Corrosion Resistance of Nanosized Silicon Carbide-Rich Composite

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Corrosion resistance of nano-sized SiC-rich composite coatings produced by noble gas ion mixing Adel Sarolta Racz, Zsolt Kerner, Attila Németh, Peter Panjan, László Péter, Attila Sulyok, Gábor Vértesy, Zsolt Zolnai, and Miklos Menyhard ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b14236 • Publication Date (Web): 05 Dec 2017 Downloaded from http://pubs.acs.org on December 11, 2017

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ACS Applied Materials & Interfaces

Corrosion resistance of nano-sized SiC-rich composite coatings produced by noble gas ion mixing A.S. Racz†, Z. Kerner†, A. Nemeth§, P. Panjan‡, L. Peter§, A. Sulyok†, G. Vertesy†, Z. Zolnai†, M. Menyhard†* †

Institute for Technical Physics and Materials Science, Centre for Energy Research, Hungarian

Academy of Sciences, Konkoly Thege M. út 29-33, H-1121 Budapest, Hungary §

Wigner Research Centre for Physics, Hungarian Academy of Sciences, Konkoly-Thege M. út 29-33,

Hungary ‡

Jožef Stefan Institute, Jamova 39, 1000 Ljubljana, Slovenia

ABSTRACT Ion beam mixing has been used to produce silicon carbide (SiC) rich nano layer for protective coating. Different C/Si/C/Si/C/Si(substrate) multilayer structures (with individual layer thicknesses falling in the range of 10-20nm) have been irradiated by Ar+ and Xe+ ions at room temperature in the energy and fluence ranges of 40-120 keV and 1-6x1016 ion/cm2, respectively. The effects of ion irradiation including the in-depth distribution of the SiC produced was determined by Auger electron spectroscopy depth profiling. The thickness of the SiC-rich region was only some nanometers and it could be tailored by changing the layer structure and the ion irradiation conditions. The corrosion resistance of the layers was investigated by potentiodynamic electrochemical test in 4M KOH solution. The measured corrosion resistance of the SiC-rich layers was orders of magnitude better than that of pure silicon and a correlation was found between the corrosion current density and the effective areal density of the SiC. KEYWORDS: room temperature SiC formation, SiC nano coating, composite coating, corrosion protection, potentiodynamic test, ion beam mixing 1 ACS Paragon Plus Environment

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1. INTRODUCTION An increasing demand for sensors that are able to operate at temperatures well above 300°C and often in harsh environments (e.g., in automotive and aerospace applications such as combustion processes, gas turbine control, oil industry etc.) has stimulated the search for silicon derivatives applicable as protective layers. Silicon carbide (SiC) is a material that has attracted considerable interest for a long time, particularly due to its high-temperature strength, thermal shock resistance, good thermal conductivity and its inertness when exposed to corrosive environments 1-5. SiC thin layers can be produced by physical vapor deposition (PVD) and atmospheric pressure chemical vapor deposition (CVD) 3. Xu et al.6 used RF plasma sputtering and a compound SiC target to deposit SiC nanoparticles on Si(100) substrates at the substrate temperature of 350 °C. Chung et al.7 demonstrated the formation of SiC nanoparticles (np) from Si/C/Si multilayer using thermal annealing at 700– 900 °C. The size and density of np–SiC were strongly influenced by the annealing temperature, the Si thickness and layer number. The size of the particles varied from tens to hundreds of nanometers. Daves et al.8 used plasma-enhanced chemical vapor deposition at 400 °C to deposit amorphous SiC thin films and studied the deposition parameters on the film properties. Mukherjee et al.9 deposited liquid polycarbosilane-derived β-SiC and α-SiC coating on silicon wafers at three moderately high temperatures by CVD method. The achieved thickness was between 650 and 1000 nm. These above mentioned methods, however, need elevated temperature, which is disadvantageous for certain substrates 3. Ion beam mixing (IBM) could be a good alternative for producing SiC thin films because it can be applied at room temperature and allows the production of compounds with high heat of formation as it operates far from equilibrium conditions

10-13

. Recently, we have demonstrated that focused ion beam

(FIB) using Ga+ ions can be applied to produce some nanometer thick SiC rich layers at room 2 ACS Paragon Plus Environment

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temperature 14. We have also presented that the produced layer shows an excellent chemical resistance against polysilicon etchants 15. When applying FIB, the irradiated area cannot be large; it is in the range of 10-2 mm2. However, for suitable protective SiC coatings macroscopic area is required. Thus, it is necessary to replace the expensive FIB with another technique that can be readily used for large macroscopic areas. IBM is conformal to this requirement since noble gas ion sources can be easily applied even for wafer-size samples. Noble gas ion irradiation has been already applied to produce SiC rich layers. Harbsmeier et al.16 irradiated a Si/C bilayer system with 350 keV Xe+ ions. 6H–SiC formation was only observed when the target temperature was at least of 600 °C. Prakash et al.17 bombarded a Si (3 nm)/C (2.5 nm) multilayer structure with 40 keV Ar+ ions. Higher fluence ion bombardment (10 × 1016 Ar+/cm2) caused a complete mixing of the layers and compound formation also occurred. Interestingly no information about the thickness of the intermixed zone achieved was provided. Kumar et al.18 exposed a C/Si interface to ion sputtering of 100–1500 eV Ar+ ions, which induced an epitaxial SiC formation. The epitaxial formation was only proved by angle dependent XPS. Battistig et al.19 applied Ar+ and Xe+ irradiation on a C (20nm)/Si (20nm) multilayer system at room temperature which led to SiC formation. Despite the detailed description of the layer formation, only one of the previous studies dealt with the chemical behavior of the produced layer. Gurban et al. 15 performed etching tests by checking the integrity of the layer by optical microscope after applying polysilicon etchant for various durations. The qualitative results show correlation between the fluence of irradiation and chemical resistance. For any practical application, however, a more precise quantitative characterization of the layer is necessary. Electrochemical polarization methods like potentiodynamic provide information about the mechanism and the rate of corrosion in a given environment. 3 ACS Paragon Plus Environment


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Herein, a detailed investigation of the corrosion resistance of nano-size (10-60 nm thickness) SiC rich layers made with IBM by means of potentiodynamic corrosion test is reported. The layers were produced by noble gas IBM (scalable up to wafer size) of pure C and Si layers. The concentration distributions in the sample in pristine as well as in the ion beam-mixed state were determined by AES depth profiling, which revealed SiC formation. By varying the initial layer structure and the irradiation conditions various distributions of SiC was achieved. The potentiodynamic corrosion test was performed in 4M KOH, which is known as a polysilicon etchant. The corrosion resistance measured by the inverse of corrosion current density increased with the amount of SiC produced by the IBM and was much better for any studied sample than that of pure silicon.

2. EXPERIMENTAL 2.1 Producing SiC-rich layers

Two types of C/Si/C/Si/C multilayer structures were made by magnetron sputtering on Si single crystal substrate. The initial structure of the specimens determined by XTEM were C 20 nm/Si 20 nm/C 19nm/Si 23 nm/C 17.5 nm/Si substrate and C 11 nm/Si 22 nm/C 11 nm/Si 22 nm/C 11 nm/Si substrate, respectively

19

. For easier reference the former sample will be denoted as sample 20-20,

while the latter one as sample 10-20. To produce the SiC-rich layers, Ar+ or Xe+ ion irradiation was applied at room temperature. The energy of the Ar+ projectiles was 40 keV and the applied fluences were 1, 3 and 6x1016 Ar+/cm2 while the angle of incidence was 7° with respect to the surface normal. The irradiation took place in a Variantype ion implanter.

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In the case of xenon irradiation the energy of the projectile was 120 keV and the applied fluences were 1 and 3x1016 Xe+ /cm2. The bombardment was performed in the Heavy Ion Cascade Implanter of the Institute for Particle and Nuclear Physics of the Wigner Research Centre for Physics in Budapest. In both cases the ion beam with typically millimeter dimensions was x–y scanned across the full sample surface (1-2 cm2) in order to achieve good irradiation homogeneity within the exposed area. The irradiation conditions, projectiles and energies, have been chosen based on TRIDYN simulation 20 to achieve the desired ion mixing, SiC productions

19-21

. Obviously other projectiles might be also

applied and the proper energies can be determined by TRIDYN simulation. 2.2 AES depth profiling

AES depth profiling was carried out to obtain the depth distributions of the target components and the projectiles before and after the ion irradiation. 1 keV Ar+ ions were used for AES depth profiling with an angle of incidence of 80º with respect to the surface normal. The ion current was kept constant during sputtering. The sample was rotated (6 rev/min) during ion bombardment. These parameters were chosen to minimize the ion bombardment-induced surface and interface morphology changes. The Auger spectra were recorded by a STAIB DESA 105 pre-retarded Cylindrical Mirror Analyzer (CMA) in direct current mode.

2.2.1 Calculation of the concentration and the depth

Generally AES provides elemental information. However, the change of the shape of C Auger peak due to chemical interaction can be utilized to monitor the appearance of compounds, too. The shape and energy of the C (KLL) Auger peak in carbide and graphite phases are strongly different; 5 ACS Paragon Plus Environment


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based on these differences the measured C Auger peak can be decomposed into graphitic and carbide components

14

. Thus, the AES analysis provided the depth distributions of C, Si, SiC, and that of the

bombarding projectile. The relative sensitivity factor method was used for the calculation of the atomic concentrations 22

. The relative sensitivity factors for Si and C were determined from the intensities measured on pure

Si or C layers. For noble gas ions, we chose such a sensitivity factor which properly provides the implanted fluence in case when all the bombarding projectiles remain in the sample. The relative sensitivity factors for C and Si emitted from SiC were determined in a sample after irradiation with Ga+ ions with energy of 30 keV to a fluence of 12 × 1016 ion /cm2. In this case a region was obtained which contained only SiC and Ga; since the relative sensitivity factor of Ga is known, the sensitivity factor for C in SiC form could be derived

14

. The same relative sensitivity factors were applied for each

evaluation of all measured AES spectra. The transformation of the sputtering time to removed thickness (called depth) is described elsewhere

23

. Its brief summary is as follows: The relative sputtering yields of pure C and Si can be

determined from the sputtering times required to remove the known layer thicknesses of the asdeposited Si and C layers. Afterwards, for samples exposed to IBM, in the regions exhibiting mixed C and Si components (independently from their chemical state) we suppose that the total sputtering yield Y is given as 1/Y=XC/YC+XSi/YSi where Yi and Xi are the sputtering yield and concentration of component i, respectively. A single fitting parameter was used to obtain the thickness of the untouched layer.

2.3 Corrosion test


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The growth of the SiC rich layer obviously started at the C/Si interface and its thickness depended on the conditions of IBM. In the majority of the cases only a part of the top C layer was consumed by IBM-induced compound formation and thus, the sample surface contained a carbon layer of various thicknesses. Therefore, before the corrosion test we had to remove the carbon layer which covered the intermixed region. This was achieved by oxidation of the intermixed sample in microwave plasma at 500 °C for 10 minutes. The oxidation procedure removed the upper carbon layer, but stopped at the SiC rich layer. This was confirmed by a subsequent AES depth profiling 15. The electrochemical measurements were carried out at room temperature in a common electrochemical three-electrode glass cell housed in a Faraday cage. The potentiodynamic experiments were performed with a computer-controlled Autolab PGSTAT 30 potentiostat. A saturated calomel electrode (SCE) was used as reference electrode and immersed with the help of a Luggin capillary. All potentials throughout this paper refer to this SCE reference electrode. A platinum mesh was used as the counter electrode. The electrolyte solution was de-aerated 4M KOH prepared from analytical grade chemicals. The samples (with an effective area of 0.8 cm2) were fixed with carbon glue to a stainless steel holder. The contact of the samples with the electrolyte was established by employing the hanging meniscus technique. The solution was purged with argon during 30 min before immersing the sample to reduce the dissolved-oxygen content in the solution. Purging was continued for an additional 30 min after the sample was contacted with the solution. The polarization curves were recorded at constant sweep rate of 5 mV/min in a potential range of (-1.5) – (-0.4) V vs. SCE. The required time for reaching open circuit potential (OCP) was 30 min. The polarization curves were characterized by the Tafel extrapolation method.



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3. RESULTS 3.1 AES depth profiles Figure 1a shows the depth profile of the non-irradiated 20-20 sample. It can be seen that in the pristine sample all interfaces are sharp. The broadening (16%-84%) of the interfaces is in the range of 2-3 nm. This means that the combined effect of the ion bombardment-induced broadening (during depth profiling) and the intrinsic broadening of the interfaces (due to sample production) is in this range. Thus, any higher value of the interface broadening observed after IBM is due to ion bombardment-induced mixing. Figure 1b presents a typical result of IBM; it shows the measured AES depth profile obtained on the same sample as depicted in Figure 1a after an irradiation of 6x1016 Ar+/cm2. It is clear that due to the irradiation serious changes occurred in the sample; namely, an intermixing took place. The first Si layer (below the topmost C layer) practically disappeared; it was consumed by the SiC production. On the other hand, only a part of the second Si layer was converted to SiC. This can be understood considering the projected range of the 40 keV Ar+ being about 40 nm (20-20 sample 40 keV Ar+).

24

The last carbon layer remained more or less untouched. Considering the shape of the SiC distribution e.g at 50 and 70 nm one can conclude that the SiC is growing from the Si/C interface.


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Figure 1. AES depth profiles of the 20-20 sample (a) pristine and (b) irradiated (6x1016 Ar+/cm2, 40keV)

The in-depth distribution of SiC certainly varies with the irradiation conditions. The projectile and its energy determine the projected range, while the fluence determines the strength of the mixing, as it is demonstrated by Figs. 2a and 2b which show the SiC in-depth distributions obtained from 20-20 samples after irradiations by Ar+ (Figure 2 a) and Xe+ (Figure 2 b) of various fluences. The total thickness of the irradiated sample changes due to the irradiation conditions since sputtering and mixing affect the thickness. These changes make the comparison of the as-received depth profile difficult. To make the comparison more straightforward the substrate/layer system boundaries are fitted to be at a nominal 100 nm depth. Comparing the SiC in-depth distributions belonging to 3x1016 Ar+/cm2 irradiation to the one produced by irradiation of 1x1016 Ar+/cm2, a significant increase of the amount of SiC is easily discernible. This can be explained if we consider that the irradiation of 1x1016 Ar+/cm2 consumes only half of the available Si from the first Si layer for making SiC; the additional irradiation continues this process. Similarly, there is some additional layer growth at the deeper Si/C interfaces with the fluence increase. On the other hand, a further increase of fluence from 3x1016 Ar+/cm2 to



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6x1016 Ar+/cm2 does not result in a considerable increase of the amount of SiC. This can be accepted considering that practically all Si has been consumed from the most exposed topmost Si layer by the irradiation of 3x1016 Ar+/cm2. In case of the xenon irradiation (20-20 sample 120 keV Xe+ Figure 2b), more SiC is produced than in the case of argon irradiation with the same fluence due to the higher projected range which is about 55nm 24. The higher energy and projected range cause an almost continuous distribution of SiC because the intermixed regions growing from neighboring interfaces begin to overlap. This overlap is due to the fact that the second interface is more affected by the ion mixing than in case of argon again because of the higher projected range of Xe+.

Figure 2. In-depth distribution of SiC of different irradiations for the 20-20 samples (a) argon 40 keV (b) xenon irradiation 120 keV

In case of sample 20-20, the amount of C (112.8 atoms/nm3) is evidently higher than that of Si (49.9 atoms/nm3); consequently, if all the Si is converted to SiC, still half of the C remains unreacted. If we want to have an intermixed layer without the presence of C, evidently its amount should be decreased; as in sample 10-20. Figure 3a and b show AES depth profiles obtained from the 10-20 10 ACS Paragon Plus Environment

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samples. The projected ranges for 40 keV Ar+ and 120 keV Xe+ in this structure are slightly higher than those in the 20/20 sample being 42 and 57 nm, respectively. Figure 3a presents the in-depth distributions of SiC for the 1x1016 Ar+/cm2, 3x1016 Ar+/cm2 and 1x1016 Xe+/cm2 irradiated 10-20 samples. The substrate/layer system boundaries are fitted to be at the nominal 70 nm depth. It can be seen that comparing the 3x1016 Ar+/cm2 irradiated sample to the 1x1016 Ar+/cm2 irradiated one, a significant increase in the amount of SiC can be observed. Moreover, the distribution is getting continuous. The in-depth distribution of the 1x1016 Xe+/cm2 sample is more or less similar to the 3x1016 Ar+/cm2 one, slight increase in the amount of SiC is noticed. Figure 3b shows that by applying 3x1016 Xe+/cm2 irradiation an almost completely continuous, uniform in-depth distributed SiC layer was achieved. All the Si has been used up to form SiC and only a slight amount of carbon has not been reacted. This means that by applying the appropriate irradiation conditions (projectile, energy, fluence and layer structure), a quasi-continuous SiC layer can be produced at room temperature.

Figure 3. AES measurements of different irradiations for the 10-20 samples (a) SiC in-depth distributions of the 0.5/1/3x1016 Ar+/cm2;1x1016 Xe+/cm2 irradiations (b) Depth profile of 3x1016 Xe+/cm2 120 keV irradiation.



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3.2. Corrosion resistance The potentiodynamic corrosion tests were performed in 4M KOH. The polarization curves were recorded at constant sweep rate of 5 mV/min in a potential range of (-1.5) – (-0.4) V vs. SCE. Figure 4 shows as an example the dependencies of the j=f (E) behavior, namely the Tafel plots, for argon- and xenon-irradiated 20-20 samples. Quantitative information on electrochemical kinetic parameters such as corrosion current densities (jcorr) were extracted from the extrapolation of Tafel lines. The corrosion current density gives information about the degradation degree of the sample 25. The primary result of the Tafel plots is the corrosion rate as expressed in current density, which can be used for calculating the loss rate in other units by using the appropriate form of Faraday’s law as follows:

It M = Ah zF d

(1)

where I is the corrosion current, t is the time period while corrosion takes place, z is the number of electrons involved in the elementary reaction of the corrosion process, F is the Faraday constant (96485 Cmol-1), M and d are the molar weight and the density of the corroding material, respectively; A is the surface area exposed to the corrosion and h is the thickness change during the corrosion process. It can be seen that both sides of Equation 1 indicates the volume of the material involved in the corrosion process. A trivial rearrangement of Eq. 1 leads to the relationship of the corrosion rate (CR) as expressed in current density and the thickness change rate (vcorr):

jcorr M = vcorr zF d

(2)

Eq. 2 is equivalent to the form of CR often used in various studies dealing with the corrosion of silicon carbide 26:


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vcorr = CR = k

jcorr EW d

(3)

where the equivalent weight (EW) is used to replace the M/z ratio and the k constant is adjusted to the unit of the other quantities, comprising also F. It was supposed that the free Si was being preferentially attacked, thus the corrosion reaction is assumed as follows 26: Si = Si4+ + 4e-

(4)

Thus, the other relevant quantities are: d = 2330 kgm-3, and M = 0.028 kgmol-1 and z =4. The polarization resistance (Rp) of different specimens can be calculated by the Stern-Geary equation 27

:

=

. ( )

(5)

where Rp is the corrosion resistance; βa is the anodic Tafel; βc is the cathodic Tafel slope. The Stern– Geary equation illustrates the corrosion current as a value inversely proportional to the polarization resistance. The resulting corrosion current densities and the corrosion rate values measured under the different conditions are given in Table 1. The obtained polarization curves can be seen in Figure 4. First of all, there is little doubt that the positive shift in corrosion potential, the decrease of the current densities and the increase of polarization resistances by orders of magnitude in the case of irradiated samples compared to pure silicon clearly demonstrates the protective character of any of the irradiated system.

Table 1. Summary of corrosion current densities, corrosion rates and polarization resistance values measured on variously irradiated samples of different structures



Layer

Irradiation

Corrosion

Corrosion rate

structure

conditions

current density,

vcorr

jcorr ( µA/cm2)

(µm/year)

20-20

10-20

bulk Si

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Rp (Ω cm2)

1x1016 Ar+/cm2

3.09x10

-1

3.1

2.83x10

5

3x1016 Ar+/cm2

1.02x10

-1

1.0

7.96x10

5

6x1016 Ar+/cm2

1.04x10

-1

1.0

8.55x10

5

1x1016 Xe+/cm2

5.73x10

-2

0.57

1.30x10

6

3x1016 Xe+/cm2

4.36x10

-2

0.44

2.20x10

6

0.5x1016 Ar+/cm2

2.30x10

-1

2.3

2.05x10

5

1x1016 Ar+/cm2

5.93x10

-2

0.60

8.97x10

5

3x1016 Ar+/cm2

2.01x10

-2

0.20

1.85x10

6

1x1016 Xe+/cm2

1.40x10

-2

0.14

2.11x10

6

3x1016 Xe+/cm2

1.31x10

-2

0.13

2.64x10

6

37.6

2.38x10

4

non-irradiated

3.74


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Figure 4. Polarization curves obtained for the irradiated a. 20-20 samples and b. 10-20 samples

We have checked the impact of the corrosion tests on the samples by measuring the Auger depth profiles after the test. In Figure 5 we show the Auger depth profile of 20-20 sample irradiated 6x1016 Ar+/cm2, before and after corrosion test. The profiles were shifted to have the third carbon/Si substrate interface at the same depth. It can be seen that only that part of the sample is missing where the SiC concentration was lower than 20%; the rest of the sample remained unhurt, however.

Figure 5. AES depth profile of the 20-20 sample irradiated by 6x1016 Ar+/cm2 before (lines) and after (symbols) corrosion test 15 ACS Paragon Plus Environment


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4. DISCUSSION Two independent measurements (corrosion test and AES depth profiles) provided two data sets; namely, corrosion resistance and SiC distribution. Our aim is to find the connection between these two data sets. The potentiodynamic corrosion test data clearly show that the noble gas irradiation-induced intermixing considerably improves the corrosion resistance of the sample against KOH compared with the pure silicon. Even in the case of the irradiation with the lowest fluence, the corrosion resistance of the sample was an order of magnitude better than that of pure Si. These findings agree with the results of previous studies, where much more aggressive etching was applied 15. Considering the details, it can be seen that the growth of the SiC rich layer decreases the electrochemical activity of the samples in general; i.e, both the cathodic and the anodic current densities are diminished as the SiC rich layer gets thicker and thicker. This suggests that not the SiC surface determines the corrosion properties, since in this case, the layer thickness above a full monoatomic coverage should exhibit very similar corrosion properties. In addition, Figure 5 shows that the topmost SiC layer (and hence, all layers beneath) remain intact during the corrosion test. This is a clear evidence for that the anodic process cannot be the dissolution of the SiC layer. For explaining the corrosion properties found, it was assumed that some imperfections are responsible for it. The ion mixed layers produced by noble gas irradiation have been carefully studied by XTEM applying resolution of 0.17 nm. These studies failed to show any structural defects, cavities etc. 19. The same was concluded when gallium ion irradiation (which even accumulated in the sample) was applied for making the ion mixed layer even after heavy Ga+ irradiation (4*1016 Ga+/cm2). Thus the appearance of structural defects, cavities due to the irradiations we applied can be ignored. 16 ACS Paragon Plus Environment

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The TEM studies are not sensitive to defects like unreacted Si, "atomistic" structural defects, vacancies etc. On the other hand these "atomistic" defects strongly affects the electrochemical activity. In such a case, the cathodic process (hydrogen evolution) takes place only at such defect spots, and the current density decreases with the decrease of the surface density of the imperfection. The density of the extended defects (i.e., those penetrating throughout the layer) is roughly inversely proportional to the layer thickness. The intact nature of the SiC rich layer (see Figure 5) also shows that its dissolution cannot be responsible for the anodic current in the voltammograms at the positive side of the corrosion potential. If the anodic process is governed by the reaction at the "atomistic" defects, the current cannot grow in accord with the exponential dependence of the electrochemical reaction rate constants since a cavity must grow beneath the protecting layer. The evolution of the current with potential is determined by a large number of factors such the composition of the solution in the cavity and its resistance, passivation of the cavity surface etc. It is common that in such cases 28, a negative deviation of the current density from the exponential function can be seen or the current falls with the potential increase, as observed in Figure 4. In the following we correlate these findings with the SiC distributions obtained by the AES depth profiling. If we consider the corrosion resistance values belonging to samples of similar structure we can conclude that longer irradiation causes higher amounts of SiC, which result in a better corrosion resistance before reaching a threshold value, after which a saturation behavior has been observed. For instance, in the case of the 20-20 samples the argon irradiation caused a significant increase of the amount of SiC when the fluence increased from 1x1016 to 3x1016 Ar+/cm2 but slightly increased for the further increase of the fluence (see Figure 2a). In good agreement with this result, the corrosion current 17 ACS Paragon Plus Environment


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density decreases by a factor of 3 for the samples of 1 x1016 to 3x1016 Ar+/cm2 and only about 10 % from 3x1016 to 6x1016Ar+/cm2 . If we apply Xe+ irradiation with higher energy but with same fluence on the same type of sample (20-20), the SiC distribution gets wider and the samples show even better corrosion resistance than the argon-irradiated ones. On the other hand, despite of the difference between the SiC distributions of the samples irradiated by the 1x1016 Xe+/cm2 and 3x1016 Xe+/cm2, corrosion properties are only slightly different. Thus, it seems that the corrosion resistance before a threshold value is determined by the amount of SiC produced; later a saturation value is reached while the amount of SiC still increases. In case of the 10-20 samples we have observed a very similar behavior. Irradiation by 3x1016 Ar+/cm2 caused a decrease in the corrosion current by a factor of 3 with respect to 1x1016 Ar+/cm2 irradiation, which resulted in a significant increase in the amount of SiC. In case of 1x1016 Xe+/cm2 there is a slight decrease in the corrosion current densities, which can be explained by the similar indepth distribution of SiC. Even though a 3x1016 Xe+/cm2 irradiation leads to a uniform SiC layer, this fact is not reflected in the measured corrosion current densities, as they are practically the same. The rough summary of the experimental results obtained by AES depth profiling is that the corrosion resistance - until a given threshold - increases with increasing amount of SiC; after the threshold the corrosion resistance seems to be saturated and it does not depend on the further increase of SiC. On the other hand it should be emphasized the surface concentration of the irradiated samples seem to be rather similar for all irradiations and thus do not show any correlation with the corrosion resistance. All these findings correlate well with the picture suggested by the electrochemical studies. Based on this agreement it was also supposed that the various imperfections determined the corrosion current.


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From Figure 3, it is clear that the SiC distribution depends on the irradiation conditions and the initial structure of the sample, and that the intermixed layer contains also unreacted Si and C atoms besides the SiC. The unreacted Si and the structural defects offer a lower resistance corrosion path. It is also clear that with increasing amount of SiC, the number of unreacted Si atoms decreases and so does the number of percolating defects, hence the layer´s resistance against corrosion increases. To find a correlation between the corrosion resistance and the amount of SiC is not straightforward since the SiC forms various distributions, the handling of which is rather difficult. Instead taking the SiC distributions we considered the areal densities. This term, which is the integral of the element, compound etc. along the depth (that is the total amount of the chosen element) is frequently used for the characterization of the layers produced by IBM. We know that if the SiC concentration is lower than 20% then the layer cannot resist to the chemical attack (see Figure 5). Moreover, the distributions, see Figure 3, are not continuous except in case of the high fluence Xe irradiation. To account for these facts we define the effective areal density as follows: the integral of SiC from that depth where its amount is higher than 20% until that depth where its concentration goes below 20% again.

Figure 6. Corrosion rate (a) and polarization resistance (b) vs effective areal density for all measured data.



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Figure 6 a and b shows the corrosion rates and the inverse of polarization resistance vs. effective areal density, respectively for all measured data obtained from the 20-20 and 10-20 samples which were irradiated by various ions, energies and fluences. The characteristics of the curves are very similar. It can be seen that they are in good correlation with the effective areal densities even though the SiC distributions are strongly different. The correlation means that the average density of the imperfections can be measured by the effective areal density. The saturation type of behavior (above 1500 mole/nm2) shows that after the threshold value the intrinsic corrosion resistance of the SiC layer is measured and the coating exhibit an impervious nature 29. It is clear that the effective areal density can be used to measure the average number of imperfections in a wide range of irradiation and sample structure parameters. The effective areal density can be calculated, however. First, by applying the TRIDYN code 20, the intermixing is determined. Having the intermixed layer our simple method can calculate the distribution of SiC 19,21. In practical case the required corrosion resistance of the coating layer is given. Knowing this value (based on Figure 6) one can get the desired effective areal density. Using the simulation techniques we can design the layer structure and the necessary irradiation conditions which will result in the necessary effective areal density. This means that we have elaborated a method to produce a properly tailored coating layer.

5. CONCLUSIONS C/Si multilayer systems of various thicknesses were irradiated at room temperature by 40 keV Ar+ and 120 keV Xe+ ions in the fluence range of 0.5-6x1016 ion/cm2. In the intermixed region SiC production was observed. The in-depth distribution of the irradiated samples was determined by AES depth 20 ACS Paragon Plus Environment

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profiling. The amount and distributions of the SiC rich layer could be tailored by varying the energy, fluence, projectile and layer structure. The chemical resistance of the layers was measured by potentiodynamic corrosion test. It was found that all irradiated samples are more resistant to corrosion in alkaline solutions than silicon. The corrosion resistance depended on the thickness of the SiC rich layer. This correlation was explained by supposing that the intermixed layer contains defects, which determine the corrosion rate. Correlation was found between the effective areal density of SiC and the corrosion rate. By selecting the irradiation parameters and layer structure, the chemical resistance of the sample can be tailored.

AUTHOR INFORMATION Corresponding Author *E-mail: [email protected]

ACKNOWLEDGEMENT The work was supported by the GINOP-2.3.2-15-2016-00041, NN 114422 M-ERANET “GRACE” project. Thanks for the discussions with Prof. G. Láng and Prof. R. Schiller. Thanks for Károlyné Payer for oxidation in microwave plasma.

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Corrosion Resistance of Nanosized Silicon Carbide-Rich Composite

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