Correlating the magnetism and gas sensing properties of Mn-doped ZnO films enhanced by UV irradiation

Abstract

In this study, we report on the correlation between the magnetism and gas sensing properties of Mn-doped ZnO films grown via aerosol spray pyrolysis. The evolution of the structure, morphology, optical properties, and chemical state of ZnO with the Mn concentration was also investigated. ZnO doped with Mn (0.1 at%) demonstrated room-temperature ferromagnetism (RTFM) due to the uncompensated surface spins primarily originating from structural defects and oxygen vacancies (V_O) on the surface, which act as active sites for the adsorption of oxygen species. The undoped ZnO structure revealed both FM and paramagnetism (PM) at the near surface of the film. Increased Mn doping destroyed the RTFM ordering due to improved PM features induced by Mn clusters on the ZnO surface and reduced amount of V_O on the surface. However, ZnO films doped with Mn (0.1 at%) exhibited an improved sensing response to oxidizing gases compared to their counterparts, showing that films with RTFM with no PM contribution exhibit improved sensing properties. These analyses revealed that the nature of the film surface plays a substantial role in both their magnetic and sensing behaviors.

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1. Introduction

The development of gas sensors capable of reliably monitoring toxic and combustible gases has become very important due to concerns regarding environmental pollution and safety requirements in industry. It is well-known that NO produced in combustion processes is harmful to human health and the environment. In addition, NO is not only a precursor of acid rain but it also causes depletion of the ozone.¹ In the presence of excess O₂, NO is easily oxidized to nitrogen dioxide (NO₂). Qureshi et al.² reported that frequent exposure to NO and NO₂ gases can cause pulmonary edema and fatality. In contrast, CO₂ is a colorless, odorless, non-flammable gas. CO₂ is regarded as the most prominent greenhouse gas, which is emitted into the air as humans exhale and released during industrial and domestic burning of fossil fuels for energy production and deforestation activities around the planet. In 2011, CO₂ accounted for approximately 84% of all United States (U.S.) greenhouse gas emissions from human activities.³ Therefore, reliable sensors for detecting lower concentration of NO and CO₂ gases at room temperature are required. Metal-oxide semiconductors (MOS) have stimulated significant research interest in various fields including spintronics, solar cells, and gas sensors. Among the MOS gas sensors, ZnO-based gas sensors have attracted much attention for a long time because of their advantageously high sensitivity and simplicity in fabrication.^4,5 This material possesses a wide and direct band gap (3.4 eV), and a large exciton binding energy of 60 meV at room temperature. ZnO crystallizes in a wurtzite structure and exhibits n-type electric conductivity. In addition, ZnO is a highly sensitive material for flammable or toxic gas detection.^6,7 In the recent past, there has been an increasing degree of awareness of the use of transitional metals (TMs) as dopants for ZnO to customize its electrical, optical, and sensing properties.^8–11 Among all the TM-doped ZnO materials, Mn-doped ZnO has stimulated considerable attention primarily due to its high thermal solubility in ZnO.¹² However, few studies have been carried out to improve the performance of ZnO by TM doping, which would enable this material to act alone as a high-performance gas sensor. A significant drawback of these traditional metal-doped semiconductor gas sensors is that they are typically activated at a high temperature. Therefore, it is imperative to overcome this requirement for heating and the corresponding power consumption, which also leads to undesirable long-term drift problems caused by sintering effects at the MOS grain boundaries. Researchers have reported practical approaches, such as doping with noble metals and exposure of sensors to UV light, to reduce the operating temperature.^13–16

Additionally, among diluted magnetic semiconductors (DMS), Mn-doped ZnO was the most attractive because Mn has the highest possible magnetic moment at a valence state of 2⁺.¹⁷ However, local spin density approximation (LSDA) calculations suggested that Mn³⁺ may be more favored in the ground state of the system.¹⁸ Theoretical calculations predicted that Mn-doped ZnO exhibits ferromagnetism (FM) with a Curie temperature well above room temperature (RT).^19–22 In DMS, the non-magnetic host ions are partially substituted by magnetic ions, which exploit the spin in magnetic materials along with the charge of electrons in semiconductors. Nonetheless, the origin of FM in TM-doped ZnO systems is not yet understood, and it is debated whether this is caused by defects other than the metal dopants. Recently, Straumal et al.^23,24 analyzed a large number of experimental publications concerning the specific grain boundary area, which is the ratio of the grain boundary (GB) area to the grain volume, s_GB. They observed that GBs are the key factor that determines the presence or absence of FM in ZnO. According to Straumal et al.,²³ FM only appears if s_GB surpasses a certain threshold value, s_th.

However, both the gas sensing mechanism and the magnetization depend primarily on defects. It is well known that oxygen vacancies (V_O) play a vital role in the gas sensing mechanism,^25,26 and similarly, it is also understood that the presence of V_O may lead to RTFM in MOS. Currently; there are only a few reports on the correlation between magnetism and sensing. Since V_O is considered a common parameter for gas sensing and magnetism, in this work, we correlate the effect of magnetism on the sensing properties of the undoped and Mn-doped ZnO films towards oxidizing gases, such as NO and CO₂. We show in detail that the nature of the magnetism greatly influences the gas sensing response of both undoped and Mn-doped films grown via aerosol spray pyrolysis. In addition, a correlation between the optical constants, crystal quality, and magnetism was also investigated using spectroscopic ellipsometry (SE). The origin of the defects responsible for mediating ferromagnetism and gas sensing properties in Mn-doped ZnO films was investigated using electron spin resonance (ESR), X-ray photo-electron spectroscopy (XPS) and photoluminescence (PL) spectroscopy. Additionally, the time-of-flight secondary ion mass spectrometry (TOF-SIMS) was also used to investigate the presence of Mn dopants and the homogeneity in the crystal lattice of ZnO in detail.

2. Experimental details

2.1. Synthesis of un-doped and Mn-doped ZnO films

Undoped and Mn-doped ZnO films were deposited on the surface of ultrasonically cleaned Corning 1737 F glass substrates (25.4 × 25.4 mm²) using an in-house-developed aerosol spray pyrolysis unit. The 0.1 M dehydrate Zn(NO₃)² (purity > 99%) and Mn(CH₃COO)₂·4H₂O (purity > 99.9%) purchased from Sigma-Aldrich, were used as precursors to grow Mn-doped ZnO films at various Mn concentrations (0.1–2.0 at%). Prior to deposition, the precursors were dissolved in distilled water and stirred thoroughly for 30 min to yield a clear and homogeneous solution. The undoped and Mn-doped films were homogenously deposited at a constant flow rate of 300 ml h⁻¹ using a nozzle assisted by a nitrogen carrier gas at 0.5 bar, over a corning glass substrate heated at 300 °C. The nozzle to substrate plane distance was maintained at the optimal value of about 28 cm.²⁷ The films were grown for 30 min.

2.2. Characterization methods

The surface morphology and cross-sectional analyses were carried out using a scanning electron microscope (SEM, Zeiss Auriga, Germany). The structural properties of the undoped and Mn-doped ZnO films were characterized using a PanAnalytical X′ pert PRO PW 3040/60 X-ray diffractometer (The Netherlands) with a Cu Kα (λ = 0.154 nm) radiation source. The measurements were recorded at 45.0 kV and 40.0 mA from 2θ = 25 to 90°. The film thicknesses (Table 1) and optical constants were determined by fitting the model function to the measured data using a WVASE32 program developed by J. A. Woollam Corporation (USA).²⁸ The electrical properties were measured using a Hall Effect instrument (Ecopia HMS-3000) at room temperature. The optical properties were characterized using room temperature PL (Jobin-Yvon NanoLog spectrometer) at an excitation wavelength of 325 nm. TOF-SIMS⁵ (IONTOF Gmbh instrument) analyses were carried out using a pulsed 30 kV bismuth primary ion beam operated at a DC current of 1 pA, and pulse repetition rate of 10 KHz (100 μs) was used to acquire chemical images of the Mn–ZnO in the positive secondary ion polarities. The elemental mapping was performed using a Cs⁺ 0.4 pA current that was rastered over an area of 100 × 100 μm of the surface to prevent edge effects. The resolution employed to record the surface images was 512 × 512 square pixels. The base pressure of the analysis chamber was maintained at ∼2 × 10⁻⁹ mbar during the experiment. XPS analyses were carried out using a PHI 5000 Versaprobe-Scanning ESCA Microprobe. A low-energy Ar⁺ ion gun and low-energy neutralizer electron gun were used to minimize charging on the surface. Monochromatic Al Kα radiation (hν = 1486.6 eV) was used as the excitation source. A 25 W, 15 kV electron beam was used to excite the X-ray beam, which was 100 μm in diameter. For the higher-resolution spectra, the pass energy of the hemispherical analyzer was maintained at 11.8 eV (C 1s, O 1s, and Zn 2p) for 50 cycles. The measurements were performed using either 1 eV per step and 45 min acquisition time (binding energies ranging from 0–1400 eV) for survey scans or 0.1 eV per step and 20–30 min acquisition times for the high-resolution scans. The pressure during acquisition was typically below 1 × 10⁻⁸ Torr. The same measurements were repeated after the surfaces were sputter cleaned for 30 s using an Ar⁺ ion gun (2 kV energy ions). The microwave absorption measurements were carried out using an X-band ESR spectrometer (JES FA 200, JEOL, Japan) equipped with an Oxford ESR900 gas-flow cryostat and a temperature controller (Scientific instruments 9700, USA). The films were mounted in the cavity centre at the position where the microwave magnetic field was at a maximum. The DC static field H_DC was slowly swept between 0 and 500 mT. The microwave power was maintained at 5 mW to prevent saturation. The DC field was modulated with a superposed AC field whose amplitude was varied between 1 mT and 6 mT at a 100 kHz frequency. The microwave response was measured as a derivative of the microwave absorption signal. The measurements were carried out at 298 K.

Table 1 Summary of the crystallite sizes and lattice parameters for the (002) orientation

Films	2θ(002) (°)	Crystallites size (nm)	Lattice strain, ε (×10^−3)	a (Å)	c (Å)	V (Å^3)
Undoped ZnO	34.45	29.09	1.19	3.2489	5.2063	47.5904
0.1 at% Mn	34.67	26.41	1.31	3.2501	5.1743	47.3329
0.5 at% Mn	34.54	22.97	1.51	3.2549	5.1940	47.6535
2.0 at% Mn	34.53	21.05	1.64	3.2729	5.1853	48.1013

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2.3. Measurement of the gas-sensing performance of the Mn–ZnO films

The undoped and Mn-doped ZnO gas sensing films were tested for their response to either NO or CO₂ using a lab-built computer-controlled characterization system. The gas chamber was connected to the gas inlet line from the mass flow controllers. A gas exhaust tube was located on the opposite side of the chamber. The flow-through technique was used to test the gas sensing properties of the undoped and Mn-doped gas sensing films. All of the measurements were conducted in a temperature-stabilized sealed chamber at 25 °C under a controlled relative humidity (RH) of approximately 20%. The operation temperature of the sensor was measured by a calibrated K-type thermocouple mounted on the device. The test chamber, which had a volume of 5 l, was composed of a glass. The ppm levels (10–100 ppm) of the target gases were controlled by the flow rate ratio between the synthetic air and target gas. A constant flow rate of 500 sccm was maintained during the measurements. The resistances of various sensors were continuously monitored with a computer-controlled system utilizing the voltage-amperometric technique at a DC bias of 10 V with current measurements through a picoammeter. National Instruments Labview™ v6.0 software was used to control the mass flow controllers and record the gas concentration. A Keithley sourcemeter-2420 was used to acquire the resistance data. Photo-reduction (activation) measurements were carried out using a UV mercury pencil lamp with an average intensity of 4 mW cm⁻² at a wavelength of 254 nm. The UV lamp was placed at a distance of 30 mm from the surface of the sensing films.

3. Results and discussion

SEM images of the undoped and Mn-doped ZnO products are shown in Fig. 1. The undoped ZnO films revealed nanorods grown perpendicular to the substrate entangled to one another to form “spaghetti-like” or coarse-grain (large-grains) structures containing voids and interspaces. These structures, which had an average diameter in the range of 20–35 nm, were previously observed by SEM.²⁸ Upon introducing Mn ions (0.1 at%) into the ZnO host, nanoparticles were observed across the film, leading to more voids and interspaces (see Fig. 1b). These voids and interspaces may be responsible for the improved gas sensing response by providing more positions for gas molecules. The low magnification (results not shown) showed “doughnut-like” structures which were made from nanoparticles or nanograins, and were slightly finer than those of pure ZnO. By increasing the Mn doping level (0.5–2 at%), the size of the nanoparticles decreased (see Fig. 1c and d).

Fig. 1 SEM images of (a) undoped ZnO, (b) 0.1 at%, (c) 0.5 at% and (d) 2 at% Mn-doped ZnO films.

Fig. 2 depicts the cross-sectional SEM analysis of the undoped and Mn-doped films. The cross-sectional view in Fig. 2a shows that columnar structures indeed grew perpendicular to the substrate with a parallel c-axis orientation. As shown in Fig. 2a, the shape of the columnar structures or grains appears to be elongated. The typical column size was estimated to be approximately 20–40 nm. It is important to note that, by doping the ZnO with 0.1 at% of Mn, the columnar structure was maintained and was visible across the film edge (Fig. 2b). At a higher doping level (0.5–2.0 at%), no columnar structure was observed.

Fig. 2 Cross-sectional SEM micrographs of (a) undoped ZnO, (b) 0.1 at%, (c) 0.5 at%, and (d) 2 at% Mn-doped ZnO films.

The crystal structure and orientation of the as-deposited and Mn-doped ZnO films grown for 30 min were investigated using XRD, and the results are shown in Fig. 3. These XRD patterns indicate that the films are polycrystalline in nature with a hexagonal wurtzite structure. The positions were indexed to the (100), (002), (101), and (102) planes, and the peak positions were determine to be consistent with the JCPDS [card 36-1451] of ZnO. No traces of manganese oxides or any binary zinc manganese phase were observed in any of the samples up to 2.0 at% Mn doping. This result clearly indicates that the speculated substitution of Zn²⁺ ions by Mn²⁺ ions did not alter the inherent Wurtzite structure of ZnO. A slight shift in the lattice constant of doped ZnO films was observed compared to undoped ZnO film, as shown in Table 1. This result is due to the larger ionic radius (0.83 Å) of the Mn²⁺ ions that replaced the smaller Zn²⁺ ions, which have a smaller ionic radius (0.74 Å). The shifting of the peak to a lower angle indicates substitution of larger Mn²⁺ in the Zn²⁺ site. As the Mn content increased, the diffraction intensity from the ZnO (002) peaks decreased, which indicates that the film crystallinity deteriorated with Mn incorporation and this in agreement with the SEM results. As shown in Fig. 3, small peak broadening also occurred with an increase in the Mn content. This small degree of broadening is the result of increasing strain (see Table 1) in the film due to Mn incorporation in the Zn lattice site²⁹ and a corresponding increase in the volume of the unit cell with an increase in the Mn doping level. The average crystallite size of the Mn-doped ZnO films was determined using the Scherrer equation,³⁰ and the lattice strain (ε) was calculated using the following relation:

where, β_2θ is the full width at half maximum intensity (FWHM) of the corresponding peak and θ is the diffraction angle. As the Mn-doping concentration increased, the crystallite sizes related to the (002) peak decreased and ε increased, as shown in Table 1.

Fig. 3 XRD patterns of the undoped and Mn-doped ZnO films.

The typical XPS wide survey spectra of undoped and Mn-doped ZnO films showed that the as-prepared materials are pure without any impurities, showing only three main elements (i.e., O, Zn, and Mn, see ESI, Fig. S1†). High-resolution XPS analyses of Zn 2p were carried out to investigate how the electronic structure of Zn changes upon Mn doping of the Zn site. The Zn 2p_3/2 XPS spectra of undoped and Mn-doped ZnO films are shown in Fig. 4. Based on the results presented in Fig. 4a, the undoped ZnO exhibited two symmetric peaks located at 1022.14 and 1047.61 eV. However, upon doping with Mn, the ZnO symmetric peaks shifted to 1021.8 and 1045.24 eV after fitting with a Gaussian, which is attributed to the partial substitution of Zn in the ZnO lattice by Mn²⁺ ions in the Zn–Mn bonding structure.³¹ Parts (d) and (e) of Fig. 4 indicate that the oxygen 1s (O 1s) spectra were asymmetric, indicating that multi-component oxygen species cannot be ruled out in the near-surface region of all these films. The peaks of the spectra fitted by two Gaussians were located at approximately 531.78 and 532.72 eV. The high-energy peak centered at approximately 532.72 was due to either the chemisorbed oxygen of the surface in the form of –OH, –CO₃, absorbed H₂O, or absorbed O₂ or primarily due to surface contamination. The medium binding energy component centered at approximately 531.78 eV (O_b) corresponded to O²⁻ ions in the oxygen deficient regions within the ZnO matrix and/or Zn–OH groups.³² Therefore, the changes in the intensity of the O_b relative area may be related to the concentration of oxygen vacancies (V_O). Based on the results in Fig. 4 and Table 2, the O_b peak was present in all of the samples and substantially decreased with Mn dopants up to 2.0 at%.

Fig. 4 XPS deconvoluted spectra of (a–c) Zn 2p_3/2, and (d −e) O 1s core level of the undoped and Mn-doped ZnO films.

Table 2 Chemical state of oxygen and zinc in undoped and Mn-doped ZnO films

Films	Oa position (eV)	Ob position (eV)	Relative intensity of Oa (%)	Relative intensity of Ob (%)
Undoped ZnO	532.72	531.47	79.35	21.65
(0.1 at%) Mn–ZnO	532.40	531.18	63.13	36.87
(2.0 at%) Mn–ZnO	532.27	531.15	94.47	5.53

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The valence band of Mn was investigated by measuring the Mn 2p XPS spectra shown in Fig. 5. These two weak XPS peaks at binding energies of 641.7 eV and 657.8 eV correspond to Mn 2p_3/2 and Mn 2p_1/2, respectively. The peak centered at approximately 641.7 eV was due to the presence of Mn in its Mn²⁺ oxidation state, which is in good agreement with the value reported in the literature.³³ Therefore, these results indicate that the Mn dopant ions were indeed in the divalent state, which is evident that Mn as a dopant substitutes for part of Zn in the ZnO lattice.³⁴ This result implies that, with the addition of Mn²⁺ ions, the bond configuration and lattice structure in the ZnO crystal lattice deteriorated, which is consistent with the XRD results shown in Fig. 3. Therefore, at higher doping levels (2.0 at% Mn), the intensity of the Mn²⁺ peak decreased probably because of segregation, and a satellite peak appeared at approximately 663.10 eV.

Fig. 5 Deconvoluted core level XPS spectra of Mn 2p of the (a) 0.1 at% and (b) 2.0 at% Mn-doped ZnO films.

To investigate the presence of Mn dopants and their homogeneity distribution in the crystal lattice of ZnO, TOF-SIMS chemical imaging analyses were performed, as shown in Fig. 6. Only elements of interest (e.g., Zn and Mn) are shown. Parts (a) to (d) of Fig. 6 show the positive ion mapping of pure ZnO film during depth profiling after 67 scans while Fig. 6d is the overlay of Si⁺, Zn⁺, and ⁶⁶Zn⁺, wherein each particle is clearly observed. The Si⁺ signal arose from the Corning glass substrate. The micrographs depict the Zn⁺ and ⁶⁶Zn⁺ elemental or chemical distribution across the surface. The results also confirm the formation of “spaghetti-like” or nanograins structures as observed by SEM. The 0.1 at% Mn-doped sample in Fig. 6e and f also showed the Mn⁺ signal, confirming that Mn was homogeneously incorporated in the ZnO surface. Fig. 6g shows that the Mn was evenly distributed across the surface within an area of 100 μm × 100 μm, although the signal was weak for the lower doping concentration. In addition, the “doughnut-like” structures were also visible, as previously observed in SEM results. The 0.5 at% Mn-doped sample also confirmed the homogenous distribution of Mn in the ZnO matrix. The overlay image shows that, the Mn ions were visible at 0.5 at% doping level. At a higher doping level (1.0 at%, parts (m) to (p) of Fig. 6), the Mn⁺ ions were not uniformly distributed; they were mostly observed on the “doughnut-like” structures surface, which are more visible in the 3D images shown in the ESI (Fig. S2†). The arrows on the overlay image, Fig. 6m, indicate the uncovered Corning substrate.

Fig. 6 TOF-SIMS images of (a–d) pure ZnO, (e–h) 0.1 at% Mn–ZnO, (i–l) 0.5 at% Mn–ZnO, (m–p) 1.0 at% Mn–ZnO.

The absorption coefficient (α) of the Mn-doped ZnO films with different compositions was calculated from the obtained k values using the following relation:

where k is obtained from the absorption curve and λ is the wavelength of light. Therefore, the optical energy gap (E_g) was derived by assuming a direct transition between the edges of the valence and conduction bands, where the variation in the absorption coefficient with the photon energy (hv) is given by Tauc's relation:³⁵

(αhν)² = A(hν − E_g) (3)

where A is a constant. Therefore, the obtained results show a direct electronic transition across the band gap of the films (Fig. 7). The extrapolation of the linear part of the curves to the photon energy axis indicates that the optical band gap varied in the range of 3.1–3.2 eV as a function of the Mn concentration. The results in Table 3 indicate that the E_g decreased (except for 0.5 at%) with an increase in the Mn concentration. A decrease in the band gap is typically caused by a shift in the energy of the valence and conduction bands resulting from electron–impurity and electron–electron scattering effects.³⁶ The estimated E_g values from the spectroscopic ellipsometer are consistent with those extracted from the UV-Vis absorbance spectrum (see the Table 3 and ESI, Fig. S10†). A band gap of 3.2 eV was observed by Kulandaisamy et al.³⁷ on pure ZnO. Upon doping with cobalt, their band gap decreases up to 2.79 eV.

Fig. 7 Plot of (αhv)² as a function of the photon energy (hv) for Mn–ZnO films.

Table 3 Summary of the surface roughness, carrier concentration and optical band (E_g) of the Mn-doped ZnO films^a

Films	Thickness (nm)	Roughness (nm)	Carrier concentration (cm^−3)	Mobility (cm^2 V^−1 s^−1)	E^*g (eV)	E^**g (eV)
a E^*g was estimated from spectroscopic ellipsometer measurements. E^**g was estimated from UV-Vis measurements.
Undoped ZnO	173.01	11.63	5.10 × 10^16	6.58	3.23	3.24
(0.1 at%) Mn	97.60	3.99	2.91 × 10^17	48.16	3.14	3.23
(0.25 at%) Mn	96.51	3.68	4.90 × 10^15	1.79	3.12	3.22
(0.5 at%) Mn	90.92	2.61	3.71 × 10^15	1.15	3.22	3.18
(2.0 at%) Mn	90.70	2.42	4.40 × 10^14	0.12	3.11	3.21

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The electrical properties of the undoped and Mn doped ZnO films with different Mn concentrations measured using Hall effect system are listed in Table 3. As listed in Table 3, the undoped shows a carrier concentration of 5.10 × 10¹⁶ cm³. However, upon doping with 0.1 at% Mn, a slight increase in carrier concentration is observed. As the Mn doping increases, the carrier concentration and mobility decrease. At lower concentration (0.1 at% Mn), oxygen vacancies increase (see Fig. 8a, PL analyses), which results to an increase in carrier concentration. The oxygen vacancies, as a shallow n-type dopant play an important role in ZnO conductivity.³⁸ Since, the mobility of the charge carriers depends on the scattering mechanisms in the film; therefore, the decrease in electron mobility in Table 3 is probably due to the increased impurity scattering centers by doping more Mn.³⁹

Fig. 8 Room temperature PL spectra of (a) the un-doped and Mn-doped ZnO films with various Mn contents and (b) corresponds to an inset of (a), (c and d) Gaussian fits of each PL spectrum. The black lines and red solid dots in (c and d) correspond to the experimental results and Gaussian fits, respectively.

To investigate the presence of defects in Mn–ZnO films, PL measurements were performed as shown in Fig. 8a. The undoped ZnO film exhibited a broad emission peak at approximately 380–550 nm. The absence of the sharp near band edge (NBE) at 385–390 nm may be due to higher defects possessed by these structures. In general, the reported defects present in ZnO-based nanostructures are oxygen vacancies with different charged states, Zn vacancies, Zn interstitials, and adsorbed molecules.^40,41 Previous studies indicated that the resulting defect-related emissions for these defects typically occur near the blue-green (approximately 480–550 nm), yellow (approximately 550–610 nm), and orange-red (approximately 610–750 nm) regions.⁴¹ Therefore, to obtain a detailed understanding of the effect of individual defects on the magnetic properties of the ZnO nanostructures, a three-peak (Gaussian) fitting method of the broad visible emission was adopted. The de-convoluted spectra of the undoped ZnO in Fig. 8b exhibit three peaks at approximately 390, 441, and 511 nm. Therefore, the UV emission at approximately 390 nm was due to exciton recombination through an exciton–exciton collision process. The emission peak at 441 nm may be due to an electron transition from the Zn interstitial energy level to the valence band.⁴² In addition, the green PL emission band at 511 nm may be assigned to oxygen vacancies (V_O) in ZnO. This luminescence may be due to the recombination of photogenerated holes with singly ionized oxygen vacancies .^42,43 Based on the results from parts (a) to (d) of Fig. 8, the intensity of the emission band related to was higher for the 0.1 at% doped material than its counterparts, which indicates that the 0.1 at% doped material contained more . The 588 nm emission band is typically attributed to the ⁴T₁(4G) → ⁶A₁(6S) transition within the 3d shell of the Mn²⁺ ion.⁴⁴ Here, the ⁴T₁(4G) → ⁶A₁(6S) emission intensity increased and reached a maximum at a Mn²⁺ concentration of 0.5 at% (Fig. 8b). In addition, the observed sharp decrease in intensity of the Mn²⁺ emission for ZnO films doped at a concentration above 0.5 at% Mn may be attributed to the concentration quenching effect due to the pairing or coagulation of the Mn ions, which is consistent with the XPS and TOF-SIMS results (Fig. 5a and 6).

To further confirm the presence of defects, EPR analyses of the Mn–ZnO films were carried out at room temperature and the results are shown in Fig. 9. The spectrum in Fig. 9a exhibits two absorption peaks associated with paramagnetism (PM) and ferromagnetic resonant (FMR) fields at 150 and 320 mT, respectively. The ESR results indicate that all the films possessed PM features at 150 mT except for the 0.1 at% Mn-doped material. However, at higher magnetic fields (>320 mT), a FM signal was observed for both the undoped and Mn-doped films up to 1.0 at%, which is consistent with the magnetization measurements shown in the ESI (Fig. S8†). This feature decreased as the PM features (see inset, Fig. 9b) increased with Mn concentration. Mn-related oxides, such as MnO, MnO₂, and Mn₂O₃, are known to be antiferromagnetic while Mn₃O₄ is ferromagnetic below 45 K.⁴⁵ Therefore, the observed room temperature (RT) FM in the current work excludes the presence of the above-mentioned oxides, and the RTFM is only attributed only to the incorporation of Mn²⁺ ions into the ZnO host lattice and the high V_O content, as depicted in the PL analyses (Fig. 8), which is in good agreement with our previous results.⁴² Furthermore, Straumal et al.^23,24 have shown that the topology of a grain boundary network in the material and the shape of the grain boundaries have a significant effect on FM. They found that saturation magnetization changed drastically with the topology of the grain boundary network in Mn-doped ZnO. Therefore, the observed FM for undoped and 0.1 at% Mn-doped ZnO film may also be due to the shape of nanograins, which appeared to be elongated according to SEM. Additionally, the lower FMR signal observed for the undoped sample compared to 0.1 at% may be due to larger nanograins (coarser grain). Tietze et al.⁴⁶ have reported that a film with an average grain size of 31 nm (i.e. finer grains) has larger magnetization compared to the film with an average grain size of 65 nm (coarser grains) because the value of s_GB was 5.32 × 10⁷ m⁻¹ for the fine-grain sample, while the value of s_GB was 2.65 × 10⁷ m⁻¹ for the coarser-grain sample. Additionally, Kumar et al.⁴⁷ observed FM on Fe (0.06)-doped ZnO showing a s_GB value of 41 × 10⁶ m⁻² m⁻³. Thus, for our system (0.1 at% Mn–ZnO), s_GB value is well above the threshold value, viz. s_GB = 6.24 × 10⁷ m⁻¹. Moreover, the FM spectra for both the undoped and Mn-doped films showed sextet hyperfine structures, which are related to interactions between the electronic state and the nuclear spin angular momenta in Zn²⁺. A g-factor of about 1.9965–2.0013 is observed for all the samples, which was related to V_O and V_Zn.^{7,42,43,48,49}

Fig. 9 ESR spectra of the undoped and Mn-doped ZnO films and (b) corresponding spectra of the inset in (a).

At higher doping level (i.e., 2.0 at% Mn), a magnetic transition occurs, and a signal related to FM transforms to a sharper peak related to PM (see the circle, Fig. 9a). This indicates that the film is dominated by the same spins, which are randomly oriented clusters with some ferromagnetic features that are still present even at relatively high Mn concentration. The observed PM was most likely caused by tiny deposited Mn particles that tend to agglomerate among themselves on the surface of ZnO at higher doping levels and suppress the number V_O (see XPS and PL analyses, Fig. 4 and 8), which reduces the FM nature of the sample. Wang et al.⁵⁰ reported that Mn-doped ZnO films with a higher Mn content resulted in a shorter magnetic coupling distance between Mn atoms, which is energetically favorable for Mn²⁺ ions to obtain antiferromagnetism.

Fig. 10 illustrates the variations in the response of Mn-doped ZnO gas sensing films exposed to CO₂ gas at room temperature. In stage A, the films are photo-reduced using a UV light. The increase in conductance (stage A) was due to the generation of free carriers within the film and photodesorption of surface species with a subsequent thinning of the electron depletion layer near the film surface. In stage B, when the UV lamp was turned off, the conductivity decreased due to a limited adsorption process on the film surface. In stage C, when the sensing films were exposed to the oxidizing gas (i.e., CO₂) at room temperature, the conductance decreased further with a distinct exponential decay corresponding to the sensing gas concentration.

Fig. 10 Current vs. time curve of the Mn–ZnO sensing film for various Mn concentrations towards CO₂ at room temperature.

Fig. 11 shows the reproducibility of the temporal response of Mn–ZnO sensing films exposed to CO₂ at room temperature. The sensors maintained their initial response amplitude without a clear decrease up on four successive sensing tests to CO₂, demonstrating that these sensors had good repeatability. This behavior was also observed for NO gas (results not shown). Therefore, all these observations specify that Mn–ZnO sensing films are good candidates for the development of high-performance CO₂ sensors operating at room temperature.

Fig. 11 Sensing performance and repeatability of Mn–ZnO sensing films with various Mn concentrations towards CO₂ at room temperature over a long working period.

Fig. 12a displays the gas sensing responses of the undoped and Mn-doped gas sensing films exposed to various oxidation gases at room temperature. Due to its intrinsic defects, the ZnO material showed an n-type conducting feature. Although Mn doping causes an increase in film conduction owing to the sp–d exchange interaction, the addition of Mn will not change the conduction type. Therefore, in the sensing process, the adsorption of oxidizing gases (i.e., CO₂ or NO) on the surface will increase the resistance of the film. In this study, the gas sensing performances were characterized in terms of the gas-sensing response (S), which is defined as:

where R_gas is the resistance measured in the presence of the gas and R_air is the resistance measured in air (in the absence of a reactive gas). Therefore, based on the results presented in Fig. 12, the sensing response increased when doping with 0.1 at% Mn. The increase in response was due to the incorporation of Mn²⁺ ions (see XPS and PL analyses, Fig. 6 and 8), which resulted in enhancement of the surface activity. This enhancement was due to the lower ionization energy of Mn than ZnO.⁵¹ In addition, the activation energy of the gas surface chemisorption decreased, which led to an improvement in the gas absorption. However, when the doping level increased, the gas sensing response decreased for all of the gases. The reason for the decrease in the sensing response of the Mn-doped sensing films may be a decrease in the number of V_O and defects in this film compared to the 0.1 at% Mn-doped ZnO film. V_O and defects are the most favorable adsorption sites for oxygen species, which can enhance the possibility of interaction with gas molecules. Moreover, the low response may be due to that tiny deposited Mn particles tend to agglomerate among themselves at higher doping levels, which leads to rough agglomerated sediment and a lower number of specific active sites as well as poor Mn dispersion on the ZnO support.

Fig. 12 (a) Response of the Mn–ZnO sensor exposed to various gases (left side) and the FMR signal (right side) as a function of the Mn concentration, (b) the response curves of 0.1 and 0.25 at% Mn-doped ZnO films exposed to different concentrations of NO under UV-light irradiation and (c) the response versus time for the 0.1 at% Mn exposed to various concentrations of NO gas.

This result indicates that a higher concentration of Mn doping in ZnO decreases rather than increases number of reactive sites. Based on the results in Fig. 12a, we infer that the 0.1 at% Mn-doped ZnO film, which showed a higher FM signal, was more sensitive to oxidation gases than its counterparts. Therefore, the gas sensing and magnetic studies imply that Mn may form clusters on the ZnO surface and reduce the FM nature (due to contribution from PM) of the sample, and this result was observed in XPS and PL analyses where the amount of V_O decreased upon increasing the Mn concentration as dopants on the ZnO surface increased. Therefore, this is the reason for the reduced response of ZnO containing higher level of Mn dopants (0.25–2.0 at%). Additionally, the sensor response to NO gas was obviously superior to the response to CO₂ gas at lower doping concentrations (0.1 at% Mn). Based on these results, we conclude that, the selectivity of the Mn–ZnO gas sensing films toward NO gas is excellent at lower Mn doping levels.

Fig. 12b demonstrates the response magnitude versus the concentration of NO gas in the range of 10–100 ppm under UV-light. Fig. 12b shows that the 0.1 at% Mn sensing film had a great response to NO gas compared to the 0.25 at% and 2.0 at% Mn-doped ZnO films (results not shown). Its response increased rapidly with the NO gas concentration. The fabricated sensor exhibited a clear response of approximately 58.6 even for a very low NO gas concentration of 10 ppm. Additionally, at 100 ppm NO, the highest response of approximately 86% was observed without any saturation trend. As depicted in the figure, the sensor had a clear linear dependence curve in the range of 10 to 100 ppm NO gas, which confirms that the as-fabricated gas sensor can be considered a possible candidate for NO detection. Fig. 12c depicts the sensor response versus time for the 0.1 at% Mn exposed to various concentrations of NO gas. As confirmed in Fig. 12c, the sensor response increases with an increase in gas concentration. Additionally, the sensing films were also tested to H₂ gas, however, only 0.1 at% showed a response to H₂ (see ESI, Fig. S11†), disclosing poor response and recovery over a long period compared to CO₂ and NO gases. Nonetheless, previous studies have reported a higher sensing response to H₂ and ethanol gases using ZnO nanostructures.^52–55 Moreover, Rout et al.⁵⁴ have shown that the relative humidity plays an important role on the gas sensing. They observed that, at higher relative humidity concentration, the sensing response decreases.

The slope of the curves in Fig. 12b was used to study the sensitivity.⁵⁶ Fig. 12b shows that the response of the 0.1 at% Mn-doped film tested with NO increased linearly resulting in a sensitivity of 0.34 ppm⁻¹ in the range of 20–80 ppm. The 0.25 at% and 2.0 at% Mn-doped films showed a linear trend at 10–50 ppm and started to saturate at higher concentrations. At 10–50 ppm, these films had a sensitivity of 0.19 ppm⁻¹ and 0.06 ppm⁻¹ respectively. Therefore, these findings show that the 0.1 at% Mn-doped sensing film had a higher sensitivity to NO gas compared to other sensing films. This indicator is very vital for gas sensing, since it indicates how accurately the sensor can differentiate one analyte concentration from another one.

Table 4 shows a comparison of the sensing films on NO and CO₂ gases prepared in this study and those published in the literature. As shown in Table 4, the 0.1 at% Mn-doped ZnO sensing films prepared in this work had either comparable or higher responses when compared to ZnO materials prepared using different methods and tested at different operating temperatures and gas concentrations.

Table 4 Literature survey on NO and CO₂ gas sensors based on ZnO nanostructures

Sensing element	Method	Gas concentration (ppm)	Operating temperature (°C)	Response	Ref.
Ca doped ZnO	Sol–gel	2500 ppm CO2	450	117	57
50% La-loaded ZnO	Hydrothermal method	5000 ppm CO2	400	65%	58
ZnO nanowires	Chemical vapour deposition	NO 10 ppm	Under UV-light	46%	59
ZnO–In2O3 composites	Co-precipitation (powder)	19 ppm NO	200	60	60
		10 ppm NO	150	65
ZnO thin film	DC reactive magnetron sputtering	1000 pm CO2	300	90%	61
0.1 at% Mn–ZnO nanorods	Spray pyrolysis	100 ppm NO	Under UV-light	87%	This work
		100 ppm CO2		66%

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The UV-activated room-temperature gas sensing mechanism is demonstrated in Scheme 1. It is well known that the gas-sensing mechanism of undoped and Mn–ZnO-based sensors involves surface-controlled processes. According to these processes, the resistance change is governed by the types of molecules and amount of oxygen species adsorbed on the surface. Once the Mn–ZnO-sensing films are positioned in ambient air (or in the dark), atmospheric oxygen molecules adsorb onto the surface or grain boundary sites, where they become ionized by accepting electrons from the conduction band to form O⁻ and O²⁻ ions at the grain boundaries and creates a space charge at the grain boundary, as depicted in the reactions below and Scheme 1a.⁶² These are the characteristics of n-type metal oxides, such as ZnO. The adsorption of oxygen ions is the most imperative process in sensors since these ions are the most reactive surface species that make the material sensitive to the presence of oxidizing gases.

O₂(g) → O₂(ads) (5)

O₂(ads) + e⁻ → O₂⁻(ads) (6)

O₂⁻(ads) + e⁻ → 2O⁻(ads) (7)

O⁻(ads) + e⁻ → O²⁻(ads) (8)

Scheme 1 Schematic diagram showing the CO₂ sensing mechanism. Note: (a) corresponds to as-deposited Mn–ZnO sensing film relaxed in dark, at room temperature, (b) UV irradiated film, (c) oxygen species are photo-desorbed from the surface, (d) sensing film exposed to NO or CO₂ gas at room temperature and (e) NO or CO₂ gas reacts with adsorbed oxygen species and NO⁻ (or N₂O) or CO₂⁻ forms and the sensor resistance increases.

A depletion layer forms in the surface region of the Mn-doped ZnO films due to the consumption of electrons in the surface region of the films, which results in a decrease in electrical resistance. When the Mn-doped ZnO film is exposed to UV light with photon energy larger than the energy band gap of ZnO, electron–hole pairs generated. Upon reaching the surface of the films, some of the photo-generated electrons and holes recombine. As a result, most of the photo-generated holes react with the adsorbed oxygen ions on the surface, as shown in the following reactions and Scheme 1b:⁶³

h⁺(hν) + O²⁻(ads) → O₂(g), (9)

Or

Consequently, the surface depletion layer width is reduced, and oxygen species are photo-desorbed from the surface (Scheme 1c). However, the photo-generated electrons also contribute to a decrease in the depletion layer width and in the resistance of the gas sensing films.

Since NO behaves as an oxidizing gas on Mn-doped ZnO films, the resistance of n-type ZnO films increases as the NO gas chemisorbs on the surface. The increase in the resistance of the ZnO sensing film may be due to NO adsorption in the charged form (NO_ads⁻) or to the reduction of NO to form N₂O, as described by the following reactions and Scheme 1d and e:^64–67

NO + e → NO_ads⁻ (12)

2NO + e → N₂O + O_ads⁻ (13)

When the sensing films are exposed to CO₂, which is also an oxidizing gas, the following reactions occur (see Scheme 1d and e):^68,69

CO₂(gas) + e⁻ → CO₂⁻(ads) (14)

CO₂⁻(ads) + O⁻(ads) + 2e⁻ → CO(gas) + 2O²⁻(ads) (15)

The adsorption of O⁻ is a very interesting step in metal-oxide gas sensors, because the O⁻ ions assist the adsorbed oxidizing ions in transporting the electrons from the metal oxide surface. The concentration of electrons on the surface of the metal oxide decreases and the resistance of the n-type metal oxide layer increases accordingly. In contrast, the resistance of a p-type metal oxide surface decreases because the extracted electrons result in the generation of holes in the valence band.

4. Conclusions

In summary, the structural, morphological, and optical properties and the chemical state of the Mn-doped ZnO nanostructures grown via ASP on Corning glass substrates with different Mn concentrations were investigated. The effect of magnetism on gas sensing was also studied in detail. XRD and XPS revealed that Mn²⁺ ions were successfully incorporated into the ZnO lattice with no change in the wurtzite structure with a preferential orientation along the (002) direction. The optical and structural characterization indicated that the crystalline quality of ZnO deteriorated with Mn dopants. PL analyses indicated enhanced ultraviolet emission of Mn : ZnO, defect-related green emission of Mn : ZnO and an orange emission due to the ⁴T₁(4G) → ⁶A₁(6S) transition of Mn²⁺ in ZnO. The results revealed that the nature of the magnetism greatly influenced the response of both undoped and Mn-doped films. Films with lower Mn doping exhibited improved RTFM ordering, and higher doping destroyed the RTFM and resulted in a reduction in the gas sensing response. This was due to the contribution from the PM induced by Mn clusters or inhomogeneity on the surface of ZnO, as confirmed by TOF-SIMS, which decreased the number of oxygen vacancies. The findings revealed that Mn doping plays a significant role in controlling the morphology of the ZnO nanostructures and magnetism as well as the gas response of ZnO films.

Acknowledgements

This work was supported by the Department of Science and Technology, the Council for Scientific and Industrial Research (HGER27S), the WIROX project (PIRSES-GA-2011-295216), a Marie Curie International Research Staff Exchange Scheme Fellowship within the 7th European Community Framework Programme and ‘ORAMA’ Oxide Materials Towards a Matured Post-Silicon Electronics ERA FP7-NMP – contract no.: 246334.

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Footnote

† Electronic supplementary information (ESI) available: Additional results related to TOF-SIMS showing the 3D images and magnetization results on undoped and Mn doped ZnO films. See DOI: 10.1039/c5ra27154a

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