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Hydrogen storage characteristics of the nanocrystalline and amorphous Mg–Nd–Ni–Cu-based alloys prepared by melt spinning
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Hydrogen storage characteristics of the nanocrystalline and amorphous MgeNdeNieCubased alloys prepared by melt spinning Yanghuan Zhang a,b,*, Haitao Wang a,b, Tingting Zhai b, Tai Yang b, Yan Qi b, Dongliang Zhao b a
Key Laboratory of Integrated Exploitation of Baiyun Obo Multi-Metal Resources, Inner Mongolia University of Science and Technology, Baotou 014010, China b Department of Functional Material Research, Central Iron and Steel Research Institute, Beijing 100081, China
article info
abstract
Article history:
Nanocrystalline and amorphous MgeNdeNieCu-based (Mg24Ni10Cu2)100xNdx (x ¼ 0e20)
Received 18 July 2013
alloys were prepared by melt spinning and their structures as well as hydrogen storage
Received in revised form
characteristics were investigated. The analysis of XRD, TEM and SEM linked with EDS
6 December 2013
reveal that all the as-cast alloys hold a multiphase structure, containing Mg2Ni-type major
Accepted 22 December 2013
phase as well as some secondary phases Mg6Ni, Nd5Mg41 and NdNi, whose amounts clearly
Available online 20 January 2014
grow with Nd content rising. Furthermore, the as-spun Nd-free alloy displays an entire nanocrystalline structure whereas the as-spun Nd-added alloys have a mixed structure of
Keywords:
nanocrystalline and amorphous, moreover, the amorphization degree of the alloys visibly
Mg2Ni-type alloys
increases with Nd content rising, implying that the addition of Nd facilitates the glass
Nd addition
forming in the Mg2Ni-type alloy. The addition of Nd results in a slight decrease in the
Melt spinning
hydrogen absorption capacity of the as-cast and spun alloys, but it significantly enhances
Nanocrystalline and amorphous
their hydrogen storage kinetics and hydriding/dehydriding cycle stability of the alloy. In
Hydrogen storage kinetics
order to reveal the capacity degradation mechanism of the as-spun alloy, the structure evolution of the nanocrystalline and amorphous alloys during the hydridingedehydriding cycles was investigated. It is found that the root causes of leading to the capacity degradation of the nanocrystalline and amorphous alloys are nanocrystalline coarsening, crystal defect decreasing and amorphous phase crystallizing. Copyright ª 2013, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved.
1.
Introduction
The successful application of fuel cell technology in portable electronic devices or vehicles is globally considered to be dependent on the availability of an economical, safe and
practical storage of hydrogen. Of the available ways of hydrogen storage, metal hydride systems are considered to be more accurate, efficient and safe, representing the frontiers of technology [1]. Among these materials, Mg and Mg-based compounds always attract the interest of many investigators due to the great abundance, the light weight of Mg, and the
* Corresponding author. Department of Functional Material Research, Central Iron and Steel Research Institute, No. 76, Xueyuannan Road, Haidian District, 100081 Beijing, China. Tel.: þ86 10 62183115; fax: þ86 10 62187102. E-mail addresses: [email protected], [email protected] (Y. Zhang). 0360-3199/$ e see front matter Copyright ª 2013, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.ijhydene.2013.12.139
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high hydrogen capacity of the hydrides, e.g. 7.6 wt.% for MgH2, 3.6 wt.% for Mg2NiH4, 4.5 wt.% for Mg2CoH5 and 5.4 wt.% for Mg2FeH6. In spite of the practical application of the Mg and Mg-based alloys being nullified by their disadvantages such as relatively high hydrogen desorption temperature, sluggish hydriding/dehydriding kinetics as hydrogen storage materials of on-board use, Mg2Ni-type metallic hydrides are still looked upon as one of the most promising hydrogen storage materials applied to hydrogen fuel cell [2]. It is universally agreed that alloying and modifying the microstructure are the effective approaches for improving the hydriding properties [3]. Particularly, the partial substitution of some elements (Y, La, Zr) for Mg and (Cu, Fe, V, Cr, Co) for Ni in Mg2Ni alloy significantly improves the hydrogen absorption/desorption properties [4,5]. Furthermore, it was documented that the hydriding and dehydriding kinetics of the Mg and Mg-based alloys are very sensitive to their structures [6]. Especially, their hydrogen storage properties are strongly affected by nanometer scale structures because of the thermodynamic and kinetic features. High energy ball-milling (HEBM) has inarguably been convinced to be quite an effective method for fabricating nanocrystalline and amorphous Mg and Mg-based alloys [7]. However, some inherent shortcomings of ball milling technology seem to be unavoidable: the milled materials are easily polluted by steel balls and air, even though in very good protection conditions. Also, the cycle stabilities of the milled Mg and Mg-based alloys are very poor owing to the vanishment of the metastable structures generated by ball milling during the multiple hydriding and dehydriding cycles [8]. Hence, it is necessary to look for other alternative methods to synthesize nanocrystalline and amorphous Mg and Mg-based alloys. The melt-spinning technique is definitely regarded as a useful method to obtain an amorphous and/or nanocrystalline structure without the inherent disadvantages of the BM process. It was ascertained that the Mg-based alloys with nanocrystalline and amorphous structure produced by meltspinning exhibit excellent hydriding characteristics, similar to the alloys produced by the HEBM process [9]. Also, the microstructure created by melt spinning displays much higher stability during the hydrogen absorbing and desorbing cycles compared with the microstructure generated by BM [10]. ic et al. [11], partial It is reported by Song et al. [8] and Simic substitution of Cu for Ni in Mg2Ni alloy makes the stability of the hydride decreased and the hydrogen desorption reaction easier, which could be very effective in preventing the corrosion of materials, so as to improve the cycle stability. What is more, the substitution of rare earth elements for Mg considerably improves the hydrogen absorption capacity and kinetics of the Mg2Ni alloy due to the catalytic effect of rare earth elements [5,12e14]. Our previous investigations also have found that the substitution of La for Mg and M (M ¼ Cu, Co, Mn) for Ni improved the electrochemical and gaseous hydrogen storage performances of the Mg2Ni-type alloys dramatically [15,16]. Therefore, we expect that the joint addition of Cu and Nd combining with a proper melt spinning technique may ameliorate the hydrogen storage characteristics of the Mg2Ni-type alloy more pronounced. To validate this, a systematical investigation about the effects of Nd content on the structures and hydrogen storage performances of the
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Fig. 1 e XRD profiles of the as-cast and spun (Mg24Ni10Cu2)100LxNdx (x [ 0e20) alloys: (a) As-cast, (b) Asspun (40 m/s).
(Mg24Ni10Cu2)100xNdx (x ¼ 0e20) electrode alloys has been performed, and some experimental results were provided.
2.
Experimental
The compositions of the experimental alloys were (Mg24Ni10Cu2)100xNdx (x ¼ 0, 5, 10, 15, 20). For convenience, the alloys were denoted with Nd content as Nd0, Nd5, Nd10, Nd15 and Nd20, respectively. The alloy ingots were prepared by using a vacuum induction furnace in a helium atmosphere at the pressure of 0.04 MPa to prevent Mg from volatilizing. A part of the as-cast alloys were re-melted and spun by melt spinning
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Fig. 2 e HRTEM micrographs and ED patterns of the as-spun (40 m/s) (Mg24Ni10Cu2)100LxNdx (x [ 0e20) alloys: (a) Nd0, (b) Nd5, (c) Nd10.
with a rotating copper roller cooled by water. The spinning rates used in the experiment were 10, 20, 30 and 40 m/s, respectively, which were approximately expressed by the linear velocity of the copper roller. The phase structures of the as-cast and spun alloys were determined by X-ray diffraction (XRD) (D/max/2400). The diffraction, with the experimental parameters of 160 mA, 40 kV and 10 ( )/min respectively, was performed with CuKa1 radiation filtered by graphite.
A Philips SEM (QUANTA 400) linked with an energy dispersive spectrometer (EDS) was used for morphological observation and chemical composition analysis of the as-cast alloys. The thin film samples of the as-spun alloys prepared by using ion etching technology were observed by high resolution transmission electron microscope (HRTEM) (JEM-2100F, operated at 200 kV) and their crystalline states were inspected by electron diffraction (ED).
Fig. 3 e SEM images together with typical EDS spectra of the as-cast (Mg24Ni10Cu2)100LxNdx (x [ 0e20) alloys: (a) Nd0, (b) Nd5, (c) Nd20.
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Thermal stability and crystallization of the as-spun alloys were studied by means of differential scanning calorimetry (DSC). Also, the DSC curve was measured in argon flow, and the samples were placed in an alumina crucible, heating temperature from 30 to 550 C with the heating rate of 10 C/ min. The hydrogen absorption and desorption kinetics of the alloys were measured by an automatically controlled Sieverts apparatus. The hydrogen absorption was conducted at 2 MPa hydrogen pressure (in fact, this pressure is the initial pressure of hydriding process) and 200 C, and the hydrogen desorption at the pressure of 1 104 MPa and 250 C.
3.
Results and discussion
3.1.
Microstructure characteristics
Demonstrated in Fig. 1 are the XRD patterns of the as-cast and spun (40 m/s) (Mg24Ni10Cu2)100xNdx (x ¼ 0e20) alloys. It is found that the addition of Nd brings on an evident change of the phase compositions of the as-cast alloys, generating secondary phases NdNi and Nd5Mg41 along with the main phase Mg2Ni, and the amounts of these secondary phases increase markedly with Nd content rising. Furthermore, we note that the as-spun (40 m/s) Nd0 alloy displays very sharp diffraction peak, characteristic for an entire crystalline structure with a grain size of about 10 nm, which is calculated by Scherrer’s equation based on the FWHM values of the major diffraction peak (203) in Fig. 1 (b). Interestingly, it is found that, even if spun at the same spinning rate, the Ndadded alloys show completely different diffraction peaks from that of the Nd0 alloy. The very broad and flat diffraction peaks indicate that the as-spun Nd-added alloys are of an amorphous structure, and the degree of the amorphization visibly increases with Nd content rising, suggesting that the addition of Nd facilitates the glass forming in the Mg2Ni-type alloy. The results are also evidenced by TEM examination, as demonstrated in Fig. 2. It can be seen that the as-spun Nd0 alloy displays an entire nanocrystalline structure with an average crystal size of about 10 nm and its electron diffraction (ED) pattern appears sharp multi-haloes, corresponding to a crystalline structure. Nevertheless, the as-spun Nd5 and Nd10 alloys differing from Nd0 alloy exhibit a clear feature of the nanocrystalline embedded in the amorphous matrix, and their electron diffraction patterns are consisted of broad and dull halo, indicating the presence of an amorphous structure, which conforms to the XRD observations as depicted in Fig. 1. Likewise, the phase components of the as-cast alloys are analyzed by SEM, as illustrated in Fig. 3. We note that the morphologies of the as-cast alloys obviously changed with the growing of Nd content. Evidently, the as-cast Nd0 alloy shows a typical dendritic structure, but it disappears completely and some secondary phases appear with Nd content increasing. The EDS patterns also reveal that all the experimental Nd-added alloys have a multiphase structure, containing major phase Mg2Ni and secondary phases Nd5Mg41 and NdNi, which is consistent with the result of the XRD detection very well. It can be seen from Fig. 1 that the Cu-rich phase has not been detected in the as-cast and spun
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alloys. Moreover, the EDS patterns show that the Cu is distributed in each phase. According our previous researches [17,18], Cu addition causes a visible enlargement in the lattice parameters and cell volume of the alloys. So, it is justifiable that element Cu gets alloyed with the main phases.
3.2.
Hydrogen storage characteristics
3.2.1.
Capacity and kinetics
Fig. 4 shows the variations of the hydrogen absorption quantity of the as-cast and spun (40 m/s) (Mg24Ni10Cu2)100xNdx (x ¼ 0e20) alloys with reaction time of hydrogenating. Evidently, all the as-cast and spun alloys display a fast hydrogen absorption rate in the initial stage, and after that the hydrogen content is almost saturated at the next quite a long hydrogenation time. We note that for all the experimental alloys, the hydrogen absorption capacity in 100 min (Ca100 ) is more than 95% of their saturated hydrogen absorption capacities. Hence, it is considered to be reasonable to take Ca100 value as hydrogen absorption capacity of the alloy. It is found that Ca100 values of the as-cast and spun alloys always decline with Nd content increasing, suggesting that the addition of Nd gives rise to an obviously negative contribution to the hydrogen absorption capacity of the alloy, probably owing to the appearing of the secondary phases (Mg6Ni, NdNi, Nd5Mg41) created by Nd adding because their hydrogen-absorption abilities are weaker than that of Mg2Ni phase under the same experimental conditions. As is well-known, in addition to the relatively high hydrogen absorption and desorption temperature, another issue by which the attempt of putting Mg2Ni-type alloy into practice application is severely frustrated by its sluggish hydriding and dehydriding kinetics. Apparently, it is very necessary to examine the influence of Nd adding on the hydriding and dehydriding kinetics of the alloys. Here, the hydriding kinetics of the alloy is characterized by its hydrogen absorption saturation ratio (Rat ), a ratio of the hydrogen absorption capacity at a fixed time to the saturated hydrogen absorption capacity of the alloy, which is defined as Rat ¼ Cat =Ca100 100%, where Cat and Ca100 are hydrogen absorption capacities at t min and 100 min, respectively. As is shown in Fig. 4, the differences of hydrogen absorption capacities are much more obvious in the 5e10 min, considering the better possibility of mutual comparison, we take the hydrogen absorption time of 5 min as a criterion. The variations of Ra5 (t ¼ 5) values of the as-cast and spun (Mg24Ni10Cu2)100xNdx (x ¼ 0e20) alloys with Nd content are presented in Fig. 5. It is derived that the Ra5 values grow from 32.3% to 43.3% for the ascast alloys and from 88.4% to 94.9% for the as-spun (40 m/s) ones with the increasing of Nd content from 0 to 20, respectively, meaning that the addition of Nd plays a beneficial role on the hydriding kinetics of the alloy. Particularly, we note that, for all the Nd contents, the as-spun alloys always exhibit much superior hydriding kinetics to the as-cast ones. Also, it is found that the minimum difference of the Ra5 values between the as-spun and as-cast alloy for the fixed Nd content is much larger than the maximum difference of the Ra5 values generated by changing the Nd content for the fixed spinning rate. So, we powerfully believe that the hydrogen absorption kinetics of the alloy is chiefly dominated by its structure.
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Fig. 4 e Evolution of hydrogen absorption quantities of the (Mg24Ni10Cu2)100LxNdx (x [ 0e20) alloys with time: (a) Ascast, (b) As-spun (40 m/s).
Fig. 6 e Evolution of hydrogen desorption quantities of the (Mg24Ni10Cu2)100LxNdx (x [ 0e20) alloys with time: (a) Ascast, (b) As-spun (40 m/s).
Described in Fig. 6 are the relationship between the hydrogen desorption quantities of the as-cast and spun (Mg24Ni10Cu2)100xNdx (x ¼ 0e20) alloys and the reaction time. An important feature of alloys emerging in the dehydriding process shows the initially fast hydrogen desorption, subsequently, the rate of hydrogen desorption declines sharply. It is evident that the H-desorbed quantities of the as-cast and spun alloys invisibly augment with Nd content growing. Specifically, the Cd100 (H-desorbed capacity in 100 min) values are
enhanced from 0.301 to 1.232 wt% for the as-cast alloys and from 1.983 to 2.296 wt% for the as-spun (40 m/s) alloys by increasing the Nd content from 0 to 20. Similarly, the hydrogen desorption kinetics of the alloy is symbolized by hydrogen desorption ratio (Rdt ), a ratio of the H-desorbed capacity at a fixed time to the saturated hydrogen absorption capacity of the alloy, which is defined as Rdt ¼ Cdt =Ca100 100%, where Ca100 is the hydrogen absorption capacity at 100 min and Cdt is the hydrogen desorption capacity at the time of t min,
Fig. 5 e Evolution of the Ra5 values of the as-cast and spun (Mg24Ni10Cu2)100LxNdx (x [ 0e20) alloys with Nd content.
Fig. 7 e Evolution of the Rd10 values of the as-cast and spun (Mg24Ni10Cu2)100LxNdx (x [ 0e20) alloys with Nd content.
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Fig. 8 e Evolution of the hydrogen absorption quantity (Ca100 ) of the as-spun (30 m/s) Nd0 and Nd10 alloys with cycle number.
respectively. In order to facilitate comparison, here, we take hydrogen absorption time of 10 min as the standard. Thus, the relationship between the Rd10 (t ¼ 10) values of the as-cast and spun (Mg24Ni10Cu2)100xNdx (x ¼ 0e20) alloys and the amount of Nd-added can be established easily, as demonstrated in Fig. 7. Obviously, the addition of Nd makes a positive contribution to the dehydriding kinetics of the as-cast and spun alloys. More specifically, increasing the Nd content from 0 to 20 gives rise to an augment of the Rd10 value from 8.2% to 36.1% for the as-cast alloy and from 38.9% to 58.2% for the as-spun (40 m/s) alloy, respectively. What is more, it is noticeable that, whatever the amount of Nd-added is, the as-spun alloy shows a much larger Rd10 value than the as-cast one, suggesting that the melt spinning facilitates the dehydriding rate of the alloy. In terms of the hydriding and dehydriding kinetics affected by Nd adding, some elucidations can be provided. As is known to all, the hydrogen absorption process of hydrogen storage alloys consists of the following steps: the rate of hydrogen molecular dissociation at the surface, the capability of hydrogen atoms penetrating from the oxide layer surface into the metal, the rate of hydrogen atoms diffusing into the bulk
Fig. 9 e Evolution of hydrogen absorption quantity of the as-spun (30 m/s) Nd0 and Nd10 alloys with reaction time at 1st and 100th cycles.
Fig. 10 e DSC profiles of the as-spun (30 m/s) (Mg24Ni10Cu2)100LxNdx (x [ 5e20) alloys.
metal and through the hydride already formed, yet, the hydrogen absorption rate is predominated by the slowest step. The positive impact of Nd-added on the hydriding kinetics is principally ascribed to the enlarged cell volume and the secondary phases created by Nd adding, which have been evidenced by our previous work [15]. Cui et al. consider that the increase of the lattice constants and cell volume facilitates to decrease the diffusion activation energy of hydrogen atoms, enhancing hydrogen diffusion rate [19]. Also, the secondary phases generated by Nd adding probably engender a catalytic action, weakening the bonding between Mg and H atoms and consequently speeding up the rate of hydrogen diffusion. The improved hydrogen desorption kinetics by Nd adding is ascribed to two reasons. Firstly, the addition of Nd considerably strengthens the glass forming ability of Mg2Ni-type alloy and the amorphous structure Mg2Ni shows excellent hydrogen desorption capability. Secondly, such substitution decreases the stability of the hydride and makes the desorption reaction easier [5]. With respect to the positive contribution of the melt spinning to the hydrogen storage kinetics, it is considered to be most likely associated with the changed structure of the alloy induced by the melt spinning. The crystalline material, when melt spun, becomes at least partially disordered and its structure turns into nanocrystalline, generating a lot of new crystallites and grain boundaries (Fig. 2). Hence, some crystal defects such as dislocations, stacking faults and grain boundaries are introduced, as evidenced by our published work [20], which may prompt the diffusion of hydrogen in materials by providing numerous sites with low diffusion activation energy [21]. As stated by Kumar et al. [22], changing the structure of the Mg2Ni alloy from polycrystalline to nanocrystalline results in a drop of about 100 K in absorbing/desorbing temperature, namely from 573 K to 473 K. It is evidenced by Zhao et al. [23] that hydrogen adheres to the surface of the nanocrystalline nickel more strongly than the polycrystalline nickel, facilitating the
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Fig. 11 e SEM observations of the granular morphologies of as-spun (30 m/s) Nd10 alloy before and after hydriding and dehydriding cycle: (a) Before cycling, (b) After 10 cycles, (c) After 20 cycles.
dissociation reaction of hydrogen molecular at the surface of the alloy particles.
3.2.2.
Cycle stability
In order to investigate the effect of Nd adding on the hydriding and dehydriding cycle stability of the as-spun alloy, the variations of the saturated hydrogen capacity (Ca100 ) with cycle number are presented in Fig. 8. It can be seen that, after 100 cycles, the Ca100 values of the as-spun Nd0 and Nd10 alloys degrade to 3.03 and 3.0 wt%, respectively. The slope of the curve qualitatively reflects the degradation rate of the hydrogen capacity during the hydriding and dehydriding cycles, namely the smaller the slope of the curve is, the better the cycle stability of the alloy will be. Clearly, the slope of the as-spun Nd10 alloy is less than that of the Nd0 one, indicating that the Nd adding plays a beneficial role on the cycle stability of the alloy. As a matter of
Fig. 12 e XRD patterns of the as-spun (30 m/s) Nd10 alloy before and after hydriding and dehydriding cycle.
fact, we found that, after 100 cycles, the Nd0 and Nd10 alloys still keep very high hydrogen absorption capacities, as demonstrated in Fig. 9. A very similar result was reported by Li et al. [24]. It is noted from Fig. 9 that increasing the cycle number results in the slope of the curves rising obviously, meaning a significant reduction of hydrogen absorption rate. Noticeably, the saturated hydrogen capacities of the alloys after different numbers of cycles are very closely if time is long enough, suggesting that increasing the cycle number has a slight impact on capacity of hydrogen. In an attempt to elucidate the possible reasons for the capacity loss of the asspun alloys, we first study the crystallization behavior of amorphous structure. As reported by Liang et al., the crystallization of amorphous phase in the Mg-based alloy caused an obvious change of hydrogen storage performances [25]. A relatively high temperature (523K) is kept during hydrogen desorption, at which the crystallization reaction might happen. The thermal stability of the as-spun (30 m/s) (Mg24Ni10Cu2)100xNdx (x ¼ 5e20) alloys were examined by means of DSC, as described in Fig. 10. The result indicates that the crystallization temperatures of the amorphous phase in the alloys are all above 456 C, which is much higher than dehydrogenation temperature (250 C). Thus, we can surmise that crystallization of the amorphous phase will not occur in the dehydrogenation temperature. As is known to all, another cause making the capacity degradation is the pulverization of the alloy particles. So, the morphologies of the particles before and after cycling are observed by SEM, as illustrated in Fig. 11. A lot of cracks can be clearly seen on the surfaces of the alloy particles after 10 cycles, and with the increasing of cycle numbers, the pulverization degree of the particles becomes more serious. It must be pointed out that the capacity degradation of the alloys can not be ascribed to the pulverization, which is supported by two aspects: Firstly, the cycling tests are curried out under nearly absolute vacuum in order to prevent the surfaces of alloy particles from oxidized. Secondly, it is generally accepted that reducing the sizes of the alloy particles facilitates to hydrogen storage capacity and kinetics of the
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Fig. 13 e TEM examination of the microstructure evolutions of as-spun (30 m/s) Nd0 and Nd10 alloys before and after hydriding and dehydriding cycle: (a) Nd0 before cycling, (b) Nd0 after 10 cycles, (c) Nd0 after 20 cycles, (d) Nd10 before cycling, (e) Nd10 after 10 cycles, (f) Nd10 after 20 cycles.
alloys. As a matter of fact, the crystallization of the amorphous phase is inevitable during repeated cycling in spite of dehydrogenation temperature being much lower than crystallization one, which is also validated by XRD data, as depicted in Fig. 12. It is found that the amorphous structure basically remains unchanged but only a little crystal Mg2Ni phase appears after ten cycles. As for the alloys after 20 cycles, the amorphous structure is found to turn to nanocrystalline almost. Obviously, the crystallization reaction of the amorphous phase gradually occurs with cycle number increasing, which is also evidenced by TEM inspection, as shown in Fig. 13. In order to observe the microstructure more clearly, the amplified TEM morphologies are also presented in Fig. 13. It can be seen that the as-spun Nd0 alloy before cycling displays an entire nanocrystalline structure. Although no any amorphous structure is detected, the partially disordered structure with numerous grain boundaries can be seen clearly from Fig. 13 a. After 10 cycles, the grains visibly grow and the disorder degree of the structure evidently decreases (Fig. 13 b). Further increasing the cycle number to 20 cycles, the disorder structure of the alloy near completely disappears and the grain size grows to about 30 nm (Fig. 13 c). It can be seen that the amount of the amorphous phase in the as-spun Nd10 alloy gradually
decreases with cycle number increasing (Fig. 13 def). Based on the above-mentioned results, it can be concluded that the capacity degradation of the as-spun alloy is principally ascribed to the crystallization of amorphous, the reduction of the crystal defects and the growth of the grains.
4.
Conclusions
The hydrogen storage characteristics of the as-cast and spun (Mg24Ni10Cu2)100xNdx (x ¼ 0e20) alloys and their structure evolution during cycling were investigated. The major conclusions are summarized as follows: 1. All the as-cast alloys hold a multiphase structure, consisting of the major phase Mg2Ni and some secondary phases. The addition of Nd brings on the forming of Nd5Mg41 and NdNi phases whose amounts markedly grow with Nd content increasing. The addition of Nd facilitates to the glass forming of the Mg2Ni-type alloy and the amorphization degree of the alloys significantly increases with Nd content rising. 2. The addition of Nd makes a negative contribution to the hydrogen absorption capacity of the as-cast and spun
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alloys, but it promotes their hydriding and dehydriding kinetics significantly. 3. The as-spun alloys display excellent cycle stability, to be attributed to high stability of the nanocrystalline and amorphous structure generated by melt spinning. The repeated cycling many times results in a gradual evolution of the nanocrystalline and amorphous structure, namely the crystallization of amorphous, the reduction of the crystal defects and the growth of the grains, which are responsible for the degradation of the cycle performances.
Acknowledgements This work is financially supported by National Natural Science Foundations of China (51161015 and 51371094) and Natural Science Foundation of Inner Mongolia, China (2011ZD10).
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ic MV, Zdujic M, Dimitrijevic R, Nikolic -Bujanovic L, [11] Simic NH. Hydrogen absorption and electrochemical Popovic properties of Mg2Ni-type alloys synthesized by mechanical alloying. J Power Sources 2006;158(1):730e4. [12] Darriet B, Pezat M, Hbika A, Hagenmuller P. Application of magnesium rich rare-earth alloys to hydrogen storage. Int J Hydrog Energy 1980;5(2):173e8. [13] Spassov T, Ko¨ster U. Hydrogenation of amorphous and nanocrystalline Mg-based alloys. J Alloys Compd 1999;287(1e2):243e50. [14] Teresiak A, Gebert A, Savyak M, Uhlemann M, Mickel C, Mattern N. In situ high temperature XRD studies of the thermal behaviour of the rapidly quenched Mg77Ni18Y5 alloy under hydrogen. J Alloys Compd 2005;398(1e2):156e64. [15] Ren HP, Zhang YH, Li BW, Zhao DL, Guo SH, Wang XL. Influence of the substitution of La for Mg on the microstructure and hydrogen storage characteristics of Mg20xLaxNi10 (x ¼ 0e6) alloys. Int J Hydrog Energy 2009;34:1429e36. [16] Zhang YH, Zhao C, Yang T, Shang HW, Xu C, Zhao DL. Comparative study of electrochemical performances of the as-melt Mg20Ni10-xMx (M ¼ None, Cu, Co, Mn; x ¼ 0, 4) alloys applied to Ni/metal hydride (MH) battery. J Alloys Compd 2013;555:131e7. [17] Zhang YH, Li BW, Ren HP, Ma ZH, Guo SH, Wang XL. An electrochemical investigation of melt-spun nanocrystalline Mg20Ni10xCux (x ¼ 0e4) electrode alloys. Int J Hydrog Energy 2010;35:2385e92. [18] Zhang YH, Li BW, Ren HP, Guo SH, Zhao DL, Wang XL. Microstructure and hydrogen storage characteristics of meltspun nanocrystalline Mg20Ni10xCux (x ¼ 0e4) alloys. Mater Chem Phys 2010;124:795e802. [19] Cui N, Luo JL. Electrochemical study of hydrogen diffusion behavior in Mg2Ni-type hydrogen storage alloy electrodes. Int J Hydrog Energy 1999;24:37e42. [20] Zhang YH, Li BW, Ren HP, Hu F, Zhang GF, Guo SH. Gaseous and electrochemical hydrogen storage kinetics of nanocrystalline Mg2Ni-type alloy prepared by rapid quenching. J Alloys Compd 2011;509:5604e10. [21] Wu Y, Han W, Zhou SX, Lototsky MV, Solberg JK, Yartys VA. Microstructure and hydrogenation behavior of ball-milled and melt-spun Mge10Nie2Mm alloys. J Alloys Compd 2008;466:176e81. [22] Kumar LH, Viswanathan B, Murthy SS. Hydrogen absorption by Mg2Ni prepared by polyol reduction. J Alloys Compd 2008;461:72e6. [23] Zhao XY, Ding Y, Ma LQ, Wang LY, Yang M, Shen XD. Electrochemical properties of MmNi3.8Co0.75Mn0.4Al0.2 hydrogen storage alloy modified with nanocrystalline nickel. Int J Hydrog Energy 2008;33:6727e33. [24] Liu Y, Li Q, Lin G, Chou KC, Xu KD. Properties of Mg2NiH4 prepared by microwave-assisted activation synthesis from micro-particles. J Alloys Compd 2009;468:455e61. [25] Liang GY, Wu DC, Li L, Huang LJ. A discussion on decay of discharge capacity for amorphous MgeNieNd hydrogen storage alloy. J Power Sources 2009;186:528e31.
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Hydrogen storage characteristics of the nanocrystalline and amorphous MgeNdeNieCubased alloys prepared by melt spinning Yanghuan Zhang a,b,*, Haitao Wang a,b, Tingting Zhai b, Tai Yang b, Yan Qi b, Dongliang Zhao b a
Key Laboratory of Integrated Exploitation of Baiyun Obo Multi-Metal Resources, Inner Mongolia University of Science and Technology, Baotou 014010, China b Department of Functional Material Research, Central Iron and Steel Research Institute, Beijing 100081, China
article info
abstract
Article history:
Nanocrystalline and amorphous MgeNdeNieCu-based (Mg24Ni10Cu2)100xNdx (x ¼ 0e20)
Received 18 July 2013
alloys were prepared by melt spinning and their structures as well as hydrogen storage
Received in revised form
characteristics were investigated. The analysis of XRD, TEM and SEM linked with EDS
6 December 2013
reveal that all the as-cast alloys hold a multiphase structure, containing Mg2Ni-type major
Accepted 22 December 2013
phase as well as some secondary phases Mg6Ni, Nd5Mg41 and NdNi, whose amounts clearly
Available online 20 January 2014
grow with Nd content rising. Furthermore, the as-spun Nd-free alloy displays an entire nanocrystalline structure whereas the as-spun Nd-added alloys have a mixed structure of
Keywords:
nanocrystalline and amorphous, moreover, the amorphization degree of the alloys visibly
Mg2Ni-type alloys
increases with Nd content rising, implying that the addition of Nd facilitates the glass
Nd addition
forming in the Mg2Ni-type alloy. The addition of Nd results in a slight decrease in the
Melt spinning
hydrogen absorption capacity of the as-cast and spun alloys, but it significantly enhances
Nanocrystalline and amorphous
their hydrogen storage kinetics and hydriding/dehydriding cycle stability of the alloy. In
Hydrogen storage kinetics
order to reveal the capacity degradation mechanism of the as-spun alloy, the structure evolution of the nanocrystalline and amorphous alloys during the hydridingedehydriding cycles was investigated. It is found that the root causes of leading to the capacity degradation of the nanocrystalline and amorphous alloys are nanocrystalline coarsening, crystal defect decreasing and amorphous phase crystallizing. Copyright ª 2013, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved.
1.
Introduction
The successful application of fuel cell technology in portable electronic devices or vehicles is globally considered to be dependent on the availability of an economical, safe and
practical storage of hydrogen. Of the available ways of hydrogen storage, metal hydride systems are considered to be more accurate, efficient and safe, representing the frontiers of technology [1]. Among these materials, Mg and Mg-based compounds always attract the interest of many investigators due to the great abundance, the light weight of Mg, and the
* Corresponding author. Department of Functional Material Research, Central Iron and Steel Research Institute, No. 76, Xueyuannan Road, Haidian District, 100081 Beijing, China. Tel.: þ86 10 62183115; fax: þ86 10 62187102. E-mail addresses: [email protected], [email protected] (Y. Zhang). 0360-3199/$ e see front matter Copyright ª 2013, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.ijhydene.2013.12.139
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high hydrogen capacity of the hydrides, e.g. 7.6 wt.% for MgH2, 3.6 wt.% for Mg2NiH4, 4.5 wt.% for Mg2CoH5 and 5.4 wt.% for Mg2FeH6. In spite of the practical application of the Mg and Mg-based alloys being nullified by their disadvantages such as relatively high hydrogen desorption temperature, sluggish hydriding/dehydriding kinetics as hydrogen storage materials of on-board use, Mg2Ni-type metallic hydrides are still looked upon as one of the most promising hydrogen storage materials applied to hydrogen fuel cell [2]. It is universally agreed that alloying and modifying the microstructure are the effective approaches for improving the hydriding properties [3]. Particularly, the partial substitution of some elements (Y, La, Zr) for Mg and (Cu, Fe, V, Cr, Co) for Ni in Mg2Ni alloy significantly improves the hydrogen absorption/desorption properties [4,5]. Furthermore, it was documented that the hydriding and dehydriding kinetics of the Mg and Mg-based alloys are very sensitive to their structures [6]. Especially, their hydrogen storage properties are strongly affected by nanometer scale structures because of the thermodynamic and kinetic features. High energy ball-milling (HEBM) has inarguably been convinced to be quite an effective method for fabricating nanocrystalline and amorphous Mg and Mg-based alloys [7]. However, some inherent shortcomings of ball milling technology seem to be unavoidable: the milled materials are easily polluted by steel balls and air, even though in very good protection conditions. Also, the cycle stabilities of the milled Mg and Mg-based alloys are very poor owing to the vanishment of the metastable structures generated by ball milling during the multiple hydriding and dehydriding cycles [8]. Hence, it is necessary to look for other alternative methods to synthesize nanocrystalline and amorphous Mg and Mg-based alloys. The melt-spinning technique is definitely regarded as a useful method to obtain an amorphous and/or nanocrystalline structure without the inherent disadvantages of the BM process. It was ascertained that the Mg-based alloys with nanocrystalline and amorphous structure produced by meltspinning exhibit excellent hydriding characteristics, similar to the alloys produced by the HEBM process [9]. Also, the microstructure created by melt spinning displays much higher stability during the hydrogen absorbing and desorbing cycles compared with the microstructure generated by BM [10]. ic et al. [11], partial It is reported by Song et al. [8] and Simic substitution of Cu for Ni in Mg2Ni alloy makes the stability of the hydride decreased and the hydrogen desorption reaction easier, which could be very effective in preventing the corrosion of materials, so as to improve the cycle stability. What is more, the substitution of rare earth elements for Mg considerably improves the hydrogen absorption capacity and kinetics of the Mg2Ni alloy due to the catalytic effect of rare earth elements [5,12e14]. Our previous investigations also have found that the substitution of La for Mg and M (M ¼ Cu, Co, Mn) for Ni improved the electrochemical and gaseous hydrogen storage performances of the Mg2Ni-type alloys dramatically [15,16]. Therefore, we expect that the joint addition of Cu and Nd combining with a proper melt spinning technique may ameliorate the hydrogen storage characteristics of the Mg2Ni-type alloy more pronounced. To validate this, a systematical investigation about the effects of Nd content on the structures and hydrogen storage performances of the
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Fig. 1 e XRD profiles of the as-cast and spun (Mg24Ni10Cu2)100LxNdx (x [ 0e20) alloys: (a) As-cast, (b) Asspun (40 m/s).
(Mg24Ni10Cu2)100xNdx (x ¼ 0e20) electrode alloys has been performed, and some experimental results were provided.
2.
Experimental
The compositions of the experimental alloys were (Mg24Ni10Cu2)100xNdx (x ¼ 0, 5, 10, 15, 20). For convenience, the alloys were denoted with Nd content as Nd0, Nd5, Nd10, Nd15 and Nd20, respectively. The alloy ingots were prepared by using a vacuum induction furnace in a helium atmosphere at the pressure of 0.04 MPa to prevent Mg from volatilizing. A part of the as-cast alloys were re-melted and spun by melt spinning
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Fig. 2 e HRTEM micrographs and ED patterns of the as-spun (40 m/s) (Mg24Ni10Cu2)100LxNdx (x [ 0e20) alloys: (a) Nd0, (b) Nd5, (c) Nd10.
with a rotating copper roller cooled by water. The spinning rates used in the experiment were 10, 20, 30 and 40 m/s, respectively, which were approximately expressed by the linear velocity of the copper roller. The phase structures of the as-cast and spun alloys were determined by X-ray diffraction (XRD) (D/max/2400). The diffraction, with the experimental parameters of 160 mA, 40 kV and 10 ( )/min respectively, was performed with CuKa1 radiation filtered by graphite.
A Philips SEM (QUANTA 400) linked with an energy dispersive spectrometer (EDS) was used for morphological observation and chemical composition analysis of the as-cast alloys. The thin film samples of the as-spun alloys prepared by using ion etching technology were observed by high resolution transmission electron microscope (HRTEM) (JEM-2100F, operated at 200 kV) and their crystalline states were inspected by electron diffraction (ED).
Fig. 3 e SEM images together with typical EDS spectra of the as-cast (Mg24Ni10Cu2)100LxNdx (x [ 0e20) alloys: (a) Nd0, (b) Nd5, (c) Nd20.
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Thermal stability and crystallization of the as-spun alloys were studied by means of differential scanning calorimetry (DSC). Also, the DSC curve was measured in argon flow, and the samples were placed in an alumina crucible, heating temperature from 30 to 550 C with the heating rate of 10 C/ min. The hydrogen absorption and desorption kinetics of the alloys were measured by an automatically controlled Sieverts apparatus. The hydrogen absorption was conducted at 2 MPa hydrogen pressure (in fact, this pressure is the initial pressure of hydriding process) and 200 C, and the hydrogen desorption at the pressure of 1 104 MPa and 250 C.
3.
Results and discussion
3.1.
Microstructure characteristics
Demonstrated in Fig. 1 are the XRD patterns of the as-cast and spun (40 m/s) (Mg24Ni10Cu2)100xNdx (x ¼ 0e20) alloys. It is found that the addition of Nd brings on an evident change of the phase compositions of the as-cast alloys, generating secondary phases NdNi and Nd5Mg41 along with the main phase Mg2Ni, and the amounts of these secondary phases increase markedly with Nd content rising. Furthermore, we note that the as-spun (40 m/s) Nd0 alloy displays very sharp diffraction peak, characteristic for an entire crystalline structure with a grain size of about 10 nm, which is calculated by Scherrer’s equation based on the FWHM values of the major diffraction peak (203) in Fig. 1 (b). Interestingly, it is found that, even if spun at the same spinning rate, the Ndadded alloys show completely different diffraction peaks from that of the Nd0 alloy. The very broad and flat diffraction peaks indicate that the as-spun Nd-added alloys are of an amorphous structure, and the degree of the amorphization visibly increases with Nd content rising, suggesting that the addition of Nd facilitates the glass forming in the Mg2Ni-type alloy. The results are also evidenced by TEM examination, as demonstrated in Fig. 2. It can be seen that the as-spun Nd0 alloy displays an entire nanocrystalline structure with an average crystal size of about 10 nm and its electron diffraction (ED) pattern appears sharp multi-haloes, corresponding to a crystalline structure. Nevertheless, the as-spun Nd5 and Nd10 alloys differing from Nd0 alloy exhibit a clear feature of the nanocrystalline embedded in the amorphous matrix, and their electron diffraction patterns are consisted of broad and dull halo, indicating the presence of an amorphous structure, which conforms to the XRD observations as depicted in Fig. 1. Likewise, the phase components of the as-cast alloys are analyzed by SEM, as illustrated in Fig. 3. We note that the morphologies of the as-cast alloys obviously changed with the growing of Nd content. Evidently, the as-cast Nd0 alloy shows a typical dendritic structure, but it disappears completely and some secondary phases appear with Nd content increasing. The EDS patterns also reveal that all the experimental Nd-added alloys have a multiphase structure, containing major phase Mg2Ni and secondary phases Nd5Mg41 and NdNi, which is consistent with the result of the XRD detection very well. It can be seen from Fig. 1 that the Cu-rich phase has not been detected in the as-cast and spun
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alloys. Moreover, the EDS patterns show that the Cu is distributed in each phase. According our previous researches [17,18], Cu addition causes a visible enlargement in the lattice parameters and cell volume of the alloys. So, it is justifiable that element Cu gets alloyed with the main phases.
3.2.
Hydrogen storage characteristics
3.2.1.
Capacity and kinetics
Fig. 4 shows the variations of the hydrogen absorption quantity of the as-cast and spun (40 m/s) (Mg24Ni10Cu2)100xNdx (x ¼ 0e20) alloys with reaction time of hydrogenating. Evidently, all the as-cast and spun alloys display a fast hydrogen absorption rate in the initial stage, and after that the hydrogen content is almost saturated at the next quite a long hydrogenation time. We note that for all the experimental alloys, the hydrogen absorption capacity in 100 min (Ca100 ) is more than 95% of their saturated hydrogen absorption capacities. Hence, it is considered to be reasonable to take Ca100 value as hydrogen absorption capacity of the alloy. It is found that Ca100 values of the as-cast and spun alloys always decline with Nd content increasing, suggesting that the addition of Nd gives rise to an obviously negative contribution to the hydrogen absorption capacity of the alloy, probably owing to the appearing of the secondary phases (Mg6Ni, NdNi, Nd5Mg41) created by Nd adding because their hydrogen-absorption abilities are weaker than that of Mg2Ni phase under the same experimental conditions. As is well-known, in addition to the relatively high hydrogen absorption and desorption temperature, another issue by which the attempt of putting Mg2Ni-type alloy into practice application is severely frustrated by its sluggish hydriding and dehydriding kinetics. Apparently, it is very necessary to examine the influence of Nd adding on the hydriding and dehydriding kinetics of the alloys. Here, the hydriding kinetics of the alloy is characterized by its hydrogen absorption saturation ratio (Rat ), a ratio of the hydrogen absorption capacity at a fixed time to the saturated hydrogen absorption capacity of the alloy, which is defined as Rat ¼ Cat =Ca100 100%, where Cat and Ca100 are hydrogen absorption capacities at t min and 100 min, respectively. As is shown in Fig. 4, the differences of hydrogen absorption capacities are much more obvious in the 5e10 min, considering the better possibility of mutual comparison, we take the hydrogen absorption time of 5 min as a criterion. The variations of Ra5 (t ¼ 5) values of the as-cast and spun (Mg24Ni10Cu2)100xNdx (x ¼ 0e20) alloys with Nd content are presented in Fig. 5. It is derived that the Ra5 values grow from 32.3% to 43.3% for the ascast alloys and from 88.4% to 94.9% for the as-spun (40 m/s) ones with the increasing of Nd content from 0 to 20, respectively, meaning that the addition of Nd plays a beneficial role on the hydriding kinetics of the alloy. Particularly, we note that, for all the Nd contents, the as-spun alloys always exhibit much superior hydriding kinetics to the as-cast ones. Also, it is found that the minimum difference of the Ra5 values between the as-spun and as-cast alloy for the fixed Nd content is much larger than the maximum difference of the Ra5 values generated by changing the Nd content for the fixed spinning rate. So, we powerfully believe that the hydrogen absorption kinetics of the alloy is chiefly dominated by its structure.
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Fig. 4 e Evolution of hydrogen absorption quantities of the (Mg24Ni10Cu2)100LxNdx (x [ 0e20) alloys with time: (a) Ascast, (b) As-spun (40 m/s).
Fig. 6 e Evolution of hydrogen desorption quantities of the (Mg24Ni10Cu2)100LxNdx (x [ 0e20) alloys with time: (a) Ascast, (b) As-spun (40 m/s).
Described in Fig. 6 are the relationship between the hydrogen desorption quantities of the as-cast and spun (Mg24Ni10Cu2)100xNdx (x ¼ 0e20) alloys and the reaction time. An important feature of alloys emerging in the dehydriding process shows the initially fast hydrogen desorption, subsequently, the rate of hydrogen desorption declines sharply. It is evident that the H-desorbed quantities of the as-cast and spun alloys invisibly augment with Nd content growing. Specifically, the Cd100 (H-desorbed capacity in 100 min) values are
enhanced from 0.301 to 1.232 wt% for the as-cast alloys and from 1.983 to 2.296 wt% for the as-spun (40 m/s) alloys by increasing the Nd content from 0 to 20. Similarly, the hydrogen desorption kinetics of the alloy is symbolized by hydrogen desorption ratio (Rdt ), a ratio of the H-desorbed capacity at a fixed time to the saturated hydrogen absorption capacity of the alloy, which is defined as Rdt ¼ Cdt =Ca100 100%, where Ca100 is the hydrogen absorption capacity at 100 min and Cdt is the hydrogen desorption capacity at the time of t min,
Fig. 5 e Evolution of the Ra5 values of the as-cast and spun (Mg24Ni10Cu2)100LxNdx (x [ 0e20) alloys with Nd content.
Fig. 7 e Evolution of the Rd10 values of the as-cast and spun (Mg24Ni10Cu2)100LxNdx (x [ 0e20) alloys with Nd content.
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Fig. 8 e Evolution of the hydrogen absorption quantity (Ca100 ) of the as-spun (30 m/s) Nd0 and Nd10 alloys with cycle number.
respectively. In order to facilitate comparison, here, we take hydrogen absorption time of 10 min as the standard. Thus, the relationship between the Rd10 (t ¼ 10) values of the as-cast and spun (Mg24Ni10Cu2)100xNdx (x ¼ 0e20) alloys and the amount of Nd-added can be established easily, as demonstrated in Fig. 7. Obviously, the addition of Nd makes a positive contribution to the dehydriding kinetics of the as-cast and spun alloys. More specifically, increasing the Nd content from 0 to 20 gives rise to an augment of the Rd10 value from 8.2% to 36.1% for the as-cast alloy and from 38.9% to 58.2% for the as-spun (40 m/s) alloy, respectively. What is more, it is noticeable that, whatever the amount of Nd-added is, the as-spun alloy shows a much larger Rd10 value than the as-cast one, suggesting that the melt spinning facilitates the dehydriding rate of the alloy. In terms of the hydriding and dehydriding kinetics affected by Nd adding, some elucidations can be provided. As is known to all, the hydrogen absorption process of hydrogen storage alloys consists of the following steps: the rate of hydrogen molecular dissociation at the surface, the capability of hydrogen atoms penetrating from the oxide layer surface into the metal, the rate of hydrogen atoms diffusing into the bulk
Fig. 9 e Evolution of hydrogen absorption quantity of the as-spun (30 m/s) Nd0 and Nd10 alloys with reaction time at 1st and 100th cycles.
Fig. 10 e DSC profiles of the as-spun (30 m/s) (Mg24Ni10Cu2)100LxNdx (x [ 5e20) alloys.
metal and through the hydride already formed, yet, the hydrogen absorption rate is predominated by the slowest step. The positive impact of Nd-added on the hydriding kinetics is principally ascribed to the enlarged cell volume and the secondary phases created by Nd adding, which have been evidenced by our previous work [15]. Cui et al. consider that the increase of the lattice constants and cell volume facilitates to decrease the diffusion activation energy of hydrogen atoms, enhancing hydrogen diffusion rate [19]. Also, the secondary phases generated by Nd adding probably engender a catalytic action, weakening the bonding between Mg and H atoms and consequently speeding up the rate of hydrogen diffusion. The improved hydrogen desorption kinetics by Nd adding is ascribed to two reasons. Firstly, the addition of Nd considerably strengthens the glass forming ability of Mg2Ni-type alloy and the amorphous structure Mg2Ni shows excellent hydrogen desorption capability. Secondly, such substitution decreases the stability of the hydride and makes the desorption reaction easier [5]. With respect to the positive contribution of the melt spinning to the hydrogen storage kinetics, it is considered to be most likely associated with the changed structure of the alloy induced by the melt spinning. The crystalline material, when melt spun, becomes at least partially disordered and its structure turns into nanocrystalline, generating a lot of new crystallites and grain boundaries (Fig. 2). Hence, some crystal defects such as dislocations, stacking faults and grain boundaries are introduced, as evidenced by our published work [20], which may prompt the diffusion of hydrogen in materials by providing numerous sites with low diffusion activation energy [21]. As stated by Kumar et al. [22], changing the structure of the Mg2Ni alloy from polycrystalline to nanocrystalline results in a drop of about 100 K in absorbing/desorbing temperature, namely from 573 K to 473 K. It is evidenced by Zhao et al. [23] that hydrogen adheres to the surface of the nanocrystalline nickel more strongly than the polycrystalline nickel, facilitating the
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Fig. 11 e SEM observations of the granular morphologies of as-spun (30 m/s) Nd10 alloy before and after hydriding and dehydriding cycle: (a) Before cycling, (b) After 10 cycles, (c) After 20 cycles.
dissociation reaction of hydrogen molecular at the surface of the alloy particles.
3.2.2.
Cycle stability
In order to investigate the effect of Nd adding on the hydriding and dehydriding cycle stability of the as-spun alloy, the variations of the saturated hydrogen capacity (Ca100 ) with cycle number are presented in Fig. 8. It can be seen that, after 100 cycles, the Ca100 values of the as-spun Nd0 and Nd10 alloys degrade to 3.03 and 3.0 wt%, respectively. The slope of the curve qualitatively reflects the degradation rate of the hydrogen capacity during the hydriding and dehydriding cycles, namely the smaller the slope of the curve is, the better the cycle stability of the alloy will be. Clearly, the slope of the as-spun Nd10 alloy is less than that of the Nd0 one, indicating that the Nd adding plays a beneficial role on the cycle stability of the alloy. As a matter of
Fig. 12 e XRD patterns of the as-spun (30 m/s) Nd10 alloy before and after hydriding and dehydriding cycle.
fact, we found that, after 100 cycles, the Nd0 and Nd10 alloys still keep very high hydrogen absorption capacities, as demonstrated in Fig. 9. A very similar result was reported by Li et al. [24]. It is noted from Fig. 9 that increasing the cycle number results in the slope of the curves rising obviously, meaning a significant reduction of hydrogen absorption rate. Noticeably, the saturated hydrogen capacities of the alloys after different numbers of cycles are very closely if time is long enough, suggesting that increasing the cycle number has a slight impact on capacity of hydrogen. In an attempt to elucidate the possible reasons for the capacity loss of the asspun alloys, we first study the crystallization behavior of amorphous structure. As reported by Liang et al., the crystallization of amorphous phase in the Mg-based alloy caused an obvious change of hydrogen storage performances [25]. A relatively high temperature (523K) is kept during hydrogen desorption, at which the crystallization reaction might happen. The thermal stability of the as-spun (30 m/s) (Mg24Ni10Cu2)100xNdx (x ¼ 5e20) alloys were examined by means of DSC, as described in Fig. 10. The result indicates that the crystallization temperatures of the amorphous phase in the alloys are all above 456 C, which is much higher than dehydrogenation temperature (250 C). Thus, we can surmise that crystallization of the amorphous phase will not occur in the dehydrogenation temperature. As is known to all, another cause making the capacity degradation is the pulverization of the alloy particles. So, the morphologies of the particles before and after cycling are observed by SEM, as illustrated in Fig. 11. A lot of cracks can be clearly seen on the surfaces of the alloy particles after 10 cycles, and with the increasing of cycle numbers, the pulverization degree of the particles becomes more serious. It must be pointed out that the capacity degradation of the alloys can not be ascribed to the pulverization, which is supported by two aspects: Firstly, the cycling tests are curried out under nearly absolute vacuum in order to prevent the surfaces of alloy particles from oxidized. Secondly, it is generally accepted that reducing the sizes of the alloy particles facilitates to hydrogen storage capacity and kinetics of the
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Fig. 13 e TEM examination of the microstructure evolutions of as-spun (30 m/s) Nd0 and Nd10 alloys before and after hydriding and dehydriding cycle: (a) Nd0 before cycling, (b) Nd0 after 10 cycles, (c) Nd0 after 20 cycles, (d) Nd10 before cycling, (e) Nd10 after 10 cycles, (f) Nd10 after 20 cycles.
alloys. As a matter of fact, the crystallization of the amorphous phase is inevitable during repeated cycling in spite of dehydrogenation temperature being much lower than crystallization one, which is also validated by XRD data, as depicted in Fig. 12. It is found that the amorphous structure basically remains unchanged but only a little crystal Mg2Ni phase appears after ten cycles. As for the alloys after 20 cycles, the amorphous structure is found to turn to nanocrystalline almost. Obviously, the crystallization reaction of the amorphous phase gradually occurs with cycle number increasing, which is also evidenced by TEM inspection, as shown in Fig. 13. In order to observe the microstructure more clearly, the amplified TEM morphologies are also presented in Fig. 13. It can be seen that the as-spun Nd0 alloy before cycling displays an entire nanocrystalline structure. Although no any amorphous structure is detected, the partially disordered structure with numerous grain boundaries can be seen clearly from Fig. 13 a. After 10 cycles, the grains visibly grow and the disorder degree of the structure evidently decreases (Fig. 13 b). Further increasing the cycle number to 20 cycles, the disorder structure of the alloy near completely disappears and the grain size grows to about 30 nm (Fig. 13 c). It can be seen that the amount of the amorphous phase in the as-spun Nd10 alloy gradually
decreases with cycle number increasing (Fig. 13 def). Based on the above-mentioned results, it can be concluded that the capacity degradation of the as-spun alloy is principally ascribed to the crystallization of amorphous, the reduction of the crystal defects and the growth of the grains.
4.
Conclusions
The hydrogen storage characteristics of the as-cast and spun (Mg24Ni10Cu2)100xNdx (x ¼ 0e20) alloys and their structure evolution during cycling were investigated. The major conclusions are summarized as follows: 1. All the as-cast alloys hold a multiphase structure, consisting of the major phase Mg2Ni and some secondary phases. The addition of Nd brings on the forming of Nd5Mg41 and NdNi phases whose amounts markedly grow with Nd content increasing. The addition of Nd facilitates to the glass forming of the Mg2Ni-type alloy and the amorphization degree of the alloys significantly increases with Nd content rising. 2. The addition of Nd makes a negative contribution to the hydrogen absorption capacity of the as-cast and spun
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alloys, but it promotes their hydriding and dehydriding kinetics significantly. 3. The as-spun alloys display excellent cycle stability, to be attributed to high stability of the nanocrystalline and amorphous structure generated by melt spinning. The repeated cycling many times results in a gradual evolution of the nanocrystalline and amorphous structure, namely the crystallization of amorphous, the reduction of the crystal defects and the growth of the grains, which are responsible for the degradation of the cycle performances.
Acknowledgements This work is financially supported by National Natural Science Foundations of China (51161015 and 51371094) and Natural Science Foundation of Inner Mongolia, China (2011ZD10).
references
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