Mechanical Study of High Resistance Silicon
Carbide Based Multi-Nano-Layers Grown by
Multifrequency PACVD
Farida Rebib, Thomas Gaudy, Audrey Soum-Glaude, Isabelle Caron,
Tony Da Silva, Christine Picard, Fabienne Cournut, Laurent Thomas*
Introduction
For ten years, hard coatings attract considerable attention in a large industrial domain in particular for thermo-mechan-ical applications where resistance to wear and high temperature is needed. Interest is then growing for the elaboration of protective coatings such as carbon or nitride-based ones on metallic parts. For those materials, the main part of the European market is actually occupied by PVD (Physical Vapor Deposition) technology . More precisely, for coatings dedicated to wear and friction, the production corresponds classically to PVD nitride/carbide or titanium-based films (TiN, TiCN, TiCrN, CrC, TiAlN . . .) and, for a minor
importance (#20% of the whole market) to DLCs (Diamond-like Carbons) and related materials that are obtained both by PACVD (Plasma-Assisted Chemical Vapor Deposition) and PVD.[1–3] Due to their high hardness and low wear in different atmospheres, DLC films find applications in aeronautics, space and high precision mechanics, but can be limited in use as their microstructure is strongly changed when the temperature is increased.[4]Doping DLC’s with elements such as Si (up to 30 at.%) can extend applications to higher temperatures,[5] and amorphous hydrogenated
silicon carbide-based films (a-SiC:H) have then been found to be a new alternative that can reach a hardness close to 30 GPa and a friction coefficient lower than 0.15 against metallic alloys (wear <2.106mm3 N1 m1).[6–8]DLC (a-C, a-C:H) coatings currently present high compressive internal residual stress responsible for both their low crack resistance and weak adhesion on metallic alloy substrates. High internal stress, classically due to intense ion bombardment during growth,[9,10] can limit the maximum acceptable normal loads in mechanical applications (friction/wear). In order to overcome this drawback and hence improve adhesion and strength, architectures composed of hard coating/soft coating periods can accommodate the com-pressive stresses at the substrate/coating interfaces and thus increase adhesion.[11,12]Unfortunately, such architectures,
F. Rebib, T. Gaudy, A. Soum-Glaude, C. Picard, L. Thomas PROMES/CNRS, Tecnosud, Rambla de la Thermodynamique, F-66100 Perpignan, France
Fax : (þ33) 4 68 68 22 65/(þ33) 4 68 68 22 13; E-mail: thomas@univ-perp.fr
I. Caron, T. Da Silva
LISMMA/ISMEP, 3 rue Fernand Hainaut, F-93407 Saint-Ouen, France
F. Cournut
EADS France – Innovation Works, 12 rue Pasteur, F-92150 SURESNES, France
Titanium alloys are commonly used to lighten aeronautical structures. Unfortunately, their
poor tribological properties need to confer onto the surface both high mechanical resistance
under load and low friction/low wear against other metals. Based on amorphous silicon
carbide (a-SiC:H), periodic multi-nano-coatings have been found to be candidates for this use.
Such coatings were deposited on Ti surfaces using multifrequency plasma CVD devices.
Scratch-testing measurements have revealed that the critical loads corresponding to
cohe-sive/adhesive cracks are significantly increased using multilayered stacks. Fretting-wear tests
also revealed the influence of multilayer architecture and preliminary analysis of wear tracks
and material transfers leaded to a first classification of the coatings.
when classically grown in capacitive devices [low- (n 500 kHz) or radio-(n ¼ 13.56 MHz) frequencies], give rise to relatively low plasma growth rates (a few mm h1) as those plasma excitations are classically limited to approximately 109 and 1010electron cm3(at P < 1 Torr). In order to improve such characteristics, plasma sources as microwave ones (n ¼ 915 MHz or 2.45 GHz) are conveni-ent.[6,13,14]When coupled to an additional plasma surfacial source, such a device allows high plasma concentrations – high deposition rates particularly when orga-nosilicon precursors are employed[15]–
independence of ions energy and flux, and possible surface temperature regula-tion. Finally, recently developed micro-wave source applicators, such as linear coaxial, modular coaxial, or dipolar ECR sources have been adapted to large-scale PACVD reactors. In this work, we present amorphous hydrogenated silicon carbide (a-SiC:H) multilayered hard coatings. These coatings have been found to be attractive for mechanical application in aeronautics, where they are believed to be excellent candidates to improve tribo-mechanical properties of titanium alloys.[16,17]Prevent studies have showed
that multifrequency PACVD of tetramethylsilane (TMS) is a powerful technique to deposit coatings of high mechanical quality. In this paper, we aim to show that coupling, to the microwave excitation, a negative biasing of the substrate using low-frequency excitation leads in a sharp control of the hardness (H) and Young modulus (E) of the deposits. Coatings with high mechanical properties contrast were obtained (10 < H < 32 GPa – 100 < E < 250 GPa), then stacked in multi-nano-layered architectures. Microwave volumic plasma was obtained by use of coaxial modular applicators connected to an appropriate laboratory reactor setup. Multilayers, elaborated within this PACVD reactor, were mechanically tested by use of a scratch-tester for fracture resistance and adhesion, and a tribological setup for high loads fretting and first wear analysis.
Experimental Part
As shown in Figure 1, the reactor setup used for this study consists in a home-made cylindrical chamber of 25 cm internal diameter. Four coaxial microwave applicators are installed around the centre plan of the reactor and connected to a 2.45 GHz–1200 W generator through a power four ways splitter. The substrate holder is a low
frequency (50 kHz) biased cathode located at the bottom part of the chamber. It can be heated up to 650 8C.
All a-SiC:H coatings were grown on TA6V substrates. The gases whose flow rates were adjusted using flow controllers, were injected into the deposition room through a stainless ring located in the upper part of the reactor. Films (mono- and multilayers) were grown using the procedure described in ref. [15] Their elemental composition was determined by means of Energy Dispersive Spectroscopy (EDS) using an electronic microscope equipped with an emission field gun (SEM-FEG Hitachi 4500 equipped with a Noran Vantage EDS analyser). Stoechiometric SiC and SiO2samples were used for the instrument calibration before
determining Si/C atomic ratio and the oxygen contamination of the deposits. Surface mechanical properties of the films were measured using a Nano-indenter II from Nano Instruments Inc. Two sequential loading, unloading cycles of a Berkovitch’s indenter, under load control (loading time of 10 s) with a low maximum load of 5 mN, have been used to prevent the influence of the substrate on the measure. Mean hardness (H) and elastic modulus (E) were calculated using 10 indentation tests. After nano-indentation tests, two monolayers with a high hardness contrast were engaged in the multi-layers [a-SiC:H (hard)/a-SiC:H (soft)] stacking. Two main parameters were varied in the architecture realization. The internal period (P), corresponding to the sum of one soft and one hard monolayer thicknesses, was of respectively 5, 25 and 50. The ratio (R) of the hard monolayer Figure 1. Schematic representation of the multifrequency PACVD experimental setup used to grow a-SiC:H/a-SiC:H periodic multilayers.
thickness on the soft one was equalled of 0.25, 1 and 4. For all the multilayer stacks, the total thickness was kept constant at 2.5 mm. Samples were named PxRy, where x and y are the period number and the thickness ratio, respectively. In order to characterise the multilayer coatings, adhesion measurements were carried out using a Micro Scratch-Tester (CSM Instruments, SA). The tests were performed with a diamond indenter (Rockwell conic-shaped diamond indenter (r ¼ 200 mm), progressive load (0–30 N at a loading rate of 30 N ( min1), track length of 5 mm at a speed of 5
mm ( min1). The averaged values of the critical loads were deduced from three tests performed for each sample and adhesion properties were estimated from the analysis of the failure modes.[11] Lc1, Lc2 and Lc3 are respectively critical loads
corresponding to cohesive, cohesive/adhesive and purely adhesive failures. For the fretting tests, a VibroCryoTriboMeter (V.C.T.M.) applies two independent loadings, a static normal one and a cyclic tangential one, to reproduce the fretting phenomenon as well as in a mechanism submitted to a static load and vibrations. The contact is usually a ball-on-plane type (a spherical pin in the present case). The pin is linked to the normal load which is applied by means of a hanging weight. The plane test specimen is set into motion by a vibrator. The two specimens are placed in a cryostat of liquid nitrogen. Oscillatory sliding motion between the plane specimen and the pin is measured by the detached sensor that allows servo-positioning. Tangential loads are controlled by a piezo-electric sensor. The V.C.T.M. apparatus is described in more detail elsewhere.[18]Fretting-wear tests were conducted at 50 Hz
under a mean pressure of 588 MPa with a displacement of 32 mm. Annealed Ti6Al4V hemispheric pins (5 mm radius) were used. This has the consequence to make test conditions more severe; the maximum pressure corresponding to Hertz pressure is close to 900 MPa (888 MPa). The factor named PV is defined as the energy inside the specimen/pin contact during test. It is representative of surface energy developed during the test. PV was initially of 3.7 W ( mm2and displacement was applied during 2.2 105cycles. Preliminary classification of the multilayers is performed.
Results and Discussion
a-SiC:H Monolayers Selection for Stacked Architectures
The first step consists in monolayers elaboration in order to choose high mechanical properties contrast, then stack functional multilayers. In this study, the DC bias (Vdc) has been chosen as the one parameter to act on chemical and mechanical characteristics of the coatings, the other process parameters being fixed (gas fluxes, pressure, microwave power).[15] In PACVD processes, among all
accessible parameters acting on film growth, DC bias voltage controls the maximum of the ion energy distribu-tion funcdistribu-tion (IEDF) in the plasma sheath. It can be used as a convenient parameter to modify composition, hence hardness and elastic modulus, of the deposited layer by controlling ion–surface interactions. Table 1 summarises the main results obtained for single a-SiC:H layers. As can be seen, independently from the applied negative bias
voltage (Vdc), oxygen contamination is very low (<2.6 at.%), probably coming both from the post-deposition oxidation of the sample surface and gas mixture impurities. Atomic ratio Si/C does not vary significantly, even with a large variation of Vdc. The slight trend to decrease when increasing Vdc has been explained in terms of selective sputtering of internal bondings in the material when submitted to energetic ion bombardment.[14–16] During growth, as Arþ are major ions in our process (mixture used Ar/few % of TMS), selective sputtering of chemical elements and film cross-linking under ion bombardment are known to act efficiently on film composition. Combined XPS, FTIR and Raman analysis has revealed that DC bias ranging in the voltage domain 0– 250 V was strongly modifying the monolayer structure (ion mean energy function of DC bias).[14–16] As Vdc
increases, hydrogen then silicon atoms first are easily sputtered leading to the slight Si/C ratio decrease. Then, another selectivity regime occurs between C–Csp2and C– Csp3bonds etching, leading when medium ion energies are reached, to an increase of sp3carbon form relatively to the sp2one. Silicon selective sputtering is then compen-sated by the sp2 carbon one, leading finally to a quite constant Si/C ratio. Hence, the mechanical properties of a-SiC:H monolayers, such as their hardness (H) and Young modulus (E) are Vdc dependent. As reported in Table 1, hardness increases from 20 to 28 GPa when Vdc increases from 40 to 150 V. It decreases for higher bias voltage. Even if films composition variation is not that much significant, the depletion of silicon, mainly bonded to carbon (SiC), leads to the creation of new CC bonds. Moreover, as C–Csp2-like bondings are more easily sputtered than C–Csp3 ones, the increase of Csp3 states (C–Csp3, sp3C–H) in the film structure is then directly linked to the increase of the harness. The elastic behavior of the hard layers can be estimated in terms of the H3/E2 ratio. This amount is related to wear resistance towards
Table 1. Evolution of the composition and the mechanical proper-ties of single a-SiC:H layers with the bias voltage (Vdc).
-Vdc Composition Mechanical properties
V Si/C O H E H3/E2
at.% GPa GPa
40 0.58 0.76 20.1 162 0.31 60 0.64 0.98 20.3 176 0.27 120 0.58 2.61 24.9 154 0.65 150 0.56 1.87 28.2 177 0.71 200 0.54 0.84 22.4 145 0.53 250 0.56 0.95 20.2 133 0.46
plastic deformation or failure (for purely fragile mate-rial)[11,15,16]and is proportional to the load which initiates plastic deformation or cracks when films are mechanically fragile. As for hardness, the H3/E2ratio is higher for films deposited at a DC bias close to 150 V. Such changes do not seem to be connected to strong stress changes as revealed in previous works.[14]Taking into account these results, two a-SiC:H monolayers that present a high hardness contrast were chosen to be stacked in periodically multi-layered architectures.
a-SiC:H-based Multilayers
As described in Table 2, two types of films were grown. Three films have a low number of period (P ¼ 5). In that case, periods are micro-scaled (0.5 mm) whatever the ratio of symmetry R. Others present a period thickness lower than 100 nm; they point out the interest of nanoscale for crack resistances. As a matter of fact, in scratch-tests, Lc1 corresponds to the occurrence of the first cracks into the material. It is significant of the film cohesive resistance. As shown in Table 2, it seems that a trend is emerging when analyzing the results for the highest critical loads obtained values: (i) for a small period number
(P ¼ 5), high Lc1 values were obtained when choosing a compromise between a thick hard layer (high resistance to cracks, H3/E2high), and a thick soft layer
that limits crack propagation in the internal periods. In that configuration, this naturally leads to a ratio of sym-metry close to R ¼ 1; (ii) for high P values (P ¼ 25 and 50) and on the contrary, it seems necessary to promote a ratio R in which the hard layer is thinner (R ¼ 0.25). That configuration delays the first cracks
occurrence to higher loads, and limits their propagation through the inner thick soft layer, and thus, flakes formation (both high Lc1 and Lc2 values). Hence, among all performed architectures, two multilayers exhibit interesting behavior (high Lc1, 2 and (3); they correspond to P5R1 and P50R0.25 architectures. Lack of value for Lc3 for P5R1 film indicates that it exceeded the maximum value obtainable by the used scratch-test apparatus (30 N). For the film P50R0.25, Lc3 is also very close to the limit of the device (Lc3 # 28 N). Figure 2 shows optical images of the tracks obtained for these two films after scratch-testing. For film P5R1, flakes corresponding to Lc2 are located around the track borders where the material piles up under the indenter load. In its case, the substrate never appears (Lc3 > apparatus limit). For P50R0.25 coating, fewer flakes are observed when Lc2 is reached. In this other case, cracks propagation appears to be more difficult over the film bulk. The substrate surface appears only around the track at the end of the test, and cohesion of the material is still kept along the whole track. Adhesion is then limited by the strong deformation of the coated substrate under load.
The achievement of fretting-wear tests allows the estimation of the friction strength required for the transition
Table 2. Description of obtained multilayers and corresponding scratch-test critical loads.
Samples Critical Loads
Reference Number of periods Ratio Lc1 Lc2 Lc3
P5R4 5 4 7.1 0.2 11.3 1.7 25.8 2.3 P5R1 5 1 11.6 1 15.8 1.1 >30 P5R0.25 5 0.25 7.2 0.2 10.9 1.0 12.8 0.8 P25R4 25 4 6.4 2.1 10.9 0.4 18.1 5.2 P25R1 25 1 10.2 1.4 13.5 0.5 24.4 1.0 P25R0.25 25 0.25 9.6 0.7 11.2 0.5 25.0 2.5 P50R4 50 4 7.1 0.2 11.1 0.5 21.5 1.0 P50R1 50 1 8.6 0.6 15.0 1.7 22.4 3.5 P50R0.25 50 0.25 11.7 0.7 14.2 0.6 27.5 0.8
between partial slip and total slip regimes. This makes possible the evaluation of a threshold friction coefficient (mthreshold) for the different stack architectures. As reported in
Table 3, a-SiC:H/a-SiC:H multilayers with symmetric ratio (R ¼ 1) exhibit the lowest friction coefficient (0.30–0.35). Note that these friction coefficient values are obtained under hard load conditions. With classical pin-on-disc tribometers (PHertz# 300–500 MPa), friction coefficients lower than 0.1–
0.15 are currently measured for those films against metals. For configurations with an asymmetric ratio, it seems that those with R ¼ 0.25 present a better friction behavior than the ones at R ¼ 4. Such a result could be directly linked to a-SiC:H monolayers. Indeed, high hardness is not absolutely necessary to obtain low friction coefficient. Huang et al.[19]
explain such behavior considering the film surface rough-ness. In our case surface roughness (Ra) is systematically lower than 100 nm. Two different types of damages were evidenced (Table 3): wear and (wear þ transfer) from the pin towards the sample. No case corresponding to transfer from the film to the pin was observed for this set of coatings. The transfer was quantified as follows: Grade 0 – localized transfer; Grade 1 – transfer on 30% of the surface; Grade 2 – important transfer. Note that even under such severe fretting-wear conditions, the a-SiC:H/a-SiC:H multilayers
wear rate was always on the order of 105–
106mm3 N1 m1. Architectures with R ¼ 1 exhibit only
a wear-type damage. Considering their low threshold friction coefficient these samples quickly reach the total slip regime which explains their highest wear rates (105mm3 N1 m1) after 2.2 105cycles. It is important
to underline that such a configuration is more interesting than that with matter transfer from the pin which could lead to systems that cannot be dismantled. Initially, the PV factor was equal to 3.7 W1 mm2, as at the end of the test, it reaches values varying from 0.14 to 0.21 W1 mm2 depending upon the damages of the plane samples and pins.
Conclusion
This work has been focused on the interest of PACVD multilayers-based on periodical a-SiC:H/a-SiC:H staked architectures for mechanical applications. Such coatings have been performed on titanium alloys in a dual frequency PACVD reactor where the whole surface pre-treatments and deposition are all realized in situ and in one run.[14]Among all process accessible parameters, this study was focused on the effect of bias voltage that can be managed to control the chemical hence the mechanical properties of films that are used for multilayered architectures. A comparison between single and periodical nano-layers (a-Si:C:H/a-Si:C:H) showed that multi-layer arrangement plays a key role in scratch-test and wear response (period number, layers thickness ratio, total thickness). As a matter of fact, coating architectures have to be adapted to each needed mechanical solicitation. Two mechanical tests have thus been employed to classify different multilayer solutions: (i) scratch-testing to qualify the performance of coatings for cracks resistance, and (ii) fretting-tests to study their behaviors under severe tribological conditions. For films mechanical strength characterization, i.e., adhesion and cracks resistance under loads, actual best architectures correspond to periodical stacks where the ratio of thickness symmetry is close to 0.25 (10 nm/40 nm, respectively). For a total thickness equal to 2.5 mm, critical load Lc3 (final substrate appearance) is close to 30 N under standard scratch-test conditions. Values as high as 100 N have been already obtained for this kind of architectures at total thicknesses close to 25 mm (see ref. [14]). For fretting tests, multilayered architectures conferred interesting wear properties under hard pressure conditions (PHertz max¼ 900 MPa). Friction
coefficients close to 0.3–0.35 against Titanium were measured for a 50 Hz test frequency – 32 mm track length
Table 3. Fretting-wear damages and threshold friction coefficient of the different a-SiC:H/aSiC:H multilayers.
Coating Damage Pm final PV final mthreshold
MPa W mm2
P5 R4 Transfer Grade 1 32.2 0.21 0.51
P5 R1 Wear 24.8 0.16 0.33
P5 R0.25 Wear Transfer Grade 0 21.8 0.14 0.46
P25 R4 Transfer Grade 2 24.7 0.16 0.42
P25 R1 Transfer Grade 0 22.5 0.14 0.29
P25 R0.25 Wear Transfer Grade 1 21.3 0.14 0.44
P50 R4 Transfer Grade 1 26.3 0.17 0.35
P50 R1 Wear 25.7 0.16 0.3
– 2.2 105 cycles number. Considering both friction
coefficient and wear behaviors, it appears that a-SiC:H/ a-SiC:H multilayers configuration with symmetric ratio and a high period number are the favorite ones for severe fretting request. Additional performed fatigue tests have to be analyzed. Finally, tribological tests have to be related to the gaseous environment where they are performed. As a matter of fact, for such materials family (Si-DLC, a-SiC:H), oxygen has a important role as it controls the nature of the material transferred onto the antagonist used for friction. First pin-on-disc experiments have shown, through comparisons between ambient air and vacuum, that an oxygen source (in those cases coming from gaseous atmosphere) was necessary to promote friction and wear. From our knowledge, due to its relatively low level (at.% O < 3%), O contamination into SiC:H films does not have a significant role in friction behavior.[17] On the opposite,
medium oxygen contents could allow the use of such films for spatial applications. We are actually developing SiC:O:H-based mono- and multi-layers for such validation.
Received: September 26, 2008; Accepted: March 7, 2009; DOI: 10.1002/ppap.200932305
Keywords: diamond-like carbon (DLC); mechanical properties; tribology; wear
[1] L. Tomasini, La Revue de Me´tallurgie-CIT 2001, 94–95, 411– 419.
[2] K. Vercammen, H. Haefe, Y. Gerbig, J. Meneve, Surf. Coat. Technol. 1999, 133–134, 466–472.
[3] G. Van der Kolk, ‘‘Expected developments of PVD and PACVD coatings for components’’, Opening Lecture of Conference PSE2006 2006, Garmisch-Partenkirchen, Germany.
[4] Y. Funada, K. Awazu, K. Shimamura, M. Iwaki, Surf. Coat. Technol. 1998, 103–104, 389–394.
[5] C. Donnet, Surf. Coat. Technol. 1998, 100–101, 180– 186.
[6] L. Thomas, L. Maille´, J. M. Badie, M. Ducarroir, Surf. Coat. Technol. 2001, 142–144, 314.
[7] L. Thomas, F. Teyssandier, M. Ducarroir, C. Boher, L. Autrique, Proc. EuroCVD 14, 2003, 359–364.
[8] A. Soum-Glaude, L. Thomas, A. Dollet, C. Picard, Plasma Pro-cess. Polym. 2007, 4, S11–S15.
[9] Y. Lifshitz, S. R. Kasi, J. W. Rabalais, Phys. Rev. Lett. 1989, 62, 1290.
[10] Y. Lifshitz, S. R. Kasi, J. W. Rabalais, W. Eckstein, Phys. Rev. B 1989, 41, 10468.
[11] S. Logothetidis, S. Kassavetis, C. Charitidis, Y. Panayiotatos, A. Laskarakis, Carbon 2004, 42, 1133.
[12] C. Strondl, G. J. Van der Kolk, T. Hurkmans, W. Fleischer, Surf. Coat. Technol. 2001, 142–144, 707.
[13] M. Joinet, O. Borrod, L. Thomas, C. Picard, V. Lucas, Diamond Relat. Mater. 2007, 16, 1254–1258.
[14] M. Joinet, S. Pouliquen, L. Thomas, F. Teyssandier, D. Aliaga, Surf. Coat. Technol. 2008, 202(11), 2252.
[15] A. Soum-Glaude, L. Thomas, E. Tomasella, J. M. Badie, Surf. Coat. Technol. 2006, 201(1–2), 174–181.
[16] A. Soum-Glaude, L. Thomas, E. Tomasella, Surf. Coat. Technol. 2006, 200(22–23), 6425–6429.
[17] G. Covarel, M. Joinet, L. Thomas, V. Lucas, Proceedings of CIP’2003 2003, 261.
[18] A. Koenen, Ph. Virmoux, R. Gras, J. Blouet, J. M. Dewulf, J. M. DeMonicault, Wear 1996, 197, 192.
[19] L.-Y. Huang, lu, K.-W. Xu, Mater. Sci. Eng. 2004, A373, 4514.
