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Journal of Alloys and Compounds 772 (2019) 164e172

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Journal of Alloys and Compounds

journal homepage: http: / /www.elsevier .com/locate/ ja lcom

Flexible supercapacitor electrodes fabricated by dealloyingnanocrystallized Al-Ni-Co-Y-Cu metallic glasses

Ayan Yao a, 1, Hao Yang a, 1, Jun-Qiang Wang a, *, Wei Xu a, Juntao Huo a, Run-Wei Li a,Huajun Qiu b, Mingwei Chen c

a CAS Key Laboratory of Magnetic Materials and Devices, Zhejiang Province Key Laboratory of Magnetic Materials and Application Technology, NingboInstitute of Materials Technology and Engineering, Chinese Academy of Sciences, Ningbo, 315201, Chinab Department of Material Science and Engineering, Harbin Institute of Technology, Shenzhen, 518055, Chinac Department of Materials Science and Engineering, Johns Hopkins University, USA

a r t i c l e i n f o

Article history:Received 20 June 2018Received in revised form5 September 2018Accepted 9 September 2018Available online 11 September 2018

Keywords:Nanocrystallized metallic glassNanoporousEnergy storageDealloying

* Corresponding author.E-mail address: jqwang@nimte.ac.cn (J.-Q. Wang).

1 These authors contribute equally.

https://doi.org/10.1016/j.jallcom.2018.09.0900925-8388/© 2018 Published by Elsevier B.V.

a b s t r a c t

Nanoporous oxides of transition metals are of great potential for applications as super-capacitor elec-trodes because of their variable valence. However, the low conductance of electrons in oxides and slowtransportation of ions in nanopores limit their performance. In this work, NiCo-contained metal/metal-oxide nanoporous composites are fabricated by one-step dealloying nanocrystallized Al82Ni6Co3Y6Cu3

(at.%) metallic glass ribbons in KOH solutions. The capacitance of the dealloyed nanoporous compositesreaches 3.35 F cm�2 (about 1522 F cm�3), which is 3e10 times higher than that dealloyed from as-spunmetallic glass and crystalline precursors. The large capacitance is attributed to the combination of effi-cient charge transportation rate and high loading amount of active materials. These results demonstratefor the first time that the nanocrystallized metallic glasses are promising easy-processing precursors forfabricating flexible free-standing nanoporous electrodes.

© 2018 Published by Elsevier B.V.

1. Introduction

Electrochemical supercapacitors have become one of thepromising energy storage devices because of their high powerdensity and fast charge/discharge rates [1e7]. According to theenergy storage mechanisms, electrochemical capacitances can bedivided to two types. One is pseudo-capacitance with Faradicprocess; the other is non-Faradic double-layer capacitance [8e11].The Faradic pseudo supercapacitors based on nanoporous metaloxides attract wide interests because of their high theoreticalcapacitance. Oxides of transition metals, such as NiO, CoO, MnO2,RuO2 and their hybrids hold promising potential for applications[12e19]. However, the low mass loading of active materials onelectrodes, small ion diffusivity in nano-pore channels and lowelectric conductivity of oxides retard their performance andapplications.

The specific surface area andmicrostructure are most importantfactors in determining capacitance, which depend strongly on

fabrication methods and precursors [20e26]. Small pores canenhance the specific surface area and thus increase the exposedamount of active materials. However, too small pores limit thediffusivity of electrolytes, while too large pores decrease the spe-cific surface area [27,28]. Thus, the hierarchical microstructurescombining small and large pores/channels are desirable. To obtainhierarchical structures, it usually requires multiple processes,which are complex and challenging for applications [29e34]. Theone-step dealloying method is preferred. However, the composi-tion of crystalline precursors cannot bewidely tuned because of thethermodynamic equilibrium nature and it is difficult to obtainhomogeneousmicrostructures as a result of the existence of defectslike grain boundaries and cast defects. Amorphous alloys exhibitpromising potentials as the precursor of nanoporous materialsbecause of their homogeneous structure and non-equilibrium na-ture that permits to tune the composition in wide range [35e38].However, the capacitance of the obtained nanoporous materials isnot very high because the single mode of pore size limits thediffusivity of electrolytes and the amount of active sites [39e42].

On the other hand, the electronic conductivity of metal oxides isusually low which limit the overall charge transfer efficiency. Toovercome this problem, metal doping in oxides [43,44] and hybrid

A. Yao et al. / Journal of Alloys and Compounds 772 (2019) 164e172 165

with carbon nanotubes/graphene/conductive polymers [45e50]are designed to enhance the electron conductivity of oxides. Themetal/oxide core/shell structure [51e54] is also demonstrated to behelpful for increasing the charge transfer efficiency. However, theinterface between oxide shell and metal ligament are difficult tocontrol, especially for chemical growth methods in solution bath.Dealloying permits in situ growth of uniform oxides on metallicligaments with perfect chemical interface betweenmetal and oxide[55,56]. Thus, nanoporous materials with combined characters ofhierarchical porosity, metal/oxide core/shell ligament structure andmetal doped oxides probably exhibit high charge conductivity andlarge capacity.

In this work, we use one-step dealloying strategy to fabricateNiCo-based nanoporous metal/oxide composites by dealloyingnanocrystallized Al82Ni6Co3Y6Cu3 (at.%) metallic glasses. Anextremely large capacitance of 3.35 F cm�2 (about 1522 F cm�3) isachieved, which is attributed to the large mass-loading of activematerials, large specific area, hierarchical porosity, metal/oxidecore/shell structure in ligaments and enhance conductivity of ox-ides. These results verify that the nanocrystallized metallic glassesare outstanding precursors for fabricating nanoporous pseudosupercapacitors.

2. Experimental details

The master alloy with nominal composition of Al82Ni6Y6Co3Cu3(at. %) was prepared by melting highly pure melts (>99.9wt %)using an induction-melting furnace under the protection of Argonatmosphere. The master alloy was then melt in quartz tube andsubsequently injected onto a spinning copper roller to preparemetallic glass ribbons. The tangent speed of the roller surface was40m/s. The thickness of the as-spun ribbon was about 20e30 mm;the width was 1e2mm; the length was as long as tens of meters.The glassy nature and phase change temperatures were measuredusing differential scanning calorimeter (DSC, NETZSCH 404C) atheating rate of 20 K/min.

The metallic glass ribbons were annealed in high vacuum(<3� 10�3 Pa) for 10min at different temperatures before chemicaltests. The preannealed metallic glass ribbons were dealloyed in 4MKOH aqueous solution for 10min. After dealloying, the sampleswere rinsed with deionized water and dehydrated alcohol. Elec-trochemical tests were carried out at room temperature using anelectrochemical workstation (Zahner Zennium) in a three-electrode cell with a Pt foil as the counter electrode and Ag/AgCl(Cl� concentration inside of the electrode was 3.5M) as the refer-ence electrode. For capacitive performance measurements, theelectrolyte was 4M KOH aqueous solution. The working electrodewas the dealloyed ribbon. The nominal area of the samplesimmersed into the electrolyte for electrochemical tests was10mm2.

The atomic structure of the precursor and de-alloyed sampleswas characterized by X-ray diffraction using a diffractometer withCu Ka radiation (Bruker D8 Advance). The microstructures of thedealloyed samples were characterized by scanning electron mi-croscopy (SEM, Hitachi S-4800) and transmission electron micro-scopy (TEM, FEI Tecnai F20). Chemical composition analysis wasperformed by energy-dispersive X-ray spectrometer (EDS, OxfordINCA x-sight). The elemental valence was investigated using X-rayphotoelectron spectrometer (XPS, AXIS Ultra DLD). The specificsurface areawas studied by gas (N2) adsorption-desorptionmethodaccording to Brunauer-Emmett-Teller (BET, ASAP 2020M) theory.The distribution of pore size was calculated using Barrett-Joyner-Halenda (BJH) method.

3. Results and discussion

Metallic glasses and their nanocrystallized composites exhibitlarger elastic strain limit (~1.5e2%), lower elastic modulus andhigher fracture strength compared to conventional crystalline al-loys. They can bear large bending and tensile elastic deformations.Fig. 1(a) shows an optical image of curved metallic glass ribbonswith different curvatures. The DSC trace in Fig. 1(b) exhibits clearendothermic glass transition character and exothermic crystalli-zation peaks, which confirm the glassy nature of the as-pun ribbon.The glass transition onset temperature Tg¼ 257 �C, crystallizationonset temperatures Tx1¼275 �C, Tx2¼ 340 �C, Tx3¼ 367 �C aremarked using red arrows in Fig. 1(b). Six annealing temperatureswere selected according to the phase transition temperatures tomodulate the microstructure of the alloy.

As shown in Fig. 1(c), the XRD curves of the as-spun Al82Ni6Y6-Co3Cu3 ribbon exhibits a broad hump without sharp Bragg peaks,which confirms the homogeneous amorphous structure. After be-ing annealed at high temperatures, crystalline phases of Al, AlY,Al3Y, (Ni,Co)3Al4 nucleate and grow consequently. The variousnanocrystalline phases are confirmed to have hierarchical crystalsizes [57e59]. For example, Al in face-centered cubic crystallinestructure grows fast andmay be larger than 50 nm. It is followed bythe precipitation of eutectic compositions including intermetallicphases in smaller sizes. After being dealloyed in 4M KOH aqueoussolution for 10min, the Al nanocrystals and Al in AlY, Al3Y and(Ni,Co)3Al4 nanocrystals have been dissolved, which is confirmedby the XRD patterns in Fig. 1(d). The crystalline phases in thedealloyed ribbons are mainly composed of (Ni,Co,Cu)O, (Ni,Co)OOH, Y2O3, little Al nanocrystal, and little Al(OH)3 precipitation.

Given the variable valence characters of Co and Ni elements(Co2þ/Co3þ and Ni2þ/Ni3þ) which are potential for application assuper-capacitor electrodes, the electrochemical capacitance wasexamined using cycling voltammetry (CV) and charge-dischargemethods. Upon scanning to higher voltage, a broad oxidationpeak is detected near 0.4 V (vs Ag/AgCl electrode). Upon scanning tolower voltage, a broad reduction peak is detected near 0.2 V, asshown in Fig. 2(a). The redox couple is associated with the pseudocapacitive behavior of the NiO and CoO components [60,61]. It re-sults from the surface Faradic oxidation and reduction reactions,where the anodic peak is due to the oxidation of NiO to NiOOH, CoOto CoOOH and the cathodic peak is for the reverse process. Theheight of the oxidation-reduction peaks increase along with theincrease of preannealing temperature and reach the maximum forthe sample preannealed at 421 �C. The galvanostatic capacitance at6mA cm�2 is evaluated, as shown in Fig. 2(b). There are obviousplateaus in the charge-discharge curves which are in accordancewith the high oxidation-reduction peaks in CV curves. The charge-discharge time reaches the maximum for the sample preannealedat 421 �C. The areal capacitance is calculated according to theequation C¼ IDt/SDV, where I is the discharge current, Dt is thedischarge time, DV is the potential range and S is the area of theelectrode. The capacitances are determined to be 0.58, 0.67, 1.75,2.05, 1.06 and 0.87 F cm�2 for samples preannealed at 257 �C,312 �C, 367 �C, 421 �C, 500 �C and 550 �C, respectively, as shown inFig. 2(c).

The capacitance of 2.05 F cm�2 for the sample preannealed at421 �C is 68% higher than that dealloyed from the as-spun amor-phous ribbon (1.22 F cm�2) [41]. The areal specific capacitance canreach 3.35 F cm�2 using cyclic electrochemical dealloying method.The thickness of the ribbon is about 22 mm (see the SEM images forthe cross section of the dealloyed ribbons in Fig. S1 in supple-mentary materials), which yields a volume capacitance1522 F cm�3. The areal capacitance or volume capacitance obtainedhere is about 2e10 times higher than that dealloyed from as-spun

Fig. 2. (a) CV curves of the dealloyed Al82Ni6Y6Co3Cu3 ribbons that are preannealed at different temperatures. The san rate is 20mV/s. (b) Charge-discharge curves of the de-alloyed samples. The charge-discharge current density is 6mA cm�2. (c) The areal/volume capacitance of dealloyed nanocrystallized Al82Ni6Y6Co3Cu3 ribbons versus preanneal-ing temperatures.

Fig. 1. (a) Optical image of flexible metallic glass ribbons. (b) DSC heat flow trace of amorphous ribbon measured at 20 K/min. The glass transition onset temperature (Tg),crystallization onset temperatures (Tx1, Tx2, Tx3) are marked by arrows. Six annealing temperatures are selected according to phase transition temperatures to fabricate nano-crystallized ribbons. XRD curves of ribbons (c) before and (d) after dealloying reaction. The preannealing temperature is marked next to the curves.

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metallic glass and crystalline precursors, as shown in Fig. 3 andTable S1. The charge current density dependence of capacitance forthe dealloyed sample preannealed at 421 �C is very stable. Asshown in Fig. 3, the capacitance decrease less than 1% when thecurrent density increases from 4mA cm�2 to 15mA cm�2. Such anexcellent stability is attributed to the good charge transferefficiency.

For understanding the evolution of capacitance along with thepreannealing temperature, the charge diffusion rate D0 (cm2 s�1) iscalculated according to the Randles-Sevcik equation [62].

Ip ¼ 2:69� 105n2=3AD1=20 v1=2c

where Ip is the peak current (in unit of A) of oxidation peaks in CVcurves; n¼ 1 is the number of electrons transferred in anelementary event; A¼ 0.1 cm2 is the electrode surface area; v isscanning rate (in unit of V s�1); c¼ 4mol L�1 is the concentration of

the electrochemically active species. The CV curves at different scanrates for the dealloyed sample preannealed at 421 �C are shown inFig. 4(a). As shown in Fig. 4(b), the Ip is plotted versus v1/2. Based onthe slope, the value of charge diffusion coefficientD0 is calculated tobe 8.43� 10�8, 1.23� 10�7, 5.53� 10�7, 1.40� 10�6, 4.37� 10�7,2.40� 10�7 cm2 s�1 for samples preannealed at 257 �C, 312 �C,367 �C, 421 �C, 500 �C and 550 �C, respectively. As shown inFig. 4(c), the value of D0 reaches the maximum for the samplepreannealed at 421 �C. The diffusivity decreases for the samplespreanealed at high temperatures, which is attributed to the growthintermetallic phases. As shown in SEM images in Fig. 6 and TEMimages in Fig. S2, the intermetallic dendrites grow too big and blockthe tunnels/pores. After being dealloyed, the intermetallics formclosed large nanopores with fine nanopores in themselves and limitthe ionic diffusivity. The fine nanopores limit the diffusivity of ions.The change of diffusivity is consistent with the evolution of arealspecific capacitance, which suggests that the increase of charge

Fig. 3. The areal/volume capacitance versus discharge current density. For the refer-ence data: the solid symbols are referred to areal capacitance, the open symbols arereferred to volume capacitance.

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diffusion plays an important role in enhancing the capacitance.As a result of large volume contraction (>20%) during deal-

loying, cracks usually form due to the large internal stress [63e71].The cracks decrease the electric conductivity of samples and result

Fig. 4. (a) CV curves at different scan rates for the dealloyed sample that is preannealed at 4scan rate (v1/2, v in mV/s). (c) The charge diffusion coefficient D0 as a function of preannea

Fig. 5. The SEM images of the dealloyed samples that are preannealed at different temperathigher temperatures.

in low capacitance. The number of cracks is reduced substantiallyfor preannealed samples, as shown in Fig. 5. When the annealingtemperature is above 367 �C, there are almost no cracks, whichbenefit the charge conductivity. The disappearance of the crackscan be attributed to the enhanced strength of ligaments as a resultof the growth of crystals in precursor alloys.

The morphology of nanoporosity plays an important role ininfluencing the pseudocapacity. Small porosity provides largespecific surface area but limit the ionic diffusivity, while largeporosity allows fast ionic diffusivity but has smaller specific surfacearea. To gain insights into effects of the preannealing temperatureon morphology of nanoporosity, we studied the micro-structure ofthe de-alloyed samples, as shown in Fig. 6. The size of the mainpores increases from 30 nm to 400 nm, when the preannealingtemperature increases from 256 �C to 550 �C. This is attributed tothe growth of nanocrystals. For samples preannealed at highertemperatures, large intermetallic crystals with length of severalmicrons and width of about 100 nm form networks, which increasethe strength of dealloyed sample. The specific surface area fordealloyed samples is studied using BET method, as shown inFig. 6(g). It is determined to be 13.8m2 g�1, 37.9m2 g�1, 7.6m2 g�1

for samples that are dealloyed from as-spun amorphous ribbon,ribbon preannealed at 421 �C and ribbon preannealed at 550 �C,respectively. The distribution of pore size for the sample

21 �C. (b) Peak current (Ip) of oxidation reaction in CV curves against the square of theling temperature.

ures. Micro-cracks (marked by arrows) are avoided for samples that are preannealed at

Fig. 6. The SEM images of the dealloyed ribbons which are preannealed at different temperatures (a) 257 �C, (b) 312 �C, (c) 367 �C, (d) 421 �C, (e) 500 �C and (f) 550 �C. (g) Therepresentative BET gas (N2) adsorptionedesorption isotherm loop for the dealloyed sample that is preannealed at 421 �C. It yields a specific surface area of 37.9m2/g. (h) The BJHpore size distribution curve calculated using desorption curve for the dealloyed sample preannealed at 421 �C. It confirms the existence of hierarchical nanoporous structures withthree typical pore sizes as marked by the arrows.

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preannealed at 421 �C is calculated according to BJH method, asshown in Fig. 6(h). There are three typical sizes of nanopores withsize distribution peaks at about 50 nm, 10 nm and 3.5 nm.

To study the formation mechanism of the pores with differentsizes, the TEM and EDS mapping experiments are performed, asshown in Fig. 7. Before dealloying, there are Al nanocrystals in sizeof about 50e100 nm, as shown in Fig. 7(a)-(f). After dealloying,there are no Al nanocrystals, as shown in Fig. 7(g)-(l). This suggeststhat the formation of large nanopores (30e100 nm) should derivefrom the dissolution of Al nanocrystals. The medium nanopores(8e20 nm) are surrounded by Y oxides, which should derive fromthe dissolution/dealloying of intermetallic nanocrystals (e.g. Al3Y).The fine nanopores (3e4 nm) distribute in the NiCo oxides, whichshould derive from the dealloying of Al-Ni-Co intermetallic nano-crystals. Such hierarchical structure composed of the coarse, me-dium and fine nanopores are suitable for ions to diffuse fast andefficiently. A line-scan of EDS is performed on the nano-ligament, asindicated by the green line in Fig. 7(m). Fig. 7(n) shows the EDSintensity of elements. From the change of O element, it is confirmed

that the ligament has a metallic core which is surrounded by a shellof metal oxides. The Al in the shell has been dealloyed completely,while the active elements Ni and Co together with Cu and Y left inthe shell. The oxide shell and metallic core has a good chemicallybonded interface. During redox reactions, the electrons can transfereasily from oxide shell to metallic core, which benefits the fastcharge/discharge progresses. For the sample preannealed at 550 �C(see data in Fig. S2), it has similar metal/oxide core/shell structure.But the crystals are much larger and some of the pores are closedwhich decrease the exposure of active oxides and lead to smallercapacitance.

The TEM images in Fig. 7 also confirm that the loading amountof the active oxides is very large. This is attributed to thenonequilibrium nature of metallic glasses that allows alloyingvarious elements in wide range but not limited by the thermody-namic equilibrium phase diagrams. Much active elements, e.g. Niand Co and other beneficial elements can be added into the pre-cursor alloy, which can enhance the areal capacitance of thenanoporous materials.

Fig. 7. The TEM and EDS images of the alloy preannealed at 421 �C. (a)e(f) before dealloying and (g)e(m) after dealloying. The element is marked at the upper-right corner of eachpanel. The arrows in (a) show the Al nanocrystals. The white arrow in (g) shows the nanoporous CoNi oxides with pore size of several nanometers; the black arrow in (g) shows thenanoporous Y oxides with pore size of about 10 nm. (n) and (o) The TEM image of dealloyed sample and the line-scan EDS curve shown in (m). The EDS curves for various elementsare shifted by a constant in y-axis direction.

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Copper has excellent electric conductivity and can formmetallicnanoporous structures after dealloying reactions [72,73]. Theaddition of Cu in precursor alloy is expected to form metallicclusters after dealloying which will increase the electron transferrate in oxides. To study the valence state of the elements, XPSexperiment is performed on the surface of the nanoporous ribbonsand sub-surface layers by polishing the ribbon using Ar-sputtering,as shown in Fig. 8. On the surface, most Ni and Co are in oxide statesconfirming that the ligaments of the nanoporous structure arecoated by a shell of active Co/Ni oxides. Most of the Cu elementremains in metallic state on the surface which verifies thedesigning strategy that precious metal forms metallic clusters andcan enhance the electric conductivity. Y is added to enhance theglass-forming ability of the precursor alloy. After dealloying, Y isoxidized into Y2O3 which is most stable oxides and should behelpful for stabilizing the microstructure of nanoporous structure.After the surface layer is polished by Ar-sputtering for 5min,metallic states of Co, Ni, Cu, Yand Al are observed. This confirms themetallic core in ligaments. It is worth noting that the oxide peaksshift to lower binding energy states for the sub-surface oxides. Thissuggests that the oxides gain more electrons which are helpful forenhancing electric conductivity.

Metallic glass ribbons can be fabricated easily using single rollerspinning method. The thickness of the ribbons can vary from 15 mmto 50 mm; the width can vary from 0.5mm to 150mm; the lengthcan be as long as hundreds of meters. Thus, it is easy for massproduction. Metallic glasses exhibit large elastic strain limitation ofabout 2% that is 3e5 times larger than crystalline alloys and lowelastic modulus that is 30% lower than their crystalline

Fig. 8. The XPS curves of the dealloyed ribbon that is preannealed at 421 �C. The sub-surfaceare the signals of metallic states of elements. The blue curves are the signals of oxide states oof mainly oxides of CoO (2p3/2: 781.1 eV, 2p1/2: 796.3 eV), Co(OH)2 (2p3/2: 783.5 eV, 2p954.7 eV), Y2O3 (3d5/2: 157.9 eV, 3d3/2: 159.9 eV); Al(OH)3 (2p: 74.6 eV), NiOOH (Ni 3p: 68.3metal state of Co (2p3/2: 777.7 eV, 2p1/2: 792.7 eV), Ni (2P3/2: 852.4 eV, 2p1/2: 869.7 eV) anfigure legend, the reader is referred to the Web version of this article.)

counterparts. Such flexible characters are favored in industrial ap-plications. Owing to the nonequilibrium metastable nature ofmetallic glasses, they will crystallize into multiple nanocrystallinephases with hierarchical crystal sizes when they are annealed atelevated temperatures. It would form hierarchical nanoporousstructure after being dealloyed. Thus, nanocrystallized Al-Ni-Co-Y-Cu metallic glasses are ideal precursors for fabricating free-standing electrodes by one-step dealloying strategy.

4. Conclusions

Binder-free nanoporous ribbons containing Ni/Co metal/oxidecomposites are fabricated by one-step dealloying nanocrystallizedAl-Ni-Co-Y-Cu metallic glasses. The microstructure of the precursoralloy is modulated bymodifying the preannealing temperature. Theareal capacitance reaches a maximum of 3.35 F cm�2 (about1522 F cm�3), which is 3e10 times higher than that dealloyed fromas-spun metallic glass and crystalline precursors. The excellentelectrochemical performance is attributed to following reasons. (1)The hierarchical continuous nanoporous structure leads to a highsurface area with more active sites and increases the chargediffusion coefficient. (2) The metal/oxide core/shell structure ofligaments can enhance the charge transfer coefficient. (3) Thealloyed Cu forms metallic clusters embedded in the active oxideswhich can increase the electric conductivity in oxides. (4) Pre-annealing the precursor alloy suppress the formation of cracksduring de-alloying process, which can increase the electric con-ductivity. (5) The nonequilibrium nature of metallic glasses allowsincreasing the alloying amount of active elements, e.g. Ni and Co,

curve is measured after being polished by 4 kV Ar-sputtering for 5min. The red curvesf elements. The green curve is the sum of red and blue curves. The surface is composed1/2: 798.7 eV), NiO (2p3/2: 855.9 eV, 2p1/2: 873.5 eV), CuO (2p3/2: 934.8 eV, 2p1/2:eV), CuO (Cu 3p: 77.3 eV) and metal Cu (2p3/2: 932.7 eV, 952.5 eV). In the sub-surface,d Al (2P3/2: 71.8 eV) are observed. (For interpretation of the references to colour in this

A. Yao et al. / Journal of Alloys and Compounds 772 (2019) 164e172 171

and other beneficial elements, e.g. Cu, Pt and Au. Thus, the nano-crystallized Al-based metallic glasses are promising candidates asprecursors for fabricating nanoporous metal/oxides compositeswith advanced electrochemical properties. The fabrication strategyis easy for mass production which has promising potential for ap-plications in energy storage and sensing.

Acknowledgement

The experimental assistance from Jingqing Feng, Zexuan Wangand Peng Zou are appreciated. The financial support from NationalNatural Science Foundation of China (NSFC 51771216, 51771217 and51701230), Zhejiang Provincial Natural Science Foundation of China(LR18E010002), One Hundred Talents Program of Chinese Academyof Sciences are acknowledged.

Appendix A. Supplementary data

Supplementary data related to this article can be found athttps://doi.org/10.1016/j.jallcom.2018.09.090.

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