Synthesis of ultra-thin nanobristles of Na-Mn-O compounds and their magnetic and structural properties

Ceramics International 42 (2016) 17059–17066

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Synthesis of ultra-thin nanobristles of Na-Mn-O compounds and their magnetic and structural properties E. Oz a, S. Demirel a, S. Altin a, A. Bayri a, S. Avci b,n a b

Physics Department, Inonu University, 44280 Malatya, Turkey Department of Engineering Physics, Istanbul Medeniyet University, 34700 Istanbul, Turkey

art ic l e i nf o

a b s t r a c t

Article history: Received 10 June 2016 Received in revised form 20 July 2016 Accepted 29 July 2016 Available online 30 July 2016

Boron substituted Na0.44MnO2 nanorods were synthesized via conventional solid state reaction technique. Optimum synthesis temperature of the nanorods is determined as 750 °C. As the temperature increases, the nanorods start merging each other. As the boron content increases, amount of nanorods decrease due to the emergence of impurity phases. 12.5% boron substitution (x¼ 0.25 in NaMn2  xBxO4 nominal composition) into Mn site striggers the formation of ultra-thin nanobristles on the surface of the nanorods. These nanobristles disappear completely in x¼ 0.75 sample. We attribute the formation of these nanobristles to the increase in the internal energy upon increase in the unit cell volume. & 2016 Elsevier Ltd and Techna Group S.r.l. All rights reserved.

Keywords: Nano-bristles Na0.44MnO2 Nano-rods

1. Introduction Nanomaterials are attractive since their physical properties differ from their bulk properties [1]. Various techniques have been used to synthesize nano-sized materials such as vapor-solid-liquid growth [2], chemical vapor deposition [3], molecular beam epitaxy [4], hydrothermal process [5], etc. Among these techniques, solid state synthesis of nano-materials is not common and it requires special conditions for uni-directional crystal growth. These conditions can be provided by a flux agent, partial melting of the main matrix [6] and/or the facets of the nuclei with different surface energies [7]. The solid state synthesis of the rods can be achieved by fast ion transfer from the main matrix to the nuclei. The main advantages of the solid state synthesis of nano-materials are low synthesis cost and easy process. Na0.44MnO2 nanorods were discovered by Sauvage et al. using stress induced splitting method [8]. The length and the width of the rods are reported as 5–10 mm and 500 nm, respectively. These nanorods are very promising for energy storage applications due to their large surface area, short diffusion distance and high crystallinity [9]. Na0.44MnO2 is an isostructure of Na4Mn9O18 and has an orthorhombic lattice structure with Pbam space group [10]. There are five different Mn sites in the lattice; one MnO5 square pyramid and four MnO6 octahedra. Two of these five Mn sites contain Mn3 þ ions (including the square pyramidal site) and the rest have Mn4 þ ions located in octahedral sites. The square n

Corresponding author. E-mail address: [email protected] (S. Avci).

http://dx.doi.org/10.1016/j.ceramint.2016.07.215 0272-8842/& 2016 Elsevier Ltd and Techna Group S.r.l. All rights reserved.

pyramids connect two double and one triple edge-linked octahedral chain forming large S-shaped tunnels which provide paths for Na-ion insertion/removal [11]. Previously, we reported the detailed growth mechanism of Na0.44MnO2 nanorods and stated that excess Na plays a crucial role in the formation of these nanorods [12,13]. The morphology of the nanorods are needle-like shape with lengths up to 500 mm and widths vary between 100 nm and 2 mm. So far Ti substitution in Mn sites has been reported in the literature [14]. It is possible to substitute up to 55% of Mn ions with Ti4 þ . This partial substitution suppresses the dissolution and the degradation of Na0.44MnO4 in electrolyte at elevated temperatures. [15]. Zhan et al. reported that Ti substitution breaks the charge ordering and significantly enhances the capacity retention of Na0.44MnO2 which is a negative electrode material in aqueous Naion batteries. The morphology of these Ti substituted samples are rod-like shape with a length of 1 mm and a diameter of 0.4–0.5 mm [11]. In this study, we investigate the synthesis of Boron substituted Na0.44MnO2 nanorods by solid state reaction method. The details of the synthesis can be found in Ref. [12]. We have studied the effects of boron substitution on structure, morphology and magnetic properties of the nanorods. We observed formation of ultrathin nanobristles on the nanorods with 12.5% Boron substitution. Above this substitution level, not only the ultra-thin nanobristles disappear but also impurity phases become dominant. Magnetic measurements show that the rods have paramagnetic behavior down to liquid helium temperature. Small hysteresis on the M-H curve for high B substitutions is attributed to the impurity phases.

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Table 1 Nanorod formation conditions and properties. Sintering temp. (°C) Width (nm) Length (lm) Number of rods (in 1 lm2) Comments

NaMn2O4

650

140–1600

1–24

8–12

700

80–500

2–16

12–14

750

45–1000

2–18

18–20

800

100–2000

2–35

9–12

850

200–2000

1–40

4–6

900

550–2000

6–40

1–2

90–1500

2–10

9–12

800

300–1500

5–40

1–2

850

200–4000

2–20

1–2

900

680–4000

50–200

1

750

750–3500

3–40

1

NaMn1.25B0.75O4 750

500–3300

12–30

1

NaMnBO4

300–1800

10–80

1

NaMn1.75B0.25O4 750

NaMn1.5B0.5O4

750

– – – – – – – – – – – – – – – – – – – – – – – – – – – – – – – – – – – – – – – – – – – – –

The nanorod formations just start at this temperature. The sample surface consist of the granular structure. Rods grow mostly on the cracked surfaces rather than on flat regions. Samples consist of  10% nanorods and  90% granular structure in per 100 mm2 surface area. Average nanorod width and length are  160 nm and  2 mm, respectively. Sample surface consists of  80% nanorods and  20% granular structure in per 100 mm2 surface area. The rods are observed on the cracked surfaces rather than the flat regions. Average nanorod width and length are  50 nm and  10 mm, respectively. Sample surface consist of 100% nanorods in per 100 mm2 surface area. Nanorods are thinner and longer than the ones sintered at 700 °C. Average nanorod width and length are  90 nm and  13 mm, respectively. Sample surface consists of 100% nanorods in per 100 mm2 surface area Nanorods look like merging each other and are thicker than the ones sintered at 750 °C. Average nanorod width and length are  180 nm and  12 mm, respectively. Sample surface consists of 100% nanorods in per 100 mm2 surface area Nanorods merge each other. Average nanorod width and length are  320 nm and  10 mm, respectively. Sample surface consists of 95% nanorods and 5% granular structure in per 100 mm2 surface area. The thickest nanorods grow at this temperature. The nanorods start melting. Average nanorods width and length are  800 nm and  8 mm, respectively. Sample surface consists of 75% nanorods and 25% granular structure in per 100 mm2 surface area There are so many ultra-thin nanobristles on the nanorods. Nanorods on the cracked surfaces are thinner and longer than the ones grow on flat regions. Average nanorod width and length are 150 nm and 10 mm, respectively. Sample surface consists of 100% nanorods in per 100 mm2 surface area Nanorods are thicker than the ones sintered at 750 °C. Nanorods look like merging with the bristles. Average nanorod width and length are  300 nm and  10 mm, respectively. Sample surface consists of 85% nanorods and 15% granular structure in per 100 mm2 surface area. At this temperature, nanorods are thicker than the ones at lower temperatures and look like merging each other. Average nanorod width and length are  650 nm and  8 mm, respectively. Sample surface consists of 40% nanorods and 60% granular structure in per 100 mm2 surface area. The longest and the thickest nanorods grow at this temperature and they look like merging each other. Average nanorod width and length are  1400 nm and  110 mm, respectively. Sample surface consists of 30% nanorods and 70% granular structure in per 100 mm2 surface area. At this composition, the nanorods look like melting. Small amount of nanobristles are observed. Average nanorod width and length are  800 nm and  30 mm, respectively. Number of nanorods decrease significantly compared to the lower B substitution. Sample surface consist of 40% nanorods and 60% granular structure in per 100 mm2 surface area. Nanobristles disappear completely with this substitution level. Average nanorod width and length are  1100 nm and  22 mm, respectively. At this doping level the nanorods are rarely observed. Samples consist of 10% nanorods and 90% granular in per 100 mm2 surface area Average nanorod width and length are  800 nm and  45 mm, respectively.

E. Oz et al. / Ceramics International 42 (2016) 17059–17066

Composition

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diameter 4 times smaller than that of Titanium substituted composition. More importantly the emergence of ultra-thin nanobristles are not reported for Titanium substituted nanorods [11]. 2.2. Characterization of the samples

Fig. 1. XRD patterns of NaMn2  xBxO4 (x ¼0, 0.25, 0.5, 0.75 and 1). Blue and black tick marks represent the Na0.44MnO2 and Mn2O3 phases, respectively. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

The micro-structural characterization and phase analysis of the materials were investigated by X-ray diffractometer. The scan speed was selected as 2° min  1 in the range of 2θ ¼ 3–80°. The automated Rigaku RadBDmax x-ray difractometer with CuKα radiation was used for the XRD analysis. The lattice parameters were calculated by Rietveld refinement technique. The strain analysis and average crystallite size of the samples were calculated using XRD data. Details of these calculation techniques can be found in the Ref. [17]. The microstructural and compositional characterization of the samples were performed with Leo EVO-40 VPX scanning electron microscope (SEM) and Bruker X-flash detector 4010 energy dispersive x-ray spectroscope. Dc-susceptibility measurements between 300 and 2 K under magnetic field H¼5 kOe were measured using Quantum Design PPMS system with VSM attachment.

2. Experimental section In this study, we use NaMn2  xBxO4 as the nominal compositions. In the rest of the paper x values denote the B content in the nominal compositions. We synthesized x ¼0, 0.25, 0.5, 0.75 and 1 samples via solid state reaction technique. High purity powders of Na2O2, Mn2O3 and B2O3 were weighed in the appropriate amounts and mixed in an Ar-filled glovebox. After sintering, the powder mixture was pressed into pellets under pressure of 5 t. To determine the optimum synthesis temperature of the nanorods, the pellets were heated between 700 and 950 °C for 16 h under oxygen atmosphere. 2.1. Synthesis of nano-rods Na0.44MnO2 structure can be obtained in the axially oriented form which are called rod in the literature [11,16]. To find the optimum synthesis temperature, we performed systematic synthesis studies. The synthesis conditions and observed features of the nanorods are listed in Table 1. We determined the optimum synthesis temperature as 750 °C based on the number, length and width of the nanorods. Detailed growth mechanism of Na0.44MnO2 can also be found in Ref. [12]. According to Table 1, optimum synthesis temperature for NaMn1.75B0.25O2 is also 750 °C. These nanorods have widths between 90–1500 nm. Some of them have

3. Result and discussion 3.1. Structural properties Laboratory x-ray diffraction (XRD) patterns of the B substituted samples after sintering 750 °C are shown in Fig. 1. XRD data are collected on the mixture of the main matrix and the nanorods, by grinding them all together. All the diffraction peaks in the XRD pattern of the parent compound are indexed by Na0.44MnO2 phase, which has orthorhombic symmetry, without any detectable impurity phases. We intended to substitute Mn with boron however even with 12.5% substitution (x ¼0.25 in the NaMn2  xBxO4 composition), impurity phases start appearing. We identified these phases as Mn2O3 and Na(BO2). With 50% B substitution (NaMnBO4), almost all the Na0.44MnO2 phase disappears and Mn2O3 phase becomes the dominant phase. Evolution of the phase fractions and Na content (obtained from Rietveld Refinements) as a function of B substitution are shown in Fig. 2. Na0.44MnO2 phase remains as the main phase up to x ¼0.5 sample. Up to this point calculated Na content also does not change significantly. On the other hand, unit cell volume and the c axis of the Na0.44MnO2 phase increase monotonically with increasing boron content for xr 0.5 and decrease for higher B content. (Table 2). One would naively expect the unit cell volume to shrink due to the small ionic

Fig. 2. Variation of phase fractions (left panel) of Na0.44MnO2, Mn2O3 and Na(BO2) phases as a function of B substitution. Na content (right panel) of the samples upon B substitution. Solid lines are guide to the eye. Error bars are smaller than the symbols.

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Table 2 NaMn2  xBxO4 cell parameters. Doping quantity

x¼0 x ¼ 0.25 x ¼ 0.5 x ¼ 0.75

Volume (Å3)

Cell parameters (Å) a

b

c

9.08908 73.58  10  4 9.11328 74.35  10  4 9.12054 7 4.06  10  4 9.10254 75.29  10  4

26.48695 7 8.24  10  4 26.4995 7 0.0012 26.481137 0.00107 26.424427 0.00147

2.82628 71.58  10  4 2.830617 1.35  10  4 2.833657 1.34  10  4 2.82733 71.73  10  4

674.55 676.95 678.03 678.24

determine the exact B content that go into octahedral Mn3 þ sites. The crystallite size and lattice strain due to imperfections, impurities and dislocations can be found from the XRD peak broadening [19,20]. The average particle size of the samples is calculated by using Scherrer equation:

D = kλ /βc cosθ

(1)

where D is the average particle size, λ is the wavelength of the x-ray ( λCuKα =1.54056 Å ), k is the Scherrer constant which is usually taken as 0.9 [21], and βc is the full width at half maximum (FWHM) of the reflection peaks located at angle θ. βc consists of instrument broadening and the width of the peaks of the sample. The instrument broadening can be calculated from polycrystalline Si standard and the width of the peaks are calculated using the equation, βc =

Fig. 3. Particle size of Na0.44Mn1  xBxO2 (x¼ 0, 0.25, 0.5, 0.75 and 1). Solid line is guide to the eye.

β2exp − β2inst [22]. The average sizes of the particles

are determined as in the range of 22–38 nm as seen in the Fig. 3 and the particle size change with boron doping shows a parabolic behavior with increasing boron content. Williamson Hall method [23] can be used for strain analysis which uses the diffraction line broadening due to crystallite size and strain. The peak broadening on the XRD can be given:

βhkl = β + βε

(2)

where β is the crystallite size contribution and βε is the strain induced broadening [24]. βhkl represents the FWHM of the XRD peaks. The crystallite size contribution on βhkl is calculated by Scherrer equation as in Eq. (1). The strain contribution on βhkl is calculated by:

βhkl =4βε tanθ

(3)

where βε is the micro-strain. Line broadening on XRD peak is a combination of crystallite size and strain. The combined equation has a more simplified form as below:

βhkl cosθhkl =

Fig. 4. Strain of Na0.44Mn1  xBxO2 (x¼ 0, 0.25, 0.5, 0.75 and 1). Solid line is guide to the eye.

size of B3 þ (41 pm) compared to that of Mn3 þ (72 pm). This increase in the volume and the c axis shows that some of the B3 þ ions occupy the interstitial sites. In the case of LiMn2O4 we were able to substitute Mn with B up to 25% without any impurity phases [18]. In the case of Na0.44MnO2, as mentioned above, there are five Mn sites and two of them have Mn3 þ , one octahedral with 6 coordination and one square pyramidal with five coordination. Considering the valance state of B3 þ , we expect B3 þ replace the Mn3 þ ions only in the octahedral site, since B3 þ cannot have five coordination. Thus, theoretically only 22% of the Mn ions can be substituted by B3 þ ions. Boron is not visible to x-rays, therefore we could not

kλ +4βε sinθ D

(4)

where D is the crystallite size, k is a constant which equal to 0.9 [25]. The value of βhklcosθhkl as a function of 4sinθ is plotted and βε is determined from the slope of the curve. Both strain and average particle size show a similar behavior with unit cell volume upon substitution. They both decrease with boron substitution up to x¼ 0.5 sample and then increase for higher boron contents (Fig. 4). 3.2. Morphology of the nanorods SEM images in Fig. 5 shows effects of boron substitution on the morphology of Na0.44MnO2 nanorods. B substitution for x r0.5 does not suppress the nanorod growth. As the B content increases further, number of the nanorods decrease due to increase of impurity phases. Details of the systematic synthesis are in Table 1. The main feature of the morphology of the B substituted Na0.44MnO2 nanorods is the ultra-thin nano birstles that grow on the surface of nanorods. These nanobristles are observed densely

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Fig. 5. SEM image of Na0.44Mn1  xBxO2 (x¼ (a) 0, (b) 0.25, (c) 0.5, (d) 0.75 and (e) 1).

in 12.5% B substituted sample (NaMn1.75B0.25O4). Increasing boron content negatively affects the formation of these nanobristles. SEM images in Fig. 5 show very little amount of these bristles in x ¼0.5 sample and they completely disappear in x ¼0.75 and 1 samples. Fig. 6 shows the SEM images of x ¼0.25 sample (NaMn1.75B0.25O4) sintered at various temperatures. We observed that optimum temperature for the formation of these nanobristles is 750 °C. Sintering the samples at higher temperature prevents the nano bristle growth. One dimensional crystal growth needs higher energy than that of the three dimensional grain growth [26]. The Gibbs free energy change of the crystals can be given as [27];

G = Gv + Gs

(5)

where Gv is the volume free energy term and Gs is the surface energy term. One dimensional growth of the crystals can increase

the free energy of the system and there are several models regarding the explanation of the surface energy term of the nanorods [28]. In our case, XRD data show that lattice volume increases with increasing boron content up x¼ 0.5 sample. Addition of boron into the interstitial regions can increase the volume energy term of the crystals which can cause the formation of new phases [29] such as nanobristles as observed in our samples. These nanobristles with their increased surface area and decreased diffusion lengths can be a good alternative for Na-based energy storage applications. 3.3. Magnetic properties As mentioned above there are five Mn sites, two of them are Mn3 þ (one octahedral and one square pyramidal) and the rest are octahedral Mn4 þ . B3 þ ions are expected to substitute octahedral

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Fig. 6. SEM images of NaMn1.75B0.25O4 sample sintered at (a) 750 °C (b) 800 °C (c) 850 °C (d) 900 °C.

Fig. 7. The spin state configurations of different environment of Mn ions.

Mn3 þ ions. 3d orbitals in an octahedral environment has five orbitals, t2g triplet and eg doublet with є(t2g)o є(eg). Since all Mn4 þ ions accommodate in an octahedral environment with three electrons, t2g gets the magnetic moments of these ions due to three unpaired electrons. On the other hand, Mn3 þ in the same environment may have either low spin (t2g4eg0) or high spin (t2g3eg1) configuration due to the energy differences of these two levels. 3d orbitals in a square pyramidal environment splits four different sets one with doubly degenerate (Fig. 7). Hence one may expect low (LS), intermediate (IS) and high spin (HS) configurations for Mn3 þ in this structure. However, due to the very small energy differences between these splitting levels (which is always the case for a square pyramidal environment) the expected magnetic moment would be the sum of four unpaired electrons without orbital contribution. The effective magnetic moment can be calculated from μeff ¼ g(S(Sþ 1))0.5μB where μB is the Bohr magneton, g is the spectroscopic splitting factor and S is the total spin quantum number. The magnetic measurements are performed under field cooling

mode under 5 kOe applied magnetic field. Fig. 8 shows the temperature dependence of magnetic susceptibility of the samples. χ values monotonically increase with decreasing temperature which corresponds to Curie–Weiss paramagnetic behavior. Magnetic susceptibility data of the parent compound Na0.44MnO2 in Fig. 8 show that the antiferromagnetic ordering starts around 25 K which is consistent with the literature [30]. The sample has Curie–Weiss type paramagnetism in the interval of 100–300 K. So the data in this interval are fitted to find the θ and C values from Curie-Weiss formula and χ  1-T. The θ value is found as   80.21 71.2 K. The large negative value of θ constant indicates strong antiferromagnetic interactions between Mn sites [30]. The effective magnetic moment is calculated from the fit and found as μeff ¼3.4μB per Mn atom. The obtained μeff values from Curie-Weiss fitting show an increase by increasing boron content as seen from the Fig. 9. The changes and the anomalies in the susceptibility and the increase in the μeff values upon B substitution are attributed to the impurity phases (α-Mn2O3 and Na(BO2)) that appear in XRD. Thus, the susceptibility data in Fig. 8 are the

E. Oz et al. / Ceramics International 42 (2016) 17059–17066

Fig. 8. χ-T of Na0.44Mn1  xBxO2 (x ¼0, 0.25, 0.5, 0.75, and 1). Solid lines are Curie Weiss fitting.

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80 K [31]. The Mn3 þ ions in α-Mn2O3 have two different electronic configurations in the octahedral environment as given in the Fig. 7. Fig. 8 shows that Neel temperature of α-Mn2O3 is quite high compared to other antiferromagnetic Mn alloys (  80 K) [32]. If the interacting magnetic centers are increased, one may expect a high Neel temperature assuming that Mn3 þ ions in this phase have HS electronic configuration. Cockayne et al. confirms that in α-Mn2O3, Mn3 þ ions are likely in high-spin configuration [33]. In our case, the increase of the effective magnetic moment by increasing boron content can be explained that the content of αMn2O3 and Na(BO2)2 phases increase by decomposition of Na0.44MnO2 phase. χ-T graph of x ¼1 sample shows a step-like feature at low temperatures. To understand this anomaly, we should focus on the different phases in this particular sample. Na0.44MnO2 phase shows paramagnetic behavior down to 25 K and the antiferromagnetic transition is observed at 25 K. α-Mn2O3show similar behavior with Na0.44MnO2 and the antiferromagnetic transition temperature is  80 K. Although, both of these two phases show no spin state transitions from LS to HS or LS to IS in their pure forms [13,33], boron substitution may induce such transitions. However, in our data, the susceptibility decreases with increasing temperature indicating that step-like features are not due to a spin state transition. The most probable explanation of these anomalies in x ¼ 1 sample is the modifications in magnetic and/or atomic structures of unidentified impurity phases. To support this idea, we measured M-H of the samples at 5 K and observed small hysteresis (Fig. 10) for x ¼1 with increasing magnetization.

4. Conclusion

Fig. 9. μeff of Na0.44Mn1  xBxO2 (x¼ 0, 0.25, 0.5, 0.75 and 1). Solid line is guide to the eye.

Boron substituted Na0.44MnO2 nanorods were synthesized via conventional solid state technique. NaMn2  xBxO4 are used as the nominal compositions. Nanorods densely grow in xr 0.5 samples, and as the Boron content increases further, number of nanorods decrease. Particle size and strain decrease and unit cell volume and c axis increase monotonically for x r0.5 samples indicating that some of the B ions may settle in the interstitial sites in the Na0.44MnO2 phase. 12.5% B substitution induces the formation of ultra-thin nanobristles on the surface of the nanorods. These nanobristles disappear completely for x4 0.5 samples. We attribute the formation of these nanobristles to the increase in the free energy of the nanorods due to increase in the unit cell volume upon B substitution. These nanobristles are promising for rechargeable Na-ion battery applications due to their increased surface area and decreased diffusion lengths. The μeff values show an increase by increasing boron content due to the emergence of impurity phases. As a conclusion B substitution affects the morphology of Na0.44MnO2 nanorods by inducing the growth of ultrathin nano birstles. At the same time, magnetic properties of Na0.44MnO2 can be controlled via B substitution.

Acknowledgment

Fig. 10. M-H of Na0.44Mn1  xBxO2 (x¼ 0, 0.25, 0.5, 0.75 and 1).

sum of magnetization of all these impurity phases. The magnetic properties of α-Mn2O3 were investigated by Meisenheimer and Cook that the susceptibility shows Curie-Weiss type behavior down to 80 K and the antiferromagnetic ordering starts below

This study was supported by TUBITAK (The Scientific and Technological Research Council of Turkey) under Grant No TUBITAK 112M487.

References [1] S. Suresh, Semiconductor nanomaterials, methods and applications: a review,

17066

E. Oz et al. / Ceramics International 42 (2016) 17059–17066

Nanosci. Nanotechnol. 3 (2013) 62–74. [2] K.W. Kolasinski, Catalytic growth of nanowires: vapor–liquid–solid, vapor– solid–solid, solution–liquid–solid and solid–liquid–solid growth, Curr. Opin. Solid State Mater. Sci. 10 (2006) 182–191. [3] N.M. Hwang, W.S. Cheong, D.Y. Yoon, D.-Y. Kim, Growth of silicon nanowires by chemical vapor deposition: approach by charged cluster model, J. Cryst. Growth 218 (2000) 33–39. [4] J. He, K. Yadavalli, Z. Zhao, N. Li, Z. Hao, K.L. Wang, A.P. Jacob, InAs/GaAs nanostructures grown on patterned Si(001) by molecular beam epitaxy, Nanotechnology 19 (2008) 455607–455613. [5] J. Yao, L. Lv, C. Shen, P. Zhang, K.-F. Agusey-Zinsou, L. Wang, Nano-sized spinel LiMn2O4 powder fabricated via modified dynamic hydrothermal synthesis, Ceram. Int. 39 (2013) 3359–3364. [6] V. Palomares, M. Casas-Cabanas, E. Castillo-Martinez, M.H. Han, T. Rojo, Update on Na-based battery materials. A growing research path, Energy Environ. Sci. 6 (2013) 2312–2337. [7] Y.S. Park, S.-H. Lee, J.-E. Oh, C.-M. Park, T.-W. Kang, Self-assembled GaN nanorods grown directly on (1 1 1) Si substrates: dependence on growth conditions, J. Cryst. Growth 282 (2005) 313–319. [8] F. Sauvage, L. Laffont, J.-M. Tarascon, E. Baudrin, Study of the insertion/deinsertion mechanism of sodium into Na0.44MnO2, Inorg. Chem. 46 (2007) 3289–3294. [9] F. Cheng, H. Wang, Z. Zhu, Y. Wang, T. Zhang, Z. Tao, J. Chen, Porous LiMn2O4 nanorods with durable high-rate capability for rechargeable Li-ion batteries, RSC Energy Environ. Sci. 4 (2011) 3668–3675. [10] M. Xu, Y. Niu, C. Chen, J. Song, S. Bao, C.M. Li, Synthesis and application of ultra-long Na0.44MnO2 submicron slabs as a cathode material for Na-ion batteries, RSC Adv. 4 (2014) 38140–38143. [11] P. Zhan, K. Jiao, J. Wang, Z. Hu, R. Ma, H. Zhu, S. Jiao, Titanium-substituted Na0.44MnO2 nanorods as cathode materials for high performance sodium-ion batteries, J. Electrochem. Soc. 162 (2015) A2296–A2301. [12] S. Demirel, E. Oz, E. Altin, S. Altin, A. Bayri, P. Kaya, S. Turan, S. Avci, Growth mechanism and magnetic and electrochemical properties of Na0.44MnO2 nanorods as cathode material for Na-ion batteries, Mater. Charact. 105 (2015) 104–112. [13] Y. Wang, J. Liu, B. Lee, R. Qiao, Z. Yang, S. Xu, X. Yu, L. Gu, Y.-S. Hu, W. Yang, K. Kang, H. Li, X.-Q. Yang, L. Chen, X. Huang, Ti-substituted tunnel-type Na0.44MnO2 oxide as a negative electrode for aqueous sodium-ion batteries, Nat. Commun. 6 (2015) 6401–6411. [14] J.A. Saint, M.M. Doeff, J. Wilcox, Electrode materials with the Na0.44MnO2 structure: effect of titanium substitution on physical and electrochemical properties, Chem. Mater. 20 (2008) 3404–3411. [15] Y.J. Park, M.M. Doeff, Effect of structure on the storage characteristics of manganese oxide electrode materials, J. Power Sources 165 (2007) 573–580. [16] C. Liu, J. Li, P. Zhao, W. Guo, X. Yang, Fast preparation of Na0.44MnO2 nanorods via a high NaOH concentration hydrothermal soft chemical reaction and their lithium storage properties, J. Nanopart. Res. 17 (2015) 142–150. [17] K.A. Aly, N.M. Khalil, Y. Algamal, Q.M.A. Saleem, Lattice strain estimation for

[18]

[19]

[20]

[21] [22]

[23]

[24]

[25]

[26]

[27] [28] [29]

[30] [31] [32]

[33]

CoAl2O4 nano particles using Williamson-Hall analysis, J. Alloy. Compd. 676 (2016) 606–612. S. Demirel, E. Oz, S. Altin, A. Bayri, E. Altin, S. Avci, Enhancement of battery performance of LiMn2O4: correlations between electrochemical and magnetic properties, RSC Adv. 6 (2016) 43823–43831. M.A. Amer, T. Meaz, M. Yehia, S.S. Attalah, F. Fakhry, Characterization, structural and magnetic properties of the as-prepared Mg-substituted Cu-nanoferrites, J. Alloy. Compd. 633 (2015) 448–455. S.N. Anitha, I. Jayakumari, Synthesis and analysis of nanocrystalline Fe2Mn2Ni0.5Zn1.5O9 at different treating temperatures, J. Nanosci. Nanotechnol. 1 (2015) 26–31. S.G. Elsharkawy, R. Awad, Thermal expansion measurements of (Cu0.25Tl0.75)1234 added by MgO-nanoparticles, J. Alloy. Compd. 478 (2009) 642–647. K. Omri, O.M. Lemine, J. El Ghoul, L. El Mir, Sol–gel synthesis and room temperature ferromagnetism in Mn doped ZnO nanocrystals, J. Mater. Sci: Mater. Electron. 26 (2015) 5930–5936. G. Rajender, P.K. Giri, Strain induced phase formation, microstructural evolution and bandgap narrowing in strained TiO2 nanocrystals grown by ball milling, J. Alloy. Compd. 676 (2016) 591–600. S.A. Kahani, F. Mashhadian, Preparation of bimetallic CoeAg and CoeCu nanoparticles by transmetallation of tetrakis(pyridine)silver(II) peroxydisulfate and tetrakis(pyridine)sulfatocopper(II) monohydrate complexes, J. Alloy. Compd. 660 (2016) 310–315. T.S. Shyju, S. Anandhi, R. Indirajith, R. Gopalakrishnan, Effects of annealing on cadmium selenide nanocrystalline thin films prepared by chemical bath deposition, J. Alloy. Compd. 506 (2010) 892–897. M. Lucas, Z.L. Wang, E. Riedo, Growth direction and morphology of ZnO nanobelts revealed by combining in situ atomic force microscopy and polarized Raman spectroscopy, Phys. Rev. B 81 (2010) 045415–045420. B.S. Mitchell, An Introduction to Materials Engineering and Science for Chemical and Materials Engineers, Wiley, New Jersey, 2004. X. Wang, Z. Li, J. Shi, Y. Yu, One-dimensional titanium dioxide nanomaterials: nanowires, nanorods, and nanobelts, RSC Chem. Rev. 114 (2014) 9346–9384. K. Jacobs, D. Zaziski, E.C. Scher, A.B. Herhold, A. Paul Alivisatos, Activation volumes for solid-solid transformations in nanocrystals, Science 293 (2001) 1803–1806. Q. Chu, X. Wang, B. Li, H. Jin, X. Cao, X. Zhao, X. Liu, Flux synthesis and growth mechanism of Na0.5MnO2 whiskers, J. Cryst. Growth 322 (2011) 103–108. R.G. Meisenheimer, D.L. Cook, Magnetic susceptibility of manganese sesquioxide at temperatures from 4 to 700 °K, J. Chem. Phys. 30 (1959) 605. V.M.T.S. Barthem, C.V. Colin, H. Mayaffre, M.-H. Julien, D. Givord, Revealing the properties of Mn2Au for antiferromagnetic spintronics, Nature Communication. doi: http://dx.doi.org/10.1038/ncomms3892http://dx.doi.org/10.1038/ ncomms3892. E. Cockayne, I. Levin, H. Wu, A. Llobet, Magnetic structure of bixbyite αMn2O3: a combined DFTþ U and neutron diffraction study, Phys. Rev. B 87 (2013) 184413–184424.