Structure and transport properties of titanium oxide (Ti2O, TiO1+δ, and Ti3O5) thin films

Journal of Alloys and Compounds 786 (2019) 607e613

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Structure and transport properties of titanium oxide (Ti2O, TiO1þd, and Ti3O5) thin films Yunjie Fan a, Chao Zhang a, Xiang Liu a, Yue Lin a, Guanyin Gao a, Chao Ma a, Yuewei Yin a, **, Xiaoguang Li a, b, * a Hefei National Laboratory for Physical Sciences at the Microscale, Department of Physics, and CAS Key Laboratory of Strongly-Coupled Quantum Matter Physics, University of Science and Technology of China, Hefei 230026, China b Collaborative Innovation Center of Advanced Microstructures, Nanjing 210093, China

a r t i c l e i n f o

a b s t r a c t

Article history: Received 30 November 2018 Received in revised form 23 December 2018 Accepted 31 January 2019 Available online 2 February 2019

Titanium oxides have partially filled or empty d orbital and are stable at various oxidation states with different structures and unique properties. Here, three kinds of titanium oxide thin films of hexagonal Ti2O metal, cubic TiO1þd superconductor, and monoclinic g-Ti3O5 semiconductor, were successfully grown on a-Al2O3 substrates by a pulsed laser deposition technique, through ablating a pure titanium target under different oxygen pressures. The electrical resistivities of these films increase with increasing oxygen content. The metallic behaviors of Ti bulk and Ti2O film can be described by the BlocheGrüneisen formula, and the semiconducting behaviors of TiO1þd films in normal state and g-Ti3O5 film obey the variable range hopping and the small polaron hopping conduction mechanisms, respectively. For titanium monoxide TiO1þd (1.05  1þd  1.17) films, increasing oxygen content is accompanied by an increase of disorder, a decrease of electron density of states at the Fermi level, and an enhancement of carrier localization, leading to a suppression of superconductivity. © 2019 Elsevier B.V. All rights reserved.

Keywords: Titanium oxide thin films Pulsed laser deposition Oxygen content Electron energy-loss spectroscopy Superconductivity Transport properties

1. Introduction Titanium and its oxides have attracted significant interest due to their abundant physical properties. For example, titanium can be used in biomedical applications and aerospace industries, because of its high wear resistance, good biocompatibility, and excellent corrosion resistance [1e3]. Titanium monoxide TiO1þd has been proposed as an electrical conductor in microelectronics, due to its low electrical resistivity and good barrier property [4e6]. And titanium dioxide TiO2 is widely used in many fields such as antibacterial applications, water purification, and solar cell applications because of its non-toxic nature, high chemical stability, and outstanding optical and electrical properties [7e9]. Recently, the superconductivities in titanium oxides have drawn much attention. The bulk TiO1þd shows a much higher onset superconducting

* Corresponding author. Hefei National Laboratory for Physical Sciences at the Microscale, Department of Physics, and CAS Key Laboratory of Strongly-Coupled Quantum Matter Physics, University of Science and Technology of China, Hefei 230026, China. ** Corresponding author. E-mail addresses: [email protected] (Y. Yin), [email protected] (X. Li). https://doi.org/10.1016/j.jallcom.2019.01.381 0925-8388/© 2019 Elsevier B.V. All rights reserved.

transition temperature Tc ~5.5 K [10] compared with titanium metal which becomes superconducting below ~0.5 K [11]. Our group recently succeeded in growing cubic TiO epitaxial thin films on aAl2O3 (0001) single crystalline substrates and enhanced its Tc to ~7.4 K [12e14]. The superconductivity was also obtained in Ti2O3 (Tc ~8.0 K) [15], g-Ti3O5 (Tc ~7.1 K), and Ti4O7 (Tc ~3.0 K) films [16,17]. Therefore, the effects of oxygen content on the physical properties of titanium oxide films, such as structure, electricity, especially the exotic superconductivity, are fascinating. It is intriguing and necessary to investigate titanium oxide materials with different oxygen contents. Although single crystalline higher titanium oxides (e.g. Ti2O3, Ti3O5, Ti4O7, TiO2, etc.) have been studied extensively [18e21], little is known about lower titanium oxides (such as Ti2O and TiO1þd) due to their polymorphism and unstable natures during single crystal growth. Besides bulk samples, single crystalline titanium oxide films can be grown by a series of methods including magnetron sputtering [22,23], pulsed laser deposition (PLD) [24,25], chemical vapor deposition [26], and low-energy ion bombardment [27]. Among all deposition techniques, PLD has the capability to grow simple metal oxide films with alterable oxidation states, such as FeOx (FeO, Fe2O3, Fe3O4) [28,29] and TiOx (TiO, Ti2O3, g-Ti3O5, Ti4O7)

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[17,30], and the fine control on stoichiometry was achieved by tuning the oxygen pressure. Aforementioned titanium oxide films prepared by PLD were all using titanium oxide targets (Ti2O3 or TiOx), which is difficult to get lower titanium oxide films because of the difficult control of low oxygen content. Furthermore, TieO bond formation is exothermic, and thus titanium oxide films can be easily produced starting from metallic target [31]. In this work, a pure titanium target was selected for PLD film deposition to obtain titanium oxide films (Ti2O, TiO1þd, and g-Ti3O5) with different oxygen contents. It was found that the oxygen content played a central role in influencing the structure, morphology, and electrical transport properties of titanium oxide films. 2. Experimental Titanium oxide thin films with different oxygen contents were grown on commercial (0001)-oriented a-Al2O3 single crystalline substrates by a PLD technique under different oxygen pressures PO2 as 4  104 Pa (labelled as P-4), 6  104 Pa (P-6), 7  104 Pa (P-7), 8  104 Pa (P-8), and 1  103 Pa (P-10), respectively. The target was titanium metal with a purity of 99.99%. The chamber was evacuated to a base pressure below 5  105 Pa and purged three times with high purity oxygen gas. A KrF excimer laser (l ¼ 248 nm) was used for the ablation. The laser energy per pulse, ablation area, laser energy density, repetition rate, and target-substrate distance were 50 mJ, 2.5 mm2, 2.0 J cm2, 5 Hz, and 4.5 cm, respectively. The films were deposited at 850  C and cooled down to room temperature naturally at the oxygen pressure for growth. X-ray diffraction measurements were carried out using a commercial Panalytical X'pert X-ray diffract meter with the Cu-Ka1 radiation at a wavelength of 1.5406 Å. The thicknesses and surface morphologies of films were measured by cryo field emission scanning electron microscopy (cryo FESEM JSM-6700F) and atomic force scanning probe microscope (AFM MultiMode V). For the structure and chemical composition characterizations, a highresolution transmission electron microscopy (TEM Talos F200X), equipped with a spherical aberration corrector on the condenser lens system, was used to obtain the cross-sectional high angle annular dark field transmission electron microscopy (HAADF-TEM) images and core-level electron energy-loss spectroscopy (EELS). In our measurements, the accelerating voltage and energy resolution of EELS spectra were 200 kV and 0.25 eV, respectively. A four-probe method was used to measure the resistivity in a Physical Property Measurement System (PPMS-9T, Quantum Design). 3. Results and discussion Fig. 1 shows the X-ray diffraction (XRD) patterns of titanium oxide films (P-4, P-6, P-7, P-8, P-10), and the full widths at half maximum (FWHM) values of their rocking curves (see Fig. S1 of the supplementary material) are displayed in the inset of Fig. 1. The diffraction peaks for the film deposited at 4  104 Pa (P-4) correspond to (002), (003), and (004) diffraction peaks of hexagonal Ti2O. According to the Ti2O (002) reflection, d002 ¼ 2.4295 Å is calculated by Bragg's law. With PO2 between 6  104 and 8  104 Pa, the films (P-6, P-7, P-8) are face-centered cubic (FCC) TiO1þd showing (111) and (222) peaks. There is a slight shift of the (111) peak toward a higher diffraction angle with increasing PO2 (see Fig. S2) and d111 values are 2.3939 Å, 2.3921 Å, and 2.3884 Å for P-6, P-7, and P-8, respectively. The magnifications of XRD patterns around Al2O3 (0006) reflections are depicted in Fig. S3, which shows an additional small peak located at the lower detection angle of the (111) peak of P-6 and P-7. This small peak may come from Ti2O (002) reflection and shifts slightly to a higher diffraction angle with increasing oxygen content. Fig. S4 shows the XRD f-scans of

Fig. 1. X-ray diffraction patterns of Ti2O (P-4), TiO1þd (P-6, P-7, P-8), and g-Ti3O5 (P-10) films. Inset: FWHM values as a function of oxygen content.

Ti2O (011) and TiO1þd (200) planes. Upon further increasing PO2 to 1  103 Pa, the (202) and (204) diffraction peaks of monoclinic gTi3O5 (P-10) were observed, indicating the multiple orientation growth of the g-Ti3O5 film with d202 ¼ 2.3878 Å and d204 ¼ 1.4434 Å. Here, using pure titanium target we didn't obtain Ti2O3 film under 850  C, in which the O/Ti ratio is between TiO1þd and Ti3O5. This may be due to a narrow PO2 range at 850  C for Ti2O3 film. Of course, it is easier to obtain Ti2O3 film with TiOx target [15,17,32]. To obtain the in-plane epitaxial relationship between the titanium oxide films and a-Al2O3 substrate, the cross-sectional HAADFTEM images of titanium oxide films and a-Al2O3 substrate were depicted in Fig. 2. It is observed that the out-of-plane (in-plane) growth directions of Ti2O [001] ([110]) and TiO1þd [111] ([112]) are parallel to a-Al2O3 [0001] ([1120]), consistent with the XRD results and previous reports [13,17]. As for g-Ti3O5, the (101) orientated phase was observed in the TEM image with its [101] ([010]) crystal direction parallel to a-Al2O3 [0001] ([1120]), corresponding to the previous reports [16,17]. Based on the XRD and TEM results, the lattice parameters of the obtained titanium oxide films were listed in Table 1, similar to the standard JCPDS cards. To further detect the composition of titanium oxide films, the quantitative EELS analyses were carried out, as shown in Fig. 3(a). For titanium oxides, the fine structures of the Ti-L3,2 and O-K edges reflect the covalent bonding states resulting from the strong hybridization between Ti-3d and O-2p electronic states. The Ti-L3,2 edges of samples P-6 to P-8 consist of two peaks (the lower energy Ti-L3 edge and higher energy Ti-L2 edge) in agreement with the features of TiO1þd [33]. The Ti-L3,2 and O-K edges further split for P10 sample, whereby the number of the additional peaks depend on the valence state, coordination, and site symmetry, corresponding to the Ti-L3,2 and O-K edges of g-Ti3O5 [33,34]. Furthermore, a comparison of the energy positions of Ti-L3,2 peaks reveals that the

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Fig. 2. Cross-sectional HAADF-TEM images of (a) Ti2O film (P-4), (b) TiO1þd film (P-6), and (c) g-Ti3O5 film (P-10), respectively. Insets illustrate the schematic atom arrangements of Ti2O (110), TiO1þd (110), and g-Ti3O5 (101) planes, respectively. The large blue and small red spheres represent titanium and oxygen atoms, respectively. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)

Table 1 Calculated lattice parameters of Ti2O (P-4), TiO1þd (P-6, P-7, P-8), and g-Ti3O5 (P-10) films, similar to the standard JCPDS cards. Chemical composition

ID

a (Å)

b (Å)

c (Å)

Ti2O

P-4 JCPDS No. 73-1116 P-6 P-7 P-8 JCPDS No. 89-3660 P-10 JCPDS No. 76-1066

2.8800 2.9600 4.1464 4.1401 4.1316 4.1735 9.9906 9.9701

2.8800 2.9600 4.1464 4.1401 4.1316 4.1735 5.1000 5.0747

4.8590 4.8300 4.1464 4.1401 4.1316 4.1735 7.1958 7.1810

TiO1þd

g-Ti3O5

Ti-L3,2 onset shifts to higher energy with increasing oxygen content, and the spacing between Ti-L3 and Ti-L2 is about 5 eV, consistent with previous report [33]. Through the quantitative analyses of EELS results (detailed analyses see supplementary material P2) [35e37], the average O/Ti ratios are about 0.49, 1.05, 1.13, 1.17, and 1.62 for the five samples P-4, P-6, P-7, P-8, and P-10, respectively, and the relationship between oxygen content and deposition

oxygen pressure is shown in Fig. 3(b). Though there exists a certain error, the obtained O/Ti ratios are consistent with the oxygen contents of titanium oxide phases indicated by the XRD results. Fig. 4 shows the relationship between Ti-L3,2 peak shift and O/Ti ratio, indicating that the peak shift increases with increasing oxygen content [33]. It should be noted that the oxygen content is not so uniform at different areas, as shown in Fig. S7. Fig. 5 shows the scanning electron microscopy (SEM) and atomic force microscope (AFM) images. The thicknesses of titanium oxide films determined by cross-sectional SEM are 70e80 nm, and the root mean square (RMS) roughnesses measured by AFM are about 1.8 nm (P-4), 1.9 nm (P-6), 3.2 nm (P-7), 4.4 nm (P-8), and 3.4 nm (P10), respectively. In Fig. 5(a), the Ti2O film (P-4) prepared at low PO2 (4  104 Pa) is composed of small and long grains which are compact and dense. The rod-shaped grains are arranged irregularly in the film. As PO2 increased to 6  104 Pa (Fig. 5(b)), the surface morphology of TiO1þd film (P-6) changed obviously by showing much larger irregular grains on the surface. By increasing PO2 to 7  104 (Fig. 5(c)) and 8  104 Pa (Fig. 5(d)), although the irregular grains gradually merge, the grooves appear on the surface,

Fig. 3. (a) EELS spectra of the titanium oxide films at Ti-L3,2 and O-K edges. (b) Oxygen content as a function of deposition oxygen pressure PO2. The error bar represents the calculated deviation based on different standard samples.

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Fig. 4. Ti-L3,2 peaks shift relative to Ti-L3,2 peaks position of Ti2O as a function of oxygen content.

indicating that the crystallization is not uniform and the lattice distortion is growing. In TiO1þd films (P-6, P-7, P-8), the higher the PO2, the more the distortions, and thus the RMS roughness increases as a result. With further increasing PO2 to 1  103 Pa (Fig. 5(e)) for g-Ti3O5 film (P-10), the surface morphology shows different sizes of crystal blocks with well-defined grain boundaries.

Our results indicate that the formed grains of these films become larger with increasing oxygen content, and crystallinity is getting worse and worse, which may be due to those larger dislocations and distortions between grains. Fig. 6(a) shows the temperature dependent resistivities (r-T) for the Ti bulk, Ti2O (P-4), TiO1þd (P-6, P-7, P-8), and g-Ti3O5 (P-10) films. Ti enters a superconducting state below ~0.5 K [11]. Ti2O is one of materials with low temperature coefficient of resistivity, which are of great importance for highly accurate electronic measurement instruments and microelectronic integrated circuits. It is worth noting that there is no sign of superconductivity in g-Ti3O5 film as previously reported [16]. The inset of Fig. 6(a) displays an enlarged view around the superconducting transitions of Ti, P-6, P7, and P-8 specimens. The onset superconducting transition temperature Tc (defined by the resistivity dropping to 90% of the normal state resistivity) gradually decreases from 6.12 K for P-6 to 2.5 K for P-8, and the resistivity at 300 K increases with increasing oxygen content, as shown in Fig. 6(b). The transport behaviors of the titanium oxide films could be described by different transport models, including the BlocheGrüneisen model [38], variable range hopping (VRH) model [39], thermally activated (TA) model [40], and small polaron hopping (SPH) model [41]. For metallic Ti bulk and Ti2O film (P-4), as shown in Fig. 7(a), the resistivities increase with increasing temperature, and follow the BlocheGrüneisen formula [38]:

Fig. 5. SEM and AFM images of titanium oxide films. (a), (bed), and (e) represent the Ti2O (P-4), TiO1þd (P-6, P-7, P-8), and g-Ti3O5 (P-10) films, respectively.

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Fig. 6. (a) Temperature dependent resistivities of Ti bulk and titanium oxide films (P-4, P-6, P-7, P-8, P-10), and the inset shows the enlarged view around the superconducting transitions of Ti bulk and TiO1þd films (P-6, P-7, P-8). (b) Oxygen pressure PO2 dependent resistivity at 300 K and Tc of Ti bulk and titanium oxide films. Squares and stars represent resistivity at 300 K and Tc, respectively.

rðTÞ ¼ rð0Þ þ AðT=QÞm

Qð=T 0

xm   dx; ðex  1Þ 1  ex

(1)

where r(0) and A are temperature-independent constants, m is a fitting parameter, and Q is the Debye temperature defined as Q ¼ 2yphkF/kB (yph is the phonon velocity, kF is the fermi wave vector, and kB is the Boltzmann's constant). The best fitting parameters are m ¼ 4, Q ¼ 421 K for Ti2O and m ¼ 4, Q ¼ 402 K for Ti [42]. While for the semiconductor-like behaviors of TiO1þd films (P-6, P-7, P-8) and g-Ti3O5 film (P-10), we used the variable range

hopping (VRH) model, thermally activated (TA) model, and small polaron hopping (SPH) model to fit the r-T curves from 300 K to 100 K. Fig. 7(b) shows the fitting results of the VRH model [39]:

rðTÞ ¼ r0 expðT0 =TÞ1=4 :

(2)

Here, r0 is temperature-independent, and the i characteristic h temperature T0 is given by T0 ¼ 24= pkB NðEF Þx3l , where N(EF) is the electron density of states at the Fermi level, and xl is the localization length. Fig. 7(c) shows the fitting results of the TA model [40]:

Fig. 7. The fitting results of r-T curves with (a) the BlocheGrüneisen formula for Ti bulk and Ti2O film (P-4) from 1.9 K to 300 K, and (b) the variable range hopping model, (c) the thermally activated model, (d) the small polaron hopping model for TiO1þd films (P-6, P-7, P-8) and g-Ti3O5 film (P-10) from 300 K to 100 K. The points are the experimental data, and the dark gray lines are the fitting results.

612

rðTÞ ¼ r0 expðEa =kB TÞ;

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(3) [3]

where Ea is the activation energy representing the energy to cross the band-gap, which is about Eg/2 (Eg is the band-gap energy) for intrinsic semiconductors [40]. Fig. 7(d) shows the fitting results of the SPH model [41]:

   rðTÞ ¼ r0 T exp Ep kB T ;

(4)

where Ep is the activation energy to overcome the barrier. Comparing the fitting results of semiconductor-like behaviors in different models, the normal states of TiO1þd films (P-6, P-7, P-8) follow the VRH model well, where the fitted values of T0 are 0.6  104 K for P-6, 1.2  104 K for P-7, and 2.8  104 K for P-8, respectively, indicating the decrease of NðEF Þx3l with increasing oxygen content. The decrease of N(EF) should contribute to the increase of resistivity, and the observed decrease of N(EF) is only about 8% as increasing oxygen content from 1.05 to 1.17 in nonstoichiometric titanium monoxide [43]. Because the T0 increases by almost five times, a decrease of localization length xl is mainly responsible for the increase of T0, and leads to an enhancement of carrier localization with increasing oxygen content, reflecting the increase of disorder strength at high oxygen content [44,45]. Moreover, it is clear that the decrease of xl can inevitably localize the Cooper pairs and suppress their formations [46], thus reduce the superconducting transition temperature significantly. The r-T behavior of g-Ti3O5 film (P-10) is more consistent with the SPH conduction, corresponding to the formation of polaron which may result in high resistivity related to the decreasing electron mobility and increasing electron effective mass [47,48].

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4. Conclusions

[17]

In summary, we have successfully grown thin films of different titanium oxides including Ti2O, TiO1þd, and g-Ti3O5 by pulsed laser deposition technique, and studied the influence of oxygen content on the structure and electric transport properties. The Ti2O, TiO1þd, and g-Ti3O5 films show metallic, superconducting, semiconducting properties, respectively. The metallic behaviors in Ti2O can be described by the BlocheGrüneisen formula, and the semiconducting behaviors of TiO1þd in normal state and g-Ti3O5 follow the variable range hopping and the small polaron hopping mechanisms, respectively. For the superconducting TiO1þd films, the increasing oxygen content leads to an increase of disorder, a decrease of N(EF), an enhancement of carrier localization, and thus an increase of resistivity and a decrease of superconducting transition temperature.

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Acknowledgments

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This work was supported by the National Natural Science Foundation of China (51790491, 21521001, and 51622209) and the National Key Research and Development Program of China (2016YFA0300103 and 2015CB921201).

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