estimated from the relationship between peak position qmax and correlation length ξ,31 it is ∼40 nm, quite consistent with the average wavelength (the total sizes of Zr-rich and Y-rich glassy domains) of spinodal decomposition in the MG precursor imaged by SEM−EDS mapping. Morphology of the Nanoporous Metallic Glass. By selective phase dissolution and passivation in a nitric acid solution, the nanoporous structure can be obtained via open circuit corrosion or electrochemical corrosion at a constant potential that is lower than the breakdown potential of passivity Vbp of the precursor (Figure S2). As shown in the SEM images of panels a and b of Figure 3, three-dimensional bicontinuous nanoporous structure uniformly spans the thickness of the ribbons. The average pore size is ∼25 nm, as determined by the fast Fourier transform analysis of the SEM images32 and Brunauer−Emmett−Teller measurements, which are in line with the domain size of the Zr-rich glassy phase (Figure S3). With the formation of the nanoscale pores, a large specific surface area (SSA) of 22.7 m2/g can be obtained. The chemical composition of the solid skeletons in the nanoporous structure is measured to be Zr43Cu40Al5Y7 by SEM−EDS. Compared with the composition of the precursor, the content of Y and Cu decreases greatly, indicating that the pore structure comes from the selectively dissolved Y-rich glassy phase. The absence of the Y-rich glassy phase in the nanoporous alloy is further supported by the disappearance of the crystallization peaks of the Y-rich glass phase in the DSC curve as shown below. These results demonstrate the porous structure is determined by the morphology of the spinodally decomposed precursor alloy. Thus, the porous structure and porosity can be tailored by manipulating the spinodal decomposition processes. According
to the spinodal decomposition theory,33 critical wavelength Λc, closely related to the porous size, is dependent on composition according to the relationship Λc = [−8π2k/(∂2f/∂c2 + 2η2γ)]1/2, where ∂2f/∂c2 is the second derivative of free energy with composition c.34 Hence, the nanopore size can to be tailored by changing the ratio of the Zr47Cu46Al7 and Y47Cu46Al7 phases. Apparently, the porosity of the nanoporous MG also changes with the fraction of the sacrificed Y47Cu46Al7 phase. Moreover, on the basis of spinodal decomposition kinetics, wavelength Λ evolves with time at a constant temperature in the fashion of t1/3.34 Accordingly, the pore size can be further tailored by controlling the cooling rate, holding time, and injection temperature within the metastable miscibility gap. Experimentally, we noticed that the control of kinetic processes is an effective approach for adjusting the pore sizes. As shown in Figure 3c−f, the pore sizes can be modified in a controllable mode from ∼25 nm to 1 μm by simply decreasing the cooling rates that are characterized by the rotation rate of copper roller from 4000 to 500 rpm during melt spinning. Associated with the decrease in pore size, the SSA increases with cooling rate, as shown in Figure 3f. Glassy Nature of the Porous Alloy. As one can expect from the glassy precursor and selective dissolution of the active Y-rich glassy phase for nanopore formation, the residual porous product remains glassy in nature as verified by the XRD pattern in Figure 4a, the SAED in Figure 4b, and the high-resolution transmission electron microscopy image in Figure 4c and the associated fast Fourier transformation. The glassy state can be further supported by the DSC curve of the nanoporous alloy (Figure 4d). Compared with the pristine precursor, the crystallization peak corresponding to the Zr-rich glassy phase 4480
DOI: 10.1021/acs.chemmater.7b01038 Chem. Mater. 2017, 29, 4478−4483
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Figure 4. Structural characterization of the nanoporous alloy. (a) XRD pattern. (b) SAED pattern. (c) High-resolution transmission electron microscopy image. The inset shows the fast Fourier transformation of the whole image. No diffraction spot can be found. (d) DSC curve of the residue alloy. The crystallization peak corresponding to the Zr-rich phase still exists, indicating the residue alloy is still in the glassy state, while the Y-rich phase is corroded. Scale bars are 5 nm for panel and 2 nm−1 for panel c.
Figure 5. Hydrogenation performance of nanoporous MG. (a) Comparison of the hydrogen absorption performance of nanoporous MG with solid MG and its nanoporous crystalline counterpart. Points are the experimental data. Solid lines are the fitting curves based on the Avarami−Erofeeev function. (b) Comparison of the m index for various samples. The values of ideal diffusion and 3D interface reaction process were also added. For the sake of clarity, the underlying hydrogen absorption procedure is schematically illustrated.
still exists. The peak position shifts to a higher temperature, which is possibly associated with the kinetic stabilization of supercooled liquids and glasses due to the surface effect. Associated with the dissolution of the Y-rich phase, the interface between the Y-rich and Zr-rich phase disappears, and more of the Zr-rich phase surface appears. The passivation layer on the surface may also affect the crystallization process of the ligament. Except for the passivation character shown in the linear scanning voltammetry curve (Figure S2), the surface passivation of the Zr-rich phase is further reflected in the existence of more Zr in contrast to Cu on the outmost surface of the glassy ligaments as shown in STEM−EDS mapping (Figure S4). The thin passivation layer inhibits the further dissolution of Cu and Al from the ligaments and thus stabilizes the residual glassy alloy as the skeletons of the nanoporous structure. Hydrogen Uptake Performance. The combination of nanoscale pores and a prominent disordered nature permits the nanoporous MG suitable for catalysis, gas permeability, gas absorption, and gas purification. Considering that Zr has a strong hydrogen absorption ability and Cu can also assist in the dissociation of hydrogen gas,35 a low-temperature hydrogen uptake experiment was conducted to demonstrate the uniqueness of the nanoporous MG. As shown in Figure 5a, the absorption process of the nanoporous MG takes only ∼300 s to reach a high saturated capacity of ∼0.56 wt % at 200 °C. In contrast, the solid MG ribbons and the nanoporous crystalline counterpart prepared by crystallization of the nanoporous MG absorb only 0.12 and 0.18 wt % hydrogen, respectively, in 1 h at the same temperature. With a similar amount of hydrogen absorption, for example, 0.18 wt %, the nanoporous MG only need 40 s, which is ∼2 orders of magnitude faster than the crystalline counterpart, highlighting the significant advantage of the nanoporous MG in gas absorption. To understand the underlying hydrogenation mechanism, the kinetic process of
absorption is further analyzed on the basis of the Avrami− Erofeev equation: α = w(t)/w∞ = 1 − exp(−Btm), where α represents the ratio of the absorbed hydrogen content to the saturated hydrogen absorption capability under the same experimental condition, m is the reaction exponent, and B is the reaction rate. As m can be utilized to distinguish the types of reactions,36 fitted index m for each sample is plotted in Figure 5b. For the solid MG ribbon with the same composition of nanoporous Zr43Cu40Al5Y7 MG, index m is close to 0.62, suggesting that the hydrogen absorption behavior is dominated by a bulk diffusion process. For the nanoporous crystallized MG, a slightly higher index m value of ∼0.76 is obtained, benefiting from the enhanced surface diffusion caused by the large surface area of the nanoporous structure. However, index m is still closer to the value characterizing bulk diffusion, suggesting that the slow hydrogen diffusion in crystalline ligaments is the ratecontrolling step. For the 3D bicontinuous nanoporous MG, the m value is as high as 0.92, close to that of the full phase boundary-controlled reaction,36 indicating the hydrogen uptake process is managed by a fast diffusion process with the help of the large surface area of the nanoporous material and disordered structure of the glass. It is worth noting that solid Zr-based crystalline alloys, particularly AB2-type compounds, face the bottleneck of slow activation and low rate capabilities at lower temperatures.35,37,38 Via the introduction of nanoscale pores with a large surface area and reserved disordered glass structure, the low-temperature absorption rate of nanoporous Zr-based MG can be dramatically improved. However, similar to solid Zr-based MGs, the high affinity between H and Zr prevents the dehydrogenation at temperatures below the 4481
DOI: 10.1021/acs.chemmater.7b01038 Chem. Mater. 2017, 29, 4478−4483
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Chemistry of Materials crystallization temperature of the nanoporous MG. Therefore, the nanoporous MGs fabricated in this study may not be appreciated as a recyclable hydrogen storage material. Nevertheless, the ultrafast hydrogen uptake and large capacity achieved with the nanoporous MGs have important implications for the development of new hydrogen storage materials and absorption materials. Compared with conventional dealloying routes for nanoporous crystalline metals and alloy,39 the selective dissolution and passivation method in this study includes two separated steps for nanoporous MGs: preparation of a spinodally decomposed MG precursor and preferential dissolution of one active glass phase. For non-noble-metal-based MGs, the surface passivation of ligaments is pivotal for retaining the integrity of the glassy phase, while the standard electrode potentials of individual glassy components are not the critical factors for selective phase dissolution, which is different from the conventional dealloying of noble-metal-based alloys,40−42 the selective phase dissolution of transition metal compounds.43,44 To obtain a nanopore structure, a considerable difference in corrosion rates between component phases in a precursor is necessary. Via heat treatment of a homogeneous metallic glass above Tg, chemical partitioning was achieved, and the porous structure was formed, associated with the breakdown of the passivation layer.11 Here, the phase separation process is achieved during the cooling process and can be rationally designed on the basis of spinodal decomposition theory.
Mingwei Chen: 0000-0002-2850-8872 Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS This work is sponsored by JST-CREST “Phase Interface Science for Highly Efficient Energy Utilization”, the World Premier International Research Center Initiative for Atoms, Molecules and Materials, MEXT, Japan, MOST 973 of China (Grant 2015CB856800), and the National Natural Science Foundation of China (Grants 11327902 and 51271113). W.J. is supported by the Japan Society for the Promotion of Science postdoctoral fellowship program (ID P15373).
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CONCLUSIONS In summary, 3D bicontinuous nanoporous MGs were successfully fabricated by selectively dissolving one active glassy phase out of a spinodally decomposed precursor and retaining the other glassy component by passivation. By controlling the phase separation kinetics of the precursors, one can effectively tune the pore sizes of resultant nanoporous MGs. The unique combination of a nanoscale pore with a disordered glassy phase results in an ultrahigh hydrogen uptake rate, compared to those of both the solid glassy alloy and its porous crystalline counterpart. The facile route of selective phase dissolution coupled with passivation opens a new possibility for fabricating nanoporous MGs, which will greatly enrich the family of nanoporous materials. Because of the synergy effect of alloying and structural disorder, nanoporous MGs show great potential in catalysis, gas separation, and absorption.
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ASSOCIATED CONTENT
S Supporting Information *
The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.7b01038. Small-angle X-ray scattering (Figure S1) and linear scanning voltammetry (Figure S2) of the precursor and pore size distribution (Figure S3) and STEM image and EDS elemental mapping (Figure S4) of nanoporous metallic glass (PDF)
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REFERENCES
AUTHOR INFORMATION
Corresponding Author
*E-mail:
[email protected]. ORCID
Pan Liu: 0000-0002-4063-9605 4482
DOI: 10.1021/acs.chemmater.7b01038 Chem. Mater. 2017, 29, 4478−4483
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DOI: 10.1021/acs.chemmater.7b01038 Chem. Mater. 2017, 29, 4478−4483
