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Novel synthesis of superparamagnetic Ni–Co–B nanoparticles and their effect on superconductor properties of MgB2
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ScienceDirect Acta Materialia 70 (2014) 298–306 www.elsevier.com/locate/actamat
Novel synthesis of superparamagnetic Ni–Co–B nanoparticles and their effect on superconductor properties of MgB2 Mislav Mustapic´ a,⇑, Josip Horvat a, Md Shariar Hossain a, Ziqi Sun a, Zˇeljko Skoko b, David R.G. Mitchell a, Shi Xue Dou a a
Institute for Superconducting and Electronic Materials, AIIM, University of Wollongong, Squires Way, North Wollongong, NSW 2500, Australia b Department of Physics, Faculty of Science, University of Zagreb, Bijenicˇka c. 32, 10000 Zagreb, Croatia Received 12 December 2013; received in revised form 24 February 2014; accepted 24 February 2014 Available online 31 March 2014
Abstract A new procedure for the preparation of amorphous Ni–Co–B nanoparticles is reported, with a detailed investigation of their morphology by X-ray diffraction and transmission electron microscopy, as well as their magnetic properties. Many factors, such as chemical composition, anisotropy, size and shape of the particles, were controlled through chemical synthesis, resulting in the control of morphological and magnetic properties of the nanoparticles. Controlling pH values with ethylenediamine and using sodium dodecyl sulfate surfactant lowered the size of the nanoparticles to below 10 nm. Such a small structure and chemical disorder in nanocrystalline materials lead to magnetic properties that are different from those in their bulk-sized counterparts. The obtained nanoparticles can be used for different purposes, from pharmaceutical applications to implementations in different materials technology. The focus of this research is the synthesis of Ni–Co–B nanoparticles in a new way and studying the reaction of Ni–Co–B nanoparticles with Mg and B precursors and their effect on MgB2 properties. New nanostructures are formed in the reaction of Ni–Co–B nanoparticles with Mg: Mg2Ni, Co2Mg and possibly Mg2Co. Ó 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
Keywords: Chemical synthesis; Nanocrystalline nanostructure; Magnesium diboride; Differential thermal analysis; X-ray diffraction
1. Introduction Ni–Co–B nanoparticles have been widely prepared by the method of chemical reduction of metallic salts, as reported by several groups [1–5]. Ni–Co–B nanoparticles have been mostly used as catalysts for hydrogenation reaction and dechlorination [2]. The exact preparation procedure for the nanoparticles determines the nanoparticle size, structure and their physico-chemical properties. Another potential use of Ni–Co–B nanoparticles is the enhancement of the superconducting properties of a MgB2 superconductor. Incorporating Ni–Co–B nanoparticles of ⇑ Corresponding author.
E-mail address: [email protected] (M. Mustapic´).
10 nm size into the MgB2 superconductor is expected to improve its vortex pinning. This is because the coherence length of MgB2 is of the order of 10 nm. The magnetic moment of these nanoparticles is also deemed useful for further improvement of superconducting properties of MgB2. However, the interaction between the magnetic nanoparticles can result in their agglomeration, preventing their homogenous incorporation into MgB2 matrix. Making Ni–Co–B nanoparticles of small enough size and without agglomeration has been a challenge. Electron d-band occupancy in Ni–Co–B and strong electronegativity may lead to applications of Ni–Co–B nanoparticles in organic chemistry synthesis as well as in pharmacy. The great advantage of Ni–Co–B is its low cost of starting materials, as compared to palladium [6–8],
http://dx.doi.org/10.1016/j.actamat.2014.02.040 1359-6454/Ó 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
M. Mustapic´ et al. / Acta Materialia 70 (2014) 298–306
platinum [6–8] and rhodium [9] catalysts. An additional advantage is its magnetic properties, which can be manipulated through the size, shape and composition of nanoparticles [10]. This paper reports a new technique for the chemical synthesis of Ni–Co–B nanoparticles with a size of 5 nm, very narrow size distribution and little agglomeration. This technique relies on controlling the alkalinity of the reaction and the use of appropriate surfactant to prevent agglomeration. We used Ni(NO3)2 and Co(NO3)2 nitrate as the source of Ni and Co and NaBH4 as reducing agents and the source of B for Ni–Co–B nanoparticles [11]. Ethylenediamine was also added in this synthesis. The ethylenediamine kept the synthesis under alkaline pH condition, which is crucial for maintaining the reduction of metal salts and providing a high yield of the reaction. Ethylenediamine has catalytic properties, as well [12]. The reducing agent NaBH4 in neutral and acidic environments reduces hydrogen from water: NaBH4 ðsÞ þ 2H2 OðlÞ ! NaBO2 ðaqÞ þ 4H2 ðgÞ
ð1Þ
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SDS surfactant in aqueous solutions has a weak alkaline behaviour, and its surfactant ability is reduced in an acidic environment due to neutralization reaction between the dodecyl sulfate ion and H+ ions. The preparation procedure in which Ni2+ and Co2+ ions were dropped from a funnel into the three-necked flask with NaBH4 (aq) (Fig. 1) played an important role in providing the maximum yield of chemical reaction. Experimental work has shown less agglomeration of Ni–Co–B nanoparticles with described apparatus, probably due to less competition between Co2+ and Ni2+ ions. The emphasis of this report is on the materials preparation and chemistry of magnetic nanoparticles and their incorporation into MgB2, starting from the chemical synthesis of nanoparticles, solving problems of reaction between precursors and nanoparticles and finally measuring the change in properties due to those reactions. A detailed description of reaction mechanism and formation of new phases provides a crucial explanation of enhancement critical current density in a MgB2 superconductor.
Further, dissociated NaBH4 in water can reduce the hydronium ion to elemental hydrogen:
2. Experimental details
þ BH 4 ðaqÞ þ 4H ðaqÞ ! 2H2 ðgÞ þ BðOHÞ3 ðaqÞ
Ni–Co–B nanoparticles were prepared by the wet technique by reduction of cobalt and nickel salts with sodium boron hydride (NaBH4). Ni(NO3)2 and Co(NO3)2 were dissolved together (no reaction between them) in deionized (DI) water and placed in an additional funnel above the three-necked flask. On the other hand, NaBH4 was previously dissolved in DI water and placed into the threenecked flask. Ethylenediamine was added into the three-necked flask to maintain the reaction under alkaline pH conditions (pH 8–10) and also to catalyse the reaction. The
ð2Þ
All this would lead to low efficiency in producing Ni– Co–B nanoparticles and contamination of the resulting powder. In an alkaline pH environment, Reaction (2) does not occur, due to the negligible concentration of existing H+ ions. Ethylenediamine is soluble in water and reacts with water: H2 NCH2 CH2 NH2 ðlÞ þ H2 OðlÞ ! H2 NCH2 CH2 NH3 þ OH ðaqÞ
ð3Þ
On the other hand, excessive OH in the environment can be counterproductive due to another reaction between the metal salt (Co(NO3)2) and OH: Co2þ ðaqÞ þ 2OH ðaqÞ ! CoðOHÞ2 ðsÞ
ð4Þ
To provide maximum chemical reduction of metal salts, the (Co, Ni) reaction should be maintained in a medium alkaline environment and according to this experimental work, this is between pH = 8 and pH = 10. In such environment, the Reactions (1), (2) and (4) will not be significant and the reduction of metallic salts into metal borides will have the highest efficiency. Therefore, the adjustment such that 8 < pH < 10 directs the chemical reaction to reduce Ni2+ and Co2+ into Ni– Co–borides: NiðNO3 Þ2 þ CoðNO3 Þ2 þ NaBH4 þ OH ! Ni–Co–B þ NaNO3 þ H2 O
ð5Þ
Another important feature of the new synthesis is the addition of the surfactant sodium dodecyl sulfate (SDS), which prevents agglomeration of nanoparticles. Generally,
Fig. 1. Scheme of air-sensitive chemical equipment for producing nanoparticles.
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chemical reaction started with Co and Ni aqueous solution being added drop-wise into the three-necked flask with NBH4 solution. Immediately before the end of the dropwise introduction of Co and Ni nitrates from the funnel, SDS was added. Firstly, SDS was dissolved in DI water, separately bubbled for 2 h and finally added to the three-necked flask with a syringe. SDS has the role of surfactant, to prevent the agglomeration of nanoparticles. A surplus [1–4] of NaBH4 was added in order to completely reduce Ni2+ and Co2+ ions (the molar ratio of B/(Ni + Co) = 5). The synthesis was performed in a closed system under argon atmosphere, and all solutions were bubbled with argon for 2 h prior to the reaction. At the exhaust was set a chloride–calcium tube filled with CaCl2(s), to prevent entry of air and moisture retention (Fig. 1). The reaction was maintained at 273 K, and the duration of the reaction was less than 2 min. Black powder was formed instantly, and it was later extracted from the solution by centrifugation. The black powder was washed with DI water and ethanol several times to remove residual ions. Finally, the powder was dried for 2 h in a desiccator. X-ray powder diffraction (XRD) was performed at room temperature using an automatic Philips diffractometer, model PW1820 (Cu Ka radiation), in Bragg–Brentano geometry. XRD patterns for all samples were measured for samples at different stages of differential thermal analysis/ thermogravimetric analysis (DTA/TGA), after the heating of nanoparticles was stopped at 650, 770 and 850 °C. In mixing boron and magnesium with Ni–Co–B, it was important to ensure that any signature of the reaction between the nanoparticles and Mg is detectable in XRD analysis; for this reason 10 wt.% of Ni–Co–B was used. Powders were examined using a JEOL 2011 transmission electron microscope (TEM) operating at 200 kV. TEM specimens were made by dispersing the powder in ethanol and placing a drop of the suspension on a carbon support film on a copper grid. The microstructure of MgB2 samples were studied with a JEOL JSM-6400 scanning electron microscope (SEM). Magnetic measurements on nanoparticles were performed using a Quantum Design MPMS-5T SQUID magnetometer. Magnetic hysteresis loops M(H) were measured in the applied field within a range of 5 T at a few temperatures. Temperature dependence of magnetization M(T) for the temperature interval 10–300 K in a constant magnetic field H was measured in two modes: as zero field cooling (ZFC curves) and field cooling (FC curves). A ZFC magnetization curve was obtained by cooling the sample in an applied zero field from room temperature to 10 K, after which the required field was applied and the magnetization was measured during subsequent warming of the sample. The FC magnetization curve was obtained by measuring the magnetization as the sample was cooled down in the required field from room temperature to 10 K.
The memory effects of ZFC magnetization were tested by stopping the initial cool-down of the sample (10 K min1) at a set temperature (TS) for 104 s. After that, the sample was cooled down to 10 K, the magnetic field applied and magnetization measured as the sample was warmed up at a rate of 2 K min1. Thus obtained M(T) was compared with the ZFC measurements without stopping on initial cool-down. Any anomaly at TS shows a presence of strong interaction between the magnetic nanoparticles, creating a superspin glass system [13,14]. Magnetic measurements on MgB2 samples were performed using a Quantum Design PPMS-9T extraction magnetometer. The magnetization was measured at 5 K and 20 K in a time-varying magnetic field up to 9 T. DTA was performed by mixing obtained nanoparticles in different combinations with Mg, boron and Ni–Co–B (10 wt.%) pressed in small bars (200 mg) and finally measured by DTA/TGA (Mettler Toledo 1) up to 1100 °C. 2.1. Preparation of Ni–Co–B nanoparticles with boron and magnesium precursors Ni–Co–B nanoparticles described in this paper have been made with a view to improving the superconducting properties of MgB2. The nanoparticles described in this paper have a different size and size distribution than the Ni–Co–B nanoparticles described elsewhere. Because of that, they will have different chemical properties relevant to the growth of MgB2. Chemical reactions of these nanoparticles with the precursor powders for MgB2 were investigated using DTA and XRD techniques, in order to assess their suitability for this purpose. DTA results for the Ni– Co–B and boron mixture did not show any reaction between the boron and the Ni–Co–B (Fig. 2, inset). This result is consistent with the properties of elemental boron, which is quite inert and has a high melting point, 2300 °C. On the other hand, important information was observed on the DTA curve relating to the reaction of the Ni–Co–B nanoparticles and magnesium (Fig. 2). The small exothermal peak close to 500 °C (peak A) suggests a first reaction of the nanoparticles with magnesium. This small peak, which is accompanied by a broad plateau, can be described as a slow local solid–solid reaction between Ni–Co–B and Mg. The nanoparticles react with Mg and form a eutectic system [15,16] before Mg melts at 650 °C. Melted Mg can easily diffuse into and react with boron in Ni–Co–B [17]. Furthermore, with a size of less than 10 nm, the nanoparticles are extremely reactive, even at lower temperatures, than bulk Ni–Co–B, due to their large surface area. The melting point of magnesium can be clearly observed at 650 °C as the endothermic peak B in the DTA curve. With increasing temperature, two further distinct peaks appear, peak C at 700 °C and peak D at 770 °C (Fig. 2). Peak C represents the formation of the MgxBy phase [18], since 30 at.% of the atoms in Ni–Co–B are boron. The exothermal peak D at 770 °C represents the
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Fig. 2. DTA curve for mixed powders of magnesium and Ni–Co–B nanoparticles (10 wt.%). Inset: DTA curve for mixed powders of boron and Ni–Co–B nanoparticles (10 wt.%).
formation of the stable phases between Co–Mg and Ni–Mg [19]. The different shape of DTA curves (Figs. 2 and 3) is associated with the difference in the heat capacity of elements in experiments. The broad plateau in a sample of Ni–Co–B + Mg (Fig. 2) is due to the high molar capacity of magnesium (24.869 J mol1 K1), while in experimental sample with Mg + B + Ni–Co–B (Fig. 3), a substantial amount of magnesium is replaced with boron (molar ratio Mg:B = 1:2), which has considerably lower molar heat capacity (11.087 J mol1 K1). Fig. 3 shows the reactions between Mg, B and Ni–Co–B nanoparticles in the temperature range from 50 °C to 1000 °C. The DTA image clearly shows three distinguishable peaks marked with A, B, C. Small peaks at 500 °C can suggest the initial reaction between Ni–Co–B nanoparticles and magnesium in forming a eutectic liquid. The Fig. 4. XRD patterns: Mg + Ni–Co–B nanoparticles sintered at different temperatures (650 °C (black), 770 °C (red) and 850 °C (blue)). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
Fig. 3. DTA curve for mixed powders of Mg + 2B + Ni–Co–B nanoparticles (10 wt.%).
following exothermic peak A at 570 °C corresponds to the creation of MgB2 in solid–solid reaction. Ni–Co–B nanoparticles act as catalysts in MgB2 formation at a lower temperature of sintering. A similar mechanism has been reported by Zhao et al. [16], where nanoparticles form a local eutectic liquid with Mg and promote a faster local formation of MgB2 in the eutectic liquid phase. The melting point of magnesium is observed as the endothermic peak B at 650 °C. The large exothermal peak C above 730 °C corresponds to the completion of the reaction between Mg and B in forming MgB2 and the formation of new phases Mg2Ni and Mg2Co/Co2Mg [11,20–25]. The broad peak C in Fig. 3 is related to twin peaks C and D in Fig. 2.
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Fig. 5. SEM images for 2.5% Ni–Co–B doped MgB2, heat treated at 650 (a) and 850 °C (b).
In this work we have expanded X-ray diffraction analysis as compared to the previously reported study (see Fig. 4) [11]. Using the XRD technique, for Mg + 10 wt.% Ni–Co–B samples, we have analysed phase evolution as a function of different sintering temperatures, i.e. 650, 770 and 850 °C, respectively. All XRD patterns have three distinctive peaks at 32, 34 and 37°, which belong to pure magnesium. This is expected due to the large amount of magnesium as a starting material. Research into the phase diagrams revealed that Mg–Ni forms Mg2Ni as a stable phase [20–24], while the Mg–Co system is stable only as Co2Mg [20–25]. Although Mg2Co is not an equilibrium phase, trace amounts of this phase were identified in the XRD patterns (peaks at 13.4, 41.1 and 47.7°), which is further supported by similar observations for hydrogen storage materials [22–25]. During the heat treatment, small nanoparticles can react easily with magnesium, forming new phases such as Mg2Ni and Co2Mg/Mg2Co. However, such reactions can occur
only in localized areas with non-consistent ratio and composition. Therefore, we believe that Mg2Ni and Co2Mg are the most likely phases to be formed between 650 and 850 °C. As the heat treatment temperature changes, one can observe noticeable differences in the XRD patterns (Fig. 4). In particular, this applies to the Mg2Ni phase, which is much more prominent after sintering at higher temperatures (770 and 850 °C). Mg2Ni peaks are very difficult to be distinguished due to overlapping with other phases. Several peaks belonging to this phase were identified in the XRD patterns: 20, 21, 37, 40, 53, 57.5 and 72.5°. The intensity of these peaks increases with increasing sintering temperature. Cobalt based compounds (Co2Mg/ Mg2Co), on the other hand, can be identified in all of the XRD patterns (see Fig. 4), with a considerably lower intensity of the peaks in the 650 °C sintered sample. The peaks for Co2Mg are detected at 21.08, 22.45, 43.6 and 44.45°. However, all of these XRD features are quite faint, most likely because these new phases are in the form of nanoparticles, which have poor crystallinity and appear in very discrete localized reaction areas [11]. Boron from our Ni–Co–B nanoparticles could also react with magnesium and form Mg–B compounds such as MgB2 and MgB4. It is extremely difficult to prove their existence in the samples via XRD analysis, as the characteristic peaks for these compounds overlap with those for Mg and Co2Mg. However, if one assumes that such compounds are formed during the heat treatment process, the aforementioned MgB2 and MgB4 would be most likely. SEM images of doped MgB2 with 2.5 wt.% of Ni–Co–B are presented in Fig. 5a and b, for samples sintered at 650 and 850 °C, respectively. The sample heat treated at 650 °C has a large number of small-sized MgB2 crystallites. A considerable number of small crystals create more boundaries between grains, which usually have a negative impact on critical current density. On the other hand, the sample treated at 850 °C shows a larger grain size with flake-like crystallites. According to energy dispersive X-ray spectroscopy results [11], grain connectivity is improved with Ni–Co–B doping. Nanoparticles could have an impact through the hindrance of MgO formation on MgB2 grain boundaries, with a negative effect on critical current flow. In the formation of new Mg2Ni and Co2Mg phases, nickel and cobalt ions may compete with oxygen at higher temperature of sintering, preventing substantial formation of MgO. At the same time, newly formed Ni–Mg and Co–Mg phases provide good pinning centres for magnetic vortices. For samples sintered at 650 °C (Fig. 5a), distinguished spherical particles positioned at the grain boundaries are noticeable (yellow1 circles), and most likely belong to MgO or pure Mg. The sample sintered at 650 °C shows a significant amount of unreacted Mg and MgO in contrast to the
1 For interpretation of colour in Fig. 5, the reader is referred to the web version of this article.
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Fig. 6. XRD patterns of MgB2 samples doped with 2.5 wt.% of Ni–Co–B and sintered at 650 and 850 °C.
sample sintered at 850 °C, as can be seen in XRD patterns of MgB2 doped samples sintered at 650 and 850 °C (Fig. 6). The porosity shown in these figures is a common feature of MgB2. Because the density of MgB2 is higher than that of the precursor powders, the solid volume decreases upon the growth of MgB2, creating pores between the crystallites. Improvement in critical current density (JC) with the addition of Ni–Co–B nanoparticles can be seen in Fig. 7a and b. Heat treatment at high temperatures (850 °C) was required to obtain the improvement. Low temperature heat treatment (650 °C) resulted in deterioration of JC. JC was enhanced through better connectivity between the MgB2 grains, as well as through an increase of the vortex pinning [11]. Improvement in grain connectivity results in a higher affective cross-sectional area for the super-current flow through the sample. This shifts the JC(H) curves upward in Fig. 7a. On the other hand, increased vortex pinning results in a weaker field dependence of JC and tilting of the JC(H) curves, in addition to improved JC for H = 0. This tilting is not obvious in Fig. 7a, but it can be seen in Fig. 7b, showing the rate of decrease of ln(JC) with the field. The lower value of d[ln(JC)]/dH is indicative of improved vortex pinning. While the heat treatment at low temperatures did not improve the vortex pinning, the 850 °C heat treatment improved it, as shown in detail in Ref. [11].
Fig. 7. (a) Field dependence of JC for 2.5 wt.% Ni–Co–B doped and pure samples sintered at two different temperatures, 650 and 850 °C. (b) Magnetic field dependence of critical current density for pure MgB2 and 2.5 wt.% Ni–Co–B doped samples heat treated at 650 and 850 °C. (b) Rate of change of ln(JC) with field for samples heat treated at 850 °C, showing the improvement of vortex pinning with the addition of Ni–Co–B nanoparticles at 5 K.
2.2. Characterization of Ni–Co–B nanoparticles Fig. 8 shows XRD patterns for the as prepared Ni–Co– B nanoparticles. There are two clearly identifiable peaks, i.e. a diffused peak at 2h 35° and a relatively sharp peak at 2h 62°, in the diffraction pattern. This is in sharp contrast to the usually observed amorphous morphology for Ni–Co–B nanoparticles with a broad peak at 45° [1–5]. The presence of the peaks in our XRD pattern would indicate that Ni–Co–B nanoparticles in this case have some sort of nanocrystalline structure. Unfortunately, the
Fig. 8. XRD pattern of Ni–Co–B nanoparticles.
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Fig. 9. (a) Ni–Co–B nanoparticles precipitated in sheets with a high degree of agglomeration. HRTEM image of the Ni–Co–B phase showing it to be highly crystalline with a mean particle size of 6 nm (inset). FFT of the lattice spacings showing a limited number of spacings, suggesting a high crystal symmetry structure, such as cubic. (b) Mostly spherical shape of nanoparticles and size less than 5 nm.
observed peaks could not be attributed to a ternary Ni–Co–B compound, as there are no reference records available in the ICSD or ICDD databases. On the other hand, these peaks also do not belong to any of the known binary Ni–B or Co–B compounds, eliminating the possibility that these peaks originate from such impurities. It is quite difficult to explain such observations, but our hypothesis lies in the unique morphology of the nanoparticles, their size and surface layer. The appearance of the peaks might be correlated to the thickness of a nanoparticle’s surface layer. As the thickness of this surface layer becomes comparable to the diameter of the nanoparticle, a pseudo-superstructure can be formed, which results in the appearance of a diffraction peak. This is supported by the fact that XRD diffraction patterns are distinctly different when atomic ratios of Ni–Co–B nanoparticles and amounts of surfactants are changed and/or doped with other elements. It is critical to note, though, that in an alloy structure, such as that of Ni–Co–B, it is unusual and difficult to distinguish a particular crystalline phase. Further crystallographic investigations for this compound are necessary, in order to clarify observed crystal structure phenomenology. Crystallite size was also determined by the program XBroad for size-strain analysis [26]. The basis for XBroad analysis is the Warren–Averbach–Bertaut method, which is implemented in the program, and details can be found in Skoko et al. [26].
The analysis was done on the maximum located around 2h = 62°, which was firstly smoothed in order to obtain the best possible results. The results showed the area-weighted crystallite size (length of coherent diffraction columns) of 6 nm with negligible strain. Crystallite size obtained with XBroad is in strong agreement with the TEM analysis. The nanoparticulate phase precipitated in sheets in which the nanoparticles were highly agglomerated (Fig. 9a and b). Examination by high resolution TEM (HRTEM) confirmed that the particles were highly crystalline (Fig. 9a), with nearly all particles exhibiting strong lattice fringe contrast. Particles were typically rounded with no marked faceting. The mean particle diameter was estimated by measuring the long axis of 60 particles as 5.3 ± 1.9 nm. This dense agglomeration may help protect the particles from oxidation, since no surface membrane of amorphous oxide was found on any of the particles. The nanoscale size of the particles resulted in shape transform effects in diffraction, specifically highly diffuse diffraction peaks. XRD data showed weak and diffuse peaks, with only two faint peaks detected above background. Analysis of the spacings found in the HRTEM images (fast Fourier transform (FFT) Fig. 9a is shown as inset) revealed four diffraction spacings (0.215, 0.24, 0.29 and 0.46 nm). The limited number of spacings suggested that the phase had a high crystal symmetry, most likely cubic. The target stoichiometry of the phase based on mass balance was Ni1Co1B1. X-ray microanalysis lacked the sensitivity to detect B. However, electron energy loss spectroscopy confirmed the presence of B within the phase. Meaningful quantitative analysis was not possible with this technique. 2.3. Magnetic properties of Ni–Co–B nanoparticles Temperature dependence of the magnetization M(T) for Ni–Co–B nanoparticles measured under different applied magnetic fields is presented in Fig. 8. A maximum magnetization can be observed in the ZFC curve (Fig. 10) at
Fig. 10. ZFC and FC curves of Ni–Co–B nanoparticles.
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90 K, indicating that 90 K is the blocking temperature (TB) of Ni–Co–B nanoparticles [10,27]. Below TB, the particles are blocked and in the ferromagnetic state with an irreversible magnetization. Above TB, the nanoparticles are gradually unblocked and are allowed to reorient in the magnetic field. They are characterized by superparamagnetic behaviour. The sample magnetization is still irreversible up to 250 K, due to the interaction between some of the nanoparticles. This gives the irreversibility temperature (Tirr) as 250 K. Splitting between the ZFC and FC curves is a consequence of magnetic moment blocking in the studied systems of nanoparticles, which causes slow relaxation of magnetization of the system. The difference between Tirr and TB is connected to the size distribution of the particles and possibly the magnetic interaction between aggregated particles. This difference between TB and Tirr is quite large, which indicates the agglomeration of nanoparticles into clusters. The FC branch of magnetization is a continuously decreasing function of temperature. This is expected for superparamagnetic nanoparticles with weak or no interaction between them. Sasaki et al. [13] showed that the system of strongly interacting superparamagnetic nanoparticles (superspin glasses) will exhibit FC magnetization that is constant at low temperatures below TB, and it even decreases with cooling below TB. The ZFC branch of the magnetization is also expected to give characteristic memory effects for superspin glasses [13]. No memory effects of this kind [28] were observed for our Ni–Co–B nanoparticles. This shows that our samples consist of non-interacting, or weakly interacting, nanoparticles. The curved shape of the hysteresis loops with an applied magnetic field up to 2 T also indicates the superparamagnetic behaviour of the nanoparticles. Irreversible M(H) curves were obtained at low temperatures (Fig. 10). This
Fig. 11. Hysteresis loops of magnetic moment vs. field at different temperatures. Inset: expanded view of the low field part of the M(H) loops.
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irreversibility (i.e. magnetic hysteresis loop) is a consequence of slow relaxation of the magnetization of the system of blocked nanoparticles. At high temperatures (P100 K), M(H) curves are reversible (Fig. 11). The temperature dependence of the low field magnetization for magnetic nanoparticles depends on the type of the distribution function of the anisotropy energy barriers, which governs the relationship between the temperature of the maximum in MZFC(T), TMAX and the average blocking temperature, TB [10,29]. According to the Stoner–Wohlfarth model [30] magnetic hysteresis loops can be used to estimate the anisotropy energy density K = MH/2h, where M is close to saturation magnetization. If taking H HC at the same temperature, parameter h = 0.5 is used for the system of randomly oriented spherical magnetic particles. Magnetization saturation at 5 K can be estimated from magnetic moment (Fig. 11) and theoretical density of Ni–Co–B nanoparticles as 8743.2 kg m3 [31]. The value of saturation magnetization of Ni–Co–B is 349,720 A m1, and extrapolated coercive field HC is 79.60 A m1 at 5 K and 1000 Oe. Obtained anisotropy energy K, calculated from the Stoner–Wohlfarth model, is 34,955.4 J m3. From derived anisotropy energy using the Neel relation the relaxation time is s = s0 exp(DEB/kBTB), where kB is Boltzmann’s constant, T is the absolute temperature and s0 is an attempt time on the order of 109–1013 s. DEB = KV is the anisotropy energy barrier, where K is the effective anisotropy constant and V is the particle volume. The temperature at which this relaxation time s0 equals the measurement time sm and at which the particle system goes into the superparamagnetic region is the blocking temperature TB [32]. The volume of the particle can be calculated from V = 25kBTB/K, where the factor 25 is a typical ln(sm/s0) value for conventional magnetometry [33]. Obtained value of volume of Ni–Co–B nanoparticles calculated from Neel’s relation is 8.889 1025 m3, (diameter of particle 11.9 nm). On the other hand the volume of Ni–Co–B nanoparticles observed directly from TEM images (Fig. 9a and b) and calculated from V = p(D)3/6, (average diameter D 5 nm, Fig. 8b) is 8.71 1026 m3. In comparing the values of D obtained by different methods, a difference by a factor of two was obtained. The reasons for this are most likely the estimates used in obtaining the saturation magnetization and relaxation times. According to Zysler et al. [29] similar results of their M(H) measurements give evidence of a significant contribution to magnetization of uncompensated spins at nanoparticle surfaces. This is responsible for a non-saturated component in the M(H) curves. The results have been interpreted by a simple model, in which each single-domain nanoparticle is considered as a core–shell system, with uniaxial anisotropy of the core and surface anisotropy of the shell. The surface contribution is more evident in the absence of interparticle interactions.
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3. Conclusions
References
Due to being a promising candidate for doping of MgB2 superconductor, the properties of Ni–Co–B nanoparticles have been improved using a new synthesis technique. Further modification of the chemical treatment achieved nanoparticles with a narrowed distribution of 5 nm and nanocrystalline morphology. Furthermore, Ni–Co–B particles exhibit superparamagnetic behaviour, which can be well described within the model of relaxation of magnetic moments of particles. Magnetic Ni–Co–B nanoparticles added to Mg and B precursors in the formation of MgB2 react only with Mg, forming mostly paramagnetic Mg2Ni and MgCo2 phases. This process also promotes a slow solid–solid reaction (eutectic system) between Mg and B below the melting point of magnesium. After melting Mg at 650 °C, these reactions between the two main precursors, boron and magnesium, accelerate. Finally, above 750 °C, boron and Mg complete the formation of MgB2 and the incorporation of new stable phases Ni2Mg, Co2Mg/(Mg2Co) into the matrix of MgB2. Improvement of JC by Ni–Co–B doping is associated with the creation and incorporation of new particles into the matrix of MgB2, which enhances the grain connectivity as well as the pinning of magnetic vortices [11].
[1] Wang H, Yu Z, Chen H, Yang J, Deng J. Appl Catal A General 1995;129:143–9. [2] Shen J-H, Chen Y-W. J Mol Catal A Chem 2007;273:265–76. [3] Liawa B-J, Chenb C-H, Chenb Y-Z. Chem Eng J 2010;157:140–5. [4] Zhao B, Chen Y-W. J Non-Cryst Solids 2010;356:839–47. [5] Zhao B, Chen Y-W. Mater Chem Phys 2011;125:763–8. [6] Baltzly R, Phillips AR. J Am Chem Soc 1946;68:261. [7] Streltsova AA, Zelinskii ND. Bull Acad Sci USSR Div Chem Sci 1943;56. [8] Rylander PN, Kilroy M, Coven V. Engelhard Ind Tech Bull 1965;6: 11. [9] Dunworth WP, Nord FF. J Am Chem Soc 1952;74:1459. [10] Novosel N, Galic´ S, Pajic´ D, Skoko Zˇ, Loncˇarek I, Mustapic´ M, Zadro K, Babic´ E. Supercond Sci Technol 2013;26:065004. [11] Mustapic´ M, Horvat J, Hossain MS, Skoko Zˇ, Dou SX. Supercond Sci Technol 2013;26:075013. [12] Brown CA, Ahuja VK. J Chem Soc Chem Commun 1973:553–4. [13] Sasaki M, Jo¨nsson PE, Takayama H. Phys Rev B 2005;71:104405. [14] Suzuki M, Fullem ShI, Suzuki IS. Phys Rev B 2009;79:024418. [15] Ma ZQ, Liu YC, Shi QZ, Zhao Q, Gao ZM. Supercond Sci Technol 2008;21:065004. [16] Zhao Q, Liu Y, Zhao N, Penner S, Ma Z. IEEE Trans Nanotechnol 2011;10:2. [17] Ku¨hberger M, Gritzner G. Physica C 2002;370:39–43. [18] Chen SK, Wei M, MacManus-Driscoll JL. Appl Phys Lett 2006;88: 192512. [19] Tatsuhiko A, Hasehira K-I, Nishimura C. Mater Trans 2003;44(4): 601–10. [20] Chen C-Y, Lin S-K, Lin C-K. Mater Trans 2007;48(4):842–6. [21] Chiba M, Hotta H, Nobuki T, Sotoma A, Kuji T. J Magn Magn Mater 2007;316:454–7. [22] Kirkpatrick EM, Leslie-Pelecky D, Kim S-H, Rieke RD. J Appl Phys 1999;85:8–15. [23] Norek M et al. Int J Hydrogen Energy 2011;36:10760. [24] Bobet J-L, Akiba E, Darriet B. J Alloys Compd 2000;297:192–8. [25] Aizawa T, Hasehira K, Nishimura C. Mater Trans 2003;44:601. [26] Skoko Z, Popovic J, Dekanic K, Kolbas V, Popovic S. J Appl Crystallogr 2012;45:594–7. [27] Mustapic´ M, Pajic´ D, Novosel N, Babic´ E, Zadro K, Cindric´ M, ˇ , Bijelic´ M, Shcherbakovc A. Croatica Chem Acta Horvat J, Skoko Z 2010;83(3):275–82. [28] Su D, Horvat J, Munroe P, Ahn H, Ranjbartoreh AR, Wang G. Chem Eur J 2012;18:488–97. [29] Zysler RD, Romero H, Ramos CA, De Biasi E, Fiorani D. J Magn Matter 2003;266:233–42. [30] Stoner EC, Wohlfarth EP. Phil Trans R Soc 1948;A240:599–642. reprinted in IEEE Trans Magn 1991; 27: 3475–518. [31] US 5213907 A. Nickel–cobalt–boron-alloy deposited on a substrate. [32] Yoon S, Krishnan KM. J Appl Phys 2011;109:07B534. [33] Yoon S. J Korean Phys Soc 2011;59(5):3069–73.
Acknowledgments The lead author (M.M.) was a recipient of University of Wollongong PhD scholarship for the duration of the work presented in this paper. This work builds upon the outcomes of a project UKF 1B 01/07, which was supported through the Unity Through Knowledge Fund, designed by the Croatian government, represented by the Ministry of Science, Education and Sport. Financial support from the Australian Research Council through project LP120100175 is thankfully acknowledged. This work was partially supported by the Australian Research Council (Grant No. DE130101247) and a 2013 AIIM-CRG grant. The authors acknowledge the use of facilities and the assistance of Mr. Darren Attard at the University of Wollongong Electron Microscopy Centre. The help of Dr. Germanas Peleckis and Dr. Zong Qing Ma at ISEM is thankfully acknowledged.
ScienceDirect Acta Materialia 70 (2014) 298–306 www.elsevier.com/locate/actamat
Novel synthesis of superparamagnetic Ni–Co–B nanoparticles and their effect on superconductor properties of MgB2 Mislav Mustapic´ a,⇑, Josip Horvat a, Md Shariar Hossain a, Ziqi Sun a, Zˇeljko Skoko b, David R.G. Mitchell a, Shi Xue Dou a a
Institute for Superconducting and Electronic Materials, AIIM, University of Wollongong, Squires Way, North Wollongong, NSW 2500, Australia b Department of Physics, Faculty of Science, University of Zagreb, Bijenicˇka c. 32, 10000 Zagreb, Croatia Received 12 December 2013; received in revised form 24 February 2014; accepted 24 February 2014 Available online 31 March 2014
Abstract A new procedure for the preparation of amorphous Ni–Co–B nanoparticles is reported, with a detailed investigation of their morphology by X-ray diffraction and transmission electron microscopy, as well as their magnetic properties. Many factors, such as chemical composition, anisotropy, size and shape of the particles, were controlled through chemical synthesis, resulting in the control of morphological and magnetic properties of the nanoparticles. Controlling pH values with ethylenediamine and using sodium dodecyl sulfate surfactant lowered the size of the nanoparticles to below 10 nm. Such a small structure and chemical disorder in nanocrystalline materials lead to magnetic properties that are different from those in their bulk-sized counterparts. The obtained nanoparticles can be used for different purposes, from pharmaceutical applications to implementations in different materials technology. The focus of this research is the synthesis of Ni–Co–B nanoparticles in a new way and studying the reaction of Ni–Co–B nanoparticles with Mg and B precursors and their effect on MgB2 properties. New nanostructures are formed in the reaction of Ni–Co–B nanoparticles with Mg: Mg2Ni, Co2Mg and possibly Mg2Co. Ó 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
Keywords: Chemical synthesis; Nanocrystalline nanostructure; Magnesium diboride; Differential thermal analysis; X-ray diffraction
1. Introduction Ni–Co–B nanoparticles have been widely prepared by the method of chemical reduction of metallic salts, as reported by several groups [1–5]. Ni–Co–B nanoparticles have been mostly used as catalysts for hydrogenation reaction and dechlorination [2]. The exact preparation procedure for the nanoparticles determines the nanoparticle size, structure and their physico-chemical properties. Another potential use of Ni–Co–B nanoparticles is the enhancement of the superconducting properties of a MgB2 superconductor. Incorporating Ni–Co–B nanoparticles of ⇑ Corresponding author.
E-mail address: [email protected] (M. Mustapic´).
10 nm size into the MgB2 superconductor is expected to improve its vortex pinning. This is because the coherence length of MgB2 is of the order of 10 nm. The magnetic moment of these nanoparticles is also deemed useful for further improvement of superconducting properties of MgB2. However, the interaction between the magnetic nanoparticles can result in their agglomeration, preventing their homogenous incorporation into MgB2 matrix. Making Ni–Co–B nanoparticles of small enough size and without agglomeration has been a challenge. Electron d-band occupancy in Ni–Co–B and strong electronegativity may lead to applications of Ni–Co–B nanoparticles in organic chemistry synthesis as well as in pharmacy. The great advantage of Ni–Co–B is its low cost of starting materials, as compared to palladium [6–8],
http://dx.doi.org/10.1016/j.actamat.2014.02.040 1359-6454/Ó 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
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platinum [6–8] and rhodium [9] catalysts. An additional advantage is its magnetic properties, which can be manipulated through the size, shape and composition of nanoparticles [10]. This paper reports a new technique for the chemical synthesis of Ni–Co–B nanoparticles with a size of 5 nm, very narrow size distribution and little agglomeration. This technique relies on controlling the alkalinity of the reaction and the use of appropriate surfactant to prevent agglomeration. We used Ni(NO3)2 and Co(NO3)2 nitrate as the source of Ni and Co and NaBH4 as reducing agents and the source of B for Ni–Co–B nanoparticles [11]. Ethylenediamine was also added in this synthesis. The ethylenediamine kept the synthesis under alkaline pH condition, which is crucial for maintaining the reduction of metal salts and providing a high yield of the reaction. Ethylenediamine has catalytic properties, as well [12]. The reducing agent NaBH4 in neutral and acidic environments reduces hydrogen from water: NaBH4 ðsÞ þ 2H2 OðlÞ ! NaBO2 ðaqÞ þ 4H2 ðgÞ
ð1Þ
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SDS surfactant in aqueous solutions has a weak alkaline behaviour, and its surfactant ability is reduced in an acidic environment due to neutralization reaction between the dodecyl sulfate ion and H+ ions. The preparation procedure in which Ni2+ and Co2+ ions were dropped from a funnel into the three-necked flask with NaBH4 (aq) (Fig. 1) played an important role in providing the maximum yield of chemical reaction. Experimental work has shown less agglomeration of Ni–Co–B nanoparticles with described apparatus, probably due to less competition between Co2+ and Ni2+ ions. The emphasis of this report is on the materials preparation and chemistry of magnetic nanoparticles and their incorporation into MgB2, starting from the chemical synthesis of nanoparticles, solving problems of reaction between precursors and nanoparticles and finally measuring the change in properties due to those reactions. A detailed description of reaction mechanism and formation of new phases provides a crucial explanation of enhancement critical current density in a MgB2 superconductor.
Further, dissociated NaBH4 in water can reduce the hydronium ion to elemental hydrogen:
2. Experimental details
þ BH 4 ðaqÞ þ 4H ðaqÞ ! 2H2 ðgÞ þ BðOHÞ3 ðaqÞ
Ni–Co–B nanoparticles were prepared by the wet technique by reduction of cobalt and nickel salts with sodium boron hydride (NaBH4). Ni(NO3)2 and Co(NO3)2 were dissolved together (no reaction between them) in deionized (DI) water and placed in an additional funnel above the three-necked flask. On the other hand, NaBH4 was previously dissolved in DI water and placed into the threenecked flask. Ethylenediamine was added into the three-necked flask to maintain the reaction under alkaline pH conditions (pH 8–10) and also to catalyse the reaction. The
ð2Þ
All this would lead to low efficiency in producing Ni– Co–B nanoparticles and contamination of the resulting powder. In an alkaline pH environment, Reaction (2) does not occur, due to the negligible concentration of existing H+ ions. Ethylenediamine is soluble in water and reacts with water: H2 NCH2 CH2 NH2 ðlÞ þ H2 OðlÞ ! H2 NCH2 CH2 NH3 þ OH ðaqÞ
ð3Þ
On the other hand, excessive OH in the environment can be counterproductive due to another reaction between the metal salt (Co(NO3)2) and OH: Co2þ ðaqÞ þ 2OH ðaqÞ ! CoðOHÞ2 ðsÞ
ð4Þ
To provide maximum chemical reduction of metal salts, the (Co, Ni) reaction should be maintained in a medium alkaline environment and according to this experimental work, this is between pH = 8 and pH = 10. In such environment, the Reactions (1), (2) and (4) will not be significant and the reduction of metallic salts into metal borides will have the highest efficiency. Therefore, the adjustment such that 8 < pH < 10 directs the chemical reaction to reduce Ni2+ and Co2+ into Ni– Co–borides: NiðNO3 Þ2 þ CoðNO3 Þ2 þ NaBH4 þ OH ! Ni–Co–B þ NaNO3 þ H2 O
ð5Þ
Another important feature of the new synthesis is the addition of the surfactant sodium dodecyl sulfate (SDS), which prevents agglomeration of nanoparticles. Generally,
Fig. 1. Scheme of air-sensitive chemical equipment for producing nanoparticles.
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chemical reaction started with Co and Ni aqueous solution being added drop-wise into the three-necked flask with NBH4 solution. Immediately before the end of the dropwise introduction of Co and Ni nitrates from the funnel, SDS was added. Firstly, SDS was dissolved in DI water, separately bubbled for 2 h and finally added to the three-necked flask with a syringe. SDS has the role of surfactant, to prevent the agglomeration of nanoparticles. A surplus [1–4] of NaBH4 was added in order to completely reduce Ni2+ and Co2+ ions (the molar ratio of B/(Ni + Co) = 5). The synthesis was performed in a closed system under argon atmosphere, and all solutions were bubbled with argon for 2 h prior to the reaction. At the exhaust was set a chloride–calcium tube filled with CaCl2(s), to prevent entry of air and moisture retention (Fig. 1). The reaction was maintained at 273 K, and the duration of the reaction was less than 2 min. Black powder was formed instantly, and it was later extracted from the solution by centrifugation. The black powder was washed with DI water and ethanol several times to remove residual ions. Finally, the powder was dried for 2 h in a desiccator. X-ray powder diffraction (XRD) was performed at room temperature using an automatic Philips diffractometer, model PW1820 (Cu Ka radiation), in Bragg–Brentano geometry. XRD patterns for all samples were measured for samples at different stages of differential thermal analysis/ thermogravimetric analysis (DTA/TGA), after the heating of nanoparticles was stopped at 650, 770 and 850 °C. In mixing boron and magnesium with Ni–Co–B, it was important to ensure that any signature of the reaction between the nanoparticles and Mg is detectable in XRD analysis; for this reason 10 wt.% of Ni–Co–B was used. Powders were examined using a JEOL 2011 transmission electron microscope (TEM) operating at 200 kV. TEM specimens were made by dispersing the powder in ethanol and placing a drop of the suspension on a carbon support film on a copper grid. The microstructure of MgB2 samples were studied with a JEOL JSM-6400 scanning electron microscope (SEM). Magnetic measurements on nanoparticles were performed using a Quantum Design MPMS-5T SQUID magnetometer. Magnetic hysteresis loops M(H) were measured in the applied field within a range of 5 T at a few temperatures. Temperature dependence of magnetization M(T) for the temperature interval 10–300 K in a constant magnetic field H was measured in two modes: as zero field cooling (ZFC curves) and field cooling (FC curves). A ZFC magnetization curve was obtained by cooling the sample in an applied zero field from room temperature to 10 K, after which the required field was applied and the magnetization was measured during subsequent warming of the sample. The FC magnetization curve was obtained by measuring the magnetization as the sample was cooled down in the required field from room temperature to 10 K.
The memory effects of ZFC magnetization were tested by stopping the initial cool-down of the sample (10 K min1) at a set temperature (TS) for 104 s. After that, the sample was cooled down to 10 K, the magnetic field applied and magnetization measured as the sample was warmed up at a rate of 2 K min1. Thus obtained M(T) was compared with the ZFC measurements without stopping on initial cool-down. Any anomaly at TS shows a presence of strong interaction between the magnetic nanoparticles, creating a superspin glass system [13,14]. Magnetic measurements on MgB2 samples were performed using a Quantum Design PPMS-9T extraction magnetometer. The magnetization was measured at 5 K and 20 K in a time-varying magnetic field up to 9 T. DTA was performed by mixing obtained nanoparticles in different combinations with Mg, boron and Ni–Co–B (10 wt.%) pressed in small bars (200 mg) and finally measured by DTA/TGA (Mettler Toledo 1) up to 1100 °C. 2.1. Preparation of Ni–Co–B nanoparticles with boron and magnesium precursors Ni–Co–B nanoparticles described in this paper have been made with a view to improving the superconducting properties of MgB2. The nanoparticles described in this paper have a different size and size distribution than the Ni–Co–B nanoparticles described elsewhere. Because of that, they will have different chemical properties relevant to the growth of MgB2. Chemical reactions of these nanoparticles with the precursor powders for MgB2 were investigated using DTA and XRD techniques, in order to assess their suitability for this purpose. DTA results for the Ni– Co–B and boron mixture did not show any reaction between the boron and the Ni–Co–B (Fig. 2, inset). This result is consistent with the properties of elemental boron, which is quite inert and has a high melting point, 2300 °C. On the other hand, important information was observed on the DTA curve relating to the reaction of the Ni–Co–B nanoparticles and magnesium (Fig. 2). The small exothermal peak close to 500 °C (peak A) suggests a first reaction of the nanoparticles with magnesium. This small peak, which is accompanied by a broad plateau, can be described as a slow local solid–solid reaction between Ni–Co–B and Mg. The nanoparticles react with Mg and form a eutectic system [15,16] before Mg melts at 650 °C. Melted Mg can easily diffuse into and react with boron in Ni–Co–B [17]. Furthermore, with a size of less than 10 nm, the nanoparticles are extremely reactive, even at lower temperatures, than bulk Ni–Co–B, due to their large surface area. The melting point of magnesium can be clearly observed at 650 °C as the endothermic peak B in the DTA curve. With increasing temperature, two further distinct peaks appear, peak C at 700 °C and peak D at 770 °C (Fig. 2). Peak C represents the formation of the MgxBy phase [18], since 30 at.% of the atoms in Ni–Co–B are boron. The exothermal peak D at 770 °C represents the
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Fig. 2. DTA curve for mixed powders of magnesium and Ni–Co–B nanoparticles (10 wt.%). Inset: DTA curve for mixed powders of boron and Ni–Co–B nanoparticles (10 wt.%).
formation of the stable phases between Co–Mg and Ni–Mg [19]. The different shape of DTA curves (Figs. 2 and 3) is associated with the difference in the heat capacity of elements in experiments. The broad plateau in a sample of Ni–Co–B + Mg (Fig. 2) is due to the high molar capacity of magnesium (24.869 J mol1 K1), while in experimental sample with Mg + B + Ni–Co–B (Fig. 3), a substantial amount of magnesium is replaced with boron (molar ratio Mg:B = 1:2), which has considerably lower molar heat capacity (11.087 J mol1 K1). Fig. 3 shows the reactions between Mg, B and Ni–Co–B nanoparticles in the temperature range from 50 °C to 1000 °C. The DTA image clearly shows three distinguishable peaks marked with A, B, C. Small peaks at 500 °C can suggest the initial reaction between Ni–Co–B nanoparticles and magnesium in forming a eutectic liquid. The Fig. 4. XRD patterns: Mg + Ni–Co–B nanoparticles sintered at different temperatures (650 °C (black), 770 °C (red) and 850 °C (blue)). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
Fig. 3. DTA curve for mixed powders of Mg + 2B + Ni–Co–B nanoparticles (10 wt.%).
following exothermic peak A at 570 °C corresponds to the creation of MgB2 in solid–solid reaction. Ni–Co–B nanoparticles act as catalysts in MgB2 formation at a lower temperature of sintering. A similar mechanism has been reported by Zhao et al. [16], where nanoparticles form a local eutectic liquid with Mg and promote a faster local formation of MgB2 in the eutectic liquid phase. The melting point of magnesium is observed as the endothermic peak B at 650 °C. The large exothermal peak C above 730 °C corresponds to the completion of the reaction between Mg and B in forming MgB2 and the formation of new phases Mg2Ni and Mg2Co/Co2Mg [11,20–25]. The broad peak C in Fig. 3 is related to twin peaks C and D in Fig. 2.
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Fig. 5. SEM images for 2.5% Ni–Co–B doped MgB2, heat treated at 650 (a) and 850 °C (b).
In this work we have expanded X-ray diffraction analysis as compared to the previously reported study (see Fig. 4) [11]. Using the XRD technique, for Mg + 10 wt.% Ni–Co–B samples, we have analysed phase evolution as a function of different sintering temperatures, i.e. 650, 770 and 850 °C, respectively. All XRD patterns have three distinctive peaks at 32, 34 and 37°, which belong to pure magnesium. This is expected due to the large amount of magnesium as a starting material. Research into the phase diagrams revealed that Mg–Ni forms Mg2Ni as a stable phase [20–24], while the Mg–Co system is stable only as Co2Mg [20–25]. Although Mg2Co is not an equilibrium phase, trace amounts of this phase were identified in the XRD patterns (peaks at 13.4, 41.1 and 47.7°), which is further supported by similar observations for hydrogen storage materials [22–25]. During the heat treatment, small nanoparticles can react easily with magnesium, forming new phases such as Mg2Ni and Co2Mg/Mg2Co. However, such reactions can occur
only in localized areas with non-consistent ratio and composition. Therefore, we believe that Mg2Ni and Co2Mg are the most likely phases to be formed between 650 and 850 °C. As the heat treatment temperature changes, one can observe noticeable differences in the XRD patterns (Fig. 4). In particular, this applies to the Mg2Ni phase, which is much more prominent after sintering at higher temperatures (770 and 850 °C). Mg2Ni peaks are very difficult to be distinguished due to overlapping with other phases. Several peaks belonging to this phase were identified in the XRD patterns: 20, 21, 37, 40, 53, 57.5 and 72.5°. The intensity of these peaks increases with increasing sintering temperature. Cobalt based compounds (Co2Mg/ Mg2Co), on the other hand, can be identified in all of the XRD patterns (see Fig. 4), with a considerably lower intensity of the peaks in the 650 °C sintered sample. The peaks for Co2Mg are detected at 21.08, 22.45, 43.6 and 44.45°. However, all of these XRD features are quite faint, most likely because these new phases are in the form of nanoparticles, which have poor crystallinity and appear in very discrete localized reaction areas [11]. Boron from our Ni–Co–B nanoparticles could also react with magnesium and form Mg–B compounds such as MgB2 and MgB4. It is extremely difficult to prove their existence in the samples via XRD analysis, as the characteristic peaks for these compounds overlap with those for Mg and Co2Mg. However, if one assumes that such compounds are formed during the heat treatment process, the aforementioned MgB2 and MgB4 would be most likely. SEM images of doped MgB2 with 2.5 wt.% of Ni–Co–B are presented in Fig. 5a and b, for samples sintered at 650 and 850 °C, respectively. The sample heat treated at 650 °C has a large number of small-sized MgB2 crystallites. A considerable number of small crystals create more boundaries between grains, which usually have a negative impact on critical current density. On the other hand, the sample treated at 850 °C shows a larger grain size with flake-like crystallites. According to energy dispersive X-ray spectroscopy results [11], grain connectivity is improved with Ni–Co–B doping. Nanoparticles could have an impact through the hindrance of MgO formation on MgB2 grain boundaries, with a negative effect on critical current flow. In the formation of new Mg2Ni and Co2Mg phases, nickel and cobalt ions may compete with oxygen at higher temperature of sintering, preventing substantial formation of MgO. At the same time, newly formed Ni–Mg and Co–Mg phases provide good pinning centres for magnetic vortices. For samples sintered at 650 °C (Fig. 5a), distinguished spherical particles positioned at the grain boundaries are noticeable (yellow1 circles), and most likely belong to MgO or pure Mg. The sample sintered at 650 °C shows a significant amount of unreacted Mg and MgO in contrast to the
1 For interpretation of colour in Fig. 5, the reader is referred to the web version of this article.
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Fig. 6. XRD patterns of MgB2 samples doped with 2.5 wt.% of Ni–Co–B and sintered at 650 and 850 °C.
sample sintered at 850 °C, as can be seen in XRD patterns of MgB2 doped samples sintered at 650 and 850 °C (Fig. 6). The porosity shown in these figures is a common feature of MgB2. Because the density of MgB2 is higher than that of the precursor powders, the solid volume decreases upon the growth of MgB2, creating pores between the crystallites. Improvement in critical current density (JC) with the addition of Ni–Co–B nanoparticles can be seen in Fig. 7a and b. Heat treatment at high temperatures (850 °C) was required to obtain the improvement. Low temperature heat treatment (650 °C) resulted in deterioration of JC. JC was enhanced through better connectivity between the MgB2 grains, as well as through an increase of the vortex pinning [11]. Improvement in grain connectivity results in a higher affective cross-sectional area for the super-current flow through the sample. This shifts the JC(H) curves upward in Fig. 7a. On the other hand, increased vortex pinning results in a weaker field dependence of JC and tilting of the JC(H) curves, in addition to improved JC for H = 0. This tilting is not obvious in Fig. 7a, but it can be seen in Fig. 7b, showing the rate of decrease of ln(JC) with the field. The lower value of d[ln(JC)]/dH is indicative of improved vortex pinning. While the heat treatment at low temperatures did not improve the vortex pinning, the 850 °C heat treatment improved it, as shown in detail in Ref. [11].
Fig. 7. (a) Field dependence of JC for 2.5 wt.% Ni–Co–B doped and pure samples sintered at two different temperatures, 650 and 850 °C. (b) Magnetic field dependence of critical current density for pure MgB2 and 2.5 wt.% Ni–Co–B doped samples heat treated at 650 and 850 °C. (b) Rate of change of ln(JC) with field for samples heat treated at 850 °C, showing the improvement of vortex pinning with the addition of Ni–Co–B nanoparticles at 5 K.
2.2. Characterization of Ni–Co–B nanoparticles Fig. 8 shows XRD patterns for the as prepared Ni–Co– B nanoparticles. There are two clearly identifiable peaks, i.e. a diffused peak at 2h 35° and a relatively sharp peak at 2h 62°, in the diffraction pattern. This is in sharp contrast to the usually observed amorphous morphology for Ni–Co–B nanoparticles with a broad peak at 45° [1–5]. The presence of the peaks in our XRD pattern would indicate that Ni–Co–B nanoparticles in this case have some sort of nanocrystalline structure. Unfortunately, the
Fig. 8. XRD pattern of Ni–Co–B nanoparticles.
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Fig. 9. (a) Ni–Co–B nanoparticles precipitated in sheets with a high degree of agglomeration. HRTEM image of the Ni–Co–B phase showing it to be highly crystalline with a mean particle size of 6 nm (inset). FFT of the lattice spacings showing a limited number of spacings, suggesting a high crystal symmetry structure, such as cubic. (b) Mostly spherical shape of nanoparticles and size less than 5 nm.
observed peaks could not be attributed to a ternary Ni–Co–B compound, as there are no reference records available in the ICSD or ICDD databases. On the other hand, these peaks also do not belong to any of the known binary Ni–B or Co–B compounds, eliminating the possibility that these peaks originate from such impurities. It is quite difficult to explain such observations, but our hypothesis lies in the unique morphology of the nanoparticles, their size and surface layer. The appearance of the peaks might be correlated to the thickness of a nanoparticle’s surface layer. As the thickness of this surface layer becomes comparable to the diameter of the nanoparticle, a pseudo-superstructure can be formed, which results in the appearance of a diffraction peak. This is supported by the fact that XRD diffraction patterns are distinctly different when atomic ratios of Ni–Co–B nanoparticles and amounts of surfactants are changed and/or doped with other elements. It is critical to note, though, that in an alloy structure, such as that of Ni–Co–B, it is unusual and difficult to distinguish a particular crystalline phase. Further crystallographic investigations for this compound are necessary, in order to clarify observed crystal structure phenomenology. Crystallite size was also determined by the program XBroad for size-strain analysis [26]. The basis for XBroad analysis is the Warren–Averbach–Bertaut method, which is implemented in the program, and details can be found in Skoko et al. [26].
The analysis was done on the maximum located around 2h = 62°, which was firstly smoothed in order to obtain the best possible results. The results showed the area-weighted crystallite size (length of coherent diffraction columns) of 6 nm with negligible strain. Crystallite size obtained with XBroad is in strong agreement with the TEM analysis. The nanoparticulate phase precipitated in sheets in which the nanoparticles were highly agglomerated (Fig. 9a and b). Examination by high resolution TEM (HRTEM) confirmed that the particles were highly crystalline (Fig. 9a), with nearly all particles exhibiting strong lattice fringe contrast. Particles were typically rounded with no marked faceting. The mean particle diameter was estimated by measuring the long axis of 60 particles as 5.3 ± 1.9 nm. This dense agglomeration may help protect the particles from oxidation, since no surface membrane of amorphous oxide was found on any of the particles. The nanoscale size of the particles resulted in shape transform effects in diffraction, specifically highly diffuse diffraction peaks. XRD data showed weak and diffuse peaks, with only two faint peaks detected above background. Analysis of the spacings found in the HRTEM images (fast Fourier transform (FFT) Fig. 9a is shown as inset) revealed four diffraction spacings (0.215, 0.24, 0.29 and 0.46 nm). The limited number of spacings suggested that the phase had a high crystal symmetry, most likely cubic. The target stoichiometry of the phase based on mass balance was Ni1Co1B1. X-ray microanalysis lacked the sensitivity to detect B. However, electron energy loss spectroscopy confirmed the presence of B within the phase. Meaningful quantitative analysis was not possible with this technique. 2.3. Magnetic properties of Ni–Co–B nanoparticles Temperature dependence of the magnetization M(T) for Ni–Co–B nanoparticles measured under different applied magnetic fields is presented in Fig. 8. A maximum magnetization can be observed in the ZFC curve (Fig. 10) at
Fig. 10. ZFC and FC curves of Ni–Co–B nanoparticles.
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90 K, indicating that 90 K is the blocking temperature (TB) of Ni–Co–B nanoparticles [10,27]. Below TB, the particles are blocked and in the ferromagnetic state with an irreversible magnetization. Above TB, the nanoparticles are gradually unblocked and are allowed to reorient in the magnetic field. They are characterized by superparamagnetic behaviour. The sample magnetization is still irreversible up to 250 K, due to the interaction between some of the nanoparticles. This gives the irreversibility temperature (Tirr) as 250 K. Splitting between the ZFC and FC curves is a consequence of magnetic moment blocking in the studied systems of nanoparticles, which causes slow relaxation of magnetization of the system. The difference between Tirr and TB is connected to the size distribution of the particles and possibly the magnetic interaction between aggregated particles. This difference between TB and Tirr is quite large, which indicates the agglomeration of nanoparticles into clusters. The FC branch of magnetization is a continuously decreasing function of temperature. This is expected for superparamagnetic nanoparticles with weak or no interaction between them. Sasaki et al. [13] showed that the system of strongly interacting superparamagnetic nanoparticles (superspin glasses) will exhibit FC magnetization that is constant at low temperatures below TB, and it even decreases with cooling below TB. The ZFC branch of the magnetization is also expected to give characteristic memory effects for superspin glasses [13]. No memory effects of this kind [28] were observed for our Ni–Co–B nanoparticles. This shows that our samples consist of non-interacting, or weakly interacting, nanoparticles. The curved shape of the hysteresis loops with an applied magnetic field up to 2 T also indicates the superparamagnetic behaviour of the nanoparticles. Irreversible M(H) curves were obtained at low temperatures (Fig. 10). This
Fig. 11. Hysteresis loops of magnetic moment vs. field at different temperatures. Inset: expanded view of the low field part of the M(H) loops.
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irreversibility (i.e. magnetic hysteresis loop) is a consequence of slow relaxation of the magnetization of the system of blocked nanoparticles. At high temperatures (P100 K), M(H) curves are reversible (Fig. 11). The temperature dependence of the low field magnetization for magnetic nanoparticles depends on the type of the distribution function of the anisotropy energy barriers, which governs the relationship between the temperature of the maximum in MZFC(T), TMAX and the average blocking temperature, TB [10,29]. According to the Stoner–Wohlfarth model [30] magnetic hysteresis loops can be used to estimate the anisotropy energy density K = MH/2h, where M is close to saturation magnetization. If taking H HC at the same temperature, parameter h = 0.5 is used for the system of randomly oriented spherical magnetic particles. Magnetization saturation at 5 K can be estimated from magnetic moment (Fig. 11) and theoretical density of Ni–Co–B nanoparticles as 8743.2 kg m3 [31]. The value of saturation magnetization of Ni–Co–B is 349,720 A m1, and extrapolated coercive field HC is 79.60 A m1 at 5 K and 1000 Oe. Obtained anisotropy energy K, calculated from the Stoner–Wohlfarth model, is 34,955.4 J m3. From derived anisotropy energy using the Neel relation the relaxation time is s = s0 exp(DEB/kBTB), where kB is Boltzmann’s constant, T is the absolute temperature and s0 is an attempt time on the order of 109–1013 s. DEB = KV is the anisotropy energy barrier, where K is the effective anisotropy constant and V is the particle volume. The temperature at which this relaxation time s0 equals the measurement time sm and at which the particle system goes into the superparamagnetic region is the blocking temperature TB [32]. The volume of the particle can be calculated from V = 25kBTB/K, where the factor 25 is a typical ln(sm/s0) value for conventional magnetometry [33]. Obtained value of volume of Ni–Co–B nanoparticles calculated from Neel’s relation is 8.889 1025 m3, (diameter of particle 11.9 nm). On the other hand the volume of Ni–Co–B nanoparticles observed directly from TEM images (Fig. 9a and b) and calculated from V = p(D)3/6, (average diameter D 5 nm, Fig. 8b) is 8.71 1026 m3. In comparing the values of D obtained by different methods, a difference by a factor of two was obtained. The reasons for this are most likely the estimates used in obtaining the saturation magnetization and relaxation times. According to Zysler et al. [29] similar results of their M(H) measurements give evidence of a significant contribution to magnetization of uncompensated spins at nanoparticle surfaces. This is responsible for a non-saturated component in the M(H) curves. The results have been interpreted by a simple model, in which each single-domain nanoparticle is considered as a core–shell system, with uniaxial anisotropy of the core and surface anisotropy of the shell. The surface contribution is more evident in the absence of interparticle interactions.
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3. Conclusions
References
Due to being a promising candidate for doping of MgB2 superconductor, the properties of Ni–Co–B nanoparticles have been improved using a new synthesis technique. Further modification of the chemical treatment achieved nanoparticles with a narrowed distribution of 5 nm and nanocrystalline morphology. Furthermore, Ni–Co–B particles exhibit superparamagnetic behaviour, which can be well described within the model of relaxation of magnetic moments of particles. Magnetic Ni–Co–B nanoparticles added to Mg and B precursors in the formation of MgB2 react only with Mg, forming mostly paramagnetic Mg2Ni and MgCo2 phases. This process also promotes a slow solid–solid reaction (eutectic system) between Mg and B below the melting point of magnesium. After melting Mg at 650 °C, these reactions between the two main precursors, boron and magnesium, accelerate. Finally, above 750 °C, boron and Mg complete the formation of MgB2 and the incorporation of new stable phases Ni2Mg, Co2Mg/(Mg2Co) into the matrix of MgB2. Improvement of JC by Ni–Co–B doping is associated with the creation and incorporation of new particles into the matrix of MgB2, which enhances the grain connectivity as well as the pinning of magnetic vortices [11].
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Acknowledgments The lead author (M.M.) was a recipient of University of Wollongong PhD scholarship for the duration of the work presented in this paper. This work builds upon the outcomes of a project UKF 1B 01/07, which was supported through the Unity Through Knowledge Fund, designed by the Croatian government, represented by the Ministry of Science, Education and Sport. Financial support from the Australian Research Council through project LP120100175 is thankfully acknowledged. This work was partially supported by the Australian Research Council (Grant No. DE130101247) and a 2013 AIIM-CRG grant. The authors acknowledge the use of facilities and the assistance of Mr. Darren Attard at the University of Wollongong Electron Microscopy Centre. The help of Dr. Germanas Peleckis and Dr. Zong Qing Ma at ISEM is thankfully acknowledged.