Nanostructured hydrogen storage materials prepared by high-energy reactive ball milling of magnesium and ferrovanadium

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Nanostructured hydrogen storage materials prepared by high-energy reactive ball milling of magnesium and ferrovanadium Mykhaylo Lototskyy a,*, Jonathan Goh a, Moegamat Wafeeq Davids a, Vladimir Linkov a, Lindiwe Khotseng b, Bulelwa Ntsendwana b, Roman Denys c, Volodymyr A. Yartys c,** a

HySA Systems Competence Centre, South African Institute for Advanced Materials Chemistry (SAIAMC), University of the Western Cape, South Africa b Department of Chemistry, University of the Western Cape, South Africa c Institute for Energy Technology, Kjeller, Norway

article info

abstract

Article history:

Hydrogen storage nanocomposites prepared by high energy reactive ball milling of mag-

Received 11 November 2018

nesium and vanadium alloys in hydrogen (HRBM) are characterised by exceptionally fast

Received in revised form

hydrogenation rates and a significantly decreased hydride decomposition temperature.

13 January 2019

Replacement of vanadium in these materials with vanadium-rich Ferrovanadium (FeV,

Accepted 15 January 2019

V80Fe20) is very cost efficient and is suggested as a durable way towards large scale ap-

Available online xxx

plications of Mg-based hydrogen storage materials. The current work presents the results of the experimental study of Mge(FeV) hydrogen storage nanocomposites prepared by

Keywords:

HRBM of Mg powder and FeV (0e50 mol.%). The additives of FeV were shown to improve

Magnesium hydride

hydrogen sorption performance of Mg including facilitation of the hydrogenation during

Hydrogen ball milling

the HRBM and improvements of the dehydrogenation/re-hydrogenation kinetics. The im-

Ferrovanadium

provements resemble the behaviour of pure vanadium metal, and the Mge(FeV) nano-

Carbon additives

composites exhibited a good stability of the hydrogen sorption performance during

Kinetics

hydrogen absorption e desorption cycling at T ¼ 350  C caused by a stability of the cycling

Cycle stability

performance of the nanostructured FeV acting as a catalyst. Further improvement of the cycle stability including the increase of the reversible hydrogen storage capacity and acceleration of H2 absorption kinetics during the cycling was observed for the composites containing carbon additives (activated carbon, graphite or multi-walled carbon nanotubes; 5 wt%), with the best performance achieved for activated carbon. © 2019 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved.

Introduction Magnesium based hydride is in a focus of studies of solid hydrogen storage materials due to its attractive properties

[1e4]. These include high abundance, low cost and low density of Mg resulting in high gravimetric (7.66 wt% H) and volumetric (110 kg H m3) hydrogen storage densities. Further to hydrogen storage, magnesium-based hydrides are very

* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (M. Lototskyy), [email protected] (V.A. Yartys). https://doi.org/10.1016/j.ijhydene.2019.01.135 0360-3199/© 2019 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved. Please cite this article as: Lototskyy M et al., Nanostructured hydrogen storage materials prepared by high-energy reactive ball milling of magnesium and ferrovanadium, International Journal of Hydrogen Energy, https://doi.org/10.1016/j.ijhydene.2019.01.135

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efficient in heat management applications operating at T ¼ 250e550  C [5e7]. Mg, however, exhibits sluggish hydrogenation/dehydrogenation kinetics and requires use of high H2 desorption temperatures exceeding 300  C. The most efficient way to overcome these drawbacks is in a use of nanocomposite materials by employing reactive ball milling of Mg in hydrogen atmosphere (HRBM; see Refs. [8e14] and references therein) in presence of catalytic additives. Particularly strong catalytic influence was observed for the additives of vanadium metal/ V-based alloys. MgeV based hydrogen storage composite materials prepared by HRBM demonstrated fast hydrogenation rates already at room temperature, and showed lowered hydrogen desorption temperatures. The improved performance was observed for the materials having a significantly decreased during the HRBM particle size, thus decreasing the diffusion length for hydrogen atoms in the material, and, also by achieving synergy in hydrogenation/dehydrogenation of V and Mg [9]. Ferrovanadium (FeV) alloys containing 80e85 wt % V are commercial products of high demand from the metallurgy and automotive industry [15] as they are less expensive and are considered as a convenient and easy available alternative to pure vanadium thus promoting applications of MgeV based hydrogen storage materials. However, iron substitution significantly alters hydrogen sorption characteristics of V. It has been shown [16] that even a minor increase in iron content in V1exFex (x ¼ 0.03e0.1) results in a two-fold drop of the hydrogen sorption capacity, as well as in slowing down the diffusion of hydrogen because of the significant changes in the phase-structural composition of the hydride phases. Along with possible influence of typical impurities in FeV (Si, Al, C, P, S) it may alter its catalytic properties in the processes of hydrogenation/dehydrogenation of Mg, as compared to pure vanadium and alloys on its basis. To the best of our knowledge, there are no reference data concerning the influence of industrial-grade Ferrovanadium additive on the hydrogen storage performance of magnesiumbased hydrogen storage composites. However, a similar study carried out for the commercial Ferroniobium (FeNb; 66 wt% Nb) showed that high purity Nb and Fe taken in the same amounts as in the commercial FeNb are having only slightly better catalytic properties as compared with that of the commercial alloy [17] and thus demonstrated interesting opportunities in using vanadium base commercial alloys. This work presents results of the experimental study of phase-structural characteristics, morphology and hydrogen sorption performance of MgeVe(Fe) hydrogen storage nanocomposites prepared by HRBM from magnesium and commercially available Ferrovanadium (80 wt% V) and a beneficial effect of carbon additives on this performance.

Experimental Methods HRBM was performed for various combinations of starting materials, Mg and FeV, to produce the Mg-based samples of the desired compositions. The composite materials were

prepared using 8 g of Mg powder mixed with 0.16e1.6 g of Ferrovanadium (80 wt% V, V0.81Fe0.19) that corresponds to FeV/Mg weight ratio of x ¼ 2e20%. Depending on the content of FeV in the composite, the samples will be denoted in the paper as Mg-x(FeV); x ¼ 2, 5, 10, 20. One sample, Mg-10(FeV), was in addition modified with 0.4 g (5 wt% with respect to Mg) of various carbon additives, including activated carbon (AC), multi-walled carbon nanotubes (MWCNT) and graphite (G) which were added to the mixture of Mg (8 g) and FeV (0.8 g). These samples are further denoted as Mge10(FeV)e5C where C ¼ AC, MWCNT, or G. Finally, three samples were prepared using 8 g of Mg alone (further denoted as Mg), or a mixture of Mg with pure vanadium or Ferrovanadium taken in a molar ratio 1:1 which corresponds to ~68 wt% of V (or FeV) in the mixture. These samples are denoted as Mge68 V and Mg68(FeV). The compositions of the studied samples are listed in Table 1; further details describing the starting materials are presented in Section Materials. The starting materials, along with 82 hardened steel balls (diameter 10 mm), were loaded under Ar atmosphere into a hardened steel vial equipped with an in-built pressureetemperature monitoring system (Evico Magnetics GmbH). Depending on the total weight of the charge, the ball to powder weight ratio (BPR) varied between 33.3:1 and 40:1 (Table 1). The vial was then evacuated and subsequently filled with hydrogen to reach a pressure of close to 30 bar. HRBM was performed in Retsch PM100 planetary mill at a rotation speed of 500 rpm, for a total time of 5 h. Milling was interrupted every 30 min to allow the vial to cool down to room temperature followed by the monitoring of the current state of hydrogen absorption based on the readings of the pressure and temperature in the vial,1 as well as re-filling the vial with H2 to the pressure of 30 bar. The as-milled samples were unloaded in a glovebox under a purified Ar atmosphere. The characterisation of the as-milled, dehydrogenated and re-hydrogenated samples was performed using the techniques described below. The details related to the SEM, TEM and XRD studies are presented in the Supplementary Information file. The XRD data (Cu-Ka) were processed by Rietveld full-profile analysis using GSAS software [18]; the characteristics of the constituent phases were taken from CRYSTMET database [19], version 5.6.0, © Toth Information Systems, Inc. Differential scanning calorimetry (DSC) studies of the asmilled samples were performed using STA 6000 analyser (Perkin Elmer). The samples (m ¼ 20e35 mg) were placed into standard alumina crucibles (the handling was done in an argon glove box), which were transported into the instrument 1

For pure Mg hydrogen absorption during HRBM was continuously monitored starting from the pressure e temperature readings during the milling. The details are presented in Refs. [10,12]. The comparison of the results with the data obtained from the measurements using the equilibrated at room temperature vial showed that the difference is insignificant. Further determination of the hydrogen absorption was carried out by the applying the latter procedure.

Please cite this article as: Lototskyy M et al., Nanostructured hydrogen storage materials prepared by high-energy reactive ball milling of magnesium and ferrovanadium, International Journal of Hydrogen Energy, https://doi.org/10.1016/j.ijhydene.2019.01.135

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Table 1 e Compositions of the studied samples and their hydrogen sorption performances during HRBM. Notation

BPR

Mg Mg-2(FeV) Mg-5(FeV) Mg-10(FeV) Mg-10(FeV)-5ACb Mg-10(FeV)-5MWCNTb Mg-10(FeV)-5Gb Mg-20(FeV) Mg-68(FeV) Mg-68V

40 39.2 38.1 36.4 34.8 34.8 34.8 33.3 40 40

a b

Additive to Mg: V or FeV (V0.81Fe0.19) mol.%

wt.%

e 0.89 2.21 4.32

e 1.96 4.76 9.09

7.66 7.59 7.48 7.32 7.02

8.27 50 50

16.67 68.1 67.9

7.03 5.03 5.07

H storage capacity achieved during the HRBM wt.% H

H/M

7.45 7.32 7.12 7.07 6.41 6.77 6.20 6.70 3.55 4.59

1.95 1.92 1.90 1.93 1.82 1.92 1.75 1.90 1.42 1.83

Assuming H/M ¼ 2 and absence of hydrogen absorption by carbon. 5 wt% of the carbon additive.

in the sealed containers followed by quick installation onto the sample holder. The measurements were performed under a flow of purified argon (100 mL min1), heating from 25 to 600  C at several heating rates chosen between 5 and 20 K min1. The DSC data was further processed using the Kissinger method [20] to yield apparent activation energy, EA, for the endothermic process of hydride decomposition using the following equation: ln

Wt.% H (theoretical)a

b T2m

!

  ZR EA ¼ ln ;  EA R Tm

(1)

where b [K min1] is a heating rate, Tm [K] is the peak temperature, R ¼ 8.3143 J mol1 K1 is a universal gas constant, Z [min1] and EA [J mol1] are, respectively, pre-exponent and activation energy of the Arhenius equation determining the temperature dependence of the decomposition rate constant. The value of ERA was determined from a slope of a Kissinger   plot of y ¼ ln Tb2 versusx ¼ T1m , similar to the procedure used in m

the studies of the dehydrogenation performance of nanocrystalline MgH2 (see, e.g. Ref. [21]). It has to be noted that due to the very broad, up to 300 K, decomposition temperature ranges for most of the studied samples (see Section DSC below), it was difficult to accurately determine the values of the peak temperatures, Tm, that resulted in uncertainties in the values of EA calculated from the Kissinger plots. More accurate results were obtained assuming that the decomposition “peak”, Q(T), appears as a superposition of several processes resulting in the overlapping of the corresponding DSC peaks, q(T), having different parameters: QðTÞ ¼

n X

qi ðTÞ:

(2)

i¼1

Accordingly, the DSC spectra were deconvoluted, and the constituent peaks were modelled by a Gaussian function: " qðTÞ ¼ A exp 

# ðT  Tm Þ2 ; 2 w2

where A is peak amplitude, and w is the peak width parameter related to the full width at half maximum (FWHM) as: FWHM ¼ 2 w

qffiffiffiffiffiffiffiffiffiffiffiffi lnð4Þ:

(4)

The deconvolution of all observed DSC spectra was carried out assuming the same number of constituent peaks, n ¼ 32, for all the samples. The sets of peak temperatures, T(i) m, were fitted by the Kissinger method (Eq. (1)), separately for i ¼ 1, 2 and 3. Further details on the processing the DSC data are presented in Supplementary Information, section “Methods: DSC studies”. In addition to the DSC, hydrogen thermodesorption from the as-milled and re-hydrogenated samples was studied using thermal desorption spectroscopy (TDS) performed in a Sievert's-type volumetric installation. 200e400 mg of the sample powder was loaded in argon glove box into a reactor which, together with the measurement system, was then evacuated at room temperature to <104 mbar. The TDS measurements were performed by heating the reactor at a constant heating rate of 5 K min1 from room temperature to 450e470  C under dynamic vacuum conditions, where H2 desorption results in the pressure increase in the evacuated system. The vacuum sensor was calibrated by comparing the data collected at a known flow rate of H2 gas into the system (see Supplementary Information, section “Methods: Volumetric measurements (TDS/re-hydrogenation)” for the details). Re-hydrogenation (hydrogen absorption) was studied in the same Sievert's-type setup. The re-hydrogenation experiments were performed at T ¼ 250  C and initial H2 pressure of 12e13 bar (final H2 pressure 6e8 bar) followed by the TDS as described above. The collected datasets for hydrogen absorption during HRBM and re-hydrogenation were used in the analysis of the formal kinetics of the processes performed by applying AvramieErofeev equation [10]:

(3) 2

See section DSC for further details.

Please cite this article as: Lototskyy M et al., Nanostructured hydrogen storage materials prepared by high-energy reactive ball milling of magnesium and ferrovanadium, International Journal of Hydrogen Energy, https://doi.org/10.1016/j.ijhydene.2019.01.135

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   n

t C ¼ Cmax 1  exp  ; t0

(5)

where C, Cmax are the actual and maximum H absorption, respectively; t is time; t0 is characteristic hydrogenation time (reciprocal rate constant); and n is an exponential factor indirectly related to the reaction mechanism. Nucleation and growth and hydrogen diffusion are two competing processes governing the hydrogenation kinetics during the formation of the metal hydrides. We note that for the diffusion-controlled hydrogenation reactions that appears to be the case for the re-hydrogenation of Mg, accurate determination of the P(H2)-related driving force or, more precisely, of the thermodynamic contribution into pre-exponent of the Arhenius equation), can be done when assuming that ( " 12 #)   where T is it is proportional to the factor T$ 1  Peq P =

the temperature, Peq and P are the equilibrium and the actual pressures, respectively [22]. Assuming that T ¼ 250  C, Peq (250  C) ¼ 0.32 bar,3 and taking the minimum and maximum values of P observed in the experiment as 6 and 13 bar, respectively (see above), we conclude that the “pressure related driving force” for the applied experimental conditions change by less than 10% only, and thus Eq. (5) can be used for the evaluation of the reaction kinetics. The experiments on cycling stability of the materials during the hydrogen absorption/desorption were carried out at T~350  C using PCT Proe2000 automated gas sorption analyser (Hy-Energy Scientific Instruments). Hydrogen absorption and desorption studies were performed for ~30 min at final H2 pressures of ~12 bar, and ~3 bar for H absorption and desorption, respectively. Due to a significant and fast increase of the measured sample temperature, Ts, at the beginning (t < 20 s) of the hydrogenation observed in the course of the cycle stability studies (similarly to our earlier observations for HRBM MgeV [9]), the actual time dependence of the sample temperature, T*s(t), was determined by taking into account a separately measured thermocouple response time, t ¼ 5.6 ± 0.1 s, as4:   Zt u d t ½Ts ðuÞdu : exp T*s ðtÞ ¼ Ts ðtÞ exp  t t du

(6)

0

The calculated values of T*s were then compared with the temperatures, Teq, for the Mg 4 MgH2 transformation calculated as a solution of the Van't Hoff equation as respect to the temperature, starting from the actual value of hydrogen pressure, P, and the reference data for the enthalpy and entropy of phase transformation3: Teq ¼

DH : R lnðPÞ þ DS

(7)

As can be seen in Supplementary Information (Section “Materials”, Fig. S3), the Ferrovanadium used in this study had the shape of regular prismatic particles, 0.1e0.5 mm in size. According to EDS analysis (Fig. S3, Table S1), the average content of vanadium in the alloy (~77%) appeared to be close to the value specified by the supplier (79.8%). Along with Fe (~15%), the material contained noticeable amounts of: Si (~4%), Al (~3%), Ca (~1%) and Mg (~0.6%). EDS also showed a significant surface oxidation of FeV. According to the XRD analysis (Supplementary Information, Fig. S4), the as-received Ferrovanadium contained BCC phase (>90 wt%), space group Ime3m (#229), with a unit cell constant a ¼ 2.9788(1)  A as a major constituent. The unit cell parameter has an intermediate value between the reference values for a pure V (3.03  A [19, ID: 28911]) and V50Fe50 solid solution alloy (2.92  A [19, ID: 37583]). The XRD also showed the presence of a tetragonal secondary phase identified as s-(VFe), space group P42/mnm (# 136), with unit cell parameters, a ¼ 8.949(4)  A, c ¼ 4.551(4)  A. The latter values well agree with  c ¼ 4.6 A  [19, ID: 134643]). the reference data (a ¼ 9.0 A, Further to the commercial Ferrovanadium, a number of the experiments was carried out using V (80 wt%) e Fe (20 wt%) alloy prepared by arc melting of a high purity V and Fe, 99.7% and 99.98%, respectively. Measurements of hydrogen absorption properties of the commercial and arc melted Ferrovanadium showed that in both cases maximum hydrogen absorption capacity of the alloy (P ¼ 30e40 bar, T~20  C) was between 1.3 and 1.5 wt% H, or H/(V þ Fe) ¼ 0.7e0.8, indicating a formation of a monohydride (V,Fe)H1ex which is in a good agreement with the reference data for the VeFe alloys [16]. However, commercial Ferrovanadium was found to be easier to activate: H absorption started after an 1 h long vacuum heating of the sample to T~450  C followed by its cooling to room temperature and introducing H2. Conversely, the arc melted Ferrovanadium started to absorb H2 only after several absorption-desorption activation cycles which involved vacuum heating at the conditions specified above and introducing H2 to the hot sample. Other starting materials used in the study included:  Magnesium powder, 20…þ100 mesh, 99.8%, Alfa Aesar;  Vanadium metal, turnings, 99.7%, Alfa Aesar;  Graphite powder, 20 mm, 99þ%, density 1.9 g cm3, Fluka;  Activated carbon, YP-50F, coconut-based, 3…20 mm, surface area 1600…1700 m2 g1, Kuraray Chemical Co.;  Multi-wall carbon nanotubes, diameter 5…20 nm, length 10 mm, aspect ratio >500, bulk density: 0.04e0.06 g cm3, impurities: amorphous carbon and Fe (<10%), Carbon Nano-Material Technology Co., Ltd.

Materials Ferrovanadium (75% V grade) was supplied by Insimbi Alloy Suppliers (Pty), Ltd.; South Africa. The material was in a form of a coarse black powder. DH ¼ 74.5 kJ (mol H2)1 and DS ¼ 135.0 J (mol H2 K)1 [1,9]. See Supplementary Information, section “Methods: Correction of the sample temperature (Eq. (6))”, for more details. 3 4

Results and discussion HRBM behaviour Experimental data on the hydrogenation of the studied samples during HRBM are presented in Fig. 1 and Table 2. Results of formal evaluation of the hydrogenation kinetics by

Please cite this article as: Lototskyy M et al., Nanostructured hydrogen storage materials prepared by high-energy reactive ball milling of magnesium and ferrovanadium, International Journal of Hydrogen Energy, https://doi.org/10.1016/j.ijhydene.2019.01.135

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Fig. 1 e Hydrogenation of the studied samples during HRBM: A e samples with different contents of Ferrovanadium (the data for Mg and Mge68V are presented for the reference); B e influence of different carbon additives on the hydrogenation performance of Mg-10(FeV).

processing the data with Eq. (5) are presented in Table 2 and Fig. S5 of the Supplementary Information file. It is clearly seen that an addition of the V-containing catalyst and, furthermore, carbon to Mg significantly alters and at optimum content of the additives e Ferrovanadium and carbon e improves the hydrogenation kinetics. Along with the clearly observed trend of the shortening the characteristic reaction time, t0, with the increase of the content of FeV or V catalyst (Fig. S5A), the introduction of V-based catalyst also results in a gradual decrease of n, from 4 for the undoped Mg to ~2.5 for Mg-10(FeV) and Mg-20(FeV) and further to ~1 for Mg-68(FeV) (Fig. S5B). The value of the Avraami exponent n ¼ 4 corresponds to the random nucleation and 3Dgrowth reaction mechanism [23,24]. Further decrease of the Avraami exponent may have its origin in the limitations related to the possible nucleation regions when the nucleation proceeds at points (n ¼ 3), edges (n ¼ 2) or at surfaces (n ¼ 1) [25]. Furthermore, non-integer values of n can be associated with the contribution of diffusion [23], and/or the change of the nucleation mechanism during the reaction [24]. The data describing the kinetics of interaction with hydrogen can be explained by considering two phenomena taking place during a transformation of magnesium into magnesium hydride. These are (a) Nucleation and growth of MgH2 leading to a phase transformation from Mg to MgH2 and (b) Diffusion of hydrogen in the Mg-based material. These mechanisms can be easily distinguished both from the shapes of the kinetic curves and from the fitting parameter n of Eq. (5). Indeed, from Fig. 1 and Table 1, it is possible to classify the interaction mechanisms into three groups: (a) Well pronounced S-shape kinetic curve. Such type of a curve indicates that nucleation and growth (N&G) is a

predominant mechanism of interaction with H2 while n values being in a range 3.4e4.0. Such a behaviour is characteristic for Mg metal itself and for the composites containing small amounts of FeV (2 and 5%) and the sample Mg-10(FeV)e5C (C ¼ G, MWCNT). (b) When content of FeV is the highest, 68%, the mechanism of hydrogenation changes to a diffusion controlled phase transformation which is manifested by a rapid progression of the hydrogenation if to take Mg/MgH2 transformation as a reference point. Indeed, from Table 2 one can observe that 50% hydrogenation of Mg during the ball milling requires 2.7 h e in contrast with just 0.3 h for Mge68V. n values for such a mechanism are in a range 1.1e1.5 and are observed for Mg-68(FeV), Mge 68V and Mg-10(FeV)-5AC. (c) Finally, a mixed DIFFUSION CONTROLLED þ NUCLEATION AND GROWTH mechanism takes place for the composites containing 10e20 wt% of FeV. n values in such a case have intermediate values 2.4e2.56, and the kinetic curves show a shape where a delayed step of the hydrogenation is significantly suppressed and the time to reach 50% of conversion is close to 1 h. Further details can be found from the data presented in Table 2 and described in Fig. 1. From the data we conclude that the milling times to achieve the same extent of hydrogenation and, the characteristic reaction times, t0 (Table 2) are significantly shortened. Thus, we conclude that the HRBM studies showed that the addition of Ferrovanadium to Mg facilitates the hydrogenation process and changes its mechanism from nucleation and growth to diffusion controlled following increase in the content of FeV.

Please cite this article as: Lototskyy M et al., Nanostructured hydrogen storage materials prepared by high-energy reactive ball milling of magnesium and ferrovanadium, International Journal of Hydrogen Energy, https://doi.org/10.1016/j.ijhydene.2019.01.135

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Table 2 e Fitted kinetic parameters of H absorption (Eq. (5)) during HRBM (Fig. 1). Sample

Mg Mg-2(FeV) Mg-5(FeV) Mg-10(FeV) Mg-10(FeV)-5AC Mg-10(FeV)-5MWCNT Mg-10(FeV)-5G Mg-20(FeV) Mg-68(FeV) Mge68V a

b c

Milling time [hrs] to achieve hydrogenation fraction of, %

Fitted using Eq. (5) kinetic parameters of H absorption

10

50

90

Cmax [wt.% H]

t0 [min]a

n

2 1.2 0.9 0.3 0.1 0.5 1.1 0.4 <0.1 <0.1

2.7 2.0 1.5 0.9 0.5 1.0 1.8 1.0 0.3 0.3

3.9 3.2 2.5 1.6 1.2 1.6 2.7 1.8 1.1 0.8

7.42(1) 7.1(1) 6.9(1) 6.96(7) 6.38(6) 6.69(5) 6.08(7) 6.67(2) 3.46(6) 4.39(5)

182.3(3) 135(2) 101(2) 61(2) 37(1) 64(1) 119(2) 72.6(5) 27(1) 20.8(6)

4.01(3) 3.8(3) 3.7(3) 2.4(2) 1.5(1) 3.4(3) 3.7(2) 2.56(6) 1.11(7) 1.7(1)

Mechanism of hydrogenationb

R2 c

N&G N&G N&G DIFF þ N&G DIFF N&G N&G DIFF þ N&G DIFF DIFF

0.99689 0.99673 0.99684 0.99643 0.99518 0.99766 0.99781 0.99971 0.99242 0.99357

e Characteristic time of transformation t0 is defined as a time to achieve (1e1/e) ¼ ~67% transformation and equals to 1/K from the Avraami equation. e N&G ¼ Nucleation and Growth; DIFF ¼ Diffusion Controlled. e Excellent quality of the fitting has been achieved in every case as is shown by the values of R2 which are close to 1.

The increase in the content of FeV also resulted in a gradual decrease of the hydrogen absorption capacity, due to the introduction of the FeV “ballast”. However, importantly, in most cases HRBM results in almost complete hydrogenation of Mg with the calculated H/Mg  1.9. For the samples Mg-x(FeV) (x ¼ 2e10), the maximum hydrogen storage capacity achieved during HRBM (Table 1) was 95.2%e96.4% of the theoretical value (97.3% for pure Mg) while for Mg-68(FeV) it was only 70.6% which is equivalent to the average value of H/Mg~1.4. The sample Mge68V exhibited a higher hydrogen sorption capacity (90.5% of the theoretical capacity corresponding to H/M ¼ 1.83). This is probably can be related to a fact that in case of V maximum reached H sorption capacity corresponding to the formation of the dihydride VH2 while for FeV it is limited to the just monohydride with 1 at. H/at. M in maximum. The effect of carbon additives to Mg-10(FeV) (Fig. 1B, Table 2) was found to be strongly dependent on the type of carbon material used. For the activated carbon the hydrogenation kinetics was improved as compared to the carbon-free material. The addition of MWCNT did not change the kinetics of hydrogen uptake during the HRBM while graphite additive resulted in slowing the hydrogenation down due to appearance of a long, >1 h, incubation period. The observed behaviour is similar to the earlier reported data describing HRBM of Mg modified with 1e5 wt.% C (C ¼ AC, MWCNT, G) [10]. The addition of graphite and activated carbon resulted in the decrease of maximum hydrogen storage capacity (88% and 91% of the theoretical values, respectively) as compared to Mg-10(FeV) (96.6%) while for Mg-10(FeV)-5MWCNT the capacity (96.4%) remains almost the same. Surprisingly, HRBM of Mg with the arc melted FeV (see Supplementary information, Fig. S6) showed no changes in the hydrogenation performance of this composite, which was very much similar to the behaviour of individual Mg. Most probably, the origin of that is in the difficulty of the activation of the arc-melted FeV alloy that probably remains in part inactive to H2 during HRBM.

DSC Fig. 2 presents DSC curves for the selected as-milled samples measured at a heating rate set point of 10 K min1. As it can be seen, all the DSC spectra are characterised by very broad ranges of the hydride decomposition; the full width of the peaks (FW) is about 100 K for HRBM Mg, while introduction of FeV and, furthermore, carbon, results in a significant increase of FW, up to 270 K for the sample Mg-10(FeV)-5AC. Moreover, the additives cause appearance of additional peaks in the DSC curves. Deconvolution of the DSC spectra using the procedure described in Section Methods (Eqs (2) and (3)) allowed to identify individual components contributing to the overall desorption behaviour. Even though the experimental DSC data (Fig. 2) show significant deviations of b ¼ dT/dt from the instrument heating rate set point, which is particularly high in the proximity to the decomposition peaks, however we have reached a good agreement between the experimental and calculated DSC traces with a goodness of fit parameters R2>0.99 for each processed data set. The parameters describing the decomposition for the hydrides retrieved from analysis of the DSC data for the samples Mg-x(FeV) (x ¼ 2, 10, 20) and Mg-10(FeV)e5C (C ¼ AC, MWCNT) is presented in Table 3. The following conclusions can be done:  Introduction of FeV results in:  Decrease of the onset temperature of hydride decomposition by 50e100 K; the more pronounced decrease is observed at higher FeV concentrations.  Lowering of the peak temperature at maximum decomposition rate by 10e120 K; increase in a FeV content results in more pronounced shifts, particularly for x  10.  Significant broadening of the decomposition temperature range (FW), which also increases with the increase of FeV concentration.  Increase in the contribution of the lower temperature decomposition stage (peak 1) into the overall

Please cite this article as: Lototskyy M et al., Nanostructured hydrogen storage materials prepared by high-energy reactive ball milling of magnesium and ferrovanadium, International Journal of Hydrogen Energy, https://doi.org/10.1016/j.ijhydene.2019.01.135

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Fig. 2 e DSC curves for the as-milled samples measured at a heating rate of 10 K min¡1: experimental (points), calculated curves (red solid lines), constituent peaks (filled area under curves; Peak 1 e blue; Peak 2 e magenta; Peak 3 - brown). Measured values of b ¼ dT/dt are shown as well (right Y axis). (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)

decomposition process well correlates with the increase of FeV concentration. The higher temperature stage (peak 3) exhibits an opposite trend at x  10.  Introduction of carbon in Mg-10(FeV) results in  Further broadening of the decomposition temperature range and increase of the peak temperature by 60e80 K (similar results were obtained earlier for the carboncontaining HRBM Mg þ FeTiO3 [12]), and  Increase of the contribution of the medium- (peak 2; C ¼ AC) and/or higher-temperature (peak 3; C ¼ MWCNT) stages into the overall decomposition process, as it can be seen in Fig. 2 (bottom).

Fig. S8 in the Supplementary Information presents Kissinger plots for the separate peaks obtained by deconvolution of the DSC curves. The calculated activation energies, their errors and Pearson correlation coefficients are presented in Table 3. The calculated values of EA for HRBM Mg5 are in a good agreement with the published data [10,21], EA ¼ 100e140 kJ mol1. The common feature of the observed changes is in a significant decrease of the activation energies for all hydrogen 5

Straightforward application of the Kissinger method to DSC data for HRBM Mg yields EA ¼ 137.9 ± 10.6 kJ mol1; R2 ¼ 0.9884.

Please cite this article as: Lototskyy M et al., Nanostructured hydrogen storage materials prepared by high-energy reactive ball milling of magnesium and ferrovanadium, International Journal of Hydrogen Energy, https://doi.org/10.1016/j.ijhydene.2019.01.135

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Table 3 e Hydride decomposition parameters retrieved from DSC data for the as-milled samples. Parameter

Sample Mg

Decomposition temperaturesa

Peak 1b

Peak 2b

Peak 3b

a b

Onset [ C] Peak [ C] FW [K] FWHM [K] T [ C] FWHM [ C] Area [%] EA Value [kJ/mol] R2 T [ C] FWHM [ C] Area [%] EA Value [kJ/mol] R2  T [ C] FWHM [ C] Area [%] EA Value [kJ/mol] R2

302 363 100 38 341.5 13.2 0.9 141.9 ± 10.1 <0.5 374.1 23.0 27.5 131.7 ± 16.7 0.9993 394.3 40.8 71.6 106.9 ± 14.0 0.9990

Mg-2(FeV) Mg-10(FeV) Mg-10(FeV)-5AC Mg-10(FeV)-5MWCNT Mg-20(FeV) 250 351 175 53 310.0 30.2 4.2 109.0 ± 5.1 0.8693 357.5 29.2 25.5 81.0 ± 11.2 0.9968 379.8 50.7 70.3 74.6 ± 10.3 0.9949

225 272 171 67 296.5 22.1 14.5 60.1 ± 9.2 0.9691 323.1 46.3 48.6 66.4 ± 10.1 0.8858 365.9 59.5 36.9 81.5 ± 11.4 0.9539

250 331 210 53 305.7 45.7 12.0 132.0 ± 9.0 0.9994 353.3 19.4 14.7 91.6 ± 14.1 0.9799 376.6 61.1 73.3 86.0 ± 14.2 0.9988

195 243 215 39 290.9 25.7 30.0 33.3 ± 6.1 0.9521 312.3 27.4 19.0 41.1 ± 8.0 0.8970 352.6 67.3 51.0 61.3 ± 10.0 0.9995

e for heating rate set point of 5 K min1. e peak temperatures, widths and areas presented at heating rate set point of 10 K min1.

desorption events (1e3) following the increase of FeV concentration. The minimum value of the activation energy for the lower-temperature peak 1, EA ¼ 33 kJ mol1, in line with the lowest observed onset of the decomposition process (<200  C; sample Mg-20(FeV)6) may be associated with the decomposition of V-based “monohydride” phase having a similar to MgH2 stability [9]. The values of EA observed for the major peaks 2 and 3 (in total, >70% of the whole area under the DSC curve) are related to the decomposition of MgH2 (>80 wt% in the studied samples). The calculated activation energies for these peaks varied between 41 (Mg-20(FeV), peak 2) and 81 kJ mol1 (Mg-2(FeV), peak 2; Mg-10(FeV), peak 3) which concludes that introduction of FeV results in a significant improvement of dehydrogenation kinetics of MgH2 by decreasing activation energy of its decomposition. Introduction of carbon results in a certain increase of the temperatures and activation energies for peaks 2 and 3. Such behaviour can be related to (i) the delay in the H2 evolution by the passivation of the sample surface with traces of air during moving the sample to the DSC setup7 and (ii) formation of the carbon layer on the surface of MgH2 which creates a barrier for hydrogen desorption. Summarising the observations presented above, as well as comparing the data of the DSC and TDS studies of the asmilled samples (see Supplementary Information, Fig. S9; more details on the TDS studies will be presented in the next section), we can suggest the following nature of the peaks 1e3 observed in the course of the DSC studies:

6

200 352 270 61 306.0 57.2 15.1 42.4 ± 2.5 0.8988 382.1 64.6 78.1 76.3 ± 5.2 0.9953 440.9 26.5 6.8 115.1 ± 7.5 0.9879

Similar values of the decomposition onset temperature and EA for peak 1 were found for the samples Mge68V and Mg68(FeV). 7 According to our earlier results [10], the MgH2eC samples are more air sensitive than MgH2 prepared by HRBM of Mg at the same conditions.

 Peak 1 is associated with the decomposition of the least stable V-based “hydride” present in the studied system.  Peak 2 is related to a decomposition of thermodynamically unstable g-modification of MgH2.  The highest in temperature Peak 3 is associated with the decomposition of a-MgH2. We note that according to the results of in-situ synchrotron XRD studies [26], the difference in the peak temperatures for the decomposition of g- and a-modifications of MgH2 (additionally complicated by the gea transformation at T ¼ 150e310  C) is around 50 K, with >35 K lower H2 evolution onset for g-MgH2. This obviously complicates the deconvolution procedure.

TDS, re-hydrogenation and cycle stability TDS studies (see Supplementary information; Figs. S9 and S10) for the as-milled samples show their similarity with the DSC data (Section DSC). However the TDS data shows the lower temperature onsets of the hydrogen desorption and, accordingly, wider decomposition temperature ranges. The re-hydrogenated samples (TDS #2) undergo a onestage decomposition process with a single peak of H2 desorption corresponding to the decomposition of a-MgH2. The peak temperature is shifted to the higher temperatures as compared to the as-milled samples (TDS #1); the increase varies from 5 K (Mg-2(FeV)) to 24 K (Mge10Ve5AC). The differences between the TDS #1 and 2 are due to the presence of a destabilized g-modification of MgH2 in the as-milled samples (Section XRD) [26]. As can be seen from Table 4, the total amount of desorbed hydrogen estimated by the integration of time dependence of H2 desorption rates for the as-milled samples (first column) is in a good agreement with the HRBM data (Table 1, column 6). At the same time, the corresponding values for the

Please cite this article as: Lototskyy M et al., Nanostructured hydrogen storage materials prepared by high-energy reactive ball milling of magnesium and ferrovanadium, International Journal of Hydrogen Energy, https://doi.org/10.1016/j.ijhydene.2019.01.135

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Table 4 e Parameters of re-hydrogenation of the studied samples (Fig. 4). Sample

Mg-2(FeV) Mg-10(FeV) Mg-10(FeV)-5AC Mg-10(FeV)-5MWCNT Mg-20(FeV)

Estimated amount of the desorbed H2 [wt.% H]

Fitted kinetic parameters of H absorption (Eq. (5))

TDS #1

TDS #2

Cmax [wt.% H]

t0 [min]

n

R2

7.3 7.1 7.2 7.0 6.8

6.1 4.8 7.1 5.9 4.2

6.504(3) 5.169(7) 6.685(3) 7.081(4) 4.589(6)

5.38(1) 0.52(1) 2.82(1) 5.30(2) 4.05(4)

0.604(2) 0.305(4) 0.424(1) 0.817(3) 0.295(2)

0.99840 0.96669 0.99785 0.99709 0.95942

re-hydrogenated samples (Table 4) with maximum hydrogen desorption capacities obtained in the course of rehydrogenation (column 3) are lower; the drop becomes more pronounced when FeV concentration increases. However, in presence of carbon species, the deterioration of hydrogen capacity of the re-hydrogenated samples becomes much smaller. Further data describing the TDS results is presented in Supplementary information, section “Results: TDS”. The re-hydrogenation kinetic curves measured for the studied samples at T ¼ 250  C and P~10 bar are shown in Fig. 3; the fitted kinetic parameters of H absorption (Eq. (5)) are presented in Table 4. Increase of FeV concentration from ~2 to ~10 wt% results in a significant kinetic improvement with characteristic reaction time, t0, decreasing by higher than one order of magnitude. However, further increase in FeV content reverses this trend as for the sample Mg-20(FeV) the characteristic time is only 25% shorter than for Mg-2(FeV). At the same time, the samples Mg-10(FeV) and Mg20(FeV) reach lower limiting H content in the hydride phase and show a slower approach to its asymptotic limit with the value n decreases from 0.6 (Mg-2(FeV)) to ~0.3 (Mg10(FeV) and Mg-20(FeV)). Taking into account that the minimum value of the Avraami exponent with a physical meaning is n ¼ 0.5 for the diffusion-controlled process [23], it is reasonable to assume that at high concentrations of FeV the hydrogenation involves one extra, slower, step, which is likely to be controlled by H diffusion. However, the low value of n between 0.5 and 1 means that the phase

Fig. 3 e Re-hydrogenation kinetics of the studied samples after TDS #1: experimental data (points) and calculated data (lines).

transformation process is diffusion controlled for all studied contents of FeV. Introduction of carbon into Mg-10(FeV) results in slowing down the kinetics of H2 desorption, with the values of t0 are becoming close to the ones for Mg-2(FeV). At the same time, the asymptotic hydrogen limiting concentrations, Cmax, significantly increase becoming closer to the values of the asmilled samples. The rate of approaching the hydrogen concentration Cmax becomes faster, with the fitted values of n being between >0.4 and 0.8. Similar results were presented in Ref. [9] for the Mg-based HRBM composites containing 10 wt% of easily hydrogenated V alloy. Although this reference work does not include a formal kinetics analysis of the re-hydrogenation process, it showed that the addition of 5 wt% of graphite results in 5e10 times longer characteristic hydrogenation times and in higher maximum hydrogen absorption capacities. Cyclic hydrogen desorption and absorption experiments performed at T ¼ 350  C for the sample Mg-10(FeV) showed that the capacity remains stable during at least 30 hydrogen desorption/adsorption cycles; the average final absorption and desorption capacities [wt.% H] were 4.93 ± 0.02 and 4.82 ± 0.02 respectively. This indicates a better cycle stability of MgeFeV composites as compared to HRBM Mg where maximum hydrogen concentration gradually decreased from 7.3 to 6.9 wt% (5.5%) during just 10 H absorption/desorption cycles performed at the same conditions [10]. The sample Mg-10(FeV)-5G exhibits similar hydrogenation/ dehydrogenation cyclic behaviour, but with 10e12% higher observed H absorption/desorption capacities (see Fig. 4A; the experimental data are presented in Supplementary Information, Fig. S11). After initial drop in H absorption capacity (most probably, due to incomplete H desorption at P  3 bar), the H absorption capacity even slightly increases. However, after the 25th cycle, a gradual decrease in the H absorption takes place, from 5.52 to 4.82 wt% (by ~13%), when passing from cycle 25 to cycle 98. At the same time, during the first 25 hydrogenation/dehydrogenation cycles a significant improvement of the hydrogenation kinetics was observed without a noticeable deterioration of the reversible hydrogen storage capacity. This can be seen from Fig. 4B which shows as-collected kinetic data of H absorption during the variable number of cycles, as well as Fig. 5A showing the evolution of the fitted kinetic parameters during the cycling (the values of the fitting parameters are presented in the Supplementary Information, Table S2). When increasing the number of cycles from 1 to 10, the characteristic time, t0, decreases in more than 2 times and

Please cite this article as: Lototskyy M et al., Nanostructured hydrogen storage materials prepared by high-energy reactive ball milling of magnesium and ferrovanadium, International Journal of Hydrogen Energy, https://doi.org/10.1016/j.ijhydene.2019.01.135

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Fig. 4 e Hydrogen absorption at P ≥ 13 bar (A) and kinetic curves for the selected cycles (B; curve labels indicate the cycle numbers) for the sample Mg-10(FeV)-5G.

the rate constant, K ¼ 1/t0, increases accordingly. Further cycling gives similar values of the fitted rate constant, K ¼ 10 ± 1 min1. A possible origin of the reduction of the H absorption capacity during the cycling may be in a gradual sintering and recrystallization of Mg/MgH2 during the re-hydrogenation when, due to a very fast exothermic process at the beginning of the hydrogenation, the sample temperature approaches the equilibrium value, Teq, for the Mg 4 MgH2 transformation, which exceeds 400  C for the applied values of hydrogenation pressures (Fig. 5B; the spikes of the sample temperature at the beginning of H absorption are also clearly seen in Fig. S11). This observation agrees with our earlier results on the re-hydrogenation behaviour of Mg e V nanocomposites [9]. Grain growth followed by the recrystallization at T > 400  C results in elongation of the H diffusion pathways in the formed MgH2 that slows the hydrogenation down and, in turn, reduces the H capacity achieved during the rehydrogenation time. Similar effect was observed in our recent study of HRBM MgeTi composites cycled at T ¼ 350  C [27].

XRD Typical Rietveld refinements of the XRD patterns for the studied samples are presented in Fig. 6. The results of the refinements are presented in Supplementary Information, Table S3. After HRBM most of the samples show presence of nanocrystalline a- and g-modifications of MgH2 with the unit cell parameters being in a good agreement the reference data [19]. Furthermore, presence of extra phases (becoming major in the samples containing 50 mol.% of V or FeV) including vanadium and FeV hydrides was observed. From the XRD patterns for the samples with the highest (50 mol.% or 68 wt%) content of the V-based additive shown in Fig. 6 it can be seen that the replacement of vanadium by Ferrovanadium results in major changes in the phase composition of the hydrogenated material. In the as-milled sample Mge68V the V-based hydride phases give wellresolved XRD peaks corresponding to vanadium dihydride (FCC-VH2) and “monohydride” (BCT-V2H1þx) with the unit cell parameters being in a good agreement with the reference data [19] and with a crystallite size of around 45 nm.

Fig. 5 e A e change of the fitted kinetic parameters for Mg-10(FeV)-5G with the cycle number. B ere-hydrogenation kinetics of Mg-10(FeV)-5G at T0 ¼ 340  C during the tenth cycle of the cycle stability study (Fig. 4). The graphs related to the right Y axis in B show the measured sample temperature (Ts) and its corrected value (T*s; Eq. (6)) evaluated by accounting the thermocouple response time. The equilibrium van't Hoff temperature (Eq. (7)) is shown as well. Please cite this article as: Lototskyy M et al., Nanostructured hydrogen storage materials prepared by high-energy reactive ball milling of magnesium and ferrovanadium, International Journal of Hydrogen Energy, https://doi.org/10.1016/j.ijhydene.2019.01.135

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Fig. 6 e XRD patterns of the studied samples: observed (points) and calculated (line) intensities; the bottom line shows the difference, (observed e calculated); peak positions of the constituent phases are shown by ticks.

The as-milled sample Mg-68(FeV) exhibits a much more complicated XRD pattern. The refinements show that HRBM of Mg-68(FeV) results in the formation of nanocrystalline (~15 nm) vanadium-based hydrides with the lattice periods close to the values reported for the hydride phases BCT(V,Fe)2H and FCC (V,Fe)H formed in the hydrogenated Fesubstituted V [16]. The sample also contains significant amounts of non-hydrogenated BCC phase (estimated crystallite size ~7 nm) with the unit cell parameters

corresponding to the BCC phase in the starting FeV (Section Materials). Furthermore, the pattern contains peaks of another BCC phase with 3.6% higher lattice period corresponding to a solid solution of hydrogen in the parent BCC; the crystallite size for this phase is similar to the values for the BCT and FCC phases. The decrease in FeV content results in a slight increase in intensities and in narrowing of the peaks belonging to the Mg dihydrides in the as milled samples that corresponds to the

Please cite this article as: Lototskyy M et al., Nanostructured hydrogen storage materials prepared by high-energy reactive ball milling of magnesium and ferrovanadium, International Journal of Hydrogen Energy, https://doi.org/10.1016/j.ijhydene.2019.01.135

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increase of MgH2 crystallite size from 5 to 8 nm for Mg-68(FeV) to 14e16 nm for HRBM Mg without additives. The peaks corresponding to the other phases (BCT-V2H, FCC-(V,Fe)H, BCC(FeV)(Hx); see Fig. 6 (bottom labels) and Table S4) decrease in the intensity following the decrease in FeV content. The samples with x ¼ 2e10 also do not show a formation of the FCC-(V,Fe)H phase. Instead, only BCC (Da/a0 ¼ 1e2%) H solid solution is observed together with BCT-(V,Fe)2H. Introduction of carbon as AC and MWCNT into Mg-10(FeV) does not significantly change the XRD patterns of the asmilled samples. In contrast, addition of graphite results in significant changes of the pattern for Mg-10(FeV)-5G. This sample (Fig. 6) exhibits narrow lines solely belonging to aMgH2 with crystallite size around 40 nm, i.e. 2e4 times larger than the crystallite size of a-MgH2 and g-MgH2 in Mg-10(FeV) without or with other carbon additives. As expected, after the TDS experiments the dehydrogenated samples exhibited narrow peaks of individual Mg while the re-hydrogenation of Mg-10(FeV) resulted in the formation of well-crystalline a-MgH2. The crystallites size of Mg and a-MgH2 was about 80 and 110 nm, respectively. Ten hydrogen desorption e absorption cycles did not significantly change the ratio Mg/MgH2 and crystallite size of MgH2 confirming the stability of the H sorption characteristics at the beginning of the cycling (Fig. 4). At the same time, further cycling results in a gradual change of the phase-structural composition reflected in the XRD pattern (see Fig. 6). Despite the peaks of MgH2 and Mg remain narrow and the crystallite sizes insignificantly changes, the peak intensities decrease in ~10 times, and this is accompanied by the increase of the refined abundance of Mg which becomes a major phase in Mg-10(FeV)-5G cycled 100 times (Mg/MgH2 ¼ 5.5, see Table S3). Introducing of carbon (graphite) also causes a smaller grain size of MgH2 in the rehydrogenated samples: 40 nm after the first hydrogenation and 70 nm after the 100th one, against 100 nm for the Mge FeV composites without carbon additives (Table S3). Such changes agree with the data for MgeTi and MgeTieC composites recently studied by the authors [27]. For all of the studied samples, the re-hydrogenation and cycling result in the disappearance of the BCT-VH1ex. Apart from MgH2 and Mg, only nanocrystalline BCC phase (crystallite size 4e6 nm; Da/a0 ¼ 0e1.7%) was observed in the rehydrogenated and cycled Mg-10(FeV) and Mg-10(FeV)-5G. The carbon-containing samples (C ¼ AC, MWCNT) show a complete re-hydrogenation with no Mg phase left; the samples exhibit ~15% lower crystallite size of a-MgH2 than in the re-hydrogenated and cycled Mg-10(FeV) without a carbon additive.

SEM and TEM studies Typical SEM images of the as-milled samples are presented in Supplementary Information, Section “Results: SEM”; Fig. S14. As distinct from previously studied HRBM Mg, MgeC [10] and MgeFeTiO3 [12] composites exhibiting noticeable agglomeration of submicron particles, the agglomeration in the samples studied in this work was less pronounced. Introduction of carbon (see example for Mg-10(FeV)-5MWCNT in Fig. S14) results in the increased amount of smaller particles.

The EDS data collected during the SEM studies are in good agreement with the initial sample compositions; similar results were obtained during the TEM studies described below. High resolution TEM images of the as-milled samples with the highest concentration (50 mol.%) of V and FeV (Fig. 7, left) show that nanoparticles of V-based catalyst (dark inclusions) are embedded into the Mg-based matrix. Indexing of the corresponding SAD patterns (Fig. 7, right) confirmed phase composition of the materials determined by XRD (Section XRD). We note that the SAD pattern of the sample modified by Ferrovanadium (Fig. 7, right bottom) exhibits better developed and uniform rings indicating that the size of nanocrystallites in this sample is smaller than in the sample modified with individual vanadium. This observation is in agreement with the XRD data. Comparison of TEM images and SAD patterns for the asmilled, dehydrogenated and re-hydrogenated Mg-20(FeV) (see Supplementary information, Fig. S15) shows that dehydrogenation results in the coarsening/coalescence of the particles and appearance of well crystalline Mg phase. The BCC phase remains nanocrystalline (rings in the SAD patterns) in agreement with XRD data. The re-hydrogenated sample exhibits similar morphology to the de-hydrogenated one, with nanocrystalline hydrogenated BCC phase and well crystalline phases of a-MgH2 and Mg. Introduction of carbon (see examples for Mg-10(FeV)5MWCNT and Mg-10(FeV)-5G in Supplementary Information; Figs. S16eS20) results in a more uniform distribution of Vbased catalyst in the matrix of magnesium hydrides. The Vbased particles (dark inclusions) are very small, 1e10 nm, as can be seen from the high magnification images (top right). Low magnification images (bottom left) show that the dark inclusions are rather large (>50 nm), the corresponding SAD patterns (bottom right) exhibit presence of the corresponding V-based phases in nanocrystallite state, in a good agreement with the XRD data. In the as-milled Mg-10(FeV)-5G (Fig. S18), only nanocrystalline a-MgH2 was observed in the SAD pattern. TDS followed by re-hydrogenation results in the appearance of well-crystalline MgH2, however, the V-based phase remains nanocrystalline with very small particles uniformly distributed in the MgH2 matrix (Fig. S15, bottom). After 10 hydrogenation/dehydrogenation cycles of Mg10(FeV)-5G, the material shows a formation of the network structure formed by very fine (~10 nm) particles; the SAD pattern indicates presence of nanocrystallites of MgH2 and BCC phases (Fig. S19). Further cycling (Fig. S20) results in the formation of agglomerates of submicron particles, most probably, caused by sintering. The SAD pattern (bottom right) shows presence of well-crystalline Mg and MgH2, as well as nanocrystalline (wide rings) BCC phase, in a good agreement with the XRD data. The above-mentioned difference between the morphologies of Mg-10(FeV)-5G after 10 and 100 hydrogenation/ dehydrogenation cycles is visible in Fig. S21 (top). Fig. S21 (bottom) also shows elemental maps of V taken from the STEMeEDS spectral images (mid). It can be observed that apart from the morphological changes, the prolonged cycling also results in the segregation of V-based phase even though its crystallite size, according to the XRD and SAD data remains very small.

Please cite this article as: Lototskyy M et al., Nanostructured hydrogen storage materials prepared by high-energy reactive ball milling of magnesium and ferrovanadium, International Journal of Hydrogen Energy, https://doi.org/10.1016/j.ijhydene.2019.01.135

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Fig. 7 e TEM images (left) and the corresponding SAD patterns (right) of the as-milled samples containing 50 mol.% (~68 wt%) of V and FeV.

Conclusions This work shows that additives of Ferrovanadium (~80 wt% V) improve hydrogen sorption performance of magnesiumbased nanocomposites. This improvement includes facilitation of hydrogenation during reactive ball milling in H2 and causes faster dehydrogenation and re-hydrogenation kinetics. The most pronounced improvements were observed at the concentration of Ferrovanadium about 10 wt% when 2 h of ball milling yielded hydrogen storage material with H storage capacity about 7 wt%, characterised by lowered activation energy of MgH2 decomposition from original 107e130 (MgH2) to 60e80 kJ (mol H2)1, shortening of the characteristic rehydrogenation time by more than 10 times as compared to MgH2 without additives. At the same time, the improvements are less pronounced than those observed for pure vanadium or easily hydrogenated BCC vanadium alloys used as catalytic additives to magnesium. The differences are caused by incomplete

hydrogenation of FeV as compared to a pure V. If hydrogenation of MgeV composite is accompanied by the transformation of BCC-V into FCC-VH2 and BCT-V2H hydrides having a crystallite size ~50 nm, the hydrogenated Mg-(FeV) composites mostly contain nanostructured various crystallographic modifications of lower V hydrides, FCC-VH and BCTV2H with crystallite sizes 10e15 nm, together with a solid solution of H in the BCC phase, with crystallite sizes <10 nm. Cycling of hydrogen absorption and desorption resulted in recrystallization and in increase of the Mg/MgH2 crystallite size from the initial 10e20 to 80e100 nm. Ferrovanadium catalyst remained nanocrystalline under these conditions, and its fine particles were uniformly distributed in the matrix of Mg or MgH2. Most probably, this is the origin of relatively high cycle stability of the Mg-(FeV) composites exhibiting no loss of hydrogenation e dehydrogenation performance during at least 30 cycles at the working temperature of T ¼ 350  C even without carbon additives.

Please cite this article as: Lototskyy M et al., Nanostructured hydrogen storage materials prepared by high-energy reactive ball milling of magnesium and ferrovanadium, International Journal of Hydrogen Energy, https://doi.org/10.1016/j.ijhydene.2019.01.135

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Reversible H sorption capacity of Mg-(FeV) composites was found to be lower than the H sorption capacity of the as-milled material. The capacity drops from the initial ~7 wt% H to 4e6 wt.% H and this drop is proportional to the content of Ferrovanadium. The main reason for that could be in a lower catalytic activity of vanadium monohydride formed by FeV as compared to vanadium dihydride formed in MgeV composites. Carbon additives show a different effect on the performance of the composites depending on the additive to Mg-(FeV) used (graphite, activated carbon, multi wall carbon nanotubes). The best performance is observed for activated carbon, while the effect of both graphite and MWCNT is inferior resulting in slowing down the dehydrogenation and rehydrogenation of the composites. This shows absence of a direct catalytic effect of carbon on hydrogenation/dehydrogenation of Mg. At the same time, carbon additives significantly improve the cycling performance of the Mg-(FeV)eC composites resulting in a more complete hydrogenation of Mg and, in turn, in a smaller drop in the reversible hydrogen storage capacity as compared to the capacity of the as-milled material. In summary, it can be concluded that Ferrovanadium has a potential to become a cheaper and effective alternative to a pure vanadium in the development of the weight-efficient hydrogen storage materials based on nanostructured magnesium hydride.

[2]

[3]

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Acknowledgements [11]

This work was supported by ERAfrica EU FP7 program, project RE-037 ‘‘HENERGY00 , and EU Horizon 2020/RISE project “HYDRIDE4MOBILITY”. South African co-authors (ML, JG, MWD, VL) also acknowledge the Department of Science and Technology (DST), Republic of South Africa, to be co-investor of the abovementioned project ‘‘HENERGY00 , as well as the funder of the HySA Program (project KP3-S02) supporting co-authors from HySA Systems. V. Yartys acknowledges a support from Research Council of Norway (project 285146 “IEA Task Energy Storage and Conversion Based on Hydrogen”) and fruitful collaboration within Magnesium Group of International Energy Agency (IEA) Task 32 “Hydrogen based energy storage”.

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Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi.org/10.1016/j.ijhydene.2019.01.135.

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Please cite this article as: Lototskyy M et al., Nanostructured hydrogen storage materials prepared by high-energy reactive ball milling of magnesium and ferrovanadium, International Journal of Hydrogen Energy, https://doi.org/10.1016/j.ijhydene.2019.01.135

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Please cite this article as: Lototskyy M et al., Nanostructured hydrogen storage materials prepared by high-energy reactive ball milling of magnesium and ferrovanadium, International Journal of Hydrogen Energy, https://doi.org/10.1016/j.ijhydene.2019.01.135