The structure and electrical properties of PbPdO2 thin films with preferred orientation prepared by PLD

Author’s Accepted Manuscript The structure and electrical properties of PbPdO2 thin films with preferred orientation prepared by PLD Xiang Chen, Yue Chen, Yanmin Yang, Hai Jia, Jian-Min Zhang, Shuiyuan Chen, Zhigao Huang www.elsevier.com/locate/ceri

PII: DOI: Reference:

S0272-8842(17)30884-2 http://dx.doi.org/10.1016/j.ceramint.2017.05.079 CERI15239

To appear in: Ceramics International Received date: 14 February 2017 Revised date: 2 May 2017 Accepted date: 10 May 2017 Cite this article as: Xiang Chen, Yue Chen, Yanmin Yang, Hai Jia, Jian-Min Zhang, Shuiyuan Chen and Zhigao Huang, The structure and electrical properties of PbPdO2 thin films with preferred orientation prepared by PLD, Ceramics International, http://dx.doi.org/10.1016/j.ceramint.2017.05.079 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

The structure and electrical properties of PbPdO2 thin films with preferred orientation prepared by PLD

Xiang Chen,1,2 Yue Chen,1,2 Yanmin Yang, 1,2 Hai Jia,1,2 Jian-Min Zhang,1,2* Shuiyuan Chen1,2 and Zhigao Huang1,2* 1

Fujian Provincial Key Laboratory of Quantum Manipulation and New Energy Materials, College of

Physics and Energy, Fujian Normal University, Fuzhou 350117, China 2

Fujian Provincial Collaborative Innovation Center for Optoelectronic Semiconductors and Efficient

Devices, Xiamen, 361005, China

Corresponding author: Telephone/Fax: 86-591-22867577; E-mail: [email protected] (Z.Huang), [email protected](J.Zhang)

The structure and electrical properties of PbPdO2 thin films with preferred orientation prepared by PLD

Xiang Chen,1,2 Hai Jia,1,2 Yue Chen,1,2 Yanmin Yang, 1,2 Hai Jia,1,2 Jian-Min Zhang,1,2* Shuiyuan Chen1,2 and Zhigao Huang1,2* 1

Fujian Provincial Key Laboratory of Quantum Manipulation and New Energy Materials, College of

Physics and Energy, Fujian Normal University, Fuzhou 350117, China 2

Fujian Provincial Collaborative Innovation Center for Optoelectronic Semiconductors and Efficient

Devices, Xiamen, 361005, China

We have fabricated PbPdO2 nano-particles and PbPdO2 films with (211) preferred orientation (sample A) and with (002) preferred orientation (sample B), respectively. The structures and physical properties were measured by XRD, Raman, SEM, XPS and AFM. The experimental results indicate that the PbPdO2 films have granular structure with the grain size of about 200nm and some Pb vacancies exist in both samples. It is found that PbPdO2 with (002) preferred orientation has near zero-gap, while PbPdO2 with (211) preferred orientation has larger band gap. This surface anisotropy band gap is also proved by our first principles band structure calculations. Moreover, one notices that the sample B possesses lower resistivity than sample A, and both samples possess a wide insulator-metal transition temperature (TMI) with around 370K.

Key words: PbPdO2 film; Preferred orientation; Gapless semiconductor; Work function

1. Introduction Materials with a zero-energy band gap (such as graphene, HgCdTe, HgTeSe, HgZnSe, PbPdO2, topological insulator, MoS2 and so on ) have intriguing physical properties and numerous potential practical applications in spintronics, electronics, optics and sensors, which arouse the research enthusiasm [1]. Wang et al [2]. proposed firstly a new class of the spin gapless semiconductors (SGS’s) by performing density functional theory (DFT) band structure calculations for Co-doped PbPdO2. Immediately following, they studied the effects of both electrical current and magnetic field on the resistivity of Co-doped PbPdO2 thin films [3]. The usually strong colossal electric current induced electroresistance (CER) and giant magnetoresistance (GMR) effects were observed. Moreover, the giant MR exhibits the interesting temperature and magnetic field orientation dependences of CER and GMR. The pioneering studies from Wang et al. invoked a series of experimental and theoretical investigations of PbPdO2. The different prepared methods were implemented, and the various crystalline structures were obtained. Ozawa et al.[4,5] synthesized firstly single-phase polycrystalline samples of PbPdO2 by solid-state reaction. The lattice parameters obtained from powder X-ray diffraction are a = 9.4547(1) Å, b = 5.45971(6) Å and c = 4.66051(6) Å in an orthorhombic cell. For 2θ<40°, 7 diffraction peaks with (111), (211), (020), (301), (400), (002), (311) were observed. Compared to standard diffraction patterns, it is found that (200) crystal face is not easy produced [4,5]. The similar polycrystalline samples and crystalline structures were obtained in PbPdO2 and Pb(Pd0.9M0.1)O2 (M= Mn, Co) by solid state reaction [6-8]. Su et al. prepared single-phase PbPd0.81Co0.19O2 film by using the sol-gel spin-coating technique and an oxidation treatment. For

2θ<40°, 4 diffraction peaks with (211), (020), (400), (002) were observed [9-11]. Wang et al. prepared the films with the preferred orientations (002) by PLD [3]. Choo et al. prepared PbPdO2 and Pb(Pd,Co)O2 thin films oriented along the (020) direction by PLD [12,13]. From the different physical properties with polycrystalline and preferred orientation samples [3-5, 6-13], it can be found that electrical transport, magnetism, CER and GMR should be associated to the preferred orientation. Unfortunately, the relation between the physical properties and preferred orientation has not studied well. On the other hand, with varying temperature, a broad metal-insulator transition TMI from a high-temperature metallic behavior to a low-temperature insulating behavior was observed in the electrical resistivity, which is related to the thermally assisted excitation near the Fermi level due to its gapless band structure. A metal–insulator transition around 90 K and p-semiconductor for the polycrystalline samples was found, and the metal–insulator transition is suppressed by the applied pressure [4]. A TMI with about 170K was clearly seen in the films with the preferred (002) orientations. Moreover, the metal-insulator transition temperature decreases with increasing magnitude of the electric current [3]. The values of TMI for polycrystalline PbPdO2, PbPd0.9Co0.1O2 and PbPd0.9Mn0.1O2 are found to be 100, 150, 70K, which means that Co and Mn dopants can give rise different TMI shift directions [6,7]. However, the insulator-metal transition temperature of PbPd0.81Co0.19O2 film prepared by using the sol-gel spin-coating technique and an oxidation treatment was found to be 358 K [9], markedly higher than the reported values of similar material systems. The metal-insulator-like transition temperature of no-annealing PbPdO2 thin film oriented along the (020) direction is about 230K with wide minimax [12]. Moreover, with increasing

annealing time, the metal-insulator like transition temperature is evidently enhanced. The experimental results above indicate that the crystalline orientation may have important roles on TMI, and it has not yet been studied well. In this paper, PbPdO2 nano-particles and PbPdO2 films with (211) preferred orientation (sample A) and with (002) preferred orientation (sample B) were prepared, respectively. The experimental results indicate that PbPdO2 with (002) preferred orientation has near zero-gap, while it with (211) preferred orientation has larger band gap. This surface anisotropy band gap is also proved by our first principles band structure calculations. Moreover, we found that, the sample B possesses lower resistivity than sample A, and both samples possess a wide insulator-metal transition temperature (TMI) at around 370K. 2. Materials and methods 2.1 Synthesis of PbPdO2 At first, the PbPdO2 nanoparticles were prepared by sol-gel method. Pb(NO3)2, Pd(NO3)2 were used as raw materials, and the chelating agent and solvent was citric acid monohydrate and deionized water respectively. An additional 5 mol % Pb(NO3)2 was used to supplement the volatilization of Pb in the process of the subsequent heating treatment. The prepared PbPdO2 nanoparticles by sol-gel method is named sample C. Secondly, the PbPdO2 cylinder bulk sheets were synthesized by those nanoparticles. Then, the cylinder sheets were used as the PLD target. The PbPdO2 thin film was deposited on (200) oriented MgO single crystal substrate using PLD technique. A KrF excimer laser with a wavelength of 248nm was used as a source of target ablation and the repetition rate is 3Hz. The substrate temperature is 550℃. The samples were deposited at different O2 ambient pressure

(sample A: 40Pa, sample B:50Pa) and the initial vacuum condition of the chamber reached 10 -6 Torr. After deposition, the samples were ex-situ annealed in air at 650℃. 2.2 Characterization methods The powder X-ray diffraction patterns (XRD) of PbPdO2 were characterized by Rigaku MiniFlex II (Cu Kα, λ=0.15418nm). The Scanning Electron Microscopy (SEM) was characterized by Hitachi SU-8010. The Raman spectroscopy was recorded at room temperature using HORIBA Jobin Yvon Evolution with laser excitation at 532 nm. The electronic property was measured by semiconductor characterization systems (HALL8686). The surface morphologies and work functions of PbPdO2 thin films were characterized using atomic force microscopy (AFM, Bruker Dimension). X-ray photoelectron spectra (XPS) were measured by Thermo Fisher ESCALAB250Xi, and the binding energy of the XPS spectra was calibrated. 3. Results and discussion Fig. 1 shows the XRD patterns of PbPdO2 powder and films. All the samples were well confirmed with the standard PDF (No.38-1357). All samples are single phase with body-centered orthorhombic structure. The diffraction peaks of sample A observed at around 2θ=31.59  , 32.85  , 38.66  can be attributed respectively to (211), (020), and (002) planes of PbPdO2 phase. As for sample B, the corresponding diffraction peaks locate at 31.64, 32.96  , 38.73 . For the location of Intensity/a.u.

(b) Sample C

PDF(No.38-1357)

25 30 35 40 45 50 55 60

2Theta/degree

diffraction peaks, there is just a minor difference between both samples. Apparently, sample A has (211) preferred orientations which is similar to the sample C, while sample B has (002) preferred direction. The XRD patterns of samples A and C are consistent with the previous reports [14]. And up to now, (002) preferred direction in sample B was only observed in pioneering studies from Wang

et al. [3]. From the structure studies, it can be deduced that the conditions of preparation influence evidently the preferred orientation of the PbPdO2 films. The growth dynamics for lattice

mismatch interface is very interesting. MgO epitaxial growth on a Si(001) surface by ultrahigh-vacuum molecular beam epitaxy was investigated [15]. It is found that the epitaxial orientation and crystallinity were strongly dependent on the initial condition of the substrate. When MgO was deposited on a clean Si(001) surface at room temperature a MgO(001) film with a=4.21 Å grew on the Si(001) substrate with two in-plane orientations: MgO[001]//Si [001] and MgO[001]//Si[011]. Since the lattice constants of MgO and Si are 4.21 Å and 5.43 Å, respectively, there is a lattice mismatch of −22.5% for MgO[001]//Si[001]. The surface unit cell length of the Si[011] direction is 3.84 Å so that the lattice mismatch is +8.8% for this orientation. This is the first observation of MgO epitaxy with the latter orientation, which has a smaller mismatch than the former orientation. When the substrate was exposed to O

2

or thermally oxidized, the former

orientation predominantly grew on the substrate. Deposition of Mg on the substrate also produced the former orientation. These results imply that nucleation sites on the initial substrate play an important role in determining the epitaxial orientation. Similarly, the lattice constant of PbPdO2 (002) film with c = 4.66Å, which is near that of Mg (001) with a=4.21 Å. Therefore, it is reasonable that PbPdO2 (002) film can be grew on Si (001) substrate by controlling O2 pressure, substrate temperature and thermal treatment. Moreover, compared to growth of PbPdO2 (002) film grew on Si (001), PbPdO2 (002) films with c = 4.66Å can be more easily grew on the MgO(001) substrate with a=4.21 Å because of less lattice mismatch (+9.7%).

Fig. 2 shows the Raman spectra of PbPdO2 samples. From the figure, it is found that the samples with different preferred orientation have only little difference in Raman spectra, which means that their microstructures were similar. The peaks appearing at around 127cm -1 and 567 cm-1 can be related to the vibration modes in the structure of PbPdO2. But the Raman study on the microstructure of PbPdO2 is seldom, so the explanation of the Raman results needs to be further investigated. The SEM images and EDS spectra of samples A and B are displayed in Fig. 3, respectively. From the figure, the SEM images and EDS spectra of both samples are similar. Moreover, it is found that, the prepared films are dense without remarkable defects, and the grain size with granular structure is about 200nm, which is larger than the one prepared using sol-gel spin coating technique [9-11]. In the EDS spectra, Pb, Pd, O and Mg signals are observed. The Mg signal comes from the MgO substrate. Moreover, it is observed that the Pb:Pd atomic ratios for samples A and B are 0.82:1.00 and 0.81:1.00, respectively, indicating that some Pb vacancies may exist in both samples. This vacancy can result from the volatilization of Pb during the heating process. To determine the valence state of the elements in both samples, the X-ray photoelectron spectra (XPS) were measured. Figs. 4(a)(b) show the XPS spectra of PbPdO2 samples. The high-resolution scans of Pd 3d and Pb 4f are shown in the insert. The peaks for Pd 3d are symmetric and centered at around 336.7eV and 342.2eV, standing for Pd 3d5/2 and Pd 3d3/2 respectively. The two peaks can be attributed to the formation of Pd2+ [11,16,17]. The binding energy of Pb 4f contains 2 peaks, coming from the excitation of Pb 4f5/2 and Pb 4f7/2 respectively [11]. The Pb 4f5/2 spectra (the high BE line) can be separated into two characteristic peaks: a main peak at 142.69eV and a satellite peak at

142.02eV. Similarly, the Pb 4f7/2 (the lower BE line) spectra can be decomposed into two peaks: a main peak at 137.88eV and a satellite peak at 137.15eV. Both main peaks can be attributed to the lattice Pb, while both satellite peaks can be assigned to Pb vacancy. Moreover, the areas of main and satellite peaks for Pb 4f5/2 and Pb 4f7/2 in Fig.4 were calculated. The calculated results indicate that average area ratio of satellite peaks (137.15eV and 142.02eV) and to total spectra (satellite+main peaks) are about 15% and 16% for samples A and B, respectively. Thus, it is obtained that Pb vacancies are about 15% and 16% for samples A and B, respectively. From EDS, it is observed that the Pb:Pd atomic ratios for samples A and B are 0.82:1.00 and 0.81:1.00, respectively, meaning that Pb vacancies are 18% and 19% for samples A and B. By comparison, the estimated values of Pb vacancies for samples A and B from XPS and EDS are almost consistent. The shape of O 1s spectra can be decomposed into three peaks: a main peak at around 529.3 eV and two satellite peaks at around 531 eV and 534 eV. The peak at 529.3 eV can be assigned to the lattice O in PbPdO2. The satellite peaks at around 531eV and 534eV can be assigned to surface absorbed oxygen and hydroxyl, respectively [18,19]. Especially, the peak at 534 eV can be associated with Pb vacancy, and it can be explained as the following: the valence state of Pb shifts from +2 to +0 due to Pb vacancy, which makes the superfluous O2- oxidate to O1-. As a result, the O1- will absorb hydroxyl and O2. Therefore, the surfaces of the samples will absorb hydroxyl and O2 after explosion to the atmosphere. Moreover, XPS spectra of the cleaned surfaces of the samples A and B were measured. It is found that after cleaning, the peak height at around 531eV is evidently reduced, which means that it should corresponds oxygen physical absorption. However, the peaks at 533.6 and 529.29eV are hardly changed. Pb vacancy can give rise to O-1, and O-1 does chemically adsorpt hydrogen. Therefore, the

XPS satellite peaks at 533.6eV can be assigned to surface hydroxyl or O-1. It is reasonable that the hydroxyl with chemical adsorption or O-1 is not easily cleaned. Fig. 5 depicts the AFM images of cleaned PbPdO2 samples. From the figure, it is found that, the grain size is about 200nm and the surface roughness is relatively high, which is similar to the results in Fig. 3 and Ref.[11]. Scanning Kelvin Probe Microscopy (SKPM), a modified version of Atomic Force Microscopy (AFM), is a non-destructive non-contact surface technique that allows imaging two-dimensional profiles of contact potential difference (VCPD) [20-25]. The VCPD is defined as

VCPD 

tip  sample q

(1)

where tip and samle stand for the work function of the conductive tip and the sample, while the q is the electronic charge. Using the equation, we can obtain the work function of the sample by measuring VCPD. Based on Eq.(1) and the measured VCPD, the work functions samle of samples A and B at room temperature are obtained to be 5.338eV and 5.396eV, respectively. As we know, the work function stands for the difference between the vacuum and Fermi level. So the difference of the work function between both samples indicates that the samples with (211) and (002) orientations have different Fermi level, and the Fermi level of the sample A is higher than that of the sample B. Figs. 6(a)(b)(c) show the work function of samples A and B at different temperature, and temperature dependence of the work function for Samples A and B, respectively. From the figure, it is found that the work functions of both samples decrease with increasing temperature. This temperature dependence of the work function reflects the carrier type (n-type or p-type) and the value of the band gap, which can be explained by the mechanism diagram of the energy band picture shown in Fig. 7 [24]. For the intrinsic semiconductor, the Fermi level lies nearly in the middle of

energy gap at T=0; for n-type semiconductor, it does close to conduction band at T=0; while for p-type semiconductor, it does close to valence band at T=0. Doping with donors (or acceptor) can change the position of the Fermi level within the energy gap. The donor (acceptor) doping can shift the Fermi level from the middle of the energy gap toward the edge of the conduction (valence) band. We define Td to be the “ionization temperature” , here kBTd=εd, εd is the binding energy of the donor (acceptor) level). As T<Td, with increasing temperature, the work function decreases, and it will approach saturated. Especially, it can be deduced that, the larger the changed value of work function is, the wider the band gap is. As found in Fig.6(c), the work function decreases with increasing temperature, which displays a typical characteristic of p-type semiconductor. This is consistent with the experimental results from Hall measurements [4,6]. Moreover, the changed values of the work

function in the measured temperature range (30-80℃) for samples A and B are 0.20eV and 0.03eV, respectively, which indicates that sample B possesses a tiny band gap (nearly zero-gap) , while sample A has a relatively big one. Therefore, it is concluded that PbPdO2 with (002) preferred orientation has near zero-gap, while it with (211) preferred orientation has larger band gap. This surface anisotropy band gap is also proved by our first principles band structure calculations [26]. The calculated results show that, PbPdO2 with (002) preferred orientation has near zero-gap (0.030eV), while it with (211) preferred orientation has 0.324eV band gap. Fig. 8 depicts the temperature dependence of the electrical resistivity (  ). From the figure, it is found that the electrical resistivities for samples A, B at room temperature are 23.7Ωcm and 5.6Ω cm, respectively, which has the same order of magnitude in the previous reports [4,5,9,14]. Moreover, one notices that the sample B possesses a lower  than sample A within the temperature range, indicating that PbPdO2 with (002) preferred orientation possesses better electrical conductivity. On the other hand, a wide insulator-metal transition (TMI) at around 370K was found in both samples. The high value of TMI are similar to the experimental that in [9], but much higher than that in other reports [3-7,14]. This clear difference of TMI may be associated with the various microstructures resulting from the different prepared conditions. The underlying mechanism needs further study. 4. Conclusions We prepared PbPdO2 film with different preferred orientation (along (211) and (002) respectively) using PLD technique. SEM, EDS and XPS measured results indicate that, both samples have all granular structure with the grain size of about 200nm and some Pb vacancies exist in both samples. AFM and SKPM measured results mean that PbPdO2 with (002) preferred orientation has

near zero-gap, while it with (211) preferred orientation has larger band gap, which is consistent to calculated those based on first principles. Moreover, the electrical resistivity measured results show that the preferred orientation has an evident influence on the electrical conductivity of PbPdO2, and both sample possess a wide TMI at around 370K.

Acknowledgements This work is supported by the National Science Foundation of China (61574037, 61404029, 11404058, 11274064), Science and Technology Major Projects of Fujian Province (2013HZ0003), Project of Fujian Development and Reform Commission (2013-577).

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Figure Captions:

Fig. 1 XRD patterns of (a) PbPdO2 films, (b) PbPdO2 nano-particles. Fig. 2 Raman spectroscopy of PbPdO2 samples. Fig. 3 SEM images of samples A and B. The inset (upper left) is the film’s EDS spectra. Fig. 4 XPS spectroscopy of PbPdO2 samples. Fig. 5 AFM images of PbPdO2 (a) Sample A and (b) Sample B. Fig. 6. (a) The work function of sample A at different temperature; (b) The work function of sample B at different temperature; (c) Temperature dependence of the work function for Samples A and B. Fig. 7 The mechanism diagram of the energy band picture for PbPdO2. Fig. 8 Temperature dependence of electrical resistivity for samples A and B.

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Pd 3d3/2

Intensity/a.u.

340

Pb 4f 7/2

(j)

Pd 3d5/2

Pb 4f 5/2

Pb 4f 7/2

342

Binding Energy/eV

(i) Intensity/a.u.

Pd 3d3/2

344

Pb 4f 7/2 Pb 4f 5/2

Intensity/a.u.

Pd 3d5/2

Intensity/a.u.

Intensity/a.u.

Intensity/a.u.

(f)

528

Binding Energy/eV

538

530.97 533.59

536

534

532

Binding Energy/eV

Fig.4

530

528

(b)

(a)

Fig.5

Counts/a.u.

30℃ 40℃ 50℃ 60℃ 70℃ 80℃

5.0

Work Function/eV

(a)Sample A

T↑

5.4

Sample A

5.3 5.2 5.1 5.0

30

40

50

60

70

80

Temperature/℃

5.1

5.2

5.3

5.4

5.5

5.6

(b) Sample B

Counts/a.u.

30℃ 40℃ 50℃ 60℃ 70℃ 80℃

5.30

Work Function/eV

Work Function/eV

5.48

Sample B

5.44 5.40 5.36 5.32 5.28

30

40

50

60

70

80

Temperature/℃

5.35

5.40

5.45

5.50

5.55

Work Function/eV

Work Function/eV

5.6

(c)

Sample A Sample B

5.5 5.4 5.3 5.2 5.1 5.0

30

40

50

60

70

Temperature/℃ Fig.6

80

T1






n-type

EC

Ef

Ei

Eg Ef

p-type

EV

Fig.7

80 Sample B(002) Sample A(211)

70

cm

60 50 40 30 20 10 0

150

200

250

300

350

Temperature/K Fig.8

400

450