ZnO2 nanocomposites and hexagonal prism shaped ZnO nanoparticles synthesized by pulse laser ablation in Heptanes

Accepted Manuscript Structural, optical, electronic and magnetic properties of multiphase ZnO/Zn(OH)2 /ZnO2 nanocomposites and hexagonal prism shaped ZnO nanoparticles synthesized by pulse laser ablation in Heptanes

Saif Ullah Awan, S.K. Hasanain, Jamshaid Rashid, Shahzad Hussain, Saqlain A. Shah, Mian Zahid Hussain, Mohsin Rafique, M. Aftab, Rashid Khan PII:

S0254-0584(18)30155-X

DOI:

10.1016/j.matchemphys.2018.02.051

Reference:

MAC 20402

To appear in:

Materials Chemistry and Physics

Received Date:

12 September 2017

Revised Date:

06 February 2018

Accepted Date:

26 February 2018

Please cite this article as: Saif Ullah Awan, S.K. Hasanain, Jamshaid Rashid, Shahzad Hussain, Saqlain A. Shah, Mian Zahid Hussain, Mohsin Rafique, M. Aftab, Rashid Khan, Structural, optical, electronic and magnetic properties of multiphase ZnO/Zn(OH)2/ZnO2 nanocomposites and hexagonal prism shaped ZnO nanoparticles synthesized by pulse laser ablation in Heptanes, Materials Chemistry and Physics (2018), doi: 10.1016/j.matchemphys.2018.02.051

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Graphical Abstract

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Structural, optical, electronic and magnetic properties of multiphase ZnO/Zn(OH)2/ZnO2 nanocomposites and hexagonal prism shaped ZnO nanoparticles synthesized by pulse laser ablation in Heptanes Saif Ullah Awan1,2*, S. K. Hasanain2, Jamshaid Rashid3, Shahzad Hussain2,4, Saqlain A. Shah5, Mian Zahid Hussain2,6, Mohsin Rafique4,7 M.Aftab2, Rashid Khan2 1. 2. 3. 4. 5. 6. 7.

Department of Electrical Engineering, NUST College of Electrical and Mechanical Engineering, National University of Sciences and Technology (NUST) Islamabad, Pakistan. Department of Physics, Quaid-i-Azam University, Islamabad 45320, Pakistan. Department of Environmental Sciences, Quaid-i-Azam University, Islamabad 45320, Pakistan. Department of Physics, COMSATS Institute of Information Technology, Islamabad 44000, Pakistan. Department of Physics, Forman Christian College University, Lahore, Pakistan. College of Engineering, Mathematics and Physical Sciences, University of Exeter, EX4 4QF, Exeter, UK. State Key Laboratory of Low-Dimensional Quantum Physics, Department of Physics, Tsinghua University, Beijing100084, China. *[email protected]

*[email protected]

Abstract: We studied structural, optical, electronic and magnetic properties of undoped well crystalline hexagonal and non-hexagonal ZnO system. X-ray diffraction (XRD) and high resolution transmission electron micrographs (HRTEM) of air annealed at 550⁰C (sample-d) confirmed the presence of hexagonal (wurtzite) single phase ZnO. Raman analysis detected the vibrations of fundamental and second order phonons of Zn and oxygen related species. In Photoluminescence (PL) spectra, we observed that the intensity of UV peak decreases as annealing temperature increases. Broad PL visible band of samples shifted towards lower wavelength due to annealing effects. Fitting of broad PL spectra confirmed the presence of Zinc interstitial (Zni), Zinc vacancy (VZn) and oxygen vacancy (Vo) defects. To determine the Zn interstitial and vacancy defects Auger peaks were de-convoluted from X-ray photoelectron spectroscopy (XPS) survey scan. Oxygen vacancies are estimated by fitting of asymmetric XPS O-1s spectra. Incomplete oxidation of magnetization versus temperature in the presence of applied field showed the presence of irreversibility in 550⁰C annealed-sample. Zn clusters and secondary phases are may be reason of room temperature ferromagnetism in non-hexagonal samples (a-c). While, complex defect (zinc and oxygen vacancies) as observed in electronic and PL data may be responsible for inducing, promoting and stabilizing room temperature ferromagnetism in well crystalline hexagonal ZnO nanoparticles (sample-d).

1. Introduction: Spintronics is one of the hottest frontiers in science and technology since the last few decades1. The search for novel room-temperature ferromagnetic (RTFM) materials with high 1/20

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Curie temperature (Tc) is very crucial for spintronics applications, e.g., spin-valve transistors, spin light emitting diodes, and nonvolatile storage2. The role of certain structural and electronic defects such as cationic vacancies3 and oxygen vacancies4 are very important for RTFM in oxides. This type of magnetism is referred to as “defect ferromagnetism”5. It is being recognized as an intriguing new area of research in which the ferromagnetism (FM) does not originate from the parent materials, but is caused by certain types of defects6 that affect the electronic band structure and hence the magnetic properties. Efficiently controlling defects and understanding their role makes these materials suitable for applications in spintronics. Experimentally RTFM has been observed in a wide range of oxides, which do not contain ions with partially filled d or f bands. The term d0 ferromagnetism7, 8 is also used to describe this type of unexpected magnetism. The advanced discoveries of these d0 ferromagnetic materials in undoped HfO28-10, ZnO11-14, TiO215, In2O316, MgO17, 18, CeO219 and SnO220, and their different synthesis techniques have attracted an enormous attentions of researchers worldwide. The origin of ferromagnetism in these systems is still under debate. Recently d0 ferromagnetism has been reported in carbon21, graphite22-24, graphene25, alkaline-earth hexaborides7, 26 and Au, Ag, Cu27 nanosystems. In the light of these discoveries, there is an emerging consensus that defects in these systems play critical roles for mediating FM28. In particular, ZnO based systems have been predicted by both, the mean-field Zener model and first-principles calculations, to have a Tc well above room temperature, provided that the combination of carrier concentration and defects are optimized29, 30. Additionally, the optical properties of ZnO that exhibits UV and visible luminescence, make it a very promising wide band gap material with a wide range of potential applications in optics31. In pure ZnO systems, various intrinsic defects have been predicted exclusively to be the cause of the RTFM, and to a certain degree, it helps to settle the controversies related to RTFM due to the absence of any externally doped metallic element. Defects7 are believed to be responsible for initiating hybridization of orbital at the Fermi level and establishing a long-range ferromagnetism28, 32. The first-principles studies suggested that the origin of the magnetism and related properties in nonmagnetic oxide materials is the cationic defects, induced by the magnetic ground state of the neutral cationic vacancies that are surrounded by oxygen ion wave functions33-40. It may not be very surprising though since the defects have been recognized for a long time to play important roles in dictating the electrical and optical properties of most wide band gap oxide semiconductors. It was also proposed computationally that defects like Zn vacancies11-13, 41, oxygen vacancies42, 43, Zn interstitials44, Zn and O vacancies13, 45-47, grain boundaries 48, 49, Zn interstitials and Oxygen vacancies 50, Oxygen interstitials and zinc vacancies 51-53 and lattice distortions54 could induce ferromagnetism and establish the correlation between physical properties in terms of the defect mediated FM45, 55, 56. Therefore, it is very important to investigate the roles of intrinsic defects in the origin of ferromagnetism in wide band gap oxides57, 58. 2/20

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Ferromagnetism has been observed in a variety of transition metal(TM) doped oxides systems59-63. Experimental studies on TM-doped ZnO have produced inconsistent results and the mechanism of FM remains unclear64. Experimental artifacts, segregation of secondary ferromagnetic phases, magnetic clusters, and indirect exchange mediated by charge carriers have been invoked to explain the RTFM in TM-doped oxides systems65, 66. Since the recent discovery of RTFM in non-TM-doped ZnO67, the magnetic properties are not exclusively related to the magnetic ions but strongly determined by the defects. It is important to understand whether the undoped ZnO is ferromagnetic or not. If it is ferromagnetic, what kind of defects cause this d0 ferromagnetism? How to establish the longrange magnetic coupling of local moments in an oxide host? Since the dimensionality affects the band structure and density of states of the systems, we want to investigate the types of defects and defect densities that promote strong ferromagnetism in well crystalline nanoparticles system. Why ferromagnetism is produced in as synthesized ZnO material (i.e., not a single phase) and enhanced due to annealing (i.e., hexagonal phase achieved). These are some of the important questions that will be addressed in this article. To avoid the influence of substrate and the interface between film and substrate, we have synthesized ZnO nanoparticles by a pulsed laser ablation (PLA) process. Nanoparticles synthesis by PLA in liquids does not require any chemicals and thus it eliminates possibilities of magnetic impurities. And there is no need of further purifications to remove harmful and undesired species introduced by wet chemistry methods. This report presents structural, electronic and magnetic properties of un-doped ZnO nano-crystalline prepared by PLA.

2. Methods and Measurements: ZnO nanoparticles were synthesized by pulsed laser ablation of Zn metal foil (Alfa Aesar, purity 99.90%). The target was irradiated with a wavelength of 1064nm pulsed Nd:YAG laser (Continuum Powerlite Precision 8000) operating at 10Hz with 150ns pulse width in the environment of Heptanes. We tried following different liquids that were available in laboratory e.g. distilled water, deionized water, acetone, ethanol, methanol, Isopropanol and cyclohexane other than Heptanes. But fortunately, we get desire results (sufficient amount of powder) at optimized parameters using Heptanes. There was no ablation of Zn metal foil except using Heptanes. This was the reason we used this liquid. The other reason was, it is cheaper and easily available in market. The laser focused on the target vertically with the spot size of about 1mm in diameter using a prism and a quartz lens with the focal length of 10cm. A laser intensity 130mJ/pulse was employed to ablate a target. The ablation rate determined by weighing the Zn foil before and after exposure was 0.3mg/h. Ablation was carried out for 12hours to obtain the required quantity for different characterization techniques. The solid powder was obtained after centrifuging the solution. The un-annealed sample (sample-a), air annealed samples at 400⁰C (sample-b), 450⁰C (sample-c) and 550⁰C (sample-d) were examined by using different characterization techniques. Sample-b was annealed for 6 hrs while sample-c and sample-d were annealed for 12 hours. 3/20

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The structural properties of these samples were studied by PANalytical X’pert PRO x-ray diffraction (XRD) with Cu Kα radiation (λ=0.154nm). Microstructure properties were studied by using JEOL-2010 field-emission, transmission electron microscope (TEM). Raman scattering spectroscopy was measured using a diode-pumped solid state laser of wavelength 532nm with a power of 50mW. He–Cd laser with 50mW power and photon energy 3.81eV (325nm) for Photoluminescence (PL) study. X-ray Photoelectron Spectroscopy (XPS) was performed on a Thermo Scientific ESCALAB-250 spectrometer with a focused monochromatic AlKα(hν=1486.6eV) source. Binding energies of photoelectrons were correlated with aliphatic hydrocarbon C-1s peak at 285eV. Vibrating Sample Magnetometer (VSM) was used to investigate the magnetic properties in the temperature range 50K-300K and maximum field strength of±3 Tesla.

3. Results Analysis and Discussion: 3.1 Structure and Microstructure: The Bragg-Brentano scan of as-grown sample from Zn metal foil (sample-a) is shown in the Fig.1(a) with many peaks indicating that as prepared sample is not completely amorphous. The analysis of different diffracted peaks confirmed that as-grown sample is not a single hexagonal phase, but contains different impurity phases like Zn metal, Zn carbide, ZnO and Zn(OH)2. To remove these impurities, as grown sample was annealed at 400⁰C (sample-b) in a furnace in air atmosphere for 6-hours. The XRD spectrum of sample-b shows that 95% hexagonal crystalline phase structure of ZnO was achieved with impurity phases of Zn metal at 39.0° and 43.2° respectively. Then the as grown sample was air annealed at 450⁰C for 12-hours. Again, we could not achieve the 100% hexagonal structure of ZnO due to extra peak at 39.0° with the shoulder at 39.1° and another extra peak at 43.2° of Zn metal. The XRD pattern of sample-b and sample-c are presented in Fig.1(b), while the extra phases of sample-b and sample-c are shown in the inset of Fig.1(b). Sample a-c are knows as non-hexagonal crystalline samples. However, the hexagonal structure of ZnO was achieved, when as grown material was air annealed at 550⁰C in the furnace for 12-hours (called sample-d). XRD pattern of sample-d exhibited peaks consistent with a hexagonal structure as demonstrated in the Fig.1(b). The XRD pattern of sample-d shows the polycrystalline orientation of diffracted planes without detection of any other secondary phases. The lattice parameters measured for sample-d are a=b=0.32(±0.05)nm and c=0.52(±0.07)nm, consistent with those in literature3. The average crystallite size calculated by using Scherrer formula for sample-d was ~27nm . Low-magnification bright field TEM image as acquired for sample-d is shown in Fig.2(a). We noticed the agglomeration of nanoparticles as the sample was post annealed for longer time (12 hrs) to improve the phase of sample. The sample-d exhibited a crystalline structure with the size range of about 20nm to 30nm. High resolution TEM images of sample-d were acquired at different scales (20nm, 10nm and 5nm) as shown in Fig.2(b-d). We found the homogeneous crystalline structure. The hexagonal morphology is shown in Fig.2(b)-(d), the crystalline ZnO granules are well separated by grain boundaries. The fast Fourier transform 4/20

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(FFT) is shown in Fig.2(e). Hough Transforms (selected area electron diffraction pattern) were applied to the FFT to extract the d-spacing values as shown in Fig.2(f). The inter-planar spacing (d-values) determined from the diffraction-ring diameters are clearly associated with the Braggreflections peaks of ZnO having hexagonal (wurtzite) type structure with space group of SGP63mc, lattice-parameters a=0.32(±0.03)nm and c=0.52(±0.04)nm. These d-values are in good agreement with the XRD data. From the XRD and TEM measurements, we found that all dvalues and diffracted planes corresponded to ZnO crystalline hexagonal structure (JCPDF 00036-1451). Hence, neither contamination nor intermediate compounds have been detected in sample-d. We will discuss optical, electronic and magnetic properties of the few selected samples.

3.2

Raman Spectroscopy:

Fig.3(a-b) depicts room temperature micro-Raman scattering spectra in the range of 701200cm-1 of the samples (a, b, c and d). The bands at 102cm-1 and 438cm-1 are attributed to ZnO non-polar optical phonons of low E2L and high E2H modes, respectively. The band at 408cm-1 and 574cm-1 are E1 symmetry with TO and LO modes, respectively. The spectrum consists of three peaks located at 102cm-1, 438cm-1 and 574cm-1, which correspond to the E2L, E2H, and A1(LO) fundamental phonons modes of hexagonal ZnO, respectively. The Raman peaks with frequencies at 202cm-1, 333cm-1 and 660cm-1 were assigned to 2E2L and E2H-E2L and TA+LO, respectively68, 69. These additional modes cannot be explained within the frame work of bulk phonons modes, which are attributed to multiple phonons process in the literature. The bands at 202cm-1 and 333cm-1 have been reported to occur under resonance conditions and can be interpreted as the second order phonons 2-TA(M) and 2-E2(M), respectively70, 71. The mode at 667cm-1, does not belong to first or second order structure of ZnO. The modes at 498cm-1 and 996cm-1 are due to multi phonon scattering process. Overall the positions of phonon vibrational frequencies are according to literature72. Normally, Raman peaks observed in between 570cm-1 and 590cm-1 are considered to be associated with structural disorders, such as oxygen vacancies, Zn interstitial and their combinations, due to strong dependence on the oxygen stoichiometric72, 73. So, the presence of a small band at 576cm-1 is due to the A1(LO) mode showing the presence of oxygen vacancies or Zn interstitial74. The peak at about 574cm-1 can be assigned to A1 longitudinal optical (LO) mode A1(LO) and 1156cm-1 mode is just a second order vibration (2LO) of LO modes respectively, which is caused by the defects such as oxygen vacancies, Zn interstitial defects or the complex defects47. We observed that the Raman intensity of sample-d is lowest and sample-a is the highest. An important observation: the intensities of Raman peaks decreased as the annealing temperature was increased, implying that the concentration of defects decreases as the annealing temperature increases. Similar observations have been reported for other materials as well47. It is interesting to note that E2L phonon mode has undergone broadening of the line width. Although this may be due to strong coupling of free carriers with E2L mode, the influence of change in particle size due to annealing cannot be ruled out. The Raman spectra showed that 5/20

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the only possible defects are either Zn interstitials (corresponding with Zn vacancies) or Oxygen vacancies in these samples. We will discuss the effect of these intrinsic defects on the magnetic behavior of these samples in section 3.5.

3.3

Photoluminescence (PL) Spectroscopy:

In n-type ZnO, Zn vacancy (VZn) acts as an acceptor, while dominant oxygen vacancy (VO) behaves as donor and its concentration is high75, 76. It is hard to distinguish VZn and VO from the PL spectrum only due to the slight difference of their emission centers, and researchers always combined the annealing environment to recognize their emission. The increase of green emission intensity is normally assigned to the enhancement of VO concentration under oxygendeficient and reducing annealing conditions61,65,66, while it is induced by VZn in an oxygen rich environment77. In the previous studies78, 79 on ZnO system, accumulated evidence indicates that defects11-13, 41-52, 54 play a crucial role in generating the ferromagnetism. However, the detailed mechanism is unclear yet. One of the important obstacles is that the chemistry of ZnO defects is sophisticated, including the main defect types of interstitial Zn and oxygen, Zn and oxygen vacancies. It is well known that ZnO is an ideal ultraviolet light emitter and could emit visible lights when defects are introduced80, 81. Many groups have done a great favor to reveal the relationship among defects in ZnO and the emissions in the visible light region82, 83. For example, the blue, blue-green, green-yellow, and red emissions have been assigned to Zni84, VZn85, 86, VO82, 87-89 and O 89, 90 respectively. Therefore, it should be quite helpful in understanding the origin of i ferromagnetism of ZnO induced by defects by study of the visible emission of ZnO induced by defects. Fig.4(a) shows the room temperature PL spectra of sample-a and sample-d, which contained the ultraviolet (UV) and visible emission defects. The UV and visible emissions are generally ascribed to near-band-edge (NBE) and deep-level (DL) recombination, respectively. We noticed that the intensity of UV emission is almost comparable with the broad band visible emission of sample-a. Intensity of UV peak decreased very significantly as compared to the intense visible broad band spectrum ranging from blue to red of sample-d. The decrease in the intensity of UV peak indicates that the optical quality of the sample-d deteriorates as compared to sample-d due to increase in annealing temperature. We measured the band gap (3.26eV) of sample-d from the UV emission at ∼380 nm. The shift in the peak position of UV emission in both samples is not considerable, while the wavelength shift (toward lower number) of the broad band has been clearly observed in Fig.4(a) of sample-d as compared to un-annealed sample. This shift reflects that concentration and type of defects vary with annealing temperature. As there is still no definite agreement on the origin and positions of the visible emission of ZnO, different groups have reported different positions of defects. The broad visible band has been deconvoluted into three (for sample-a) and four (for sample-d) different nicely fitted Gaussian peaks along the entire emission range as shown in Fig.4(b-c) respectively. Similar deconvolution of the broad visible band for doped ZnO nanowires has been reported91. In Fig.4(b-c), the peak “1” located at 480±2.25nm can be attributed to the electron transition from 6/20

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the shallow donor level of zinc interstitial and oxygen interstitial to the valence band92-94. It has been suggested that Zn interstitial is a shallow donor level defect, which emits blue light and has high formation energy95, 96. We observed blue light emission corresponding to Zn interstitial in the PL spectra that’s corresponding to interstitial Zn defects. The deconvoluted peak “2” in Fig.4(b-c) centered at 520±2.5nm is attributed to the transition from the conduction band to VZn defect level93. Peak “3” located at 560±3.2nm is typically related to oxygen defects94, 97. The peak “4” observed at 665nm in sample-a and at 700nm in sample-d lies in the yellow-red emissions of ZnO, and is commonly attributed to the complex of Zinc85, 86 and oxygen82, 87-89 vacancies, respectively. These defects extracted from PL measurements are further verified using electronic properties of these samples.

3.4

Electronic Properties:

In order to identify the elements and the other possible magnetic contaminations and their chemical oxidation states, X-ray photoelectron spectroscopy (XPS) measurements were performed. Fig.5 shows a typical wide scan survey spectrum of two selected samples (a and d). The photoelectrons peaks of the survey spectra reflect the main elements, Zn and O and C, and Auger Zn-LMM and O-KLL were obtained. No other contamination was detected in both samples. Lorentzian-Gaussian fittings were employed for the deconvolution of XPS spectra into separate peaks using XPSPEAK4.1 software after subtracting the Shirley background98. High resolution Zn-2p core level XPS spectra were measured to examine how the electronic structure of Zn modifies upon air annealing. The Zn 2p3/2 and 2p1/2 XPS spectra of the sample-a and sample-d are displayed in Fig.6. The spectra of the 550°C annealed sample-d show symmetric single peaks located at 1021.88eV and 1044.98eV corresponding to Zn 2p3/2 and 2p1/2 levels, respectively. The single peak rules out the existence of multiple components of Zn in the samples. The Lorentzian-Gaussian fitting of each Zn 2p3/2 and 2p1/2 peaks were nicely fit to these spectra (fitting not shown here). Both Zn 2p3/2 and Zn 2p1/2 peaks of sample a and d showed a slight shift in their positions after annealing at higher temperatures. However, the peak positions of Zn-2p3/2 and Zn-2p1/2 in both samples match closely with the standard values of ZnO3, 99, indicating that Zn species are in +2 oxidation state. The difference of Zn-2p3/2 and Zn-2p1/2 peak positions was 23.1eV100 and this difference remained the same in all samples. From the observed spin-split (∆Zn) values and the binding energy positions, it may be concluded that the Zn atoms are in +2 oxidation states. Thus we did not observe any electronic structural changes in the Zn site from Zn-2p core level XPS spectra. So, we will further concentrate on Auger peaks, since these peaks show larger shape changes rather than XPS spectra to identify the different oxidation states of species, because mostly three electrons are involved with many body effects in a single Auger transition101. We obtained Auger peaks of Zn from the XPS survey spectra to measure the presence of interstitial Zn (Zni) defects. We observed that high resolution XPS spectra of Zn 2p peaks have symmetric features, whereas Auger peaks usually showed asymmetric shape. Hence, to determine the chemical oxidation states of Zn atoms, the analysis of Zn LMM Auger peaks were 7/20

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utilized to quantify Zn interstitial versus Zn-O. Typically measured Auger peak Zn L3M45M45 of sample-a and sample-d was deconvoluted. The two Lorentzian-Gaussian fitted Auger peaks of each sample that are attributed to the Zn-O bond and Zni, which are located centrally at higher binding energy 497.6eV and lower binding energy 493.95eV respectively101, are plotted in Fig.7(a-b). The relative normalized area obtained from the deconvoluted peaks was estimated to extract Zni defects concentration and Zn-O bond. In ZnO hexagonal crystal structure, Zn interstitials are usually found to be located between O2- and Zn2+ positions101 accompanied by Zn vacancies (Vzn)101. These Vzn defects may be responsible for ferromagnetic behavior in ZnO systems as has been reported13. In our samples, we found that with the increased annealing temperature up to 550°C, the concentration of Zni decreased (at higher annealing, we achieved single hexagonal phase structure). It is understood that the generation of each Zni defect would in general be accompanied by the generation of a Zn vacancy. However, we noticed that the relative concentration of interstitial Zn defects (Corresponding Zn vacancies) in sample-a are greater than sample-d. In other words, we may suggest that the estimated Zn vacancy defect concentration is lower in sample-d. Our XPS studies were carried out at higher resolution O-1s core level spectra were measured for samples-a and sample-d to investigate how the oxygen content is varied upon annealing and whether it has some correlation to the observed ferromagnetic properties. We observed that measured spectrum is asymmetric indicating the presence of multi-component of the oxygen species. In Fig.8(a)-(b), a typical O-1s spectrum is shown and fitted with two Lorentzian-Gaussian peaks having the different binding energy components. However, the most intense first peak at 530.5±0.3eV is due to the bulk oxygen 1s contribution from the samples, ensuring that the main signal is eminent from the oxygen bonded with Zn ions. This lower binding energy peak marked as O-Zn is attributed to the O2− ions in the wurtzite structure. The peak marked as Vo for higher binding energy 532.52±0.2eV is associated with oxygen ions in oxygen deficient regions (Oxygen vacancies) within ZnO matrix40,60. The intensity of O-Zn Gaussian peak enhanced strongly for sample-d as compared to a sample-a. This intensity being a direct measure of bulk oxygen content clearly shows a very large depletion of oxygen in sample annealed at 550°C. The relative variation in the fitted peak area for both samples is within an error limit of 5%. We do not claim these values to be an absolute quantitative estimate of the bulk oxygen (O-Zn) content, but we are very confident that these values present a reliable qualitative estimation on the relative variation of bulk oxygen in the samples. However, we may argue that the hexagonal structure of annealed sample-d has less oxygen vacancies as compared to un-anneal sample-a. Now, after confirming the defects from electronic data, we will discuss how these defects play a role on magnetic properties of laser ablated materials.

3.5

Magnetic Measurements:

Room temperature (300K) and low temperature (50K) magnetic hysteresis loops were measured for all samples. The M-H (magnetic moment versus applied field) curve of sample (a, b and d) at ~300K has been demonstrated in the Fig.9. The raw data of the sample holder and 8/20

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sample (a, b and d) have been presented in the upper inset Fig.9. After the necessary sample holder background diamagnetic subtraction, we observed S-shaped hysteresis curve for all samples. Due to the limitation of our magnetometer system (±3 Tesla field), we noticed the nonsaturation magnetization signals showing the contribution of paramagnetic behavior in sample-b and sample-d while diamagnetic response was present in sample-a. From M/H curves (at 300K), we calculated the saturation magnetization (Ms)of sample a, b and d to be 0.008emu/g, 0.015emu/g and 0.011emu/g respectively. The coercivity and remanent magnetization is more clearly seen in lower inset Fig.9. The coercivity of sample-d is 250Oe, while the remanent magnetizations 0.003emu/g respectively. The small coercivity value is an indication of soft magnetic material at 300K. Similarly, we performed the M/H measurements for sample (a, b and d) at 50K, the raw data are shown in upper inset Fig.10. The raw data looks like diamagnetic behavior due to sample holder effects. After subtracting the background signals of sample holder from measured raw data, the non-saturated ferromagnetic hysteresis signals of sample (a, b and d) are presented in Fig.10. The measured values are 0.013 emu/g, 0.022emu/g and 0.028emu/g for sample a, b and d respectively at 50K. The coercivity and remanent magnetization from the hysteresis curves of samples at 50K are presented in lower inset of Fig.10. Overall, we found that the Ms of samples at 50K is higher than 300K measurements, which is a usual phenomenon in weak ferromagnetic oxide systems due to paramagnetic contributions. Our XRD data reflects that sample-a is not showing a single phase (i.e. contained highest secondary phases and clusters) even then this sample confirming ferromagnetic response. Interestingly, we noticed that at room temperature, un-annealed sample-a attained maximum magnetization (0.012emu/g) at 0.8T field, and on further increasing the field, the curve bent downward and its magnetization decreased. We may suggest that, still the diamagnetic contribution was present in the sample-a. Similarly ferromagnetic results of Zn clusters embedded in ZnO has been presented by Yi et al102. Ferromagnetism probably originates by incomplete oxidation of Zn particles in sample-a. It indicates that ferromagnetism in sample-a may be due to the unique structure of Zn clusters embedded in ZnO matrix. In sample-b, XRD data showed 95% hexagonal structure i.e. as more and more Zn particles are oxidizing (compared to sample-a) the ferromagnetism response is enhancing. In non-hexagonal samples (ac) ferromagnetism originate due to due to clusters and secondary phases of Zn species. We mostly observe the intrinsic response of the properties of samples at low temperature. By comparing the room temperature and low temperature M/H curves of samples, it is clear that saturation magnetization of sample-d (pure hexagonal crystalline material) is higher than other samples. So at low temperature the prominent intrinsic ferromagnetic behavior of pure crystalline sample-d is obvious. By comparing magnetic properties of well crystalline hexagonal material (sample-d) with other samples (i.e. that are not fully crystalline), from M/H data, we concluded that in pure well crystalline material the ferromagnetic response may be due to intrinsic defects (cationic and anionic) in the absence of secondary phases and clusters. Similarly, complex defect (cationic and anionic) has been reported in undoped crystalline CeO2 9/20

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semiconductors103, 104. This comparison between the samples clearly indicates that the observed enhancement of ferromagnetism is not due to any magnetic impurity or clusters in the sample-d, rather it is an intrinsic effect and more pronounced due to the annealing effects and stabilizations of cationic and anionic defects. This phenomenon of enhancement of ferromagnetism upon thermal annealing is explained by measuring the magnetization at different temperatures. Magnetization versus temperature M(T) measurements were obtained with magnetic field intensities of 500Oe and 1000Oe in order to measure and determine possible irreversible effects. In field-cooling (FC) mode, the measurement was performed starting at the maximum temperature (300K) with the magnetic field applied and lowering temperature to 50K. In zero field cooling (ZFC) modes, the sample was first cooled to the minimum accessible temperature (50K) in the absence of magnetic field. Once in thermal equilibrium, the magnetic field was applied and the measurement was initiated, increasing temperature to 300K. With those two modes, the behavior related to irreversible processes could readily be observed. FC-ZFC magnetization curves of sample-a show the typical behavior like superparamagnetism with irreversibility between the FC and ZFC branches of the magnetization below the blocking temperature (TB) as demonstrated in Fig.11(a). As expected, TB decreases when the magnetic field increases and ZFC and FC branches collapse for fields larger than about 3KOe. Very similar ZFC-FC curves of Co doped ZnO nanoparticles have been reported in literature105. One of prominent features of the M(T) curves is that, the discrepancy between the FC and the ZFC curves becomes pronounced with decreasing magnetic field. In addition to this, the blocking temperature (freezing of spins), tends to decrease with increasing the magnitude of the applied magnetic field. This behavior is expected because the reduced energy barrier caused by the external magnetic field can allow the reorientation of the super-paramagnetic moments by thermal fluctuations at lower temperatures106. The room-temperature M/H curves of sample-a presented in Fig.10, clearly show hysteresis characteristics that are not normally expected in a system of uniformly sized nano-meter scale super-paramagnetism. The M/H curve of a superparamagnetic system above TB should ideally possess negligible values of remanence and coercivity but our measured values might be due to defects. Correlating, the M/H and M/T data, it is clear that saturation magnetization in sample-a increased at lower values of applied magnetic field. In other word, we observed that with increasing the magnetic field, the response of magnetic moments was smaller, which is typically a diamagnetic behavior. Fig.11(b) shows the ZFC and FC magnetization curves taken at 500Oe for the sample-b. The distinct bifurcation of the FC and ZFC starts at around 250K moving toward lower temperature. The magnetization increases with decreasing the temperature from 300K to 50K as illustrated from both FC and ZFC curves. Fig.11(c)-(d) shows the ZFC and FC magnetization curves taken at 500Oe and 1000Oe for the sample-d. Both the ZFC and the FC curves of sampled are showing similar behavior like paramagnetic materials107, but not overlapping with each. The distinct bifurcation of the FC and ZFC starts well above 230K at 1000Oe and 100K at 500Oe. The bifurcation between ZFC and FC curves indicate a large thermal irreversibility with 10/20

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the onset near 230K and 100K, which is characterized as the existence of super-paramagnetism above 230K and 100K temperatures, due to the nano-size effect107. It is obvious from this M(T) data that the thermo-remanent magnetization has been enhanced significantly upon thermal annealing. The FC and ZFC magnetization increases as a function of temperature for ~ T ≤100K, indicating that there are blocked moments which start to contribute to the magnetization when the temperature is increased. The FC and ZFC magnetization decreases as the temperature increases, since the frozen moments start randomizing due to thermal energy. This behavior is seen in most dilute magnetic semiconductor systems. It is interesting to note that there is a sharp increase in the magnetization in both the FC and the ZFC data below 100K. The steep increase in magnetization with decreasing temperatures below around 50K in the M vs. T curve is characteristic of all defects based materials, and is probably related to the defects structure and possible fraction of defects (vacancies/interstitial) which are participating in the long range ferromagnetic ordering108. We propose the possibility of complex defects formed due to intrinsic defects as the effective source of RTFM in ZnO. Understanding ferromagnetism in d0 magnetic materials is a big challenge. An open question still remains unanswered related to defect-engineered magnetism in oxides: e.g., what kinds of defects can contribute to magnetic moments? Some researchers have explained it through the Zener model2 and the bound magnetic polaron (BMP) model28, 109. But there is currently no consensus on the origin of ferromagnetism in d0 systems. It is reported and believed that the native defects introduced during sample preparation play very important roles for the magnetism observed. Although the ferromagnetism obtained in undoped oxides is conformably attributed to the existence of native defects, conflicting arguments exist. While earlier experimental groups proposed that the ferromagnetism arises from oxygen vacancies9, 15, 20, 43, theoretical work suggested that it was the cationic vacancies that introduced the ferromagnetism38, 46, 110-112. About the origin of the ferromagnetic signal, on theoretical basis, Zuo et al51 have studied the FM induced by both oxygen interstitial and zinc vacancy. Yang et al54 have explored that the Zn vacancy defect brings a spin polarized state in the nearest neighbor oxygen atoms, whereas the oxygen vacancy defect has no such influence on the magnetism. They also found that the lattice distortion is a crucial factor for the Zn vacancy induced ferromagnetism, because the ferromagnetic ground state cannot be achieved if there is no lattice distortion due to Zn vacancy defect. Along with these, the theoretically calculated x-ray absorption spectroscopy and x-ray magnetic circular dichroism (XMCD) showed the FM due to Zn vacancies. Wang et al53 calculated that magnetic moment arose from the unpaired 2p electron at O site surrounding the Zn vacancies with each nearest-neighbor oxygen atom carrying a magnetic moment. Experimentally, Mal et al113 has reported that FM in ZnO thin films might be defect-mediated. Khalid et al13 have investigated that FM in ZnO films is related to Zn vacancies. Xing et al114 suggested oxygen vacancies induced and significantly boosted the RTFM in undoped single-crystalline ZnO nanowires. Xu et al45 synthesis with the careful control of defects are a better way to obtain reproducible, intrinsic and homogeneous RTFM in undoped 11/20

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ZnO thin films. Xing et al115 observed RTFM in the as-grown Zn deficient ZnO thin films. Sundaresan et al43 has demonstrated that FM may be due to the exchange interactions between unpaired electron spins arising from oxygen vacancies at the surface of the ZnO nanoparticles. Reproducible switching between “off” and “on” ferromagnetic states using oxygen annealing in combination with either laser irradiation or vacuum annealing treatments of ZnO thin films has also been reported56. Straumal et al48 synthesized ZnO thin films with respect to the ratio of grain boundary to grain volume ratio, showing that reproducible magnetization was proportional to the film thickness, and strongly suggested that grain boundaries and related vacancies were intrinsic origin for RTFM. Kapilashrami et al116 studied the transition from FM to paramagnetism and diamagnetism as a function of film thickness. Zhang et al50 measured that Zn interstitial and oxygen vacancies caused FM. In the light of these discoveries, there has been an emerging consensus that defects in these systems play critical roles for mediating FM 112, 117. Elfimov et al118 have shown how divalent cationic vacancies can develop ferromagnetism in oxides with the rock salt structure. Bouzerar et al117 extended the above discussion to include a larger class of oxides and concluded that the vacancy creates an extended defect on neighboring sites and that there is an extended magnetic moment. Osorie et al119 however estimated that an unrealistically high cation vacancy concentration would be required to generate sufficient holes in this picture. Chan et al120 further extended the discussion and showed that including the electron correlation effects on the p-orbital near the vacancy generates localization of the hole on an oxygen site and prevents the formation of the moment, as opposed to the case of Elfimov et al118. Adeago et al121 studied the ZnO system with the addition of N and H and showed some interesting features. They121 showed that a number of magnetic defects could be generated in these systems but with the introduction of the correlation effects via a Hubbard type potential on the oxygen atoms, the holes tended to become more localized and the magnetization was moved to oxygen atoms situated further from the cationic vacancy. They121 concluded that the moment was generated by the Zinc vacancy, and the FM was related to the delocalization of the holes, and that introduction of electron correlation caused hole localization and destruction of FM. They thus concluded that the observed FM could not be explained solely on the basis of these models. Peng et al122 in contrast argued that d0 FM in the ZnO type systems was a feature related to the localized nature of the 2p states in such systems. They argued that large exchange interactions could occur in the valence bands of these materials when the Fermi level was made to move away from the band maximum as by hole doping due to intrinsic defects. They122 argued that the threshold doping level required for FM could be made smaller by acceptor doping and also interestingly by quantum confinement effects. We thus see that d0 and related effects on ZnO system are intimately related to the nature of the defects and the grain or particle size of the system. Different models have been presented in literature suggesting the origin of magnetism in oxide based semiconductor systems. Fernandes et al103, 104 proposed a model using density function calculations and our results are closer to this group. This group has demonstrated that 12/20

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both cationic (cerium) vacancies and oxygen vacancies played a role in inducing ferromagnetism in oxide based ferromagnetic system, e.g., CeO. We focus our attention on this model as they103, 104 discussed that band-structure calculations indicated that both the cationic and anionic vacancies were consistent with a ferromagnetic state for all the investigated vacancy concentrations. In the light of above discussion, we emphasized that the role of complex defect (Zn and oxygen vacancies) is more important for promoting and stabilizing the ferromagnetism in our well crystalline ZnO nanoparticles. Similarly, in our case, cationic vacancies (Zn vacancies) and oxygen vacancies mutually interacted with each other and the net effect of charge carriers (holes in case of cationic vacancies and electrons in case of oxygen vacancies) played a major role in inducing a long range order ferromagnetism in un-doped ZnO crystalline systems. Over all, on the basis of above mentioned literature and our observations, it may be suggested that the complex defects of cationic and anionic vacancies play a major role in net spin and observed ferromagnetism in fully crystalline oxide based semiconducting systems.

4. Summary and Conclusion: Un-doped ZnO nanoparticles have been synthesized via pulse laser ablation of Zn metal foil in Heptanes. The as-synthesized, 400⁰C and 450⁰C air annealed samples showed secondary phases but exhibited the M/H hysteresis response. XRD analysis confirmed the hexagonal structure of 550⁰C air annealed sample. Zinc interstitial, oxygen and Zn vacancies were observed from XPS measurements. Air annealed sample at 550⁰C showed the hysteresis curve which exhibited the presence of ferromagnetism. FC and ZFC curves data of 550⁰C annealed sample showed that magnetization increased by increasing the applied field. Form XPS and magnetic data, we inferred that both oxygen and Zn vacancies are responsible for ferromagnetism in single phase ZnO nanoparticles. We also conclude that complex defects of oxygen and Zn vacancies have a particular role in inducing ferromagnetism in ZnO nanoparticles. Thus, not only oxygen vacancies are responsible for RTFM, but zinc vacancies also contribute to the observed ferromagnetic character in non-magnetic ZnO nanoparticles systems. Overall, there is competition between oxygen and Zn vacancies i.e. intrinsic defects are responsible for promoting and stabilizing RTFM in hexagonal ZnO system. ACKNOWLEDGMENTS Saif Ullah Awan and S. K. Hasanain acknowledge the financial support from the Higher Education Commission of Pakistan under “5000 Indigenous Ph.D fellowship program” and the project “Development and Study of Magnetic Nanostructures.”

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ACCEPTED MANUSCRIPT

Sample-a (As grown) Ablated

Intensity (a.u)

(a)

20 25 30 35 40 45 50 55 60 65 70 75

2 Thetha (Degree)

(101)

(b)

b

(39.052)

38

39

(43.212)

(43.238)

40

41

(110)

d b

43

(103)

44

(200)

(102)

42

(112) (201)

c

(39.026) (39.13)

(002)

Intensity (a.u)

(100)

d

c

20 25 30 35 40 45 50 55 60 65 70 75

2 Thetha (Degree)

Fig. 1: Color online: (a) X-ray diffracted (XRD) patterns of as grown (sample-a) of Zn metal foil ablated in Heptanes. (b) XRD patterns of air annealed sample-b at 400⁰C, sample -c at 450⁰C and sample-d at 550⁰C, inset to (b): closer view of 2θ values from 38-44 degrees.

ACCEPTED MANUSCRIPT

d

0.0009

Sample Holder

0.03

0.0006

emu

0.0003

0.02

a

b

b d

0.0000

a

-0.0003

emu/g

0.01

-0.0006

-0.0009

0.00

-30000

-20000

-10000

0

10000

20000

30000

H(Oe) 0.01

-0.01 0.00

-0.02 -0.01

-0.03

-4000

-30000 -20000 -10000

0

10000

0

20000

4000

30000

H(Oe)

Fig. 10: (color online) M/H loop of sample-a (un-annealed), sample-b (400⁰C Air annealed) and sample-d (550⁰C annealed) at 50K, upper inset: Rawa data of samples (a, b and d) and sample holder, lower inset: coercivity and remanence at 300K of samples (a, b and d).

ACCEPTED MANUSCRIPT

0.0039

(b)

0.0049

FC (500 Oe) ZFC (500 Oe)

0.0048 0.0047

ZFC at 500 Oe

0.0046

M(emu/g)

emu/g

0.0036

0.0050

(a)

0.0033 FC(1000 Oe) ZFC (1000 Oe)

0.0045 0.0044 0.0043

FC at 500 Oe

0.0042

0.0030

0.0041 0.0040 0.0039

0.0027

0.0038

50

100

150

200

250

300

50

Temp (K)

250

300

(d)

Oe ZFC at 500 OeZFC at 1000 (c) 0.0032 FC at 500 Oe FC at 1000 Oe

0.0028

0.0030

M(emu/g)

M(emu/g)

200

T(K)

0.0034 0.0030

0.0024

150

0.0036

0.0032

0.0026

100

0.0028 0.0026

Fig. 11: (color online) (a) FC and ZFC of sample-a (un-annealed), (b) sample -b (400⁰C air annealed) at 500Oe (c-d) FC and ZFC of sample-d (550⁰C air annealed) at 500Oe and 1000Oe.

0.0024

0.0022

0.0022 0.0020

0.0020 0.0018

0.0018 50

100

50 150

100 200

T(K)

150 250

T(K)

200 300

250

300

ACCEPTED MANUSCRIPT

(201) (112) (200) (103) (110) (102) (101) (002) (100)

Fig.2: (a) BF-TEM micrograph at low, scale bar 100nm (b-d) high magnification TEM images, sclae bar 20nm, 10nm and 5nm respectively (e) FFT image (extracted from HR-TEM-d) and (f) Hough Transform of sample-d.

ACCEPTED MANUSCRIPT

14000

(a)

E2L

6000

(574)A1(LO)

(438) E2H

8000

(333) E2H-E2L

(202) 2E2L

Raman Intensity (a.U)

10000

(1156) [2A1(LO),E1(LO),2LO]

Sample-a (As grown) Ablated

12000

4000

2000 0

200

400

600

800

1000

1200

1400

5000

4000

d c

3000

(1156) (2LO)

6000

(b)

(996)(2TO)

E2L

(202)2E2L

Raman Intensity (a.U)

7000

(333) E2H-E2L (408)E1(TO) (438)E2H (498) 2LA (574)A1(LO) (660)(TA+LO)

Raman Shift(cm-1)

*4 *2

b 2000

0

200

400

600

800

1000

1200

1400

Raman Shift(cm-1)

Fig. 3: Color online (a) Backscattering Raman spectroscopic spectra of (a) sample -a(as grown) without annealing (b) air annealed sample-b at 400⁰C, sample-c at 450⁰Cº, and sample-d at 550⁰C.

6 550C [PL Intensity (arb.unit.)*10 ]

8 10

(a)

sample-a sample-d 6

8 6

4 4 2

2 0 300

400

500

600

700

800

5 un-annealed [PL Intensity (arb.unit.)*10 ]

ACCEPTED MANUSCRIPT

0

Wavelength (nm)

Sample-a

PL Intensity

(b)

1 3

400

500

600

700

Wavelength (nm)

800

Sample-d

(c)

PL Intesnity

4

2 1

400

3

4

500 600 700 Wavelength (nm)

800

Fig.4: (Color online): Room temperature PL emission spectra of sample-a (un-annealed) and sample-d (air annealed at 550⁰C).

ACCEPTED MANUSCRIPT

12 8 4

Zn-3p Zn-3s

C-1s

Zn-2s C-KLL

Zn-2p3/2 Zn-2p1/2

14 12 10 8 6 4

Zn-L2M22M23

2

550C annealed [Counts *105]

16

Zn-L2M45M45

20

16

O-KLL

24

Zn-L3M45M45 O-1s

28

Zn-L3M23M45 Zn-L3M23M23

32

Zn-3d

un-annealed [Counts *104]

36

0

0 0

200

400

600

800

1000

1200

1400

B.E (eV)

Fig.5: (color online) X-ray Photoelectron Spectroscopy (XPS) survey spectra of sample-a (unannealed) and sample-d (550⁰C air annealed) were recorded at room temperature.

ACCEPTED MANUSCRIPT

4

un-annealed (Counts)*10

18

3.5

16 3.0

14

Zn2p1/2

12

2.5

10 2.0

8 6

1.5 1.0

4

20

Zn2p3/2

550C annealed (Counts)*10

4.0

4 1020

1025

1030

1035

1040

1045

1050

1055

B.E (eV)

Fig.6 : (Color online) X-ray Photoelectron Spectroscopy (XPS) of Zn (2p3/2 and 2p1/2) core level spectra of the sample-a (un-annealed) and sample-d (550⁰C air annealed) were measured at room temperature.

ACCEPTED MANUSCRIPT

(a)

Counts/s

Orignal data Fitted data Background

490

Zn-O

Zn-O Zni

Zni

492

494

496

Orignal data Fitted data Background

Counts/s

498

B.E (eV)

500

502

504

(b) Zn-O

Zn-O Zni

Zni

490

492

494

496

498

B.E (eV)

500

502

504

Fig. 7: (color online) The original and deconvoluted Zn L3M4,5M4,5 Auger peaks of (a) sample-a and (b) sample-d.

ACCEPTED MANUSCRIPT

12000 11000

Raw data Fitted data Background O-Zn Vo

(a)

Counts

10000

O-Zn

9000

Vo

8000 7000 6000

527

528

529

530

531

532

533

534

535

536

B.E (eV)

60000

Raw data Fitted data Background O-Zn Vo

(b)

55000

Counts

50000 45000 40000 35000

O-Zn

Vo

30000 25000 20000

527 528 529 530 531 532 533 534 535 536

B.E (eV)

Fig.8: (color online) Deconvoluted O-1s XPS core level spectra of (a) sample-a (un-annealed) and (b) sample-d (550⁰C air annealed) were recorded at room temperature (symbols are defined in text).

ACCEPTED MANUSCRIPT

0.020

0.0009

0.015

0.0006

emu/g

0.005 0.000

emu

0.0003

0.010

a d b

b

Sample Holder

d

0.0000

a

-0.0003

-0.0006

-0.0009 -30000

-20000

-10000

0

10000

20000

30000

a

H (Oe)

b 0.006

-0.005 -0.010

0.000

-0.015

-0.006

-0.020

-6000

-30000 -20000 -10000

0

d

-4000

-2000

0

2000

4000

6000

10000 20000 30000

H(Oe)

Fig. 9: (color online) M/H loop of sample-a (un-annealed), sample-b (400⁰C Air annealed) and sample-d (550⁰C annealed) at 300K, upper inset: Rawa data of samples (a, b and d) and sample holder, lower inset: coercivity and remanence at 300K of samples (a, b and d).

ACCEPTED MANUSCRIPT

Highlights 

Well controlled synthesis of nanocomposites and nanoparticles.

   

XRD identify mutiple and single phase systems. PL and XPS spectra confirmed the presence of Zni, VZn and Vo defects. Complex defects may be induce , promote and stabilize RTFM.

ZnO is suitable for Spintronics.