Journal of Alloys and Compounds 602 (2014) 294–299

Contents lists available at ScienceDirect

Journal of Alloys and Compounds

journal homepage: www.elsevier .com/locate / ja lcom

Highly conducting and wide-band transparent F-doped Zn1�xMgxO thinfilms for optoelectronic applications

http://dx.doi.org/10.1016/j.jallcom.2014.02.1810925-8388/� 2014 Elsevier B.V. All rights reserved.

⇑ Corresponding author at: State Key Laboratory of Silicon Materials, Departmentof Materials Science and Engineering, Zhejiang University, Hangzhou 310027, PRChina. Tel.: +86 571 87951958; fax: +86 571 87952625.

E-mail address: zlp1@zju.edu.cn (L.P. Zhu).

Y.M. Guo a, L.P. Zhu a,b,⇑, J. Jiang a, Y.G. Li a, L. Hu a, H.B. Xu a, Z.Z. Ye a

a State Key Laboratory of Silicon Materials, Department of Materials Science and Engineering, Zhejiang University, Hangzhou 310027, PR Chinab Cyrus Tang Center for Sensor Materials and Applications, Zhejiang University, Hangzhou 310027, PR China

a r t i c l e i n f o a b s t r a c t

Article history:Received 14 October 2013Received in revised form 27 February 2014Accepted 27 February 2014Available online 12 March 2014

Keywords:F-dopedWide-bandTransparent conducting oxideF concentrationPulsed laser depositionZn1�xMgxO thin films

Fluorine (F) doped Zn1�xMgxO thin films were deposited on quartz via pulsed laser deposition (PLD). Fdoping can decrease resistivity and broaden the bandgap of Zn1�xMgxO thin films as well when F concen-tration is less than 3%, otherwise F doping will backfire. The structural, electrical, and optical properties ofthese thin films were studied as a function of deposition temperatures. The Zn0.9Mg0.1OF0.03 thin filmsdeposited at 350 �C are optimal to be applied as transparent electrodes, taking both electrical and opticalproperties into account. Thin films have a low resistivity about 6.92 � 10�4 X cm, with a carrier concen-tration of 5.26 � 1020 cm�3, and a Hall mobility of 17.2 cm2 V�1 s�1. The average optical transmittance ishigher than 85% in the visible wavelength region.

� 2014 Elsevier B.V. All rights reserved.

1. Introduction

Transparent conducting oxides (TCO), an important class ofmaterials for various optoelectronic devices, have been studiedand applied extensively in the field of liquid crystal displays(LCDs), organic light emitting diodes, and thin solar cells [1–3].Generally speaking, indium tin oxide (ITO) is applied widely asthe transparent electrode in solar cells. However, indium (In) israre metal and reserve volume is scare, which limits its applicationin large-scale commercial production. compared to ITO or SnO2

[4–6] ZnO is one of the most promising candidates to replaceITO, owing to its wide direct energy gap, abundant raw material,environmental friendliness, high radiation resistance and relativelylow growth temperature, moreover/furthermore, the bandgap (Eg)of ZnO can be controlled by alloying with MgO, from 3.37 to4.05 eV [7–10], while that of ITO is constant, Eg = 3.75 eV. A widerbandgap TCO offers the potential for improved UV response insolar cells [11] and photoconductive detectors [9]. However,employing ZnMgO as a substitute for ZnO has limitations, due tothe decreased conductivity [12]. In order to compensate for thedecrease of conductivity, Al and Ga have been added to ZnMgOto create electronic carriers [13–17]. Matsubara et al. [18] reported

the preparation of Al-doped ZnMgO thin films and studied therelation between bandgap and resistivity in 2004. Ke et al. [19]deposited Ga-doped ZnMgO films on sapphire and investigatedthe origin of electrical property reduction with increasing Mgcontent. Recently, Park et al. [20] studied the effects of the Al con-centration and the growth variables on the optical and electricalproperties of UV-range Zn0.88�xMg0.12AlxO TCO films. They foundthat the resistivity significantly decreased by doping the film withAl, to 2–3 � 10�3 X cm, but was not critically affected by the addi-tion of a larger amount of Al. Wang et al. [21] reported the effect ofMg and Al co-doping on the structural and photoelectric propertiesof ZnO thin films and got the similar conclusions to the former. Onthe other hand, several groups prepared Mg and Ga co-doped ZnO(MGZO) thin films that maintain a low electrical resistivity ofbelow 10�4 X cm and wide band gap energy of over 3.75 eV[22–26]. Liu et al. [27] combined Ga-doped MgZnO (GMZO)transparent conductive oxide with a silver (Ag) layer to form aGMZO/Ag/GMZO composite structure, which yielded a resistivityof 5.8 � 10�5 X cm.

It has been shown that fluorine is an effective donor in F-dopedZnO thin films, because its ionic radius (1.31 Å) is similar to that ofoxygen (1.38 Å) [28]. Another advantage is that electronic pertur-bation is largely confined to the filled valence band and scatteringof conduction electrons is reduced when fluorine substitutes foroxygen, which could lead to high electron mobility as well aslow absorption loss [29]. Our previous study has successfullyprepared highly transparent and conducting fluorine-doped ZnO

Y.M. Guo et al. / Journal of Alloys and Compounds 602 (2014) 294–299 295

thin films [30], it is pretty interesting that we extended the charac-teristics from binary compounds to ternary compounds and makethe ZnO-based thin films be more widely applied. Pan et al.[31–33] prepared F and Al/Sn co-doped thin films using sol–gelmethod and studied the effect of annealing on the structures andproperties of F and Al co-doping thin films. A minimum resistivityof 1 � 10�3 X cm was obtained from the F and Sn co-doped ZnOthin film with 3% F doping, and the average optical transmittancein the entire visible wavelength region was higher than 90%, owingto its high carrier concentration and mobility, as well as goodcrystal quality.

This paper presents systematic research investigation on thegrowth and characterization of wide-band transparent and highlyconducting F-doped Zn1�xMgxO thin films deposited by pulsed la-ser deposition (PLD) for the first time. The effects of Mg content,substrate temperature on structural, electrical and optical proper-ties of the thin films were investigated and discussed in detail.Also, the influence of F doping was researched by changing the Fconcentration in Zn0.9Mg0.1OFy (y = 0, 0.03, 0.06, 0.09) thin filmsprepared under the same deposition conditions.

Fig. 1. Properties of Zn1�xMgxO0.985F0.03 thin films with different Mg contents,deposited at 350 �C.

2. Experimental details

Highly transparent and conducting F-doped Zn1�xMgxO thin films were depos-ited on quartz substrates using sintered target with different Mg, F concentrations(the atom percent when mixing the powers), prepared from spectroscopically purepowders of ZnO, MgO and ZnF2. Prior to deposition, the target was pre-sputtered foraround 10 min to remove the contaminants from the surface. The substrates, wereultrasonically cleaned with acetone and de-ionized water for 30 min, and subse-quently dried in flowing nitrogen gas before being loaded into the chamber. Beforedeposition, the deposition chamber was evacuated to a pressure of 8.9 � 10�4 Pa.During deposition, high purity (99.9995%) oxygen gas was introduced into the reac-tion chamber and the chamber pressure was kept at 1 � 10�3 Pa. A pulsed KrF laserwas operated at a wavelength of 248 nm and a repetition of 5 Hz. The laser energywas 280 mJ per pulse, and the target-substrate distance was kept at 5.0 cm. All ofthe thin films were deposited for 60 min and had an approximate thin films thick-ness of 240 nm.

X-ray diffraction (XRD) using a XPERT-PRO system with a Cu Ka (k = 1.5406 Å)source was used to characterize the crystal structure of the thin films. The depthprofile of the thin films was investigated by secondary ion mass spectroscopy(SIMS; Cameca IMS-6F). Chemical binding energies and composition of the thinfilms were investigated by the X-ray photoelectron spectroscopy (XPS; ESCALAB250Xi). The surface morphology and the thin films thickness were studied usingthe field emission scanning electron microscopy (FE-SEM; HITACHIS-4800). TheHall measurements in the van der Pauw configuration (BIO-RAD HL5500PC) wereutilized to examine the electrical properties. The optical transmission measure-ments were performed using a UV-near-IR grating spectrometer (U-4100;300–1000 nm). In addition, to make sure the reliability of the results, all thesamples were measured under the same condition.

3. Results and discussion

3.1. The effects of Mg

Fig. 1 shows effects of Mg on the properties of Zn1�xMgxO0.985F0.03

thin films, deposited at 350 �C. In Fig. 1(a), the resistivity increasesfrom 4.53 � 10�4 X cm to 5.85 � 10�2 X cm, more than two orders,when the Mg content increases to 20%, as both the carrier concentra-tion and Hall mobility steeply decrease. This is a common phenom-enon existing in Zn1�xMgxO and Al-doped Zn1�xMgxO thin films andwas explained clearly by Matsubara et al. [18]. In Fig. 1(b), it is ob-served that the Eg, estimated using Tauc’s relationship from thetransmittance spectra of the Zn1�xMgxO0.985F0.03 thin films, in-creases from 3.614 eV to 4.216 eV as the Mg increases from 0% to20%, demonstrating the bandgap of F-doped Zn1�xMgxO thin filmscan be modified just like undoped Zn1�xMgxO thin films [7–10].Besides, a compression of the c-axis with increased Mg content isobserved by the systematic shift of the (002) reflection to higher an-gles. We took the Zn1�xMgxO0.985F0.03 (x = 0.1) (ZMOF) thin films for

further research, considering their high conductivity and wide band-gap together.

3.2. Structural properties

Fig. 2(a) shows the XRD patterns of the ZMOF thin films fabri-cated at different growth temperatures. It is found that the crystal-linity increases greatly as the growth temperature increases. Onlythe peaks indexed to (002) and (004) appear in thin films, indicat-ing these thin films are preferred c-axis orientation with hexagonalwurtzite structure. The absence of extra peaks in the XRD patternexcludes the possibility of any extra phrases in the thin films, suchas MgO or ZnF2, etc., demonstrating the formation of quaternaryZMOF alloy thin films unambiguously. 2 Theta of (002) peak ofZMOF thin films deposited at different substrate temperatureswas shown in the inset of Fig. 1(a). It is obvious that, (002) peaksshift to higher angle with increasing deposition temperature, dueto the substitution of Zn2+ (0.074 nm) by Mg2+ (0.065 nm) [34].Such phenomenon could be explained by the differences in vaporpressures between ZnO and MgO [19]. The saturated vapor pres-sure of ZnO is higher than MgO in a wide temperature range [35]and results in decreasing Zn content with increasing temperature.

In order to assess the quality of the ZMOF thin films, we pre-sented the full width at half maxima (FWHM) values of (002) peakand the average crystallite size calculated according to the Scherrerformula [36] as a function of growth temperature, as shown inFig. 2(b). To make the results more precise, the instrumentalbroadening of the diffraction peaks has been subtracted from mea-sured FWHM in our study. We can see that the FWHM first de-creases from 0.366� at the substrate temperature of 200 �C to0.233� at 350 �C and then decreases slowly to 0.219� at 500 �Cwhile the crystallite size shows the opposite trend, indicating thecrystallinity increases due to the interaction and agglomerationwith each other at high temperature. This can be explained bythe residual compressive stress existing in the (002) plane in thethin films deposited at the temperature range of 200–350 �C [37].The stress was generated during the deposition process due to

Fig. 2. (a) XRD patterns of ZMOF thin films deposited at different substratetemperatures. The inset displays normalized XRD spectra of (002) peak. (b)Dependence of the FWHM (black hollow circles) of the (002) peak and crystallitesize (blue solid squares) on growth temperature for the ZMOF thin films. (Forinterpretation of the references to colour in this figure legend, the reader is referredto the web version of this article.)

296 Y.M. Guo et al. / Journal of Alloys and Compounds 602 (2014) 294–299

the freezing-in of the structural defects [38] at low growth temper-ature and it was completely relaxed when the thin films wasgrown at temperature above 350 �C. Thus all the thin films depos-ited above 350 �C have good crystalline quality.

Fig. 3(a) and (b) shows the top and cross-section SEM images ofZMOF thin films prepared at 350 �C, respectively. As can be seen,the surface is uniform and compact, thus implying good crystallin-ity, which is beneficial to the advantages of PLD device. The thinfilms exhibit a columnar crystal structure normal to the surfaceof the substrate, confirming the c-axis orientation growth, whichis consistent with the results of XRD in Fig. 1. Besides, the thinfilms have a uniform thickness of about 245 nm.

The depth profile of O, Mg, Zn, and F in the ZMOF thin filmsgrown at 350 �C was carried out by SIMS measurements, as shownin Fig. 3(c). As lack of reference samples for the SIMS measurement,quantitative results could not be obtained. The contents of Zn, Mg,and O keep constant throughout the whole thin films. Also, it is evi-dent that F has been clearly detected, and uniformly distributed inthe thin films.

To get the quantitative composition of the films, we showed theXPS spectra of ZMOF thin films grown at 350 �C in Fig. 3(d). Thetypical survey spectrum confirmed the presence of Zn, Mg, O andF from the ZMOF thin films. In the inset, F1s peak is observed at685.0 eV, representing binding energies of the Zn–F bonds. Thisindicates that F has been incorporated into the Zn1�xMgxO thinfilms as donor, which is responsible for the high conductivity.

Table 1 displays the F and Mg contents in the ZMOF thin filmsdeposited at different substrate temperatures measured usingXPS, we can find that the F contents decrease linearly with increas-ing substrate temperature, this may ascribed to the volatilizationof F during high temperature while the Mg contents increases inagreements with the results of XRD in Fig. 1.

3.3. Electrical properties

Fig. 4 shows resistivity (q), carrier concentration (n), and Hallmobility (l) as a function of substrate temperatures. It can be seenthat the resistivity decreases with the increasing substrate temper-ature, the electron concentration shows the opposite trend, and themobility displays a substantially rising trend. As the growthtemperature increases from 200 �C to 350 �C, the resistivity ofthe thin films decreases rapidly from 1.00 � 10�2 X cm to the min-imum value of about 6.92 � 10�4 X cm, and then increases to2.63 � 10�3 X cm with a further increase of 500 �C. As is wellknown, resistivity is proportional to the reciprocal of the productof carrier concentration and mobility. Therefore, the change inresistivity with substrate temperature is attributed to the changein carrier concentration and/or Hall mobility, which are character-istic parameters reflecting the thin films structure and/or theimpurity contents [39]. As shown in Fig. 4, both the carrier concen-tration and Hall mobility increase from 200 �C to 350 �C, thus theresistivity decreases in this temperature region. These tendenciescan be explained by the thin films’ structures. When deposited atlow temperature, the thin films has not only a large residual stresswhich induces lattice distortion and leads to an increase of disloca-tions, but also a plenty of small-size grains which gives rise to grainboundaries. However, with a further increase of growth tempera-ture up to 500 �C, carrier concentration decreases rapidly, but Hallmobility still increases. Therefore, the increase in resistivity ismainly attributed to the reduction in carrier concentration, whichmay be caused by the absorption of the oxygen atoms under hightemperatures oxidation environment [40]. The possible reason forthe increase in mobility is the decrease in carrier density and theresulting electron–electron scattering. In addition, the decrease ofF content and the increase of Mg content, showed in Table 1, dete-riorate the conductivity of the thin films. Thus, the thin film depos-ited at 350 �C has the best electrical properties.

Zn0.9Mg0.1OFy (y = 0, 0.03, 0.06, 0.09) thin films were preparedunder the same deposition conditions in order to explore theeffects of F concentration on the Zn0.9Mg0.1OFy thin films. Resultssuggested that F benefited the electrical performance of theF-doped thin films, due to the increase in both carrier concentra-tion and mobility, compared with the undoped one when the Fconcentration was below 3 at%, as shown in Fig. 5(a). This confirmsthat F atoms have doped into ZnMgO lattices and effectively act asdonors when they occupy O in the hexagonal lattice [29,30].However, both the carrier concentration and mobility decreasewith further F addition. We found an unknown peak in theZn0.9Mg0.1OFy (y = 0.06, 0.09) shown in Fig. 5(b), which can beattributed to the nonconducting clusters form in Zn0.9Mg0.1OFy thinfilms when F contention exceeds a critical concentration, limitingthe carrier concentration and decrease the mobility. The decreaseof carrier concentration was proposed that free electron were ini-tially generated by the F doping but further generation wasblocked by the cluster formation and the decrease of mobilitywas caused by the increasing internal strain and the number ofscattering centers [13].

3.4. Optical properties

Fig. 6(a) shows the optical transmission spectra of the ZMOFthin films deposited at different temperatures. All the thin films

Fig. 3. (a) Surface and (b) cross-sectional SEM morphologies, (c) SIMS depth profile, and (d) high-resolution XPS spectral of ZMOF thin films prepared at 350 �C.

Table 1The F and Mg contents in the ZMOF thin films deposited at different substratetemperatures measured using XPS.

Temperature (�C) 200 250 300 350 400 500

F (at. mol%) 2.71 2.54 2.16 2.04 1.67 1.24Mg (at. mol%) 20.5 22.4 23.7 25.7 27.1 28.4

Y.M. Guo et al. / Journal of Alloys and Compounds 602 (2014) 294–299 297

have sharp absorption edge in the UV range of 300–400 nm and ex-hibit an average optical transparency of higher than 85%. The insetof Fig. 6(a) shows the (aht)2 plots for ZMOF and undopedZn0.9Mg0.1O (ZMO) thin films prepared at 350 �C. As can be seen,estimated optical bandgap of the ZMOF thin films is much largerthan that of undoped ZMO due to the Burstein–Moss effect [41].

Fig. 4. Resistivity, carrier concentration, and Hall mobility as function of growth temperature for ZMOF thin films.

Fig. 5. Properties of Zn0.9Mg0.1OFy thin films with different F contents, deposited at350 �C.

298 Y.M. Guo et al. / Journal of Alloys and Compounds 602 (2014) 294–299

According to the Burstein–Moss effect, apparent bandgap as deter-mined by optical absorption increases as states near conductionband edge are populated on doping.

Fig. 6(b) shows the optical bandgap (Eg) of ZMOF thin films, cal-culated using transmittance data, as a function of the growth tem-perature. We can see all the bandgaps are much larger than bothundoped ZnO (3.2 eV) and F-doped ZnO thin films (3.63 eV) [30]due to incorporation of Mg. It is clear that the optical bandgap ofthe thin films increases from about 3.816 to 3.933 eV as the sub-

Fig. 6. (a) Transmission spectra of the ZMOF thin films deposited at different temperaturat the same parameter) and (b) the optical bandgap (Eg) of ZMOF thin films as a functio

strate temperature increased from 200 �C to 350 �C and then de-creases to 3.838 eV with a further increase in the substratetemperature. There are two possible reasons for this phenomenon.Firstly, the in-plane residual stress, as mentioned in Fig. 2 [42],leads to change in the electronic structure, resulting in bandgapborden. Secondly, the crystallinity of Zn1�xMgxO thin films grownat a low temperature is lower than that grown at a high tempera-ture, and there are a mass of disorder arrangement of thin filmsgrown at a low temperature, which induces the localized loweststates in the conduction band. Therefore, the localized states willmake the absorption edge shift to a higher energy by DE, leadingto a large bandgap [35]. As the deposition temperature increasesfrom 200 to 350 �C, both the relaxation of residual stress and therising electron concentration due to Burstein–Moss (B–M) effectresult in a bandgap borden. However, when the substrate temper-ature is above 350 �C, the effect of residual stress on the bandgapcan be neglected. Therefore, the decrease of electron concentrationvia B–M effect becomes the dominant mechanism, and induces arapid reduction of optical bandgap in this temperature region.The similar phenomena existed in former reports [37,43,44].

4. Conclusions

In summary, highly transparent and conducting F-dopedZn1�xMgxO thin films were deposited on quartz substrates bypulsed laser deposition. The effects of Mg content, F concentration,substrate temperature on structural, electrical, and optical proper-ties of these thin films have been researched. All the thin films are

es (the inset shows the (aht)2 plots for ZMOF and undoped ZMO thin films preparedn of the growth temperature.

Y.M. Guo et al. / Journal of Alloys and Compounds 602 (2014) 294–299 299

polycrystalline and have a highly c-axis preferential orientation.The thin films deposited at 350 �C has the best crystallinity, andthe lowest resistivity of 6.92 � 10�4 X cm, with a carrier concentra-tion of 5.26 � 1020 cm�3, and a Hall mobility of 17.2 cm2 V�1 s�1.The average optical transmittance in the entire visible wavelengthregion is higher than 85%. The wide optical bandgap energy com-bined with good transparency implies F-doped Zn1�xMgxO thinfilms are promising candidates for many optoelectronic applica-tions. Moreover, the reduction in resistivity and broaden bandgapdue to F doping evinces the possible enhancement of their poten-tially active role in transparent electrodes in optoelectronic devices.

Acknowledgements

This work was supported by National Natural Science Founda-tion of China 51072181, Program for Innovative Research Teamin University of Ministry of Education of China (IRT13037, and Na-tional Science and Technology Support Program (2012BAC08B08).

References

[1] Z. Pan, X. Tian, S. Wu, C. Xiao, Z. Li, J. Deng, G. Hu, Z. Wei, SuperlatticesMicrostruct. 54 (2013) 107–117.

[2] W. Beyer, J. Hüpkes, H. Stiebig, Thin Solid Films 516 (2007) 147–154.[3] D.H. Zhang, X.H. Liu, X. Wang, J. Alloys Comp. 509 (2011) 4972–4977.[4] M.M. Jumidali, M.R. Hashim, Superlattices Microstruct. 52 (2012) 33–40.[5] R. Yousefe, B. Kamaluddin, J. Alloys Comp. 479 (2009) L11–L14.[6] H. Chettah, D. Abdi, Thin Solid Films 537 (2013) 119–123.[7] S.I. Shimakawa, Y. Hashimoto, S. Hayashi, T. Satoh, T. Negami, Sol. Energy

Mater. Sol. Cells 92 (2008) 1086–1090.[8] J.Y. Cho, S.W. Shin, Y.B. Kwon, H.K. Lee, K.U. Sim, H.S. Kim, J.H. Moon, J.H. Kim,

Thin Solid Films 519 (2011) 4282–4285.[9] W. Yang, R.D. Vispute, S. Choopun, R.P. Sharma, T. Venkatesan, H. Shen, Appl.

Phys. Lett. 78 (2001) 2787–2789.[10] Y.Z. Zhang, J.H. He, Z.Z. Ye, L. Zou, J.Y. Huang, L.P. Zhu, B.H. Zhao, Thin Solid

Films 458 (2004) 161–164.[11] T. Minemoto, Y. Hashimoto, T. Satoh, T. Negami, H. Takakura, Y. Hamakawa, J.

Appl. Phys. 89 (2001) 8327–8330.[12] D.C. Olson, S.E. Shaheen, M.S. White, W.J. Mitchell, M.F.A.M. van Hest, R.T.

Collins, D.S. Ginley, Adv. Funct. Mater. 17 (2007) 264–269.[13] J.G. Lu, S. Fujita, T. Kawaharamura, H. Nishinaka, Y. Kamada, T. Ohshima, Appl.

Phys. Lett. 89 (2006) 2621071–2621073.[14] W. Yuan, L.P. Zhu, Z.Z. Ye, X.Q. Gu, Appl. Surf. Sci. 256 (2009) 1452–1454.[15] P. Prathap, N. Revathi, A.S. Reddy, Y.P.V. Subbaiah, K.T.R. Reddy, Thin Solid

Films 519 (2011) 7592–7595.[16] K. Fleischer, E. Arca, C. Smith, I.V. Shvets, Appl. Phys. Lett. 101 (2012)

1219181–1219184.

[17] C.Q. Li, Z.Z. Zhang, H.Y. Chen, X.H. Xie, D.Z. Shen, Thin Solid Films 548 (2013)456–459.

[18] K. Matsubara, H. Tampo, H. Shibata, A. Yamada, P. Fons, K. Iwata, S. Niki, Appl.Phys. Lett. 85 (2004) 1374–1376.

[19] Y. Ke, J. Berry, P. Parilla, A. Zakutayev, R. O’Hayre, D. Ginley, Thin Solid Films520 (2012) 3697–3702.

[20] J.H. Park, J.B. Lim, B.T. Lee, Semicond. Sci. Technol. 28 (2013) 0650041–0650045.

[21] C.Y. Wang, S.Y. Ma, F.M. Li, Y. Chen, X.L. Xu, T. Wang, F.C. Yang, Q. Zhao, J. Liu,X.L. Zhang, X.B. Li, X.H. Yang, J. Zhu, Mater. Sci. Semicond. Process. 17 (2014)27–32.

[22] S.W. Shin, I.Y. Kim, K.S. Jeon, J.Y. Heo, G.S. Heo, P.S. Patil, J.H. Kim, J.Y. Lee, J.Asian Ceram. Soc. 1 (2013) 262–266.

[23] S.W. Shin, G.L. Agawane, I.Y. Kim, S.H. Jo, M.S. Kim, G.S. Heo, J.H. Kim, J.Y. Lee,Surf. Coat. Technol. 231 (2013) 364–369.

[24] S.W. Shin, I.Y. Kim, G.V. Kishor, Y.Y. Yoo, Y.B. Kim, J.Y. Heo, G.S. Heo, P.S. Patil,J.H. Kim, J.Y. Lee, J. Alloys Comp. 585 (2014) 608–613.

[25] W.S. Liu, W.K. Chen, K.P. Hsueh, J. Alloys Comp. 552 (2013) 255–263.[26] J.M. Liu, X.R. Zhao, L.B. Duan, H.N. Sun, X.J. Bai, L. Chen, C.L. Chen, Appl. Surf.

Sci. 258 (2012) 6297–6301.[27] W.S. Liu, Y.H. Liu, W.K. Chen, K.P. Hsueh, J. Alloys Comp. 564 (2013) 105–113.[28] S. Pearton, Prog. Mater. Sci. 50 (2005) 293–340.[29] R.G. Gordon, MRS Bull. 25 (2000) 52–57.[30] L. Cao, L.P. Zhu, J. Jiang, R. Zhao, Z.Z. Ye, B.H. Zhao, Sol. Energy Mater. Sol. Cells

95 (2011) 894–898.[31] Z.C. Pan, P.W. Zhang, X.L. Tian, G. Cheng, Y.H. Xie, H.C. Zhang, X.F. Zeng, C.M.

Xiao, G.H. Hu, Z.G. Wei, J. Alloys Comp. 576 (2013) 31–37.[32] Z.C. Pan, Y.H. Xiao, X.L. Tian, S.K. Wu, C. Chen, J.F. Deng, C.M. Xiao, G.H. Hu, Z.G.

Wei, Mater. Sci. Semicond. Process. 17 (2014) 162–167.[33] Z.C. Pan, J.M. Luo, X.L. Tian, S.K. Wu, C. Chen, J.F. Deng, C.M. Xiao, G.H. Hu, Z.G.

Wei, J. Alloys Comp. 583 (2014) 32–38.[34] Y.J. Wang, X.Q. Wei, R.R. Zhao, C.S. Chen, S. Gao, J. Lian, J. Electroceram. 27

(2011) 203–208.[35] D.Y. Jiang, C.X. Shan, J.Y. Zhang, Y.M. Lu, B. Yao, D.X. Zhao, Z.Z. Zhang, D.Z. Shen,

C.L. Yang, J. Phys. D. Appl. Phys. 42 (2009) 0251061–0251063.[36] S. Merazga, A. Brighet, A. Keffous, K. Mirouh, L. Guerbous, Y. Belkacem, M.

Kechouane, Int. J. Nano Technol. 10 (2013) 587–596.[37] X.Q. Gu, L.P. Zhu, Z.Z. Ye, Q.B. Ma, H.P. He, Y.Z. Zhang, B.H. Zhao, Sol. Energy

Mater. Sol. Cells 92 (2008) 343–347.[38] I.K. Durukan, Y. Ozen, K. Kizilkaya, M.K. Ozturk, T. Memmedli, S. Ozcelik, J.

Mater. Sci.-Mater. El 24 (2013) 142–147.[39] Y. Tian, J.N. Dai, H. Xiong, G. Zheng, M. Ryu, Y.Y. Fang, C.Q. Chen, Chin. Phys.

Lett. 29 (2012) 0881011–0881014.[40] M.S. Kim, S. Kim, K. Gug, J.Y. Leem, D.Y. Kim, G. Nam, D.Y. Lee, J.S. Kim, J.S. Kim,

J.S. Son, J. Korean Phys. Soc. 60 (2012) 1570–1575.[41] K. Zak, M.E. Abrishami, W.H. Abd Majid, R. Yousefi, S.M. Hosseini, Ceram. Int.

37 (2011) 393–398.[42] J. Li, J.H. Huang, W.J. Song, Y.L. Zhang, R.Q. Tan, Y. Yang, J. Cryst. Growth 314

(2011) 136–140.[43] H.Y. Choi, M.S. Kim, M.Y. Cho, G.S. Kim, S.M. Jeon, K.G. Yim, D.Y. Kim, H.H. Ryu,

J.Y. Leem, D.Y. Lee, J.S. Kim, J.S. Kim, J.S. Son, J.I. Lee, J. Korean Phys. Soc. 56(2010) 1514–1518.

[44] R.K. Gupta, K. Ghosh, R. Patel, P.K. Kahol, Mater. Sci. Eng. B 156 (2009) 1–5.