Electronically controlled chemical stability of compound

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Electronically controlled chemical stability of compound semiconductor surfaces Junning Gao, Yeonbae Lee, Kinman Yu, Samuel S. Mao, and Wladek Walukiewicz ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.9b10739 • Publication Date (Web): 13 Aug 2019 Downloaded from pubs.acs.org on August 16, 2019

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ACS Applied Materials & Interfaces

Electronically Controlled Chemical Stability of Compound Semiconductor Surfaces Junning Gao,†‡∥ Yeonbae Lee, ‡ Kin Man Yu,§ Samuel S. Mao,∥and Wladek Walukiewicz*‡⊥ †School

of Materials Science and Engineering, South China University of Technology, Guangzhou 510640, China

‡Materials

Sciences Division, Lawrence Berkeley National Laboratory, Berkeley, California 94720, USA

§Department

∥Department

of Physics, City University of Hong Kong, Kowloon, Hong Kong 999077

of Mechanical Engineering, University of California, Berkeley, California 94720, USA

⊥Department

of Materials Science and Engineering, University of California, Berkeley, California 94720, USA

ABSTRACT: Effects of humid environment on the degradation of semiconductors were studied to understand the role of the surface charge on material stability. Two distinctly different semiconductors with the Fermi level stabilization energy EFS located inside the conduction (CdO) and valence band (SnTe) were selected and effects of an exposure to 85℃ and 85% relative

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humidity conditions on their electrical properties were investigated. Undoped CdO films with bulk Fermi level EF below EFS and positively charged surface are very unstable. The stability greatly improves with doping when EF shifts above EFS and the surface becomes negatively charged. This charge-controlled reactivity is further confirmed by the superior stability of undoped p-type SnTe with EF above EFS. These distinct reactivities are explained by the surface attracting either the reactive OH- or passivating H+ ions. The present results have important implications for understanding the interaction of semiconductor surfaces with water or in general with ionic solutions.

KEYWORDS: CdO; SnTe; Semiconductor stability; Surface charge determined stability; Surface charge layer; Fermi stabilization energy; Moisture degradation 1. INTRODUCTION The progressing miniaturization and increasing surface to volume ratio in modern semiconductor devices greatly emphasizes the importance of the semiconductor surface stability under various environmental conditions. In addition, in some instances, e.g., sensors or water splitting devices, the surface plays a role essential in the device operation and cannot be separated from the ambient. Therefore, there is a need to fully understand the factors affecting chemical reactions on semiconductor surfaces induced by different environments. Different factors have been identified to impact the stability of compound semiconductors upon moisture exposure. The nature and the strength of the chemical bonds are one of the main factors that determine the material stability and the degradation process. It is well known that a prolonged exposure of weakly bound perovskite-type hybrid materials to moisture results in the material decomposition1,2. Also, Hisaka et. al3 found that the degradation of AlGaAs/InGaAs


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pseudomorphic high-electron-mobility transistors started with the oxidization reaction at the drain side corner under high humidity conditions. Comizzoli et. al.4 found that the performance degradation of their avalanche photodiodes upon moisture exposure began with the corrosion of the SiN passivation layer. However, it should be noted that in some instances, the presence of moisture can be beneficial. For example, Guo et. al 5 found that the electric-double-layer oxide-based thin-film transistor gated by granular SiO2 can only realize its full performance at relatively high humidity because only these conditions assure the migration of protons that play an essential role in the device fabrication process. It is well recognized that the stability of semiconductor surfaces as well as semiconductor devices depend on the type of doping and the doping level. Profound differences in stabilities of the n- and p-type GaAs and Si have been observed, although this phenomenon was discovered when studying stability against strong oxidants, acids and alkalis, not the relatively mild humid conditions 3,6-10. It is well established that the surface of n-type GaAs is more resistant to damage than its p-type counterpart in photochemical etching. It was also found that a light illumination significantly increases the etching rate of n-type GaAs. The corrosion has been believed to be activated by the broken bonds associated with holes on the GaAs surface

6-8.

Reineke et. al.

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developed a model claiming that the reaction rates reach the same value when the Fermi level of the holes at the surface for p-type GaAs in the dark equals to the quasi Fermi level in n-type GaAs under illumination. However, in the case of photochemical wet etching of Si in HF solution, it was found that under illumination n-type materials are more reactive than p-type materials. According to Noguchi et. al 9

and Cho et. al 10, this is because the upward band bending at silicon–electrolyte interface, which


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leads to the accumulation of photoexcited holes at the interface. These results suggest that the band bending and the charge state of a semiconductor surface may play an important role in controlling the degradation process of semiconductor materials. However, the actual mechanism of the surface charge state and the surface chemical reactions is still not understood.

Figure 1. The position of

in the band line ups for a range of semiconductors17-23. For each

material, the lower bar indicates the position of the valence band maximum (VBM), the upper bar the position of the conduction band minimum (CBM), and the red dashed line the position of the .

Numerous studies11-15 have shown that surface Fermi energies of semiconductor materials are pinned at the Fermi stabilization energy

which is a common energy reference located at about

4.9 eV below the vacuum level. Although in most semiconductors gap, in some instances the

is located inside the band

can fall into either the conduction or the valence band, as shown

in Figure 1. In these extreme cases, the semiconductors show a clear proclivity for the n-type or ptype doping. The surface Fermi energy is pinned by dangling-bond-like defects whose nature,


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donor or acceptor, is determined by the location of the Fermi energy relative to

. This leads to

a surface electron or hole accumulation (depletion) depending on the location of the semiconductor bulk Fermi energy

relative to

. The charge in the accumulation (depletion) layer

originates from the surface dangling bond defects that are donor-like (acceptor-like) resulting in positively (negatively) charged surface for

). The donor-like and acceptor-like

dangling bonds compensate each other and the surface is neutral when

=

.

Consequently, the surface electrical charge can be controlled by a judicious choice of the material, doping type and the doping level. Thus a surface of undoped CdO with

located

deeply in conduction band12 is positively charged whereas the surface of undoped SnTe with located in the valence band16 is negatively charged. The surface charge can also be controlled by shifting of the Fermi energy with an intentional doping. In this paper we present results of a systematic study of the effects of electronic properties on the reactivity of semiconductor surfaces. We have selected materials with distinctly different surface Fermi level pinning energies and thus also electrostatic properties. We show that the stability of the surface in a humid environment is determined by the surface charge that either attracts or repels reactive species. 3. RESULTS The original undoped CdO layers are continuous films with grain size of a few tens of nanometers. After exposure to moisture, the film is severely damaged showing a rough and bumpy canyon-like morphology with feature sizes of a few hundreds of micrometers. SEM micrographs shown in the supporting information section (Figure S1) illustrate the non-uniform degradation of the samples.


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3.1 The degradation behavior of CdO The effects of exposure to 85℃/RH85% treatment on the structural degradation of CdO films were studied with RBS. Figure 2 shows Cd signals of Rutherford backscattering (RBS) spectra from three CdO samples with different doping levels before and after the moisture exposure. The undoped CdO thin film experienced severe changes after only 1h of exposure, while the sample with the highest electron concentration shows only small changes after much longer15h exposure. A SIMNRA analysis of the RBS spectra indicates no significant loss of Cd atoms in all treated samples, as listed in Table 1. The undoped CdO, exposed to 1 h 85 ℃/RH85% treatment shows significant changes in the RBS spectrum. The Cd signal is significantly reduced and varies with depth. These features suggest that the film becomes O rich with decreasing Cd content towards the surface. Moreover, the low energy side of the Cd signals become non-abrupt, indicating that the film becomes very rough with FWHM roughness on the same order as the film thickness. This is consistent with the SEM micrographs shown in SFigure 1.

Figure 2. The comparison of Cd RBS profiles before and after moisture exposure for the CdO films with different electron concentrations: (a) undoped CdO exposed for 1h, (b) CdO:2%In exposed for 15 h and (c) CdO:4%In exposed for 15h.


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Table 1. The total amounts of Cd in the CdO samples before and after being damaged by the 85 ℃/RH85% moisture, as well as the surface roughness of the samples after the treatment. n[cm-3]

1.08×1020

1.10×1021

1.35×1021

(Undoped)

(2%In doped)

(4%In doped)

Before

5.07x1017 (133 nm)

6.62x1017(177 nm)

8.83x1017(240 nm)

After

4.87x1017 (1 h)

6.5x1017 (15 h)

8.82x1017 (15 h)

Roughness [nm]

>100

45

20

Cd [cm2]

As is seen in Figure 2 increasing of the bulk electron concentration of the CdO with In doping makes the material drastically more stable as only a small reduction of the Cd signal is observed after a prolonged 15 h 85 ℃/RH85% treatment. Thus, less than 50 nm layer on the surface of the film is affected for the 4% In doped sample with an initial n=1.35×1021 cm-3. The effects of the treatments were also evaluated with changes in electrical properties of the films measured by Hall effect. A dramatic 18 times increase of the sheet resistance

is found in the undoped film treated

for 1 h. The In doping greatly reduces the effects of the treatments and

increases about 6 times

for the sample with n=1.1×1021 cm-3 and only 13% for the most heavily doped sample with n=1.35×1021 cm-3. The total amounts of Cd before and after 85 ℃/RH85% treatment for the three CdO samples are also listed in Table 1. The last row in the table gives the estimated roughness of the film after treatment.


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3.2 Electrical properties The Hall effect measurements were used to determine how the electrical properties of the films depend on the treatment time. Shown in Figure 3 are sheet electron density ns, mobility  and sheet resistance RS of the undoped CdO and CdO:In thin film samples as functions of the exposure time. The electron concentrations n of the as grown, undoped CdO thin films are in the range of 1 to 2×1020 cm-3 whereas In doped samples have concentrations above 1021 cm-3. Here we use the sheet electron density ns since after moisture exposure the thickness of the films is not well defined, especially for films with lower n. As is seen in Figure 3, profound decreases in both the ns and  are observed for undoped CdO samples with n in the low 1020cm-3 range when exposed even briefly (≤1 h) to moisture. On the other hand, both nS and μ become much more stable in highly doped films with n>~1021cm-3. Thus, for the sample with n=1.35x1021cm-3, degradation of the electrical properties is negligible even after 30 h moisture exposure. The mobilities of as grown films are quite similar and range from 125 cm-2·V-1·s-1 in the undoped to around 80 cm-2·V-1·s-1 for In doped CdO. The observed decrease of sheet electron concentrations and mobility leads to a distinct increase of the sheet resistance Again, as shown in Figure 3(c),

with increasing exposure time.

of the undoped CdO samples increase significantly faster than

the doped ones, indicating much faster degradation rate.


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Figure 3. The evolution of

(a), mobility  (b),

(c) and the degradation rate

(d) of

undoped and In doped CdO thin films after being exposed to the 85℃/RH85% moisture.

In order to quantify the exposure induced degradation process, we introduce the degradation rate parameter the

. Figure 3(d) shows

calculated using consecutive data points in

dependencies shown in Figure 3(c). Despite significant variations in

for each

sample there are clear trends observed in the effects of In doping. The rate parameters are drastically larger for the undoped samples and they rapidly decrease with increasing electron concentration.


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Figure 4. The evolution of

(a), mobility (b),

(c) and the degradation rate

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(d) of

CdO:Ga after being exposed to the 85℃/RH85% moisture.

In order to demonstrate that the moisture stability effects are dependent on electron concentration but are not dopant specific, we also performed similar studies on CdO thin films doped with other donor species, i.e., Ga (CdO:Ga) with n ranging from ~3.2-9.8×1020cm-3 .The results of the effects of 85℃/RH85% treatment on electrical properties of the CdO:Ga samples are shown in Figure 4. Similar to the In doped samples, exposure to the moisture decreases ns and  and increases

.The

changes depend on the initial electron density ns. Overall, the samples with higher initial n degrade slower than the ones with lower n. Thus, for the sample with n=9.81×1020 cm-3 the values of ns,  and Rs are relatively insensitive to the exposure. The

curves shown in Figure 4(d) indicate

similar variation in degradation rate as with CdO:In samples. The degradation rate parameters are large for the samples with n in the low range of 1020 cm-3 and show a profoundly smaller value for the film with n=9.81×1020 cm-3. It is seen in Figure 4 that the degradation rate is not a monotonic function of the electron concentration in the intermediate electron concentration range of 3.26.0×1020 cm-3.


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3.3 Stability of undoped SnTe The above results on CdO thin films lead to the question if the film degradation is dependent on the nature of the amphoteric native defects responsible for the proclivity of this material for the ntype conductivity24. To explore this issue, we have studied the effect of moisture on SnTe which is a material exhibiting opposite behavior to CdO. Thus, as is shown in Figure 1, SnTe has the VBE located at 0.5 eV above EFS.16 This exceptionally high location of the VBE lowers the formation energy of acceptor defects explaining propensity of SnTe for the p-type conductivity. We have investigated stability of undoped SnTe film with hole concentration p=9.54x1020 cm-3 exposed to the 85 ℃/RH85% treatment. As is shown in Figure 5, the electrical properties of SnTe show negligible changes after prolonged exposure up to 36 h, suggesting the extraordinary stability of the sample under investigation. This is in a clear contrast to CdO whereas is show in Figure 3 undoped samples are extremely unstable.

Figure 5. The evolution of sheet hole concentration degradation rate

(a), mobility (b),

(c) and the

(d) of SnTe after being exposed to the 85℃/RH85% moisture.


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4. DISCUSSION In order to understand the degradation mechanism of these semiconductors, we note that, as is shown in Figure 1, CBE (VBE) of the CdO (SnTe) is located about 1 eV (0.5 eV) below (above) . This accounts for the extreme propensity of CdO (SnTe) material for n-type (p-type) conductivity. Also, it has a profound effect on the sign and the density of the surface charge responsible for pinning of the surface Fermi energy at

as the surface charge is determined by

the location of the bulk Fermi energy relative to semiconductor is related to

in

according to12: 1

(

𝐸𝐹 𝐵𝑇

exp 𝑧 ― 𝑘

∞

𝑛(𝐸𝐹) = 3π2∫0

where

. Electron concentration

[

(

1 + exp 𝑧 ―

is the Boltzmann constant,

,

)

𝐸𝐹 𝑘𝐵𝑇

)]

𝑘3(𝑧)𝑑𝑧

(1)

2

and

are relative to the conduction

band edge. The nonparabolic dispersion relation can be derived from Kane’s two band k·p model 12,25:

𝐸𝑐(𝑘) =

ℎ2𝑘2 2𝑚0

―

𝐸𝑔 2

+

𝐸𝑔 2

( ) ( 2

+

)

𝐸𝑔ℎ2𝑘2 2𝑚𝑒∗

where the conduction band edge electron effective mass 𝑚𝑒∗ is assumed to be 0.21𝑚0

(2) 26

and Eg

is the intrinsic band gap which has been reduced by electron-electron and electron-ion interactions 12,13.

The calculated bulk Fermi energy as a function of the electron concentration plotted relative to is shown in Figure 6 (a). The

coincides with

for electron concentration of

4.9×1020 cm-3. Thus, pinning the surface Fermi energy in undoped low electron concentration CdO with Nbn(EFS)


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pinning of the Fermi energy on the surface results in an electron depletion layer compensated by negative surface charges. The case when Nb=n(EFS) corresponds to the flat band condition, i.e. the surface is neutral with the same density of positive and negative charges. The net charge density on the surface is given by27 𝑁𝑠 =

[

𝜀|𝐸𝐹 ― 𝐸𝐹𝑆|𝑁𝑏 1/2 2𝜋𝑒2

]

(3)

where 𝑁𝑠 is the surface charge density, positive when 𝐸𝐹 ― 𝐸𝐹𝑆 < 0 and negative when 𝐸𝐹 ― 𝐸𝐹𝑆 > 0, 𝑁𝑏 is the bulk concentration, 𝑒 is the electron charge, 𝜀 is the dielectric constant. The depletion (accumulation) layer thickness d is defined as 𝑁𝑠

𝑑 = 𝑁𝑏 = The surface net charge density

[

𝜀|𝐸𝐹 ― 𝐸𝐹𝑆| 1/2 2𝜋𝑒2𝑁𝑏

]

(4)

and depletion (accumulation) layer thickness d as

functions of the bulk electron concentration

of CdO are presented in Figure 6(b). As n

increases, the surface net charge changes from positive to negative and reaches 0 when n=4.9×1020 cm-3, corresponding to the transition from the surface accumulation to depletion condition. For the positively charged surface shown on the left side of the figure, the charge density first slightly increases and then decreases to 0. For the negative charged surface shown on the right side, the net charge increases almost linearly.


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Figure 6. (a) Bulk

relative to

surface net charge density

(

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) versus electron concentrations for CdO; (b) The

and depletion (accumulation) layer thickness

as functions

of the bulk electron concentration n of CdO; (c) The degradation rate rD VS 𝐸𝐹 ― 𝐸𝐹𝑆, where rD is the average of

for each curve shown in Figures 3(d), 4(d) and 5(d); (d) The rD as a function

, showing overlapping between Region I and II.

In the case of SnTe, according to Nishitani et. al.16, the bulk Fermi energy is located at the hole concentration

for

=3×1021 cm-3 whereas it lies at 0.3 eV below VBE in the investigated

sample with p=9.54x1020 cm-3. Using this energy difference and the dielectric constant ε=50 28 in Eq.s 3 and 4 yields Ns = 8×1013 cm-2 for the negative surface charge density and d= 0.83 nm for the depletion width. This negative charge density is equivalent to an intentionally doped CdO sample with bulk electron concentration of n≈8.1×1020 cm-3 that have been also proven to be very stable under the moisture treatment. To relate these calculations to the experimental data we calculate the average degradation rate (rD) which is defined as an arithmetic mean of each of the

curves shown in Figures 4(d) and

5(d). As is seen in Figure 6 (c) there are three distinct regions of rD values. In the region I, the bulk Fermi energy 𝐸𝐹 lies well below

and the rD values are consistently very large in the range


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of a few hundred Ω·sq-1·h-1. In contrast, more than two orders of magnitude lower rD values are found for samples in region III where 𝐸𝐹 is located well above where the 𝐸𝐹 lies in a closer vicinity of

. In the intermediate region II

the rD values range from 50 to 150 Ω·sq-1·h-1. It is

worth noting that there is a significantly larger variation of the rD values in region II. Undoped and lightly doped CdO samples have profoundly high degradation rate. This is because CdO has a weak bond that can be easily broken by reactive OH- ions of water. Our results show that under 85% relative humidity conditions increasing temperature greatly increases the surface reactions and reduces stability of the material. Also, it could be argued that the large density of the positive surface charge affects the water decomposition and increases concentration of OH- ions in the water layer adjacent to a surface of undoped CdO. The most important clue for understanding the degradation mechanism is provided by the results in Figure 6 (b) which clearly show three degradation regions defined by the sign and the magnitude of the surface charge. Thus, in the region I the surface is positively charged with the density of about 4.4×1013 cm-2 whereas a negative surface charge density of more than 5.8×1013 cm-2 is observed in region III. This provides a convincing argument that for a large enough surface charge density the degradation process is controlled by electrostatic interaction of the surface charge with water ions. In the intermediate region II, the surface charge density and thus also the electrostatic effects are largely reduced and the stability becomes more dependent on the quality and uniformity of the surface structure. This, as seen in Figure 6 (c), leads to larger variations in the rD values for different samples. The concept of electronically controlled stability is further supported by the SnTe data showing an excellent stability is achieved in undoped film with very high density of the negative surface charge of 8×1013 cm-2 that places it in the moisture resistant region III in Figure 6 (d). Further, as is seen in Figure 6 (d), the undoped SnTe with nS=8×1013 cm-2 and rD = 0.04


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Ω·sq-1·h-1 is firmly located in the stable region III. Note that Region I is overlapping with Region II in Figure 6 (d), this is because the net surface charge density of CdO does not vary monotonously for n smaller than the saturation concentration. Our results are consistent with previous observations of Millesi et al.

29

who reported using

nanostructured CdO thin film with high electron concentration of 7-8×1020 cm-3 as a photocatalyst for water treatment. CdO with such high electron concentration should be stable since this corresponds to a Fermi level ~0.3 eV above EFS and close to region III in Figure 6 (a). The above results have very important consequences for understanding of the interaction of semiconductor surfaces with water or in general with ionic solutions. Figure 7 shows schematics illustrating the surface band bending and the surface charge for different semiconductor materials. Three possible cases of semiconductor materials are to be considered: (i) EFS > ECBE, (ii) EFS < EVBE and (iii) EVBE< EFS 3x1021 cm-3 will result in a positively charged surface that attracts OH- ions and dissolves more rapidly. Our discovery of the charge-controlled stability of semiconductor surfaces has been developed and demonstrated on extreme cases of semiconductors with EFS located either in the conduction (CdO) or the valence band (SnTe). However, it has also very significant implications for standard semiconductors in which EFS locates in the band gap. In this case the sign of the charge on the surface is determined by the doping type. Thus, as is seen in Figure 7 (e) (Figure 7 (f)) for GaAs with EFS located at about 0.6 eV above the valence band edge p-type (n-type) doping shifts bulk Fermi energy towards VBE (CBE) producing a negatively (positively) charged subsurface layer compensated by positively (negatively) charged defects on the surface. Here, we consider two specific conditions with the bulk Fermi level set at CBM (𝐸𝐹 = 𝐸𝐶, n-type) and at VBM (𝐸𝐹 = 𝐸𝑣, p-type) which give respective ∆E of 0.8 eV and -0.6 eV. Using

(4.7×1017 cm-3) and

(9.0×1018 cm-3) and ε=12.6 33 in Eq.s 3 and 4 yield Ns =2.3×1012 cm-2 for the negative and 8.8×1012 cm-2 for the positive surface charge density, and d= 49.3 nm and 9.8 nm for the depletion widths. The difference in the surface charge for p-type and n-type doped GaAs is consistent with the experimental observations that n-type GaAs is more resistant than its p-type counterpart in the photochemical etching and that above band gap illumination is required to achieve a considerable etching rate for the n-type material 6-8. The illumination creates electrons and holes that flatten the upward band bending and neutralize the negative surface charge that attracted H+ ions stabilizing the surface. The reactive negative ions can diffuse to the neutral surface assisting in etching the material. Our present study has important implications for various applications of semiconductor materials that require a contact of semiconductor surface with ionic liquids. A most prominent


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example is the use of a semiconductor as an optically active layer in photoelectrochemical (PEC) water splitting. In the most common design of a photoelectrochemical device the surface of a standard semiconductor light absorber with the gap of 1.9 to 2 eV is in contact with water32,33. For p-type semiconductor the photoexcited electrons drift to the surface and produce hydrogen in the reaction 2H2O +2e– = 2OH– + H2, whereas in n-type semiconductor photoexcited holes are driven to the surface and split water according to, H2O +2h+ = 2H+ + 1/2 O2. In a typical analysis of these reactions, it is assumed that the semiconductor surface is neutral and does not affect distribution of water ions. However, our results indicate that the surface is positively (negatively) charged for n-type (p-type) absorber affecting distribution of water ions close to the surface. Specifically, the OH- ions formed at interface with p-type semiconductor will diffuse to, and react with the positively charged surface resulting to dissolution of the semiconductor absorber. This explains previously observed rapid corrosion of initially efficient p-type GaInP2 PEC cells34. An interesting case is represented by BiVO4 which is another promising material for PEC water splitting. It has a band gap of 2.48 eV and the CBM at 4.79 eV

35

below the vacuum level that

places the EFS level inside the gap, slightly below the CBM. This band offset configuration implies that lightly n-type BiVO4 samples should have a neutral or slightly positively charged surface susceptible to the corrosion. However, an extrinsic n-type doping of BiVO4 that shifts EF upward above EFS makes the surface negatively charged and a more corrosion resistant material. Indeed, it has been reported that increasing the n-type conductivity results in an increase the charge carrier extraction efficiency and improved performance of BiVO4 PECs35-36. 4. CONCLUSIONS We have systematically studied the degradation of semiconductor materials exposed to humid environment (85℃/85%RH) conditions. Our results show that the surface stability is


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predominantly affected by the sign and density of the surface charge. Undoped CdO with a positively charged surface decomposes very rapidly whereas heavily n-type doped CdO with large density of negative surface charge is very stable. The concept of the charge-controlled surface reactivity is confirmed by an exceptional stability of undoped p-type SnTe that is known to have a highly negatively charged surface. We show that the sign and the magnitude of the surface charge is determined by the location of the bulk Fermi energy relative to the Fermi level stabilization energy EFS and can be easily calculated for all semiconductors. These results can be understood in terms of the reactivity of water ions OH- and H+ with the semiconductor surface. The OH- ions readily react with and remove the metal component of the semiconductor whereas H+ stabilizes the surface by passivating the dangling bonds. Our work has far reaching implications for understanding the reactions at the interfaces between semiconductors and ionic liquids. Also, it offers an interesting opportunity to control semiconductor stability by doping and/or engineering of the band offsets. 5. EXPERIMENTAL SECTION Thin films studied in this report were grown on glass substrates by magnetron sputtering with the details of the deposition process presented elsewhere

23,37-39.

The CdO thin films used were

either nominally undoped or doped with different concentrations of In (CdO:In), Sc (CdO:Sc) and Ga (CdO:Ga). The sample thicknesses ranged from 130 nm to 250 nm. Typical electron concentration of undoped CdO was about 2×1020 cm-3 and that of the doped ones was in the range from 3×1020 to 1.35×1021 cm-3. The SnTe films were undoped with a thickness of 219 nm and a hole concentration of 9 ×1020 cm-3. The samples were cut into the sizes of about 10×10 mm2 square pieces and placed into an environmental chamber for different time durations at 85 ℃ and 85% relative humidity (RH)


20

Page 21 of 29 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


(85 ℃/RH85% treatment). The moisture in the chamber was produced from a demineralized water which was the mixture of 1 volume part of tap water and 5.66 volume parts of deionized water. The sample exposure and evaluation were carried out in a sequential manner. After specific exposure time the samples were taken out of the chamber for evaluation and then put back into the chamber for a consecutive exposure. The electrical properties were measured using Hall effect in the Van der Pauw configuration with 0.55 T magnetic field at room temperature. Thicknesses, compositions and degradation profiles of the films were investigated by RBS using a 3.04 MeV He2+ beam and the data were analyzed using the SIMNRA software package40. The surface morphology was evaluated by scanning electron microscopy (SEM, FEM Quanta FEG 250). The relevant information about the CdO and SnTe thin films is listed in Table 2. Table 2. The thickness d, electron density n, mobility μ and sheet resistance RS of CdO and SnTe samples used in the present study, where the superscript * represents the hole concentration of the SnTe sample. CdO

CdO:In

x [%]

0

0

0

2

2

2

4

4

d [nm]

133

133

240

170

140

215

196

235

n ×10-20 [cm-3]

1.08

1.24

1.97

11.0 11.5

12.0 10.8

13.5

μ[cm2·V-1·s-1]

125.5

89.5

80.4

84.7 77.4

81.8 75.5

82.3

RS[Ω·sq.-1]

34.6

42.4

16.5

3.9

2.9

2.4

5.0

CdO:Ga

3.9

SnTe

x[%]

0.75 0.90

2.0

3.5

6.8

0

d[nm]

200

187

265

210

199

219

n×10-20 [cm-3]

3.63 3.22

5.92

5.00

9.81

9.54*


21


Page 22 of 29

μ[cm2·V-1·s-1]

87.3 99.3

128.6 111.5

67.3

31.5

RS[Ω·sq.-1]

9.8

3.1

4.8

9.5

10.5

5.3

ASSOCIATED CONTENT Supporting Information The Supporting Information is available free of charge on the ACS Publications website at DOI:. SEM micrographs; list of acronyms and symbols. AUTHOR INFORMATION Corresponding Author *Email: [email protected]. ORCID Junning Gao: 0000-0002-7921-1438 Kin Man Yu: 0000-0003-1350-9642 Notes The authors declare no competing financial interest ACKNOWLEDGMENT The work was performed in EMAT Program at LBNL supported by the Office of Basic Energy Sciences, Materials Sciences and Engineering, Division, of the U.S. Department of Energy under Contract No. DE-AC02-05CH11231. J. Gao appreciates the financial support from the China Scholarship Council, National Natural Science Foundation of China (No. 51602105), the Natural


22

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For Table of Contents Only


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Electronically controlled chemical stability of compound

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