Magnetic and optical properties in degenerated transition metal and Ga co-substituted ZnO nanocrystals

Journal of Alloys and Compounds 805 (2019) 1191e1199

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Magnetic and optical properties in degenerated transition metal and Ga co-substituted ZnO nanocrystals €a €mbre b, Mati Kook b, Vambola Kisand b, Anzelms Zukuls a, Raivis Egl
ıtis a, Tanel Ka lix Duarte e, Martin Ja €rvekülg b, Mikhael Maiorov c, Reinis Ignatans d, Roberto Fe a, b, *  Andris Sutka a

Research Laboratory of Functional Materials Technologies, Faculty of Materials Science and Applied Chemistry, Riga Technical University, Paula Valdena 3/ 7, 1048, Riga, Latvia Institute of Physics, University of Tartu, Wilhelm Ostwald Str. 1, 50411, Tartu, Estonia c Institute of Physics, University of Latvia, Miera 32, 2169, Salaspils, Latvia d  Institute of Materials, Ecole Polytechnique F ed erale de Lausanne, CH-1015, Lausanne, Switzerland e Helmholz-Zentrum Berlin für Materialien und Energie GmbH, Berlin, 14109, Germany b

a r t i c l e i n f o

a b s t r a c t

Article history: Received 23 April 2019 Received in revised form 9 July 2019 Accepted 17 July 2019 Available online 18 July 2019

In order to study the influence of itinerant electrons on magnetic properties of transition metal substituted ZnO nanocrystals, nanopowders containing different amounts of Ga and fixed amounts of Fe, Ni and Mn ions were synthesized. The ions of different transition metals and Ga were successfully introduced into the ZnO structure using solvothermal synthesis method. X-ray diffraction, scanning electron microscopy, hard X-ray photoelectron spectroscopy and Rietveld refinement were used to characterize the synthesized nanocrystals. Optical measurements revealed that Ga substitution can change the light transmittance/absorption in the infrared part of the electromagnetic light spectrum due to itinerant electrons in the nanocrystals, as well as influence the magnetic properties of the obtained nanocrystals. © 2019 Elsevier B.V. All rights reserved.

Keywords: Diluted magnetic semiconductors Substituted ZnO Nanoparticles Solvothermal synthesis

1. Introduction Transition metal (TM) substituted ZnO is a widely studied diluted magnetic semiconductor (DMS) in which carriersubstituting cation magnetic exchange interactions can be manipulated. DMS materials have been sought after in spin-based information technologies for higher data storage and transfer efficiency [1]. Ferromagnetic behaviour at room temperature has been observed for Ni [2], Co [3], Fe [4e6], Mn [7], Cr [8] and Cu [9] substituted ZnO. In fact, ferromagnetism has been observed also in pure ZnO nanopowders. This has been interpreted as (Stoner) ferromagnetism arising from the presence of itinerant charge due to oxygen vacancies [10]. However, the magnetisation values are several orders of magnitude smaller than that of transition metal substituted ZnO [3] and too low for practical applications [2].

* Corresponding author. Research Laboratory of Functional Materials Technologies, Faculty of Materials Science and Applied Chemistry, Riga Technical University, Paula Valdena 3/7, 1048, Riga, Latvia.  E-mail address: [email protected] (A. Sutka). https://doi.org/10.1016/j.jallcom.2019.07.197 0925-8388/© 2019 Elsevier B.V. All rights reserved.

The magnitude of ferromagnetism in transition metal substituted ZnO is strongly dependent on point defects i.e. Zn interstitials or oxygen vacancies [11]. To maintain electric neutrality of the crystalline lattice, defects are compensated by itinerant charge carriers which in turn drive magnetic ordering [11]. It has been shown that Co substituted ZnO without point defects is not ferromagnetic [12], although ferromagnetism in Co substituted ZnO is widely reported [3,13]. In order to increase the charge carrier concentration, various DMSs have been co-substituted with aliovalent donor cations. This idea has been realised in Fe-Sn [14e18] and MneSn [19,20] cosubstituted In2O3 nanocrystals. However, although degenerated plasmonic ZnO nanocrystals have been widely reported where aliovalent donor substituting cations such as Al [1,21,22], In Refs. [21,23] and Ga [3,23] have been used, transition metal and donor co-substituted ZnO DMS materials have been poorly investigated. In a previous work we have co-substituted ZnO with Co (transition metal) and Ga [3] and demonstrated that by co-substituting ZnO with Ga and Co two-fold higher magnetisation values were

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observed in comparison with Co substituted ZnO at the same Co concentration. Gallium ion introduction increases the density of itinerant electron, which in turn may interact with Co magnetic spins and contribute to room temperature ferromagnetism. In this work we intend to synthesize Zn0.95-x(Me0.05Gax)O where Me ¼ Fe2þ, Ni2þ, Mn2þ to further contribute to the exploration of transition metal and Ga co-substituted ZnO DMS materials.

absorbance and transmittance spectra were measured by a UVevis spectrophotometer (UV-3700, Shimadzu Scientific Instruments Kyoto, Japan) for particle suspensions containing 0.1 mg/ml nanocrystals in tetrachloroethylene. Nanocrystals were stabilized using a mixture of oleylamine and 4-dodecylbenzenesulfonic acid (0.01 mg/ml) in 1:1 ratio. 3. Results and discussion

2. Experimental section 2.1. Synthesis of substituted zinc oxide nanoparticles The solvothermal synthesis method was used to produce different Zn0.95-x(Me0.05Gax)O (Me ¼ Fe, Ni and Mn) nanocrystals with different Ga concentration (x ¼ 0, 0.025, 0.05, 0.075 and 0.1). Zinc acetate dihydrate (99.5%), manganese (II) acetate tetrahydrate (99.0%), nickel (II) acetate tetrahydrate (99.0%), iron (II) chloride (98%), Disolol® (absolute ethanol denatured with the addition of 2% isopropyl alcohol and 2% methyl ethyl ketone), 0.5 M gallium (III) chloride (99.99%) solution in absolute ethanol and sodium hydroxide (98.0%) were used as materials for synthesis. Solution A was prepared by dissolving 2 mmol of Me precursor salts in 15 mL of ethanol (Disolol) and heated up to 60  C. Solution B was made in a two neck round bottom flask by mixing 30 mL of ethanol (Disolol) and 0.6 g of NaOH, stirred and heated until solution boiled under reflux. Then Solution A was added to solution B using syringe with a needle pushed through a perforated rubber stopper. Solution was left under reflux with stirring for 30 min, then transferred to an autoclave reactor vessel and heated at 150  C for 24 h. Next, the autoclave was cooled to room temperature. The obtained nanocrystals were washed thoroughly three times using centrifugation and methanol to remove sodium containing by-products. In the last step methanol was replaced with n-hexane for storage to eliminate ZnO degradation in time. It has been demonstrated that ZnO nanocrystals grow and change the morphology even in ambient atmosphere at room temperature [24]. For this reason, ZnO can't be stored in air or polar solvents. 2.2. Characterization Nanocrystal powders were characterized using X-ray diffraction (XRD, Rigaku Ultimaþ, Japan) with CuKa radiation and scanning electron microscope (SEM, Nova NanoSEM). Profex software was used to perform Rietveld refinement for the XRD patterns [25]. The depth uniformity of chemical composition was studied using hard X-ray photoelectron spectroscopy (HAXPES) at the KMC-1 beamline at BESSY-II (HZB) [26] at two different photon energies, 2.3 keV and 6.9 keV, corresponding to mean probe depths of approximately 30 Å and 90 Å, respectively, for the O 1s and TM 2p core level spectra presented. The calibration reference for the binding energy scales was the Au 4f7/2 photoelectron line at 84.0 eV binding energy, measured from an Au foil at both the used photon energies and at the same photoelectron analyser pass energy (200 eV) as the sample spectra. A few additional more surface sensitive X-ray absorption (XAS) and photoelectron spectroscopy (PES) measurements were carried using soft X-ray excitation (at the D1011 beamline at MAX-Lab, Lund, Sweden) on a Fe:ZnO reference sample to rationalise the role of Ga ions on the depth distribution of the Fe charge state. Magnetic properties were measured on sample powders filled in special plastic capsules using a vibrating sample magnetometer (Lake Shore Cryotronic Co., Model 7404 VSM, USA) at room temperature for fields up to 10 kOe. The sample holder was measured separately, and its signal was subtracted from the full raw magnetisation signal of each sample in the holder capsule. Light

The XRD patterns of the different Zn0.95-x(Me0.05Gax)O powders are shown in Fig. 1 (a) e (c). All sample XRD patterns show wurtzite type ZnO crystalline phase (JCPDS 36e1451), that can be assigned to corresponding (hkl) planes of diffraction maxima at 2q ¼ 31.8 (100), 34.5 (002), 36.3 (101), 47.6 (102) and 56.6 (110). The formation of an additional minor crystalline phase was observed for Ni and Mn substituted ZnO (Fig. 1 (b) and (c)). The increase of Ga ion concentration in Ni and Mn substituted ZnO nanocrystals results in an increase in the formation of metallic Ni and Mn oxide crystalline phases along with ZnO crystalline phase. Metallic Ni crystalline phase was observed at 2q values of 44.7 (111) and 52.1 (200) [27], but Mn substituted ZnO contained b-MnOOH with 2q angle of 19.2 (002) [28]. However, it is hardly possible to surely determine the crystalline phase of Mn species from one weak single peak on XRD diffractogram, thus deeper insights were pursued by HAXPES as discussed below. Gallium ion incorporation in Zn0.95-x(Me0.05Gax)O system resulted in peak broadening and loss of intensity, which could be related to reduction of crystallite size and/or non-uniform strain formation in the crystalline structure [29]. The absence of impurity phases in iron containing ZnO samples may be attributed to similar sizes of Fe2þ and Zn2þ - 0.77 Å and 0.75 Å, respectively [30]. The Rietveld refinement of the XRD data (Table 1) indicates that all the samples belong to the P63mc space group of the wurtzite structure. Substituting the ZnO with Fe, Ni or Mn ions changes the lattice parameters compared to the unsubstituted ZnO nanorods where the lattice constants are a ¼ 3.251 Å and c ¼ 5.207 Å. This suggests that the Fe, Ni and Mn ions occupy not interstitial but substitutional sites, i.e. the tetrahedral Zn sites. Unit cell volume (V0) calculations have shown that the Fe and Mn substituted ZnO have a slightly larger unit cell volume when compared to unsubstituted ZnO (47.6211 Å3) (scan step of the XRD was 0.04 ). This can be attributed to slightly larger radii of Fe2þ and Mn2þ high spin states (Fe2þ ¼ 0.77 Å, Mn2þ ¼ 0.82 Å and Zn2þ ¼ 0.75 Å) [30] that are predominant in tetrahedral complexes [31]. Similar analysis of the Ni substituted ZnO data shows a slightly smaller unit cell volume, which agrees with the smaller Ni2þ ion radius (Ni2þ ¼ 0.70 Å) [30]. However, the possibility of the change of lattice parameters due to occupation of the interstitial positions of solute ions cannot be disregarded. Rietveld refinement results show that Ga ion incorporation into the ZnO lattice result in a slight increase in the a parameter of the unit cell that can be observed in correlation with the increase in Ga ion concentration. The opposite is true for the c parameter e it gets smaller with increasing Ga ion concentration. This means that the smaller Ga ion size (Ga3þ ¼ 0.62 Å) causes a sort of “flattening” effect on the unit cell as can be seen from c/a unit cell ratios (Table 1). For all samples the c/a ratio decreases with Ga ion increase. Ga3þ ion incorporation into the tetrahedral sites of the wurtzite structure may cause formation of Zn vacancies, oxygen interstitials and/or free electrons in accordance with equations (1)e(3):

Ga2 O3

ZnO

/

00

2GaZn þ 3OxO þ V Zn

(1)

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Fig. 1. XRD patterns of different Zn0.95-x(Me0.05Gax)O powders with different Ga ion concentration: (a) Me ¼ Fe; (b) Me ¼ Ni and (c) Me ¼ Mn.

Table 1 Rietveld refinement results. Sample

Unit cell parameters

ZnO Zn0.95Fe0.05O Zn0.925Fe0.05Ga0.025O Zn0.9Fe0.05Ga0.05O Zn0.875Fe0.05Ga0.075O Zn0.85Fe0.05Ga0.1O Zn0.95Ni0.05O Zn0.925Ni0.05Ga0.025O Zn0.9Ni0.05Ga0.05O Zn0.875Ni0.05Ga0.075O Zn0.85Ni0.05Ga0.1O Zn0.95Mn0.05O Zn0.925Mn0.05Ga0.025O Zn0.9Mn0.05Ga0.05O Zn0.875Mn0.05Ga0.075O Zn0.850Mn0.05Ga0.1O

Ga2 O3 Ga2 O3

c/a

a, Å

c, Å

3.2512 3.2522 3.2520 3.2532 3.2530 3.2544 3.2490 3.2528 3.2532 3.2524 3.2545 3.2524 3.2528 3.2507 3.2528 3.2547

5.2076 5.2021 5.2013 5.1905 5.1838 5.1882 5.2071 5.2103 5.2094 5.2023 5.2046 5.2090 5.2063 5.2055 5.1992 5.2027

1.6017 1.5996 1.5994 1.5955 1.5935 1.5942 1.6027 1.6018 1.6013 1.5995 1.5992 1.6016 1.6006 1.6013 1.5984 1.5985

V0 (Å3)

47.6211 47.6487 47.6355 47.5717 47.5044 47.5857 47.6007 47.7414 47.7449 47.6564 47.7391 47.7178 47.7048 47.6359 47.6397 47.7275

pffiffiffiffiffiffi Micro strain ( k2)

0.0024 0.0029 0.0034 0.0107 0.0122 0.0095 0.0011 0.0046 0.0050 0.0082 0.0083 0.0031 0.0040 0.0048 0.0097 0.0084

ZnO

/

2GaZn þ 2OxO þ Oi

(2)

ZnO

1 2GaZn þ 2OxO þ O2ðgÞ þ 2e 2

(3)

/

00

Crystallite size calculations (Table 1) reveal, that the incorporation of Ga ion into the ZnO crystalline structure leads to particle size decrease. Whereas we did not carry out extensive high resolution electron microscopy studies on the exact particle size and morphology distribution of powders, direct comparisons cannot be drawn between the results depicting the size and shape of the particles. Nevertheless the calculated crystallite dimension values

Goodness of fit (S)

Crystallite size d(100) and d(010) (nm)

d(001) (nm)

101 ± 5.0 51 ± 2.1 40 ± 1.2 25 ± 1.4 15 ± 0.3 19 ± 0.7 97 ± 4.0 45 ± 3.9 29 ± 1.1 18 ± 0.8 21 ± 1.1 113 ± 12 38 ± 1.7 36 ± 1.2 36 ± 3.5 25 ± 1.1

431 ± 27 117 ± 9.3 75 ± 3.5 67 ± 8.8 33 ± 1.6 33 ± 1.9 258 ± 39 186 ± 53 69 ± 5.3 71 ± 9.9 63 ± 8.0 226 ± 41 99 ± 9.4 122 ± 12 140 ± 45 69 ± 6.9

1.1589 1.1374 1.2315 1.2441 1.2412 1.2503 1.1228 1.8303 1.0989 1.1671 1.1760 1.1303 1.4216 1.3183 1.1767 1.2959

exhibit similar sample to sample variation as observed in SEM studies. The incorporation of Ga ion also introduces micro strain, which increases with Ga ion concentration in the ZnO lattice due to the large size of Ga ion. This leads to limited growth of ZnO crystals. The morphology of different Zn0.95-x(Me0.05Gax)O powders were evaluated using scanning electron microscope (SEM). Fig. 2 (a)-(d) shows ZnO SEM micrographs for Zn0.95Me0.05O. Introduction of Fe ion (Fig. 2 (a)) in ZnO results in powders with a small rounded rod like particle shape with size smaller than 100 nm in all dimensions. ZnO powders substituted with Ni ions (Fig. 2. (b)) consist of nanowires up to several micrometres in length and a width of 100e300 nm. The Mn substituted (Fig. 2 (c)) ZnO nanocrystals are also dominantly in nanowire shape. However, the length of these

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Fig. 2. Comparison between SEM images of (a) ZnO nanorods and (b)e(d) Zn0.95Me0.05O particles where Me ¼ Fe (b), Ni (c), and Mn (d).

does not exceed 500 nm, while the diameter is smaller than 100 nm. The SEM micrographs of different Zn0.95-xFe0.05GaxO nanopowders with increasing gallium and decreasing zinc ion concentrations are shown in Fig. 3 (a)-(d) respectively. It can be observed that an increase in Ga ion concentration decreases the crystallite size and aspect ratio. Gallium ion introduction drastically reduces the particle size and reduces it further it as the concentration increases. Elongated particles can now be observed only in sample powders containing low Ga content. Zn0.95-xMn0.05GaxO with different ratios can be seen in Fig. 3 (i)e(l). The same as for the FeeGa co-substituted ZnO, the aspect ratio and size decreases with increasing Ga concentration. At the highest Ga concentration, particle size decreases below 50 nm. The reduction of crystallite size with addition of Ga ions can be attributed to the Coulomb repulsions, caused by free electrons introduced by compensation of Ga3þ at Zn2þ site. This repulsion makes it difficult for the zincate ions to diffuse to the crystal surface [32].

The UVeViseNIR Kubelka-Munk absorption spectra for colloids of ZnO substituted with different metal ions is presented in Fig. 4 (a). It can be seen that the Mn and Fe substituted ZnO crystals exhibit visible light absorption, which decreases with decreasing photon energy and drops to around zero at around 650 nm and 850 nm for manganese and iron, respectively. Both these samples also exhibit an additional absorption peak at around 420 nm. Contrary to this, Ni substituted ZnO exhibits the opposite effect e the absorption first drops off sharply at the UVeVis threshold and then steadily increases with the decrease in photon energy. This phenomenon could be attributed to the combined contribution of different effects. Ni cations in ZnO lattice produce energy levels in the band gap, thus providing absorption in visible range, while metallic Ni (the presence of segregated metallic nickel was indicated by the XRD results above, Fig. 1b, whereas an estimate of the relative amounts of metallic vs. oxidised nickel is given by the HAXPES results further below) should cause the surface plasmon resonance adsorption. Other point defects, such as oxygen vacancies and free electrons in conduction band gap can also induce absorption in visible and infrared range, respectively. The optical transmission spectra of the Me-Ga co-substituted ZnO nanocrystal colloids are presented in Fig. 4 (b)e(d). The transmittance measurements for nanocrystals with single transition metal ions are not demonstrated, because large crystallite sizes produce strong light scattering, which hinders direct measurement of qualitative transmittance spectra. As expected, all Me-Ga cosubstituted ZnO nanocrystal colloids exhibit light absorption in the infrared range, which can be attributed to the plasmonic resonance of the delocalised electrons in the ZnO crystal lattice [22]. The optimal Ga concentration with the highest infrared adsorption was estimated to be at x ¼ 0.05 for different compositions, although Ga concentration did not have a strong influence on the maximal infrared absorption of MneGa co-substituted ZnO. Band gap calculations of substituted ZnO nanoparticles are shown in Table 2. The optical band gap values changes from 3.12 to 3.35 eV for Fe and Ga co-substituted ZnO, from 3.21 to 3.37 eV for Ni and Ga substituted ZnO and from 3.23 to 3.32 eV for Mn and Ga substituted ZnO. The optical band gap value increases with increasing Ga concentration, which is due to BursteinMoss shift [33].

Fig. 3. SEM images of Zn0.95-x(Me0.05Gax)O nanopowders: row (a)e(d) Fe ion substituted particles, (e)e(h) Ni ion substituted particles and row (i)e(l) Mn ion substituted particles.

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Fig. 4. Optical properties of Zn0.95-x(Me0.05Gax)O products - absorbance measurements: (a) Zn0.95Me0.05O samples (Me ¼ Fe, Ni, Mn) and transmittance spectra of different Zn0.95powders: (b) Zn0.95-x(Fe0.05Gax)O; (c) Zn0.95-x(Ni0.05Gax)O and (d) Zn0.95-x(Mn0.05Gax)O.

x(Me0.05Gax)O

Table 2 Calculated optical band gaps for synthesized TM and Ga co-substituted nanocrystals. Sample

Zn0.95-xFe

x¼0 x ¼ 0.025 x ¼ 0.05 x ¼ 0.075 x ¼ 0.1

3.12 eV 3.20 eV 3.32 eV 3.33 eV 3.35 eV

0.05GaxO

Zn0.95-xNi0.05GaxO

Zn0.95-xMn0.05GaxO

3.21 eV 3.34 eV 3.35 eV 3.36 eV 3.37 eV

3.23 eV 3.24 eV 3.25 eV 3.30 eV 3.32 eV

Magnetisation versus magnetic field strength curves at room temperature for different Zn0.95-x(Me0.05Gax)O powders are shown in Fig. 5 (a)-(c) and magnetic properties are summarized in Table 3. Samples containing Fe and Ni ions exhibited ferromagnetic behaviour, but powders with Mn showed only paramagnetic response (no hysteresis). In case of the Ni and Mn co-substituted Ga:ZnO, the presence of alien impurity phases (metallic Ni or oxidised Mn, respectively) has to be taken into account. Metallic nickel should be the main reason for ferromagnetic behaviour and higher saturation magnetisation, coercive force and remanent magnetisation in Ni substituted ZnO comparing to other samples [34], while Mn species may be responsible for the paramagnetic behaviour [35]. Studies related to Mn substituted ZnO materials suggest that samples exhibiting pure ZnO phase show a very weak ferromagnetic response at room temperature (RT) which is completely suppressed at higher concentrations [36]. Similar studies have shown, that Mn substituted ZnO nanoparticles exhibit a very weak diamagnetic effect [37], while others have shown that larger particles (~1 mm) exhibit relatively strong ferromagnetism [38]. They attribute the ferromagnetic behaviour to the charge mediated magnetic coupling between Mn atoms [36]. The Fe substituted ZnO shows decrease of magnetisation, coercive force and remanent magnetisation (Fig. 5 (a), Table 3) with Ga

addition. This is a somewhat unexpected behaviour for a magnetic exchange interaction in the presence of delocalised electrons. An explanation can be that the addition of Ga produces lattice defects, which may impact and further decrease the ferromagnetic response of ferromagnetic Fe substituted ZnO. The ferromagnetic behaviour of the Fe substituted ZnO has also been observed by other authors [4e6]. It has been argued that the Mn substituting ion in ZnO requires a complementary hole co-substituting ion (e.g. Cu) in order to stabilise the ferromagnetic ground state for such a system [39]. Otherwise it would assume the lower energy antiferromagnetic configuration [40]. However, the (hole) co-substituting ion effect was found to be overcome by varying sample fabrication (and in particular, post-annealing) parameters and demonstrated that room temperature ferromagnetism was achievable with Mn:ZnO without a co-substituting ion [41]. Taking this background into consideration we take the absent ferromagnetic response of the Mn/Ga co-substituted samples as an additional indication of efficient itinerant electron doping by the Ga substituting ion, and proceed to depth sampling by HAXPES to see what effect it has on the (radial) variation of the chemical composition and the TM charge state (Fig. 6 (a)-(g)). The Mn/Ga co-substituted ZnO Mn 2p and O 1s level hard X-ray photoelectron spectroscopy data (Fig. 6 (a) and (b))) indicate Mn to be dominantly in the 3 þ charge state, but the spectrum still deviates from the available Mn:ZnO reference spectra [42] at the 2p3/2 sharp leading edge and the spectra of potentially similar model compounds [43,44]. Meanwhile, it is not trivial to find reference for the primarily assumed ligand configuration for Mn2þ substituting at the tetrahedrally oxygen coordinated Zn2þ sites, with similar model compounds, e.g. MnSO4, is with considerably stronger ligand fields, and it is not straightforward to model the photoelectron spectra in such detail without ambiguity as to be fully useful in the

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Fig. 5. Magnetic properties of Zn0.95-x(Me0.05Gax)O powders: (a) Zn0.95-x(Fe0.05Gax)O; (b) Zn0.95-x(Ni0.05Gax)O and (c) Zn0.95-x(Mn0.05Gax)O.

Table 3 Magnetic properties e comparison of the magnetic saturation (MSat), coercivity (HC) and retentivity (MR) of nanocrystals. Sample name

MSat (emu) ,103

HC (G)

MR ,103

Zn0.95Fe0.05O Zn0.925Fe0.05Ga0.025O Zn0.9Fe0.05Ga0.05O Zn0.875Fe0.05Ga0.075O Zn0.85Fe0.05Ga0.1O

75.667 73.852 40.553 17.638 48.897

57.840 47.388 13.017 19.106 11.223

2.56 2.66 0.26 0.20 0.40

Zn0.95Ni0.05O Zn0.925Ni0.05Ga0.025O Zn0.9Ni0.05Ga0.05O Zn0.875Ni0.05Ga0.075O Zn0.85Ni0.05Ga0.1O

107.82 57.041 123.64 123.20 109.15

142.72 128.43 127.69 113.84 103.84

4.51 7.64 18.03 13.84 14.83

assignment. It nevertheless can be noticed that the Mn 2p spectra are void of the typically present Mn2þ satellite at approximately 5 eV above the Mn 2p3/2 main peak [44], but on the other hand the ionicity of the bond can be quite different, which might facilitate ligand charge transfer and suppress the satellite [45]. Therefore, it is plausible that the Mn2þ satellite is suppressed, and because of the relatively lower binding energy we consider the data not to exclude a possible Mn2þ contribution. However, Mn4þ, particularly if present in segregated manganese oxide (a plausible assignment of the XRD impurity reflex below 20 in Fig. 1c), is expected to be octahedrally oxygen coordinated (as in common Mn oxides other than only the B-site in Mn3O4). This would, apart from being shifted higher in binding energy, give rise to a resolved multiplet structure in the 2p3/2 peak even in XPS. As such structure is not observed, we therefore conclude that the Mn3þ to be the most plausible dominant charge state of the Mn included. Fe substituted ZnO has been shown earlier to be prone to charge state variation dependent on fabrication, with the best

functionality not necessarily at the most closely stoichiometric substitutional inclusion of Fe2þ substituting at the Zn2þ sites [46]. While most studies have reported at least a significant Fe3þ content [47,48], the ferric (Fe3þ) component may quite often be overrepresented in the more oxidised surface layer. As seen from a comparison of the HAXPES of Fe/Ga cosubstituted versus only Fe substituted ZnO (Fig. 6 (c)e(e)), the latter has far greater abundance of iron in the Fe2þ charge state in the bulk of the nanoparticles, which would favour spin glass behaviour, whereas electron-doping was theoretically predicted early on to stabilise the ferromagnetic ground state. Experimental observations have later confirmed this prediction [46] of a strongly increased Fe3þ content, the smaller (than Fe2þ) ionic radius of which has (apart from being relevant for DMS functionality) even stronger influence on piezoelectric properties. Also, both ions are subject to feasibility of orienting the magnetic ion at the substitutional lattice site. In the Fe 2p HAXPES of the Fe/ Ga co-substituted ZnO the raised relative abundance of Fe3þ is clearly displayed. We take both this and the observation of Mn/Ga co-substituted ZnO becoming a frustrated magnet (Fig. 5c), as well as the close alignment of the plasmonic features in the optical transmission measurements (Fig. 4 b, d) as an indication of efficient electron doping using Ga co-substitution via the solvothermal synthesis route. For Ni/Ga co-substituted ZnO (Fig. 6 (f) and (g)) the TM 2p HAXPES results indicate a significant contribution from metallic nickel (the sharp peak at 852.7 eV for 2p3/2 and at 870 eV for 2p1/2), which becomes significantly more pronounced for the deeper probe (at the higher incident photon energy). The latter gives a coarse estimate of ~40% of the Ni included in the sample (in the ‘bulk’ of the nanoparticles) being in metallic (non-oxidised Ni0) charge state. We note that in earlier reports of Ni substituted ZnO the metallic Ni 2p XPS peak has been observed on some occasions [49], but not always, even for wet chemical synthesized samples

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Fig. 6. HAXPES of select samples: (a) Mn 2p and (b) O 1s XPS spectra at two different incident X-ray energies: 2.3 keV and 6.9 keV, corresponding to mean probe depths ~30 Å (solid lines) and ~85 Å (dashed lines) respectively; (c) Fe 2p XPS spectra at 2.3 keV, (d) O 1s XPS spectra at two different incident X-ray energies (same as with the (a)e(b) panels) and (e) Fe 2p TEY spectra (solid lines) and Auger PEY spectra (dashed lines) for both Fe0.05Zn0.95O and Fe2O3; (f) Ni 2p and (g) O 1s XPS spectra at two different incident X-ray energies.

[50] (the corresponding imfp~15 Å for Al-Ka excitation used in both of these examples). We find the large relative amount of metallic nickel deeper in the sub-surface of these samples as we observed it at higher incident energy. Ni 2p HAXPES are quite intriguing and perhaps even promising for some alternative paths to construct a functional material, because the magnetic response of these samples is still considerably large. From a different viewpoint, the synthesis (but also the orientational freedom of the included cosubstituting ion at the lattice sites) has an affinity towards decreasing ionic radius of the observations of (segregated) metallic clusters or minor nanoparticles. We observed this earlier for Co/Ga co-substituted ZnO [3] which also had a significant metallic component of ~8e9% of the included Co in a Co/Ga co-substituted ZnO.

4. Conclusions In summary, plasmonic transition metal and gallium cosubstituted degenerated diluted magnetic semiconductor ZnO nanocrystals were successfully synthesized by solvothermal approach. Ga ions primarily introduce delocalised electrons in the conduction band, while Ni and Fe ions give rise to the ferromagnetic properties in the ZnO nanocrystals. Mn and Ga substituted ZnO exhibit paramagnetic material behaviour. Results confirm that a considerable delocalised electron density is established, which significantly contributes to light absorption in the infrared range and magnetic exchange coupling.

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Acknowledgements This work was supported by the Latvian Council of Science in the framework of FLPP (Plasmonic oxide quantum dots for energy saving smart windows, lzp-2018/1e0187) and by the Estonian Centre of Excellence in Research project “Advanced materials and high-technology devices for sustainable energetics, sensorics and nanoelectronics” TK141 (2014-2020.4.01.15e0011). We are grateful to the staff of BESSY II for the assistance and co-operation during the synchrotron-based measurements.

[20]

[21]

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