Abstract
Zn3N2 (ZN) and ZnSnN2 (ZTN) are a promising class of nitride semiconductors for photovoltaic and light-emitting-diode applications due to their particular electrical and optical properties, elemental abundance and non-toxicity. So far, most of the experimental results show the degenerate carrier concentration. However, we find that low-temperature growth of these films in a chamber with ultra-high background vacuum can attain a non-degenerate electrical conductivity. This work provides the recent progress of the electrical properties of ZN and ZTN semiconductor thin films. The origins for the high carrier concentrations in ZN and ZTN have been discussed, demonstrating that non-intentional oxygen and hydrogen-related defects play significant roles in such high carrier concentrations. The strategies to suppress the carrier concentrations have also been addressed, such as ultra-high vacuum conditions and low temperature growth.
1 Introduction
Zn3N2 (ZN) is an attractive nitride composed of only earth-abundant elements. The small effective mass of 0.08 m o (m o is the free electron mass) indicates the potential high electron mobility [1,2]. In addition, a theoretical band gap of 0.84 eV matches the required ones to achieve a theoretically high conversion efficiency in single junction solar cells referring to the Shockley–Queisser limit [3]. These features render ZN to be a promising material for thin-film photovoltaic absorber. However, most of the experimental results show quite high carrier concentrations (1018–1021 cm−3) even without intentional doping [4–8]. Such degenerate electrical properties may prevent its semiconducting applications.
Wurtzite-derived II–IV-N2 semiconductors, such as ZnSnN2 (ZTN) and ZnGeN2, have also received increasing interest for their great potential to complement group-III nitrides (like GaN, AlN and InN) as optoelectronic materials [9–15]. From the viewpoints of both economic and ecological perspectives, II–IV-N2 semiconductors composed of Zn, Sn and Ge are highly attractive, due to the low cost of these elements than Ga and In, as well as their more mature recycling infrastructures [11]. On the other hand, these II–IV-N2 semiconductors are structurally analogous to the group-III nitrides [16,17], which is beneficial to the growth of epitaxial heterojunctions by the existing technologies and substrates. In addition, some physical properties of II–IV-N2 semiconductors are comparable or even superior to group-III nitrides. A typical example among various II–IV-N2 semiconductors is ZTN, which is composed of only earth-abundant elements. The theoretical electron and hole effective masses of ZTN are 0.11 and 1.96 m 0, respectively [2], whereas the same calculation approach gives the electron and hole effective masses of 0.18 and 1.97 m 0 for GaN, respectively [2]. Besides, the medium band gap (1.4–1.8 eV) of ZTN may make up the drawbacks of indium segregation and phase separation in In x Ga1−x N with similar band gap [10,18,19]. These fascinating properties of ZTN qualify it to be a very promising low-cost material for photovoltaic and light-emitting-diode applications [10,16].
Although various research studies on the ZTN material have been performed for nearly one decade, an obstacle that remains to hinder its semiconducting applications is the high electron concentration (1017–1021 cm−3) [9,10,20–22] regardless of ZTN in bulks [23] or in thin films grown by sputtering or molecular beam epitaxy [9,10,24,25,26]. Joint experimental and theoretical studies reveal that the high carrier concentration of ZTN comes from the large amounts of intrinsic Sn-on-Zn (
Although ZTN and ZN have been widely studied, the results are still subject to controversial issues and ambiguities. Particularly, their degenerate electrical properties seriously frustrate their applications. Hence, a critical look at their electrical properties is highly desirable. In this work, we try to summarize the recent progress of the electrical transport mechanisms of ZTN and ZN, in order to reveal their origins of the degenerate carrier concentrations.
2 Experimental details
Zn3N2 thin films were grown on Si (100) and glass substrates by reactive magnetron sputtering in an ultra-high vacuum (UHV) chamber (background pressure of 10−7 Pa, Eiko) connected with a load-lock chamber (background pressure of 10−5 Pa). High purity N2 (Grade 1, 99.99995% purity) was used as the sole reactive gas with a partial pressure of 0.78 Pa. A radio frequency (RF) power supply (Kyosan) with a power of 55 W was connected to a Zn target (2 in. diameter with a purity of 99.99%). Various substrate temperatures (near room temperature, 100 and 150°C) were employed to grow the Zn3N2 thin films. The thickness of thin films was around 400–430 nm. After the depositions, Zn3N2 thin films were transferred and stored in the load-lock chamber to prevent the pollution in atmospheric environments [31]. In the following, these Zn3N2 thin films are referred to as “ZN.”
Two chambers with background vacuum of 10−3 and 10−7 Pa were used to grow zinc tin nitride thin films. The Zn/(Zn + Sn) was determined by electron probe microanalyzer (EPMA, JEOL 8530F) and energy dispersive X-ray spectroscopy (EDS, JEOL JSM-7001F). The O anion composition (at%) is estimated by the EPMA using various standard species (like metallic Fe4N, MgO, Zn and Sn). For the thin films grown in the chamber with the poor background vacuum of 10−3 Pa, the O contents are 2.3–5.7%. In the following, these oxynitride thin films are referred to as “ZTNO.” On the contrary, the O contents of ZTN thin films deposited in the UHV chamber with background vacuum of 10−7 Pa are smaller than the detectable limit, indicating quite a little oxygen contamination. In both cases, thin films were synthesized on glass and Si (100) substrates in high purity N2 atmosphere and no intentional heating was applied to the substrates. Thus, the substrate temperature was at near room temperature. ZTNO thin films were deposited by reactive co-sputtering of Zn and Sn metal targets (2 in. diameter with a purity of 99.99%) connected to two pulsed DC supplies (Pinnacle + Advanced Energy), respectively. A fixed voltage of 440 V is applied in the Sn target, while the voltage on the Zn target is tuned from 240 to 300 V to achieve the various cation compositions of ZTNO thin films. During the growth process of ZTNO thin films, a nitrogen flow rate of 60 sccm was used and the total pressure was kept at 1 Pa. The thickness of ZTNO thin films is 1.7–2.2 μm. More details about the growth of ZTNO films are described in ref. [25]. On the other hand, two RF power supplies (Kyosan) were used to grow ZTN thin films. Zn and Sn metal targets were inclined at an angle of 45°. The input powers on the Sn and Zn targets were tuned at 20–35 and 30–80 W, respectively. ZTN thin films were deposited at various substrate temperatures, including near room temperature, 250 and 350°C. During the growth process of ZTN thin films, the N2 partial pressure and the total sputtering pressure are 1.56 and 1.59 Pa, respectively. More details concerning the growth of ZTN-U thin films can be found in ref. [28]. The phase structures of thin films have been checked by the X-ray diffraction (XRD, Bruker D8 Advance with Cu Kα1 radiation). Electrical properties were determined by resistivity and Hall-effect measurements (Toyo Corp. Resitest 8300) using a van der Pauw configuration. The microstructure was investigated by a transmission electron microscopy (TEM, JEOL 2100F). For this purpose, TEM specimens were prepared by diamond tip cleave.
A series of ZTNO thin films with various Zn/(Zn + Sn) compositions have been deposited on glass substrates by sputtering chamber with poor background vacuum of 10−3 Pa. Here it is worth noting that the intentional inlet of oxygen gas into the sputtering chamber is detrimental to form the crystallized ZTN, as the sputtered Zn and Sn atoms easily react with oxygen to produce the oxide phases. But the poor background vacuum (like 10−3 Pa) may keep the crystal structure of nitride and provide certain amount of oxygen to generate the oxynitride thin films simultaneously. On the contrary, UHV (like 10−7 Pa) will suppress the formation of oxynitride thin films due to quite a little content of oxygen.
3 Electrical and defect properties of ZN thin films
ZN thin films have been grown at different substrate temperatures (near room temperature, 100 and 150°C) by reactive magnetron sputtering in an UHV chamber. The n-type conductivity and carrier concentration (n) in these ZN thin films have been identified by standard Hall-effect measurements. However, it is found that the growth temperature has a significant influence on the magnitude of n (Figure 1). The value of n at 300 K (n 300K) of the ZN thin films deposited at near room temperature is 1.6 × 1017 cm−3. In the cases of growth temperatures of 100 and 150°C, n 300K of 8.2 × 1017 and 1.1 × 1019 cm−3 has been achieved, respectively, indicating that n increases with an increase of the growth temperature. The analyses of carrier concentration as a function of temperature give an activation energy of 0.29 eV in the ZN thin film grown at near room temperature, demonstrating the non-degenerate characteristics. These results seem to show that low temperature deposition yields non-degenerate ZN thin films, while elevating the growth temperature generates significantly larger n [31].
![Figure 1
Temperature dependence of carrier concentration of Zn3N2 thin films grown at different temperatures (near room temperature, 100 and 150°C) (reprinted from ref. [31]).](/document/doi/10.1515/htmp-2022-0028/asset/graphic/j_htmp-2022-0028_fig_001.jpg)
Temperature dependence of carrier concentration of Zn3N2 thin films grown at different temperatures (near room temperature, 100 and 150°C) (reprinted from ref. [31]).
Checking the XRD patterns of these ZN thin films grown at different substrate temperatures (Figure 2) [31], it is found that all thin films exhibit the polycrystalline diffraction features, without amorphous characteristics, even for thin films grown at near room temperature. Elevating the growth temperature from near room temperature to 150°C results in significant enhancement of (400) diffraction intensities in these ZN thin films, demonstrating a preferred growth at higher temperatures. However, the full width at half maximum (FWHM) of (400) diffraction peaks just slightly decrease from 0.30° (near room temperature), 0.28° (100°C) to 0.26° (150°C), implying that the coherent lengths of these ZN thin films slightly increase when increasing the substrate temperature by referring to the Scherrer equation. These XRD features seem to indicate that the slight increase of grain boundary effect at near room temperature growth is not the dominant source for the remarkable reduction of carrier concentration [31].
![Figure 2
XRD patterns of ZN thin films on glass substrates with temperatures of near room temperature, 100 and 150°C. The inset shows the magnified (400) diffraction peaks of the corresponding ZN thin films (reprinted from ref. [31]).](/document/doi/10.1515/htmp-2022-0028/asset/graphic/j_htmp-2022-0028_fig_002.jpg)
XRD patterns of ZN thin films on glass substrates with temperatures of near room temperature, 100 and 150°C. The inset shows the magnified (400) diffraction peaks of the corresponding ZN thin films (reprinted from ref. [31]).
Analyses of defect concentrations based on the defect formation energies estimated from the first-principle calculations indicate that all the native defects (such as
![Figure 3
Temperature dependence of defect concentrations in Zn3N2. Curve (1) shows the concentrations of
O
N
+
{\text{O}}_{\text{N}}^{+}
,
H
i
+
{\text{H}}_{i}^{+}
or both of them fitted to the experimental n
300 K of Zn3N2 thin films grown at near room temperature. Curves (2) and (3) represent the concentrations of
V
N
+
{\text{V}}_{\text{N}}^{+}
calculated by using the formation free energies depending on the temperature and the formation energies of the competing phases at 0 K, respectively. Same data as (2) but for
Zn
i
2
+
{\text{Zn}}_{i}^{2+}
is shown as curve (4) (reprinted from ref. [31]).](/document/doi/10.1515/htmp-2022-0028/asset/graphic/j_htmp-2022-0028_fig_003.jpg)
Temperature dependence of defect concentrations in Zn3N2. Curve (1) shows the concentrations of
Referring to such theoretical calculations, the mechanisms behind the reported high n and our tunable n by different growth temperatures in ZN thin films can be qualitatively explained. The background vacuum with residual gas in the chamber may be an important origin for O/H contaminations. A detailed check on the growth conditions and carrier concentration of ZN thin films shows that most of other reported ZN thin films fabricated in non-UHV chamber (background vacuum >10−5 Pa) possess relatively high n, where much more amount of O/H impurities may remain. This may be due to the low formation energies of donor-type
4 Electrical properties of ZTN thin films grown in an UHV chamber
The electrical properties of zinc tin nitride (Zn1+x
Sn1−x
N2) thin films grown by reactive sputtering in an UHV chamber have been investigated. The fractional Zn atomic percent (at%) to the total cation (at%) in ZTN thin films, expressed by Zn/(Zn + Sn), was used to describe the off-stoichiometry of Zn element. Figure 4 shows the n
300 K of ZTN thin films with various cation compositions deposited at near room temperature, 250 and 350°C. It is demonstrated that both Zn-rich content and low-temperature growth are beneficial for suppressing the carrier concentration. The lowest n
300K of 2.7 × 1017 cm−3 was achieved in thin films grown at near room temperature with a composition of Zn/(Zn + Sn) = 0.72, which are in contrast to the previously formed degenerate thin films with much higher carrier concentrations [28]. As shown in Figure 4, thin films with Zn-poor (Sn-rich) contents exhibit larger carrier concentrations. This may come from the high amount of donor-type
![Figure 4
Carrier concentration as a function of cation composition Zn/(Zn + Sn) in ZTN thin films deposited at various temperatures (near room temperature, 250 and 350°C) (reprinted from ref. [28]).](/document/doi/10.1515/htmp-2022-0028/asset/graphic/j_htmp-2022-0028_fig_004.jpg)
Carrier concentration as a function of cation composition Zn/(Zn + Sn) in ZTN thin films deposited at various temperatures (near room temperature, 250 and 350°C) (reprinted from ref. [28]).
5 A comparative study on ZTN thin films grown in UHV and non-UHV sputtering chambers
A magnetron sputtering chamber with a poor background vacuum of 10−3 Pa is also used to grow zinc tin nitride thin films, in order to perform a comparative study with thin films deposited in the UHV chamber with a background vacuum of 10−7 Pa. The purposes are to develop a comprehensive picture on tuning the electrical properties of zinc tin nitride thin films with excess Zn and O compositions.
First, we show the XRD results of ZTNO and ZTN thin films. The XRD patterns of ZTNO and ZTN thin films with different Zn/(Zn + Sn) compositions are shown in Figure 5(a) and (b), respectively. In addition, a diffraction pattern of ZnO thin film grown in the poor vacuum chamber with the O2 and Ar flow rates of 7 and 25 sccm is depicted as a reference in Figure 5(a). Referring to the theoretical electronic structures of ZTN, it is seen that orthorhombic and wurtzite phases possess the direct band gap of 1.4 and 0.4 eV, respectively [19,27]. However, our ZTNO and ZTN thin films show the optical band gap of 1.3–1.6 eV, implying that the predominant phase of these thin films should be orthorhombic, rather than wurtzite. While the theoretical diffraction patterns of orthorhombic ZTN and wurtzite ZnO are presented in Figure 5(c), which are acquired from the Materials Database of NREL (http://materials.nrel.gov) [33]. Since the energetically favorable orthorhombic ZTN derives from the wurtzite structure, the crystallographic characteristics of ZTN and ZnO are quite similar, yielding the close diffraction peaks (like ZTN (120) and ZnO (002), ZTN (200) and ZnO (100)), as shown in Figure 5(c). Looking at Figure 5(a), the (200) diffraction peaks of ZTNO shift to a lower angle when increasing the Zn/(Zn + Sn) content, whereas the (120) diffraction peaks have a tendency to move to a larger angle. At high Zn/(Zn + Sn) composition, the (120) diffraction peaks of ZTNO thin films are much more close to the (002) diffraction peak of ZnO thin film. Unlike ZTNO films, ZTN thin films with different Zn/(Zn + Sn) contents do not exhibit such shift on (120) diffraction peaks, rather nearly constant (120) peaks regardless of Zn/(Zn + Sn) compositions. These results seem to demonstrate that ZTNO thin films with high Zn content show some crystallographic features of ZnO thin films.
(a) X-ray diffraction patterns of ZTNO thin films on glass substrates with various Zn/(Zn + Sn) (0.53, 0.57, 0.61 and 0.65) compositions grown by a sputtering chamber with the background vacuum of 10−3 Pa. ZnO thin film grown by the same chamber is plotted for comparison. (b) X-ray diffraction patterns of ZTN thin films with various Zn/(Zn + Sn) (0.52, 0.59 and 0.72) compositions grown by a sputtering chamber with the background vacuum of 10−7 Pa. (c) Theoretical diffraction patterns of ZnSnN2 and ZnO (authors’ original work).
Next, dramatic difference in electrical properties between ZTNO and ZTN thin films was observed. Figure 6 shows resistivity (ρ) of ZTNO and ZTN thin films at 300 K. As shown in Figure 6, the ρ of nearly perfect cation composition ZTNO ∼0.53 Zn/(Zn + Sn) thin film is quite close to that of ZTN ∼ 0.52 Zn/(Zn + Sn) thin film. However, their mobility and carrier concentrations exhibit notable discrepancies. The Hall mobilities (μ
H) of ZTNO ∼0.53 Zn/(Zn + Sn) and ZTN ∼ 0.52 Zn/(Zn + Sn) thin films are 0.27 and 1.9 cm2·V−1·s−1, respectively. Their corresponding values of n are 5.9 × 1019 and 1.9 × 1019 cm−3, respectively. The O anion composition of ZTNO ∼ 0.53 Zn/(Zn + Sn) is around 2.3%, whereas the O content of ZTN ∼ 0.52 Zn/(Zn + Sn) is smaller than the detectable limit. Thus, these experimental results seem to show that oxygen impurities from the residual gas in the growth chamber increase n and simultaneously suppress μ
H in the nearly stoichiometric ZTNO thin film. This agrees well with the theoretical prediction that
Resistivity vs cation composition Zn/(Zn + Sn) of ZTNO and ZTN thin films on glass substrates grown by chambers with background vacuum of 10−3 and 10−7 Pa, respectively. The oxygen anion composition of ZTNO thin films is also plotted as a function of Zn/(Zn + Sn) (authors’ original work).
Increasing the Zn/(Zn + Sn) compositions, the ρ of ZTN thin films just enhance a little at Zn-rich compositions, from 0.167 (Zn/(Zn + Sn) ∼ 0.52) to 8.72 Ω·cm (Zn/(Zn + Sn) ∼ 0.72). As reported in our previous work, the slight increase of ρ of ZTN thin films at Zn-rich compositions is mainly due to the reduction of n and the enhancement of μ H [28]. Considering quite a little O content of ZTN thin films grown in the UHV chamber, this seems to indicate that the sole Zn cation engineering cannot tune the ρ of ZTN thin films notably.
On the contrary, the resistivities of ZTNO increase significantly, from 0.38 (Zn/(Zn + Sn) ∼ 0.53) to 2.91 × 105 Ω·cm (Zn/(Zn + Sn) ∼ 0.65) when increasing the Zn/(Zn + Sn) compositions, as shown in Figure 6. Taking the ZTNO ∼ 0.65 Zn/(Zn + Sn) thin film as an example, its resistivity even exceeds the 9.91 × 103 Ω·cm of ZnO thin film deposited in the same chamber. Unfortunately, its n and μ H cannot be identified precisely due to the large noise level for such ZTNO thin films with rich Zn compositions. The O content of ZTNO thin films as a function of Zn/(Zn + Sn) composition is also plotted in Figure 6, which shows similar evolution tendency to ρ of ZTNO. This means the high ρ of ZTNO with rich Zn cations is accompanied by the large amount of O anions. Comparing the electrical properties and O compositions of ZTN and ZTNO thin films, it is demonstrated that the joint incorporation of excess O and Zn compositions reduces n effectively.
Finally, microstructures and spatial distribution of ZTNO and ZTN thin films have been investigated. Figure 7(a) shows a scanning transmission electron microscopy (STEM) top-view image of highly resistive ZTNO thin film with 0.65 Zn/(Zn + Sn) composition, indicating grain sizes of less than 20 nm. Figure 6(b) exhibits EDS line profile analyses of Sn, Zn, O and N elements in the selected region marked by the red line in Figure 7(a). As seen in the bottom panel of Figure 7(b), N distributes homogeneously in ZTNO thin film, whereas the elements of Sn, Zn and O are inhomogeneous. Looking at Sn and Zn, their content evolutions as a function of position surprisingly show opposite tendency. As shown in Figure 7(b), the region with low amount of Sn contains high content of Zn. In addition, O possesses the similar evolution tendency to that of Zn, which seems to indicate that locally inhomogeneous ZnO grains also exist in the Zn- and O-rich regions in ZTNO thin film. Coincidentally, this is consistent with crystal features determined by the XRD patterns in Figure 5. On the contrary, the STEM and EDS line profile analyses of ZTN thin films do not show such non-homogeneous distributions of elements, as shown in Figure 7(c) and (d).
(a) and (b) are STEM image and EDS line profile analyses of ZTNO thin film on Si (100) substrate with 0.65 Zn/(Zn + Sn) composition, respectively. (c) and (d) are STEM image and EDS line profile analyses of ZTN thin film on Si (100) substrate with 0.72 Zn/(Zn + Sn) composition, respectively. The red lines in (a) and (c) represent the selected regions for EDS line profile analyses (authors’ original work).
The tunable electrical properties of ZTNO thin films by O and Zn composition shown in Figure 6 can be understood from the viewpoints of the theoretical defect properties and the microstructure of thin film. The defect properties of oxynitride ZTNO have been investigated using the first-principle calculations by Pan et al. [29]. It is demonstrated that the defect complexes, like (ZnSn + 2ON)0, are of great importance to tune n. These stable defect complexes are composed of defects with opposite charges, such as
6 Conclusion
In this work, we demonstrated the electrical measurements of ZN and ZTN thin films deposited at low growth temperature. It is shown that the low temperature deposition is beneficial to suppress the carrier concentrations of ZN and ZTN thin films, giving rise to non-degenerate electrical conductivity. Theoretical analyses of defect properties of ZN demonstrate that unintentional oxygen-on-nitrogen (
In addition, the electrical properties of ZTN thin films can also be tuned by combining the rich Zn and O compositions. It is found that the Zn-rich ZTN thin films incorporated with more O content grown in a chamber with the poor background vacuum of 10−3 Pa possess much higher resistivities than those deposited by a UHV chamber with background vacuum of 10−7 Pa. These results seem to demonstrate that combining the rich Zn and O compositions together in ZTN thin films can suppress the carrier concentration significantly. These electrical properties are in good accordance with the recently theoretical calculation on the defect properties of ZTN, which indicates that the defect complexes (like (ZnSn + 2ON)0) are of great importance to tune the carrier concentrations. The association between
Acknowledgments
Y. Wang would like to acknowledge the funding from NIMS-ICYS (No. QN3140), the Open Project of State Key Laboratory of Environment-friendly Energy Materials (19kfhg14) and the Qilu Young Scholar at Shandong University.
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Funding information: NIMS-ICYS (No. QN3140), the Open Project of State Key Laboratory of Environment-friendly Energy Materials (19kfhg14) and the Qilu Young Scholar at Shandong University.
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Author contributions: Y.W. carried out main parts of the work and wrote the first version of this paper. Y.W. directed the study. T.O. contributed to partial experiments and lots of analyses for the results. F.A. and J.P. contributed to the partial experiments and discussion. N.O. provided the resources and contributed to the discussion. All authors commented on the manuscript.
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Conflicts of interest: Authors state no conflict of interest.
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Data availability statement: The data that support the findings of this study are available from the corresponding author upon reasonable request.
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