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Modification of silicon nitride ceramics with high intensity pulsed ion beams
Materials Science and Engineering A253 (1998) 86 – 93
Modification of silicon nitride ceramics with high intensity pulsed ion beams F. Brenscheidt a, J. Piekoszewski b,c, E. Wieser a,*, J. Langner b, R. Gro¨tzschel a, H. Reuther a a
Forschungszentrum Rossendorf, Institut fu¨r Ionenstrahlphysik und Materialforschung, Postfach 51 01 19, D-01314 Dresden, Germany b Soltan Institute for Nuclear Studies, Pl-05 -400 Otwock/Swierk, Poland c Institute of Nuclear Chemistry and Technology, Dorodna 16, Pl-03 -145 Warszawa, Poland
Abstract Sintered Si3N4 ceramics were irradiated with intense plasma pulses containing metal (Ti, Ni) and nitrogen ions in various proportion. The energy density of the pulse was 6.5 J cm − 2 with a duration in the ms range. The samples were characterised by Rutherford backscattering, Auger electron spectroscopy, scanning electron microscopy and by tribological tests. For Ti deposition the top layer is relatively uniform and consists of TiNx, titanium silicide and possibly metallic Ti. For Ni nodules composed mainly from nickel silicide were observed with an additional thin surface layer containing nickel silicide and metallic Ni. Both for Ti and Ni, the plasma irradiation results in a considerable improvement of the wear resistance. © 1998 Elsevier Science S.A. All rights reserved. Keywords: Ceramics; Silicon nitride; Surface alloying; Plasma-pulse irradiation; Microstructure; Wear properties
1. Introduction The tribological properties of ceramic materials are known to be controlled to a large extent by the presence of surface defects and the propagation of cracks originating at the surface into the bulk. Therefore, in the past efforts have been made to alter the surface properties of ceramics using techniques such as ion implantation and ion-beam or laser-beam mixing of metallic overlayers pre-deposited on the substrate. In particular, these efforts have been aimed at the formation of surface alloys with appropriate properties to improve the surface properties of ceramic materials. In the case of ion implantation and ion-beam mixing techniques, the mixing occurs via collisional effects and/or normal and radiation-enhanced diffusion mechanisms. In the case of pulsed laser mixing, the surface layer is rapidly heated to a high temperature (well in excess of the melting point of even refractory metals) * Corresponding author Tel.: + 49 351 2603096; fax: + 49 351 2603285; e-mail: [email protected]
and remains at a high temperature for a period long enough to make it possible that diffusional mixing occurs. Several examples of successful mixing of metalceramic systems and the improvement of their mechanical properties have been reported in the literature. For instance, Ti was mixed on SiC and Si3N4 [1], Ti on ZrO2 [2] and Ni on Si3N4 [3] using pulsed-laser irradiation and Ni on SiC using both ion-beam and pulsedlaser mixing [4]. However, some systems appeared to be refractory to mixing with ion beams such as Ni on Si3N4 and Ti on ZrO2 [1]. Attempts to mix Zr and Ti on Al2O3 have been unsuccessful for both, laser and ionbeam mixing. The present work reports preliminary results of another approach to surface alloying of ceramic systems. Instead of a two step process, i.e. predeposition and subsequent pulse heating, we propose a single-step process in which deposition and heating (melting) occurs within a single plasma pulse. In this experiment plasma pulses containing nickel or titanium and nitrogen as working gas were applied. Polycrystalline Si3N4 sinter ceramics were chosen as substrate.
0921-5093/98/$ - see front matter © 1998 Elsevier Science S.A. All rights reserved. PII S0921-5093(98)00713-8
F. Brenscheidt et al. / Materials Science and Engineering A253 (1998) 86–93
2. Experimental The investigated material is a gas-pressure sintered silicon nitride ceramic with 5 wt.% Y2O3 and 3 wt.% Al2O3 as sintering additives. The samples were polished to an average roughness of Ra =0.01 mm as determined by a surface profilometer. Before ion treatment the material was cleaned with ethanol in an ultrasonic bath. The samples were irradiated with plasma pulses generated by a rod plasma injector type of accelerator IBIS described in detail elsewhere [5,6]. Here we shall briefly outline the principal operation of this machine. The plasma pulses are generated as the result of a low-pressure, high-current plasma discharge between two concentric, cylindrical sets of rod-type electrodes. The discharge is initiated by applying high voltage after some time delay t from the moment of the injection of the working gas into the inter-electrode space. If t is sufficiently long to allow the injected gas to expand over the whole inter-electrode space, the plasma contains almost exclusively the elements of the working gas. This mode of operation is referred to as pulse implantation doping (PID), typically occurring for t= 190–210 ms. For short t, when there is a steep gradient of the gas concentration in the inter-electrode space, the plasma beam contains a mixture of the elements of the working gas and the metallic electrodes chosen deliberately for a given purpose. This mode is referred to as deposition by pulsed erosion (DPE) and occurs typically for t= 160– 180 ms. In the present experiment the irradiations were performed under the following conditions: The energy density of the pulse was about 6.5 J cm − 2. The power versus time characteristic of the pulses consists of two phases. In the first one, lasting about 1 ms, the mean energy of the ions arriving at the substrate amounts to several keV, in the second phase,
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lasting several ms, the ion energy is of the order of tens of eV. As working gas nitrogen of 99.996 purity was used. The outer 32 and inner 32 electrodes were made of commercial nickel or titanium rods of 260 mm length and 2.0 mm diameter. The delay times t were varied in the range between 160 and 200 ms. Each sample was irradiated with ten pulses. The plasma beam diameter was about 100 mm. To avoid contamination of the plasma with elements other than required, the sample holder was coated either with nickel or titanium foil for Ni-N and Ti-N plasma treatment, respectively. In the case of Ti-N plasma irradiation additional samples of stainless steel AISI 321 were placed nearby the silicon nitride samples to control the experimental conditions. Before and after treatment, the samples were weighed with a microbalance to check the mass gain or loss during the treatment. The topography of the irradiated surfaces was determined by surface profilometry with a Dektak 8000 profilometer and by scanning electron microscopy (SEM) with a Zeiss DSM 962 microscope. The composition of the surface and the surface near layer was investigated by Auger electron spectroscopy (AES) combined with sputter etching using a Fisons Microlab 310 F spectrometer and EDX measurements with the SEM. Additionally, Rutherford backscattering spectroscopy was applied to determine the depth distribution of the metal ions. 1.5 or 1.7 MeV He + ions and a backscattering angle of 170° were used. The hardness of the samples was measured with a Shimadzu DUH 220 dynamical microhardness tester. A special software [7] allowed us to correct the data for the elastic deformation of the surface, for deviations of the diamond tip from the ideal geometry and to calculate a depth depending hardness. The tribological tests were carried out with a ball-on-disc type tribometer which also allowed to determine the friction coefficient. The ball was from untreated silicon nitride loaded with 2N. The tests were carried out in normal air at humidity of 60–65%. The typical sliding distance was about 40 m.
3. Results
3.1. Weight loss
Fig. 1. RBS spectra of samples irradiated with Ti in dependence on the delay time t.
The difference in weight before and after processing of the samples irradiated with Ti shows that in all cases there is a mass loss Dm, in average 359 8 mg cm − 2/pulse. No correlation was found between the amount of mass change and the delay time t. The situation is quite different for the stainless steel samples simultaneously treated with the Si3N4 substrates. The steel samples show a clear dependence between mass difference and t. A mass gain is observed for small t where the content of Ti in the plasma leads to the
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Fig. 2. SEM micrographs of samples irradiated with Ti (a −t =160 ms, b− t = 200 ms).
expected metal deposition. For longer t the plasma consists only of nitrogen and a mass loss by evaporation from the surface is detected.
3.2. Surface composition 3.2.1. Titanium deposition Fig. 1 shows the RBS spectra of the samples coated with titanium by ten metal-nitrogen plasma pulses, respectively, for different delay times. The samples with short delay times show a sharp surface peak of the deposited metal. Therefore, a thin metal rich layer exists at the surface. This peak decreases and becomes narrower
for longer delay times, indicating that less material is deposited on the surface. The area density of retained Ti is 1.0 × 1017 cm − 2 and 2.4 × 1016 cm − 2 for t=160 and 170 ms, respectively, and is below the detection limit of RBS at 200 ms. The shift of the silicon edge to higher energies for longer delay times correlates with the decreasing thickness of the titanium layer. The steep slope of the silicon edge and the narrow Ti peaks indicates a high degree of uniformity of the titanium layer. No indication of mixing between the deposited metal layer and the substrate is found. The small peaks on the high energy side of the Ti peak in Fig. 1 are probably due to the deposition of ablated components from the steel samples.
F. Brenscheidt et al. / Materials Science and Engineering A253 (1998) 86–93
The SEM micrographs presented in Fig. 2 show no significant differences comparing the samples irradiated with the shortest and the longest t. The visible structures are somewhat more smooth for t = 160 ms. This might be explained by the thin metal cover layer which is not present for t =200 ms. AES has been used to get more insight into the composition of the metallic surface layer. The formation of TiNx can be seen from the shape of the Ti(LMM) peak at 419 eV. Fig. 3 (upper part) compares this Ti(LMM) peak of Ti metal, TiN and of the sample irradiated with t =170 ms. The peak of metallic Ti is centered at about 419 eV. TiNx is characterized by an additional shoulder at about 415 eV. The energy spectrum of the irradiated ceramic sample shows clearly this shoulder characteristic for TiNx. The coexistence of metallic Ti is probable, but it cannot be confirmed without doubt because of the overlapping contributions from metal, silicide and nitride. In Fig. 3 (bottom part) the Si(KLL) peak measured on this sample is presented. The Si(KLL) peak has to be considered for any information about silicide formation [8]. The silicide peak is found at about 1617 eV, and the peak of Si3N4 is situated at about 1612 eV. In Fig. 3 the contributions of titanium silicide and Si3N4 can be clearly seen. The analysis of the measurements on Ti coated samples irradiated with different low t values leads to the
Fig. 3. Ti(LMM) AES peak at 419 eV for pure TiN, Ti metal and the ceramic sample irradiated with Ti using t = 170 ms (upper part). Si(KLL) AES peak for the sample irradiated with Ti using t =170 ms (bottom part).
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conclusion that both, TiNx and titanium silicide are formed. The relative amount of TiNx increases with decreasing thickness of the Ti containing cover layer (longer t).
3.2.2. Ni deposition The most remarkable difference in the surface modification due to irradiation with the Ni containing plasma compared to Ti is demonstrated in Fig. 4 showing SEM micrographs of samples irrradiated with different t. For the shortest t, where metal deposition is expected, droplet-like bright features can be seen. Their dimensions decrease with increasing t. For the highest t value of 200 ms they are not visible. EDX spectra of the bright nodules show that they are Ni rich, whereas between these nodules a Ni signal appears also, but with a much lower intensity. The Si(KLL) AES peak presented in Fig. 5, measured on a Ni nodule of the sample irradiated with t = 160 ms, shows a strong line at about 1617 eV characteristic for Ni silicide. These findings indicate a very inhomogeneous Ni deposition at low t values (160–170 ms), mainly as nodules. The RBS results shown in Fig. 6 confirm this assumption. The Ni distribution is characterized by the superposition of two parts. The narrow surface peak is caused by a relatively homogeneous Ni coverage on the surface. The long tail is due to the thick Ni droplets. For delay times t of 180–190 ms the tail disappears in agreement with the corresponding SEM micrographs in which nodules cannot be detected. However, the intensity of the surface peak is still remarkable. The surface concentration of Ni cannot be determined exactly for the low t values due to overlapping of the tail with the Si signal. It amounts to ] 6× 1017 cm − 2 and 1.8×1017 cm − 2 for t= 160 and 170 ms, respectively. Also for t= 200 ms a small Ni contamination of 0.3× 1017 cm − 2 is observed by RBS. The shift of the Si edge to lower energies due to the metal layer on the surface is less pronounced as in the case of the Ti samples because a large amount of Ni is concentrated in the nodules. The question as to whether mixing occurs between the Ni distributed nearly homogeneously at the surface and the substrate can not be answered unambigiously from the RBS results. The relatively steep silicon edges give no indication for mixing. On the sample irradiated with t=160 ms AES depth profiles of Ni, Si and N were measured up to a depth of about 30 nm by sputter etching using 3 keV Ar ions in an area between the noduls. As shown in Fig. 7, the coexistence of Ni and Si near the interface indicates the formation of nickel silicide while the top layer consists mainly of Ni. For larger depth’s the conducting film is etched away and charging effects prevent from performing evaluable AES measurements. The missing conductivity indicates that the mixed region, i.e. the region with a significant metal content, is limited to a depth range 530 nm.
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Fig. 4. SEM micrographs of samples irradiated with Ni (a − t =160 ms, b−t = 200 ms).
3.3. Hardness and tribological beha6ior The relative change in hardness, i.e. the hardness of the irradiated sample normalized with the hardness of the untreated substrate, has been determined for penetration depths of the indenter between 0.20 and 1.0 mm. A typical result is demonstrated in Fig. 8. All irradiated samples show a reduced surface hardness. The maximum reduction scatters between 0.5 and 0.75. Taking into account the error limits no significant dependence on t or the metal species can be stated. The most probable reason for the reduced
hardness is the degradation of the substrate surface as proven by the measured material lost. Typical wear tracks of the virgin and the processed samples are shown in Fig. 9 for a sample irradiated with Ni at t= 170 ms. Without irradiation a deep wear groove is formed whereas the treated sample exhibits essentially no wear. Instead, material is deposited on the surface, obviously from the ball. The friction coefficient is raised slightly from 0.5 to 0.6, and the wear of the ball is increased by about 20%. This behaviour was found for all Ni irradiated samples independent of the plasma conditions. The simi-
F. Brenscheidt et al. / Materials Science and Engineering A253 (1998) 86–93
Fig. 5. Si(KLL) AES peak of the sample irradiated with Ni using t = 160 ms.
lar results were also obtained for the samples coated with Ti.
4. Discussion For all irradiated ceramic samples a nearly equal material loss has been observed independent of the delay time t. The steel substrates, irradiated simultaneously with the Ti samples, show the expected dependence of the mass change on the delay time, similar to that observed in the previous investigation with copper [9]. This proves that the experimental conditions were reproducible in every pulse. The mass loss of the Si3N4 samples for all t can only be explained by an overcompensation of the metal deposition by degradation of the surface and release of constituents. That both effects, metal deposition and surface decomposition, take place is confirmed by the RBS measurements. For low t a surface coverage with Ti or Ni is detected. Decomposition of the surface of the substrate is indicated by the
Fig. 6. RBS spectra of samples irradiated with Ni in dependence on the delay time t.
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Fig. 7. AES depth profiles for Ni, Si, N, C, and O measured on the sample irradiated with Ni using t= 160 ms (depth profiling has to be stopped after 150 s sputtering time because of sample charging).
RBS spectra of the Ti coated samples, which show an enhanced silicon yield in the near surface layer, corresponding to a loss of nitrogen in this depth range. This effect is also visible but less prominent in the case of the samples treated with Ni. The different behavior of steel and ceramic substrates is concluded to arise from the different thermal conductivities. Under identical pulse energy density, the ceramic material reaches higher temperatures, resulting in an increased ablation compared to the metal. Comparing Ni and Ti irradiation, the mass loss for both species is almost the same despite the fact that the efficiency of deposition of Ni is significantly higher than that of Ti as demonstrated by comparing the metal surface concentrations given in Sections 3.2.1 and 3.2.2 in dependence on t. From results obtained for the deposition of different metals on steel substrates (J. Piekoszewski, private communication) we assume that the flux of metal ions arriving at the surface is comparable for Ni and Ti. The most plausible explanation for the different amount of retained metal is the lower sublimation enthalpy of Ti. The effect of surface alloying shows also remarkable differences comparing the morphology of Ti and Ni coating. By Ti irradiation a relatively homogeneous surface layer composed mainly from TiNx is formed and can be explained by the reaction of the deposited metal with the nitrogen plasma. The formation of titanium silicide at the interface is a competing, but slower process. The AES measurements reveal the presence of this phase. However the sharp Si edges of the RBS spectra (see insert in Fig. 1) indicate a very thin silicide layer (few nm). In the case of Ni irradiation also a thin alloyed layer is formed with mainly Ni in the top region and an increasing silicide content in the direction of the substrate. However, the higher amount of Ni deposited at low t is to a large extent concentrated in nodules
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Fig. 8. Dependence on the indenter depth of the relative hardness change due to irradiation with Ni at t = 180 ms.
composed mainly from silicide. This difference with Ti can be understood by taking into account the large differences in the melting points of nickel and of titanium silicides, respectively (NiSi: 992°C, NiSi2: 993– 1025°C, TiSi: 1570 – 1760°C, TiSi2: 1540°C [10]). Therefore, it is assumed that deposited Ni forms silicide by reaction with the substrate which melts and contracts into droplets. The higher surface tension at the melting point of Ni (1.74 N m − 1 [11]) compared to Ti (1.57 N m − 1 [12]) has also to be considered. The tribological behavior of the irradiated material is obviously independent of the amount of metal on the surface. Also Ti, which has a stronger tendency to form
lubricating oxides, shows nearly the same effect as Ni. Therefore, lubrication by the deposited metals can be excluded as a possible reason. We attribute the severe wear improvement to the following causes (i) the layer formed on the surface of the ceramics as result of the irradiation keeps its near surface region under compression, which prevents the propagation of cracks and (ii) the deposited metal and binding phase flowing out may close pre-existing flaws and also hinder their propagation. It was shown by Bolse et al. [13] that during implantation in ceramics, flowing glassy phase indeed can close surface cracks. The observed ablation of material from the substrate surface is a further effect which reduces the influence of pre-existing surface defects. This wear reduction by a reduced influence of macroscopic surface defects explains that it is observed in spite of the slightly reduced surface hardness. The hardness measurement with the low load used here is not affected by crack propagation. The increase in friction and ball wear is due to the higher roughness of the inhomogenously modified surface.
Acknowledgements The authors are grateful to Dr R. Mu¨ller for the SEM investigations, Mr J. Krolik for pulsed plasma irradiations and to Prof Z. Werner for valuable discussions.
References Fig. 9. Typical wear tracks of an untreated sample and a sample irradiated with Ni using 170 ms.
[1] T.R. Jervis, M. Nastasi, K.M. Hubbard, J.-P. Hirvonen, J. Am. Ceram. Soc. 76 (1993) 350.
F. Brenscheidt et al. / Materials Science and Engineering A253 (1998) 86–93 [2] M. Zaleski, T.R. Jervis, J.W. Mayer, J. Am. Ceram. Soc. 76 (1993) 356. [3] R.K. Singh, K. Jangannadham, J. Nerayan, J. Mater. Res. 3 (1988) 1119. [4] J. Narayan, D. Fathyy, O.W. Holland, B.R. Appleton, R.F. Davis, P.F. Becher, J. Appl. Phys. 56 (1984) 1577. [5] J. Piekoszewski, J. Langner, Nucl. Instrum. Methods B53 (1991) 148. [6] J. Langner, J. Piekoszewski, C. Pochrybniek, F. Rosatelli, S. Rizzo, J. Kucinski, A. Miotello, L. Guzman, P. Lazzeri, Surf. Coat. Technol. 66 (1994) 300. [7] T. Chudoba, PhD. Thesis, TU Dresden, 1996.
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[8] H. Reuther, M. Dobler, Surf. Interface Anal. 24 (1996) 411. [9] J. Piekoszewski, J. Langner, L. Walis, Z. Werner, Surf. Coat. Technol. 93 (1997) 209. [10] Werkstoffe der Halbleitertechnik, II. Aufl., Ed. H.F. Hadamovsky, VEB Deutchser Verlag f. Grundstoffindustrie, Leipzig, 1990, p. 256. [11] Gmelins Handbuch der Anorganischen Chemie No 57, part 1 II, Verlag Chemie, GmbH, Weinheim/Berstr. 1967. [12] D. Popp, A. Mehling, R. Wilzbach, H. Langhoff, Appl. Phys. A 55 (1992) 561. [13] W. Bolse, S.D. Peteves, W. Saris, Appl. Phys. A 58 (1994) 493.
Modification of silicon nitride ceramics with high intensity pulsed ion beams F. Brenscheidt a, J. Piekoszewski b,c, E. Wieser a,*, J. Langner b, R. Gro¨tzschel a, H. Reuther a a
Forschungszentrum Rossendorf, Institut fu¨r Ionenstrahlphysik und Materialforschung, Postfach 51 01 19, D-01314 Dresden, Germany b Soltan Institute for Nuclear Studies, Pl-05 -400 Otwock/Swierk, Poland c Institute of Nuclear Chemistry and Technology, Dorodna 16, Pl-03 -145 Warszawa, Poland
Abstract Sintered Si3N4 ceramics were irradiated with intense plasma pulses containing metal (Ti, Ni) and nitrogen ions in various proportion. The energy density of the pulse was 6.5 J cm − 2 with a duration in the ms range. The samples were characterised by Rutherford backscattering, Auger electron spectroscopy, scanning electron microscopy and by tribological tests. For Ti deposition the top layer is relatively uniform and consists of TiNx, titanium silicide and possibly metallic Ti. For Ni nodules composed mainly from nickel silicide were observed with an additional thin surface layer containing nickel silicide and metallic Ni. Both for Ti and Ni, the plasma irradiation results in a considerable improvement of the wear resistance. © 1998 Elsevier Science S.A. All rights reserved. Keywords: Ceramics; Silicon nitride; Surface alloying; Plasma-pulse irradiation; Microstructure; Wear properties
1. Introduction The tribological properties of ceramic materials are known to be controlled to a large extent by the presence of surface defects and the propagation of cracks originating at the surface into the bulk. Therefore, in the past efforts have been made to alter the surface properties of ceramics using techniques such as ion implantation and ion-beam or laser-beam mixing of metallic overlayers pre-deposited on the substrate. In particular, these efforts have been aimed at the formation of surface alloys with appropriate properties to improve the surface properties of ceramic materials. In the case of ion implantation and ion-beam mixing techniques, the mixing occurs via collisional effects and/or normal and radiation-enhanced diffusion mechanisms. In the case of pulsed laser mixing, the surface layer is rapidly heated to a high temperature (well in excess of the melting point of even refractory metals) * Corresponding author Tel.: + 49 351 2603096; fax: + 49 351 2603285; e-mail: [email protected]
and remains at a high temperature for a period long enough to make it possible that diffusional mixing occurs. Several examples of successful mixing of metalceramic systems and the improvement of their mechanical properties have been reported in the literature. For instance, Ti was mixed on SiC and Si3N4 [1], Ti on ZrO2 [2] and Ni on Si3N4 [3] using pulsed-laser irradiation and Ni on SiC using both ion-beam and pulsedlaser mixing [4]. However, some systems appeared to be refractory to mixing with ion beams such as Ni on Si3N4 and Ti on ZrO2 [1]. Attempts to mix Zr and Ti on Al2O3 have been unsuccessful for both, laser and ionbeam mixing. The present work reports preliminary results of another approach to surface alloying of ceramic systems. Instead of a two step process, i.e. predeposition and subsequent pulse heating, we propose a single-step process in which deposition and heating (melting) occurs within a single plasma pulse. In this experiment plasma pulses containing nickel or titanium and nitrogen as working gas were applied. Polycrystalline Si3N4 sinter ceramics were chosen as substrate.
0921-5093/98/$ - see front matter © 1998 Elsevier Science S.A. All rights reserved. PII S0921-5093(98)00713-8
F. Brenscheidt et al. / Materials Science and Engineering A253 (1998) 86–93
2. Experimental The investigated material is a gas-pressure sintered silicon nitride ceramic with 5 wt.% Y2O3 and 3 wt.% Al2O3 as sintering additives. The samples were polished to an average roughness of Ra =0.01 mm as determined by a surface profilometer. Before ion treatment the material was cleaned with ethanol in an ultrasonic bath. The samples were irradiated with plasma pulses generated by a rod plasma injector type of accelerator IBIS described in detail elsewhere [5,6]. Here we shall briefly outline the principal operation of this machine. The plasma pulses are generated as the result of a low-pressure, high-current plasma discharge between two concentric, cylindrical sets of rod-type electrodes. The discharge is initiated by applying high voltage after some time delay t from the moment of the injection of the working gas into the inter-electrode space. If t is sufficiently long to allow the injected gas to expand over the whole inter-electrode space, the plasma contains almost exclusively the elements of the working gas. This mode of operation is referred to as pulse implantation doping (PID), typically occurring for t= 190–210 ms. For short t, when there is a steep gradient of the gas concentration in the inter-electrode space, the plasma beam contains a mixture of the elements of the working gas and the metallic electrodes chosen deliberately for a given purpose. This mode is referred to as deposition by pulsed erosion (DPE) and occurs typically for t= 160– 180 ms. In the present experiment the irradiations were performed under the following conditions: The energy density of the pulse was about 6.5 J cm − 2. The power versus time characteristic of the pulses consists of two phases. In the first one, lasting about 1 ms, the mean energy of the ions arriving at the substrate amounts to several keV, in the second phase,
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lasting several ms, the ion energy is of the order of tens of eV. As working gas nitrogen of 99.996 purity was used. The outer 32 and inner 32 electrodes were made of commercial nickel or titanium rods of 260 mm length and 2.0 mm diameter. The delay times t were varied in the range between 160 and 200 ms. Each sample was irradiated with ten pulses. The plasma beam diameter was about 100 mm. To avoid contamination of the plasma with elements other than required, the sample holder was coated either with nickel or titanium foil for Ni-N and Ti-N plasma treatment, respectively. In the case of Ti-N plasma irradiation additional samples of stainless steel AISI 321 were placed nearby the silicon nitride samples to control the experimental conditions. Before and after treatment, the samples were weighed with a microbalance to check the mass gain or loss during the treatment. The topography of the irradiated surfaces was determined by surface profilometry with a Dektak 8000 profilometer and by scanning electron microscopy (SEM) with a Zeiss DSM 962 microscope. The composition of the surface and the surface near layer was investigated by Auger electron spectroscopy (AES) combined with sputter etching using a Fisons Microlab 310 F spectrometer and EDX measurements with the SEM. Additionally, Rutherford backscattering spectroscopy was applied to determine the depth distribution of the metal ions. 1.5 or 1.7 MeV He + ions and a backscattering angle of 170° were used. The hardness of the samples was measured with a Shimadzu DUH 220 dynamical microhardness tester. A special software [7] allowed us to correct the data for the elastic deformation of the surface, for deviations of the diamond tip from the ideal geometry and to calculate a depth depending hardness. The tribological tests were carried out with a ball-on-disc type tribometer which also allowed to determine the friction coefficient. The ball was from untreated silicon nitride loaded with 2N. The tests were carried out in normal air at humidity of 60–65%. The typical sliding distance was about 40 m.
3. Results
3.1. Weight loss
Fig. 1. RBS spectra of samples irradiated with Ti in dependence on the delay time t.
The difference in weight before and after processing of the samples irradiated with Ti shows that in all cases there is a mass loss Dm, in average 359 8 mg cm − 2/pulse. No correlation was found between the amount of mass change and the delay time t. The situation is quite different for the stainless steel samples simultaneously treated with the Si3N4 substrates. The steel samples show a clear dependence between mass difference and t. A mass gain is observed for small t where the content of Ti in the plasma leads to the
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Fig. 2. SEM micrographs of samples irradiated with Ti (a −t =160 ms, b− t = 200 ms).
expected metal deposition. For longer t the plasma consists only of nitrogen and a mass loss by evaporation from the surface is detected.
3.2. Surface composition 3.2.1. Titanium deposition Fig. 1 shows the RBS spectra of the samples coated with titanium by ten metal-nitrogen plasma pulses, respectively, for different delay times. The samples with short delay times show a sharp surface peak of the deposited metal. Therefore, a thin metal rich layer exists at the surface. This peak decreases and becomes narrower
for longer delay times, indicating that less material is deposited on the surface. The area density of retained Ti is 1.0 × 1017 cm − 2 and 2.4 × 1016 cm − 2 for t=160 and 170 ms, respectively, and is below the detection limit of RBS at 200 ms. The shift of the silicon edge to higher energies for longer delay times correlates with the decreasing thickness of the titanium layer. The steep slope of the silicon edge and the narrow Ti peaks indicates a high degree of uniformity of the titanium layer. No indication of mixing between the deposited metal layer and the substrate is found. The small peaks on the high energy side of the Ti peak in Fig. 1 are probably due to the deposition of ablated components from the steel samples.
F. Brenscheidt et al. / Materials Science and Engineering A253 (1998) 86–93
The SEM micrographs presented in Fig. 2 show no significant differences comparing the samples irradiated with the shortest and the longest t. The visible structures are somewhat more smooth for t = 160 ms. This might be explained by the thin metal cover layer which is not present for t =200 ms. AES has been used to get more insight into the composition of the metallic surface layer. The formation of TiNx can be seen from the shape of the Ti(LMM) peak at 419 eV. Fig. 3 (upper part) compares this Ti(LMM) peak of Ti metal, TiN and of the sample irradiated with t =170 ms. The peak of metallic Ti is centered at about 419 eV. TiNx is characterized by an additional shoulder at about 415 eV. The energy spectrum of the irradiated ceramic sample shows clearly this shoulder characteristic for TiNx. The coexistence of metallic Ti is probable, but it cannot be confirmed without doubt because of the overlapping contributions from metal, silicide and nitride. In Fig. 3 (bottom part) the Si(KLL) peak measured on this sample is presented. The Si(KLL) peak has to be considered for any information about silicide formation [8]. The silicide peak is found at about 1617 eV, and the peak of Si3N4 is situated at about 1612 eV. In Fig. 3 the contributions of titanium silicide and Si3N4 can be clearly seen. The analysis of the measurements on Ti coated samples irradiated with different low t values leads to the
Fig. 3. Ti(LMM) AES peak at 419 eV for pure TiN, Ti metal and the ceramic sample irradiated with Ti using t = 170 ms (upper part). Si(KLL) AES peak for the sample irradiated with Ti using t =170 ms (bottom part).
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conclusion that both, TiNx and titanium silicide are formed. The relative amount of TiNx increases with decreasing thickness of the Ti containing cover layer (longer t).
3.2.2. Ni deposition The most remarkable difference in the surface modification due to irradiation with the Ni containing plasma compared to Ti is demonstrated in Fig. 4 showing SEM micrographs of samples irrradiated with different t. For the shortest t, where metal deposition is expected, droplet-like bright features can be seen. Their dimensions decrease with increasing t. For the highest t value of 200 ms they are not visible. EDX spectra of the bright nodules show that they are Ni rich, whereas between these nodules a Ni signal appears also, but with a much lower intensity. The Si(KLL) AES peak presented in Fig. 5, measured on a Ni nodule of the sample irradiated with t = 160 ms, shows a strong line at about 1617 eV characteristic for Ni silicide. These findings indicate a very inhomogeneous Ni deposition at low t values (160–170 ms), mainly as nodules. The RBS results shown in Fig. 6 confirm this assumption. The Ni distribution is characterized by the superposition of two parts. The narrow surface peak is caused by a relatively homogeneous Ni coverage on the surface. The long tail is due to the thick Ni droplets. For delay times t of 180–190 ms the tail disappears in agreement with the corresponding SEM micrographs in which nodules cannot be detected. However, the intensity of the surface peak is still remarkable. The surface concentration of Ni cannot be determined exactly for the low t values due to overlapping of the tail with the Si signal. It amounts to ] 6× 1017 cm − 2 and 1.8×1017 cm − 2 for t= 160 and 170 ms, respectively. Also for t= 200 ms a small Ni contamination of 0.3× 1017 cm − 2 is observed by RBS. The shift of the Si edge to lower energies due to the metal layer on the surface is less pronounced as in the case of the Ti samples because a large amount of Ni is concentrated in the nodules. The question as to whether mixing occurs between the Ni distributed nearly homogeneously at the surface and the substrate can not be answered unambigiously from the RBS results. The relatively steep silicon edges give no indication for mixing. On the sample irradiated with t=160 ms AES depth profiles of Ni, Si and N were measured up to a depth of about 30 nm by sputter etching using 3 keV Ar ions in an area between the noduls. As shown in Fig. 7, the coexistence of Ni and Si near the interface indicates the formation of nickel silicide while the top layer consists mainly of Ni. For larger depth’s the conducting film is etched away and charging effects prevent from performing evaluable AES measurements. The missing conductivity indicates that the mixed region, i.e. the region with a significant metal content, is limited to a depth range 530 nm.
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Fig. 4. SEM micrographs of samples irradiated with Ni (a − t =160 ms, b−t = 200 ms).
3.3. Hardness and tribological beha6ior The relative change in hardness, i.e. the hardness of the irradiated sample normalized with the hardness of the untreated substrate, has been determined for penetration depths of the indenter between 0.20 and 1.0 mm. A typical result is demonstrated in Fig. 8. All irradiated samples show a reduced surface hardness. The maximum reduction scatters between 0.5 and 0.75. Taking into account the error limits no significant dependence on t or the metal species can be stated. The most probable reason for the reduced
hardness is the degradation of the substrate surface as proven by the measured material lost. Typical wear tracks of the virgin and the processed samples are shown in Fig. 9 for a sample irradiated with Ni at t= 170 ms. Without irradiation a deep wear groove is formed whereas the treated sample exhibits essentially no wear. Instead, material is deposited on the surface, obviously from the ball. The friction coefficient is raised slightly from 0.5 to 0.6, and the wear of the ball is increased by about 20%. This behaviour was found for all Ni irradiated samples independent of the plasma conditions. The simi-
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Fig. 5. Si(KLL) AES peak of the sample irradiated with Ni using t = 160 ms.
lar results were also obtained for the samples coated with Ti.
4. Discussion For all irradiated ceramic samples a nearly equal material loss has been observed independent of the delay time t. The steel substrates, irradiated simultaneously with the Ti samples, show the expected dependence of the mass change on the delay time, similar to that observed in the previous investigation with copper [9]. This proves that the experimental conditions were reproducible in every pulse. The mass loss of the Si3N4 samples for all t can only be explained by an overcompensation of the metal deposition by degradation of the surface and release of constituents. That both effects, metal deposition and surface decomposition, take place is confirmed by the RBS measurements. For low t a surface coverage with Ti or Ni is detected. Decomposition of the surface of the substrate is indicated by the
Fig. 6. RBS spectra of samples irradiated with Ni in dependence on the delay time t.
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Fig. 7. AES depth profiles for Ni, Si, N, C, and O measured on the sample irradiated with Ni using t= 160 ms (depth profiling has to be stopped after 150 s sputtering time because of sample charging).
RBS spectra of the Ti coated samples, which show an enhanced silicon yield in the near surface layer, corresponding to a loss of nitrogen in this depth range. This effect is also visible but less prominent in the case of the samples treated with Ni. The different behavior of steel and ceramic substrates is concluded to arise from the different thermal conductivities. Under identical pulse energy density, the ceramic material reaches higher temperatures, resulting in an increased ablation compared to the metal. Comparing Ni and Ti irradiation, the mass loss for both species is almost the same despite the fact that the efficiency of deposition of Ni is significantly higher than that of Ti as demonstrated by comparing the metal surface concentrations given in Sections 3.2.1 and 3.2.2 in dependence on t. From results obtained for the deposition of different metals on steel substrates (J. Piekoszewski, private communication) we assume that the flux of metal ions arriving at the surface is comparable for Ni and Ti. The most plausible explanation for the different amount of retained metal is the lower sublimation enthalpy of Ti. The effect of surface alloying shows also remarkable differences comparing the morphology of Ti and Ni coating. By Ti irradiation a relatively homogeneous surface layer composed mainly from TiNx is formed and can be explained by the reaction of the deposited metal with the nitrogen plasma. The formation of titanium silicide at the interface is a competing, but slower process. The AES measurements reveal the presence of this phase. However the sharp Si edges of the RBS spectra (see insert in Fig. 1) indicate a very thin silicide layer (few nm). In the case of Ni irradiation also a thin alloyed layer is formed with mainly Ni in the top region and an increasing silicide content in the direction of the substrate. However, the higher amount of Ni deposited at low t is to a large extent concentrated in nodules
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Fig. 8. Dependence on the indenter depth of the relative hardness change due to irradiation with Ni at t = 180 ms.
composed mainly from silicide. This difference with Ti can be understood by taking into account the large differences in the melting points of nickel and of titanium silicides, respectively (NiSi: 992°C, NiSi2: 993– 1025°C, TiSi: 1570 – 1760°C, TiSi2: 1540°C [10]). Therefore, it is assumed that deposited Ni forms silicide by reaction with the substrate which melts and contracts into droplets. The higher surface tension at the melting point of Ni (1.74 N m − 1 [11]) compared to Ti (1.57 N m − 1 [12]) has also to be considered. The tribological behavior of the irradiated material is obviously independent of the amount of metal on the surface. Also Ti, which has a stronger tendency to form
lubricating oxides, shows nearly the same effect as Ni. Therefore, lubrication by the deposited metals can be excluded as a possible reason. We attribute the severe wear improvement to the following causes (i) the layer formed on the surface of the ceramics as result of the irradiation keeps its near surface region under compression, which prevents the propagation of cracks and (ii) the deposited metal and binding phase flowing out may close pre-existing flaws and also hinder their propagation. It was shown by Bolse et al. [13] that during implantation in ceramics, flowing glassy phase indeed can close surface cracks. The observed ablation of material from the substrate surface is a further effect which reduces the influence of pre-existing surface defects. This wear reduction by a reduced influence of macroscopic surface defects explains that it is observed in spite of the slightly reduced surface hardness. The hardness measurement with the low load used here is not affected by crack propagation. The increase in friction and ball wear is due to the higher roughness of the inhomogenously modified surface.
Acknowledgements The authors are grateful to Dr R. Mu¨ller for the SEM investigations, Mr J. Krolik for pulsed plasma irradiations and to Prof Z. Werner for valuable discussions.
References Fig. 9. Typical wear tracks of an untreated sample and a sample irradiated with Ni using 170 ms.
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