An investigation on the microstructure, tribological and corrosion performance of AISI 321 stainless steel carbonitrided by RF plasma process

Surface & Coatings Technology 205 (2010) 674–681

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Surface & Coatings Technology j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s u r f c o a t

An investigation on the microstructure, tribological and corrosion performance of AISI 321 stainless steel carbonitrided by RF plasma process A.M. Abd El-Rahman Physics Department, Faculty of Science, Sohag University, 82524-Sohag, Egypt

a r t i c l e

i n f o

Article history: Received 8 June 2009 Accepted in revised form 6 August 2010 Available online 14 August 2010 Keywords: Carbonitriding Microstructure Friction coefficient Corrosion resistance Wear rate Activation energy

a b s t r a c t An anticorrosive and wear-resistant superficial top layer was distinctly produced in AISI 321 austenitic stainless steel by rf plasma carbonitriding. Details about the elemental distribution profiles and the microstructure of the carbonitrided samples, as a function of treatment temperature, were characterized by glow discharge optical spectroscopy (GDOS), grazing incidence X-ray diffraction (GIXRD) and optical microscopy (OM). The tribological behavior was investigated using a ball-on-disc tribometer. Further the corrosion performance of the treated samples compared to the untreated one was examined using electrochemical tests with potentiodynamic anodic polarization curves. Glow discharge optical spectroscopy indicated that, the nitrogen and carbon contents diffused into the surface during carbonitriding were found to be treatment temperature dependent. The lowest friction coefficient and superior wear resistance in the presence of excellent corrosion resistance were observed for the samples carbonitrided at treatment temperature of 818–863 K. According to Arrhenius plot analysis, the activation energy was 4 eV/atom while the effective diffusion coefficient was found to be 199.3 × 102 μm2/h. © 2010 Elsevier B.V. All rights reserved.

1. Introduction In general, austenitic stainless steel alloys with their excellent corrosion properties are commonly used in biomedical, food and chemical, pulp and paper chemical, petrochemical, heat exchange and nuclear power plant industries [1–4]. The problem of biological contamination in milk industry is much reduced after using the anticorrosive materials of austenitic stainless steels in the development of machinery and transportation tools [5]. Moreover, austenitic stainless steels are much used as an orthopedic implant material in the surgical field when the cost is of major importance of the selection [6]. Among these alloys, AISI 321 stainless steel containing a small amount of titanium is preferred for high temperature applications due to its high resistance to intergranular corrosion and relatively good mechanical properties. However, the wider applications of nearly all austenitic stainless steels are restricted by their relatively low hardness and poor tribological properties. Various surface treatment technologies are applied to improve the surface properties of these alloys and to increase their service life to widen their applications in severe chemical and physical environments. Among these, rf plasma nitriding, carburizing and carbonitriding are established previously for improving the surface properties of various types of stainless steels, iron-based alloys, titanium and titanium aluminum alloys [7–13]. It has been observed that the thickness, mechanical, tribological and corrosion properties of the modified stainless steel layers after a

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plasma surface treatment depend on the plasma processing parameters as well as on the properties of the untreated substrates [14–17]. The modified layers on austenitic stainless steel were found to be composed of complex sub-layers and have a large influence on their surface properties [18–24]. Some authors [25–28] have found a top layer with amorphous/nanocrystalline structure when nitriding austenitic steel and others have denied it [8]. Different mechanisms for plasma treatment of stainless steels have been extensively studied and discussed by many authors [29–33]. However, until today, it is still considered as an attractive subject open for debate to add valuable interpretations particularly to the extremely high growth rate and the microstructure formation. The present work was planned to improve the tribological and corrosion properties of AISI 321 stainless steel using rf plasma carbonitriding. Different characterization techniques were used to investigate the modified layer such as GDOS, GIXRD, OM and a ballon-disc tribometer. The corrosion behavior was performed by the potentiodynamic polarization technique. Carbonitriding mechanism was discussed according to the elemental depth profiles and the microstructure characterization of AISI 321 stainless steel treated at different temperatures. 2. Experimental The material used in this work is an AISI 321 stainless steel. The chemical compositions of the specimen are (in wt.%): Cr-17.7, Ni-9.8, C-0.02, Ti-0.14, Mn-1.9, Si-0.4, P-0.03, S-0.03, Mo 0.32, Co-0.11, and Fe-balance. The specimens were cut from 3 mm thick rolled sheet and

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then machined into small dimensions of 20 mm × 12 mm. The surface of the specimens was polished to surface roughness of about 5 nm investigated by a Veeco Technology, Dektak 8000 measuring system and then cleaned with acetone. The carbonitriding process was carried out using an inductively coupled rf plasma with a continuous mode of operation; details of the plasma system can be found in an earlier publication [17]. Pure nitrogen and acetylene gas mixture (1:1 gas pressure ratio) were introduced into the reactor tube from the base pressure of 0.2 Pa to about 8 Pa measured by means of a capacitance manometer. Samples were processed at different rf plasma power inputs from 350 W to 650 W and at a fixed plasma exposure time of 15 min. Then the carbonitrided sample was left in the evacuated reactor tube until it cooled down to room temperature. The sample temperature was measured during the carbonitriding process by a chromel–alumel thermocouple, which was lightly pressed on the surface of the sample. Carbonitriding process was carried out without chemical or physical pretreatment to remove the native oxide layer from the surface of the “raw” substrates. The treatment temperature of the sample was increased from about 663 K to 963 K while increasing the plasma power from 350 W to 650 W (see column 2 of Table 1). X-ray diffraction analyses were performed on the carbonitrided surface using X-Ray Diffractometer, Siemens D 5000. Cu Kα radiation (λ = 1.540560 Å) was used. The patterns of X-ray diffraction analysis were run in grazing incidence X-ray diffraction (GIXRD) between 30° and 95°, with step interval 0.1° and incidence angle of 2°. Glow discharge optical spectroscopy (GDOS) was used to determine the elemental distribution profiles of the carbonitrided layer. To investigate the cross-sectional morphology, the treated samples were cut into small work pieces using ISOMET™ low speed saw. These small parts were cold mounted as cross-sections and grinded using abrasive grit with different grades that started by 40 and ended by 400 meshes and polished to mirror like using micro polish of alumina suspensions 5 μm. After that, the cross-sectioned samples were etched using Nital etch solution to reveal the surface microstructure under optical microscopy. Finally, the layer thickness was measured by microhardness tester and visibly confirmed by the optical images. Wear and friction measurements were performed at room temperature in laboratory air with a humidity of 39 to 51% using an oscillating ballon-disc type tribometer wear tester without lubrication. A 3-mm ball of cobalt tungsten carbide was moved on the surface at a mean sliding speed of 15 mm/s with a normal load of 8 N. The friction coefficient was recorded automatically during the test. The wear volume was measured using a surface profilometer. The wear rates of all samples were calculated using the equation of K = V/SF where V is the wear volume in mm3, S is the total sliding distance in m and F is the applied load in N. The corrosion tests were performed in a 1 wt.% NaCl solution using the potentiodynamic polarization method. A threeelectrode electrochemical cell was used, the counter and reference electrode were related to Pt and saturated calomel electrode, respectively. The anodic polarization curves were recorded with potential scan rate of 10 mV/s. It is obtained to examine the passivation behavior and to determine the corrosion current density.

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The anodic and cathodic Tafel lines that intersect at Icorr, Ecorr point are tabulated in Table 2. The corrosion test was performed at an ambient laboratory temperature of 25 ± 2 °C and humidity of 42 ± 2%. 3. Results and discussion 3.1. Treatment temperature In the present experiment of carbonitriding, AISI 321 stainless steel substrates were heated up only by the electric field of the inductively coupled rf plasma coil. Carbonitriding process using inductively coupled rf plasma allows fast temperature stabilization and good plasma stability. Therefore, in the beginning of the carbonitriding process including the stage of stabilization, surface treatment of austenitic substrates is reasonably effective [18]. It was observed that, a maximum of two minutes from the beginning of the process is required to attain stabilized plasma and treatment temperature. This can be attributed to the small size of the reactor tube (4 cm diameter) and that of the treated substrates (20 mm × 12 mm × 3 mm). Furthermore, the flat geometry of the AISI 321 substrates exposed to plasma reactions is also considered. Fig. 1 shows a linear relationship of treatment temperature values as a function of plasma power input. It has been observed that the treatment temperature sharply increases as the plasma power increases. From the previous study using the same rf plasma system, it was addressed that plasma power controls the electron energy and plasma ion density which have significant influence on the treatment temperature and the nature of chemical reactions between plasma species and the surface of the immersed substrate [34]. In order to correlate the present results to other previous work, it is more convenient to study the physical and chemical properties of the carbonitrided samples as a function of the treatment temperature. 3.2. Elemental depth profiles and carbonitriding efficiency Fig. 2(a–b) shows the nitrogen and carbon depth profiles after rf plasma carbonitriding as a function of the treatment temperature. A correlation of the elemental depth profiles and the diffusion process with the treatment temperature reflects directly the effect of plasma energy and density on the sample temperature and accordingly on the surface properties. The amount of nitrogen and carbon diffused into the bulk austenitic substrate at temperature ranges of 663 K to 963 K were determined by calculating the area under their elemental profile. Nitrogen and carbon contents at different processing temperatures are calculated relative to their content at 663 K. The obtained result is plotted as a function of treatment temperature in Fig. 3(a–b). It was found that the lowest nitrogen content is observed for sample treated at relatively low treatment temperature of 663 K. Nitrogen content abruptly increases approximately 11.5 times for a sample treated at 723 K compared to the one at 663 K. It increased continuously with the increase of treatment temperature to reach a maximum value at 863 K which equals approximately 16-fold the base value. After that it is gradually decreased again with increasing

Table 1 Preparation and characterization parameters of the untreated and treated samples. Plasma power (W)

Sample temperature (K)

Carbonitrided layer thickness (μm)

Friction coefficient

Relative friction coefficient (100%)

Wear depth (μm)

Track wide (μm)

Wear volume (×105 μm3)

Untreated 350 400 450 500 550 600 650

– 663 723 773 818 863 913 963

– 7 31 38 52 53 53.5 45

1.24 0.75 0.43 0.38 0.38 0.37 0.43 0.48

100 60.48 34.68 30.65 30.65 29.84 34.68 38.71

97.2 33.74 1.26 1.23 1.00 0.72 0.9 0.92

1089 587 184 167 166 166 203 158

1499 242 2.36 2.13 1.9 1.19 1.74 1.52

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Table 2 Corrosion parameters of the untreated and treated samples studied in a 1% NaCl solution. Treatment temperature (K)

Ic (×10−6 A/cm2)

Ec (V)

Untreated 663 723 773 818 863 913 963

1.63 8.78 5.48 3.58 1.91 1.27 4.8 3.93

0.005 − 0.186 0.005 0.023 0.009 0.016 0.022 − 0.037

treatment temperature up to 963 K. The great difference in nitrogen content for the sample treated at 663 K compared to the others is mainly ascribed to the most important effect of the surface oxide layer on the carbonitriding efficiency. On the other hand, the effect of surface porous and microcracks on the carbonitriding efficiency should be considered especially when the treatment process was performed at a relatively elevated temperature [33,35]. It is proposed that the surface vacancies generated by surface instabilities move deeply into the bulk at elevated temperatures and form a highly defective layer with porous and microcracks, enhancing the diffusion process of plasma species into the bulk material [35]. The dependence of the carbon diffusion on the treatment temperature is completely different. High carbon content is detected with a deep diffusion underneath the surface (approximately 7 μm) for the sample treated at low temperature of 663 K. It decreases drastically up to approximately 34% for the sample treated at 723 K compared to that treated at 663 K. Carbon content is increasing progressively with the increase of treatment temperature, the highest content is observed for the treated sample at 963 K. The high carbon content and extraordinary diffusion behavior of carbon at relatively low treatment temperature of 663 K can be attributed to the high diffusivity of carbon compared to that of nitrogen through the oxide layer of austenitic stainless steel [36]. From another side, the nonexistence of a carbon or nitrogen suppersaturation region “plateau region” can't prevent carbon or nitrogen diffusion at low temperature which might be achieved at longer plasma processing time. Otherwise at relatively elevated temperatures and at the beginning of carbonitriding process, the species of carbon/nitrogen diffuses into the surface. Subsequently, a thick nitrided layer is formed which would prevent further diffusion from carbon species into the depth. Similar carbon/nitrogen depth profiles with second carbon peak near to the end of nitrogen suppersaturation region have been previously observed [14,37–39]. However, the gradual increase of the carbon content as increasing treatment temperature can be ascribed to the

Fig. 2. (a–b) Typical GDOS depth profiles of nitrogen and carbon as a function of treatment temperature.

mechanism of concentration gradient [32]. The concentration gradient mechanism means that plasma species diffuse from the higher concentration region (on the surface layer) to the lower one (in the bulk material), the concentration of carbon/nitrogen is found higher in the near surface region compared to underneath. The dependence of plasma reactivity and treatment temperature on plasma conditions provides different supersaturation of carbon in the near surface and different carbon contents diffused into the bulk material. Therefore, it is important to clarify that plasma carbonitriding provides high multiplication of carbon species compared to nitrogen. That is because acetylene molecule dissociates much easier than that of nitrogen which is attributed to the first ionization energy of acetylene (11.4 eV) and carbon (11.260 eV) is lower than that of nitrogen (14.534 eV) [40,41]. The oxygen concentration depth profile is not presented here because the residual oxide layer is nearly the same for all carbonitrided samples treated at a temperature of ≥ 723 K. However, the residual oxide layer is much thicker only for the treated sample at 663 K. It was previously examined that hydrogen has a good capability to eliminate the oxide layer by chemical etching [42]. 3.3. XRD analysis

Fig. 1. Treatment temperature variation as a function of plasma power input.

Fig. 4 shows the XRD pattern of the AISI 321 substrate and samples carbonitrided at different treatment temperatures. It can be seen that the untreated substrate consists of fcc austenite plus bcc ferrite expanded by Δa/a = (a − ao)/a ≈ 0.2%. The formation of ferritic phase might be attributed to the phase transformation due to the mechanical effect applied on the surface during severe polishing of the sample to be mirror-like finish. The ferritic phase was detected in

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content of CrN phase. By increasing the treatment temperature up to 723 K, the detected phases are mainly changed to nitrogen-expanded austenite phase (γN) with higher lattice expansion approximately 6.2% and high content of CrN besides minor reflections of Fe3C. The peaks shift of γN to lower angles and broaden after treatment is attributed to the compressive residual stresses [45,46] and the high density of stacking faults introduced by the incorporation of nitrogen and carbon into the surface region of stainless steels [46,47]. The values of the peak shift allow us to distinguish between the nitrogenexpanded austenite and carbon-expanded austenite as solid solution phases. Nitrogen-expanded austenite phase has higher lattice expansion compared to carbon-expanded austenite [14,39]. The detected phases of the carbonitrided sample treated at 818 K are almost the same of the treated one at 723 K. As previously studied, carbonexpanded austenite phase (γC) can be detected at the end of the nitrided layer achieved at treatment temperature ≥723 K [18–24]. Therefore, the disappearance of γC in the present diffraction patterns is due to low penetration depth of the X-ray at low incidence angle. As a result, it is important to mention that carbon depth profiles analyzed by GDOS support the XRD results. The maximum intensity of the nitrided phases was found to be at treatment temperature of 863 K. The phases of CrN and Fe3C were stable for all treated samples. However, at relatively high treatment temperature (913–963 K), the sample temperature was high enough to partly dissolve the solid solution phases to chemical compound phases; high intensities are observed for chemical compounds [48]. The same detected phases of CrN and γN besides Fe3C have been previously detected in the carbonitrided layer modified by simultaneous implantation of nitrogen and carbon into austenitic stainless steel at 673 K [38]. Mändl et al. have observed small peaks corresponding to CrN at low treatment temperature of 653 K in the nitrided austenitic stainless steel using PIII technique [49]. 3.4. Carbonitriding mechanism Fig. 3. (a–b) Relative nitrogen and carbon contents as a function of treatment temperature.

near surface of AISI 321 and AISI 304 austenitic stainless steel at lower incident angle of 2°. It disappeared at higher incident angles and was not observed after nitriding or carbonitriding which agreed well with the present results [14,43,44]. After carbonitriding at low treatment temperature of 663 K, the treated surface mainly composed of carbonexpanded austenite phase (γC) with lattice expansion of approximately 2.3% and minor reflections from the bulk material besides low

Fig. 4. Grazing incidence X-ray diffractograms of the untreated sample and samples treated at different treatment temperatures.

Carbonitriding process is carried out in order to modify the surface properties of AISI 321 through simultaneous reactions between nitrogen, carbon and carbon-containing species and the substrate surface. Nitrogen species are supplied from nitrogen gas as an only source but carbon and carbon-containing species are abounded from various sources. Acetylene gas is the main source for carbon and carbon-containing species. The hydrocarbon contamination found in the reactor tube contributes slightly in the process of carbonitriding [18]. Sequence usages of the reactor tube in carbonitriding process lead to the deposition of dense hydrogenated carbonitride film mixed with carbon powder. This is considered as another contributing source for carbon and carbon-containing species especially when the samples are treated at high carbon content (50% C2H2). Therefore, other reactions between the plasma species (carbide and nitride species) and the deposited film have to be considered during carbonitriding process. Consequently, the plasma reactivity and treatment temperature which control the elemental depth profiles and the microstructure of the modified layers can be affected. In order to further our understanding of the mechanism of carbonitriding, cross-sectional microstructure is investigated by optical microscopy (see Fig. 5). The cross-sectional micrographs show a uniform compound layer with a sharp edged interface. Three sub-layers are clearly observed corresponding to the treated samples at temperature range of 723–863 K, whereas treated samples at other temperatures are characterized only by two sub-layers. The top layer seems to be composed of amorphous carbide or carbonitride/ crystalline structure. As observed, the thickness of the top layer is treatment temperature dependent and has a maximum value of approximately 2–3 μm at 863 K. The other two sub-layers are mostly composed of nitride and carbide phases mixed with polycrystalline solid solution phases. Czerwiec et al. reported the possible

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which block the inward diffusion of carbon species and form the second carbon peak near to the end of the nitrogen plateau region [50]. Afterward, the near surface region of the carbonitrided layer becomes supersaturated with carbon and nitrogen. At later stage of the carbonitriding process, the supersaturated surface defeats the diffusion process of nitrogen/carbon species into the surface. This promotes the deposition of an amorphous carbide or carbonitride/ crystalline structure and leads to high concentration of nitrogen and carbon in the top layer compared to that investigated underneath. An amorphous carbonitrided layer was previously observed in case of duplex treatment of AISI 304 stainless steel using rf plasma nitriding and carbonitriding for processing time of 10 min for each [24]. However, it was not observed after carbonitriding as a single treatment at processing time of few minutes [9,17]. Other authors observed the formation of amorphous carbonitrided layer in the near surface region at high oxygen content. They obtained results confirming the formation of amorphous/nanocrystalline sub-layers when using gas nitriding, after plasma pretreatment with high power density or during surface treatment process [25–28]. It is probably composed of oxy-nitrides of alloying elements of austenitic steel and contains a significant quantity of graphite-like phase [28]. It was found that high content of oxygen is initially responsible for stabilizing fine crystalline structure (10–30 nm) of oxy-nitrides even at temperature higher than 600 °C [28]. The variation of treated layer thickness as a function of treatment temperature in the plasma carbonitriding process is shown in Fig. 6. It can be seen that, the layer thickness is approximately linear with temperature within the temperature range of 663–818 K. The layer thickness has approximately the value of 53 ± 0.5 μm at 818–913 K. After that, it decreases to 45 μm at 963 K. The carbonitriding rate is calculated using the formula of d2/t, where d is the thickness of the modified layer in μm, and t is the plasma processing time in hour. Fig. 7 shows the variation of carbonitriding rate as a function of the treatment temperature for AISI 321 austenitic stainless steel. It has typically the same behavior of the layer thickness as shown in Fig. 6. The obtained values (1.96 × 102–114.5 × 102 μm2/h) are much higher in comparison to the nitriding rate achieved by plasma immersion ion implantation, 6.16 μm2/h is achieved at a sample temperature of 673 K [51]. It is higher compared to the rate of 46.56 × 102 μm2/h for AISI 304 carbonitrided at sample temperature of 823 K obtained by the same rf plasma technique [14]. The high rate of carbonitriding might be ascribed to the previously discussed mechanisms of concentration gradient and surface porous and microcracks [32,33]. The activation energy should be considered in controlling the thermal diffusion process of nitrogen/carbon species into the austenitic

Fig. 5. Optical cross-sectional micrographs for carbonitrided samples at different treatment temperatures: (a) 723 K, (b) 818 K, (c) 863 K and (d) 913 K.

achievement of two different metastable solid solutions in AISI 316 [18] and Inconel 690 [22,23]. After that, they observed such distinct nitride and carbide sub-layers with a well defined interface in AISI 316L [21]. In the present work, grain boundaries and formed microcracks work well as effective channels for fast diffusion of active nitrogen/carbon plasma species into the austenitic substrate [33]. It was also discussed that adjusting the carbonitriding process using different plasma techniques allows simultaneous diffusion of nitrogen and carbon underneath the surface of austenitic stainless steel [14,37–39]. However, carbon diffusion in austenitic steel is faster than nitrogen diffusion [36]. At the beginning of the carbonitriding process, the nitrogen species start to form a nitrided layer, which pushes the carbon species toward the untreated part. The nitrided layer continues growing through the diffusion of highly active nitrogen species until forming nitrogen suppersaturation region

Fig. 6. Case depths of AISI 321 sample carbonitrided at various treatment temperatures.

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saturation region. However, energy and density of plasma species for the rest of all treated samples are efficient enough to restrict the potential barrier of the native oxide layer. It is also important to address that, the carbonitriding at relatively high treatment temperature allows more porous and microcracks to form in the surface containing the residual oxides. 3.5. Friction coefficient and wear performance The friction coefficient versus track number and sliding time curves of untreated and plasma carbonitrided specimens is shown in Fig. 9. The results of the friction test are illustrated in column 4 of Table 1. The friction coefficient is calculated as a mean value along the total sliding distance. It can be observed that, the friction coefficient of

Fig. 7. Carbonitriding rate as a function of treatment temperature.

substrate. It is calculated according to the Arrhenius equation which is expressed as follows: a d2 kT : = Do e t −E

Where Ea is the effective activation energy and T is the treatment temperature. The carbonitrided rate as a function of the treatment temperature is plotted in an Arrhenius type plot as shown in Fig. 8. A linear relationship was observed, the slope and the intersect of the line represent the effective activation energy of 4 eV/atom and effective diffusion coefficient of 199.3 × 102 μm2/h, respectively. The activation energy is higher than the values of 1.4, 1.87 and 0.44 eV/ atom for nitrogen diffusion in AISI 316 and AISI 304L [52–54]. Therefore, the higher activation energy is mostly the reason for the high carbonitriding rate. In general, the native oxide layer plays a significant role in surface treatment process of metallic alloys. It affects the surface treatment efficiency due to creating a potential barrier, reducing active sites and consequently hindering the diffusion of energetic plasma species into the bulk material [55,56]. In the present study, the oxide layer is much effective for carbonitriding at relatively low treatment temperature of 663 K. This means that, the plasma energy is still low to support fast diffusion process of plasma species into the bulk austenitic substrate. Accordingly it is low to create large-sized microcracks in the surface, achieving, carbonitrided layer with only 7 μm thick without nitrogen

Fig. 8. Arrhenius style plot of the temperature dependent carbonitriding rate. The line is a least squared fit to the data.

Fig. 9. (a–d) Friction coefficient curves for carbonitrided samples compared to the untreated one as a function of track number and sliding time measured under a load of 8 N.

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the untreated sample is higher than that of the carbonitrided samples due to the ductile and low work hardening ability of the untreated sample compared to that of the carbonitrided one. The high oscillations observed in the friction behavior of the untreated sample should be attributed to the interactions between the surface asperities and the induced dynamics of the contacting bodies [57]. Further, as depicted by Larsson et al. [58], self-lubricating of abrasive dust in the wear track is more likely to be a reason for the oscillation of the friction coefficient. However, the relatively low friction coefficient values of the carbonitrided specimens are mainly attributed to the microstructure of the top layer. It is earlier investigated that amorphous compounds such as CN are characterized with lubricating properties due to its graphite-like structure and low hardness [59]. Chromium nitride with different contents is considered for the decreasing of the friction coefficient of the nitrided sample [60]. Last and not least, the plate-like wear debris formed from the adhesive transfer could be also considered [61]. Details are shown by investigating the friction coefficient as a function of sliding time. The friction coefficient at 663 K has low values in the near surface region and slowly increased with the increase of sliding time until reaching stable values at the interface region. After that, breaking the carbonitrided layer is observed at wear depth of approximately 17 μm. This exceeds the thickness of the carbonitrided layer determined by optical micrograph and confirmed by the elemental depth profiles (see Table 1 and Fig. 2). Fig. 9(c and d) shows a friction coefficient that is oscillations free as a function of sliding time. For clarity, friction coefficient values for carbonitrided AISI 321 at various treatment temperatures are relative to the untreated one given in column 5 of Table 1. The friction coefficient of the carbonitrided sample at 663 K decreases by only 60%. A significant decrease approximately 34 ± 4% is observed for sample carbonitrided at treatment temperature ≥723 K. Fig. 10 shows wear rate values of carbonitrided AISI 321 stainless steel as a function of treatment temperature for load of 8 N. The wear rate of the untreated substrate was found to be 23.4 × 10− 6 mm3/Nm at total wear path of 80 m. The wear rate of the carbonitrided layer at 350 W has reduced by only one order compared to that of the untreated material. The carbonitrided layer at higher treatment temperature achieves maximum reduction by more than three orders for sample treated at 863 K in comparison with the untreated one. The relatively high wear rate of carbonitrided sample at low temperature compared to that at high temperature is attributed to a variation in the microstructure with the treatment temperature. Furthermore, the low thickness of the carbonitrided layer without nitrogen saturation region allows the hard ball to penetrate underneath and rapidly wear

the bulk material. The carbonitrided layer at treatment temperature higher than 723 K exhibits very low wear depth in the range of 1.1 ± 0.2 μm compared to 33.7 and 97.2 for sample treated at 663 K and untreated sample, respectively. The wear depth, track wide and wear volume as a function of treatment temperature are shown in Table 1. The superior tribological properties of the treated layer are ascribed mainly to harder nitride phases especially γN [62,63]. Furthermore, the existence of carbide phases and amorphous compounds in the near surface region at higher carbon content reduces the wear rate of the examined samples due to their lower friction coefficient [61]. Moreover, the oxide layer has a double maker role in the wear process of untreated and treated substrates. The oxide particles lead to severe wear resulting in a high friction coefficient for the untreated sample while acting as a lubricator layer, reducing friction in the case of treated samples [64]. Finally it should address the humidity variation (39%–51%) as a minor factor, affecting the wear rate trend and friction coefficient values [65].

Fig. 10. Wear rate of treated samples as a function of treatment temperature measured under a load of 8 N.

Fig. 11. Anodic polarization curves for untreated and treated samples at different treatment temperatures obtained in 1 wt.% NaCl solution.

3.6. Corrosion Fig. 11 displays the anodic polarization curves, showing current (A) as a function of applied potential (V) obtained in the environment of 1% NaCl for untreated and treated samples, at room temperature. For more information about the corrosion performance, the result of corrosion current density and corrosion potential for all tested samples is listed in Table 2. In general, a material with excellent corrosion performance displays low current density at more positive potential. The untreated AISI 321 substrate composed of austenitic phase mixed with minor ferritic provides an excellent region of passivity and superior corrosion resistance. The carbonitrided layer at 663 K with mainly γC and CrN provides the worse corrosion resistance and the highest current density in the passivation region compared to all tested substrates. The corrosion resistance increases continuously with the increase of treatment temperature and reaches a maximum value at treatment temperature of 863 K. The significant improvement of the corrosion resistance for sample carbonitrided at 863 K compared to the reference is accompanied by high current density in the active dissolution region. The variation in the corrosion performance as a function of treatment temperature is mainly attributed to the variation of the microstructure of the carbonitrided layer besides N and C contents [66,67]. It is addressed that the modified surface layers containing “expanded austenite” not adversely affect the good corrosion properties of austenitic stainless steel [68–70]. The

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superficial amorphous carbide or carbonitride/crystalline layer can play a significant role in protecting materials against corrosion. Li and Chung [71] reported a good corrosion performance of ultrasmooth CNx overcoats. Further, the a-CHx films perform better corrosion resistance than the pure carbon film [72]. Carbonitriding which is processed at extremely short treatment time (15 min) is also considered as a contributing factor. It might be affecting the microstructure especially precipitation of CrN on the grain boundaries which is mainly responsible for degradation of corrosion resistance of austenitic stainless steel. Samples treated at temperatures of 863– 963 K show some degradation in the corrosion resistance which is still much lower than that is treated at 663 K. 4. Conclusion Rf plasma carbonitriding technique succeeded to improve the surface properties of 321 AISI austenitic stainless steel. The results tell us that the treatment temperature has a significant influence on the microstructure, tribological and corrosion properties of the carbonitrided layer. Superficial top layer containing amorphous compounds was observed after carbonitriding at treatment temperature ≥723 K. An extremely high effective diffusion coefficient of 199.30 × 102 μm2/h was achieved. This is mainly owing to the high activation energy of 4 eV/ atom. The wear rate of the carbonitrided sample is decreased by more than three orders of magnitude compared to that of the untreated 321 AISI stainless steel. The friction coefficient of the carbonitrided layers is decreased to approximately 30% of the untreated material. The superior improvement of the tribological properties is ascribed to the formation of the hard nitriding phases that mixed with carbides and the existence of superficial carbonitrided top layer. Furthermore, the carbonitrided layers displayed good corrosion performance at relatively moderate treatment temperatures in sodium chloride solution. Considering the friction coefficient, wear resistance and corrosion resistance of AISI 321 stainless steel processed at 818–863 K produce the best overall performance. Acknowledgements The author thanks Prof. Dr. F. M. El-Hossary, Dr. M. Raief and Mr. M. Hammad, Faculty of Science, Sohag University for their help in the processing of the final revision. Thanks to Ion Beam Physics and Materials Research Institute, FZD, Dresden, Germany for facilities that approved for samples evaluation. Thanks for Dr. Ronghua Wei and Barrandy Theis, Southwest Research Institute, San Antonio, Texas for their English language revision and related suggestions. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12]

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