Influence of deposition conditions on the adhesion of sputter-deposited WC(Co) films

Materials Science and Engineering, A 140 ( 1991 ) 6 3 1 - 6 3 8

631

Influence of deposition conditions on the adhesion of sputterdeposited W-C-(Co) films A. Cavaleiro and M. T. Vieira Departamento de Engenharia Mec(mica, Faculdade de Ci~ncias e Tecnologia, Universidade de Coimbra, 3000 Coimbra (Portugal)

Abstract A commercially available scratch tester was used to study the response of samples of high speed steel, coated with W-C-(Co) by d.c. diode and r.f. magnetron sputtering, to the scratch test. The critical load, mode of coating removal and acoustic signals are discussed. The influence of substrate bias, film thickness and substrate hardness was studied. The films deposited by the r.f. sputtering technique show cohesive failure without presenting an extensive spalling, while those obtained by d.c. sputtering reveal characteristic adhesive failure. In some cases good correlation between acoustic emission and the events in the scratch track has been obtained. In the study of substrate bias influence, it was concluded that critical load is interrelated with the hardness and the stress level of the coatings. For a large number of films, critical load increases with the film thickness and the substrate hardness.

1. Introduction During recent years there has been considerable research into the best hard coating suitable for deposition onto cutting tools. Among different possible coatings, films deposited from the system W-C-(Co) seem to be suitable for the envisaged application based on the hardness values obtained for some deposition conditions. However, the performance of coated cutting tools depends not only on the hardness of the coating but essentially on the compromise between the hardness and the wear coefficient of the film and the adhesion of the f i l l to the substrate. The influence of deposition conditions on the hardness and wear behaviour of W-C-(Co) films has been presented elsewhere [1-3]. The evaluation of the adhesion of the coatings to the substrate is very difficult to achieve. Nevertheless, the scratch test is often used to compare film quality concerning the strength of the interface between similar coating/substrates obtained with different deposition conditions. The aim of this work was to study the behaviour of W-C-(Co) sputtered films when they are submitted to the scratch test. More specifically, different M2(AISI) heat-treated high speed steel samples have been coated by sputtering techniques from 0921-5093/91/$3.50

W-C-(Co) targets with different cobalt contents. They were then submitted to the scratch tester in order to evaluate their critical load (L c).

2. Experimental procedure The W-C-(Co) coatings were deposited on high speed steel M2(AISI) substrates by d.c. diode and r.f. magnetron sputtering. The substrates were heat treated (normally) to obtain a final hardness of 920 HV30 and polished before coating to a final roughness (Ra) better than 0.05 /zm. The detailed procedure of the coatings' deposition was described elsewhere [1-5]. For similar thicknesses of film we have studied the influence of negative substrate bias (Vs) on the critical load determined by the scratch test. Thus, substrate bias was changed from zero to 400 V in the r.f. and d.c. equipment and films were deposited from three different targets (WC, WC + 6 wt.% Co and WC + 15 wt.% Co) maintaining the other deposition conditions constant. The deposition conditions were as follows: PJep (discharge power)=6.25 W cm -2 (r.f.) or 3.5 W cm-2 (d.c.), Pdep (deposition pressure)= 1 Pa (r.f.) or 11 Pa (d.c.); dint (interelectrode distance) = 6.5 cm (r.f.) or 3 cm (d.c.). The deposition time was © Elsevier Sequoia/Printed in The Netherlands

632 TABLE 1 Deposition conditions of the W-C-(Co) films Sample number

Technique

Target

1 2 3 4 5

d.c. d.c. r.f. r.f. r.f.

WC + 6 wt.% Co W C + 15 wt.% Co WC WC + 6 wt.% Co WC + 15 wt.% Co

that necessary to obtain a film thickness (t) of 2/~m. Table 1 summarizes the deposition conditions used in the study concerning the influence of the thickness (conditions 1-5) and of the substrate hardness (conditions 2, 4; t = 2/~m) on the critical load. The variation of the substrate hardness was obtained by quenching M2 steel from different austenitizing temperatures (07) indicated in Table 2. In all the cases the steel was tempered at 550 °C (three times) for 2 h. The physical and chemical properties of films deposited using these deposition conditions were the subject of other publications [1-5]. The critical load was evaluated by a commercially available scratch testing equipment (CSEM REVETEST) fitted with an acoustic detector. The scratch tests were performed under standard conditions: diamond tip radius R = 0.2 mm, scratching speed dx/dt= 10 mm min-~ and loading rate dL/dt= 100 N min- i. The critical loads indicated were obtained by averaging the values of five different scratches. Detailed morphologies of scratch channels were viewed with optical and electron scanning microscopes (SEM). The SEM was fitted with energy dispersive X-ray analysis (EDXS) which allows us to evaluate the chemical composition and to distinguish the coating from the substrate when the scratch profile is observed. 3. Results and discussion

The main difficulty found in the study of scratch test adhesion of the W-C-(Co) films is related to the determination of the critical load (Lc) representative of the failure of the coating/ substrate during the test. In fact, beyond the known problems in adhesion scratch testing, such as the difficulty in selecting the type of morphol-

dim

Vs

t

Pd~o

Pdep

(l'a)

(cm)

(V)

(~m)

3.5 3.5 6.25 6.25 6.25

11 11 1 0.3 0.3

3 3 6.5 6.5 6.5

100 200 50 100 100

0.5-6.4 1.2-10.0 1.9-5.0 1.2-4.5 1.2-4.6

(W cm -')

TABLE 2 Substrate hardness as a function of 0y 0y (°C)

Hardness (HV30)

Fullyannealed 900 950 1050 1150 1220

250 430 460 610 710 92O

ogy which defines the critical load when the nature of the coating/substrate changes, due to the variation of coating failure, in the W-C-(Co) films we have observed even for the same type of coating/substrate different modes of failure depending on the deposition conditions.

3.1. Influence of substrate bias The failure mode observed in films obtained with different substrate bias can be classified into four types. Type I. Films presenting conformal cracking [6] for low loads which changes to lateral cracking when the load increases. For higher load, some kind of flaking can be detected (Fig. 1 ). However, no vestige of extensive spalling was identified. This behaviour was observed in films deposited by r.f. from targets without cobalt, for all the substrate biases studied, or from targets with 6 wt.% Co with high substrate bias. The failure occurs within the coating itself (cohesive failure) rather than at the steel/coating interface (adhesive failure) which is a characteristic of a brittle coating [7]. This is consistent with the analysis by the EDXS technique, where no local concentration of iron exists near the cracks. Type II. The tensile cracking appears in the track before lateral cracking, with or without flak-

633

Fig. 1. Scratch track micrographs of a W-C-(Co) film (r.f.; W C + 6 wt.% Co; p = 1 Pa; 1'=6.25 Wcm -'. V,= 100 V) (SEM): (a) conformal cracking; (b) lateral cracking; (c) end of the scratch.

ing (Fig. 2). As in type I, no sign of spalling was detected. This morphology of cracking was observed in films obtained by r.f. sputtering from a target with 6 wt.% Co with low substrate bias or from a target of 15 wt.% Co with high substrate bias. Type IlL Only lateral cracking occurs during the scratch test (Fig. 3). This failure mode is characteristic of coatings that have resulted from r.f. sputtering targets with 15 wt.% Co when the substrate is low biased.

Type IV. The films present extensive spalling, with or without previous cracking. This failure mode was detected in films obtained by d.c. sputtering. Whereas the films deposited by this technique from the target, WC + 6 wt.% Co spall off without previous cracking (Fig. 4(a)), the films obtained from the target with higher cobalt content show tensile cracking before spalling (Fig. 4(b)). Although spalling of the coatings is evident from the microscopical observations, application of the EDXS technique confirmed this morphology by detecting the iron sign in the spall regions. Type I, II and III films deposited by r.f. sputtering refer to low, medium and high cobalt content respectively. The cobalt content of the film is a function of the percentage of this element in the target and of the substrate bias value. In fact, the substrate bias is the most important deposition parameter determining the cobalt content of the films. The increase in substrate bias induces a decrease of the cobalt percentage in the films [5]. For high values of this deposition parameter the percentage of cobalt in the films can be decreased to values similar to those obtained in the films deposited without substrate bias from targets with lower cobalt content. Thus, the substrate bias has an important role on the failure mode of W-C-(Co) films deposited by r.f., owing to its contribution to the changing of cobalt content in the film. In fact, for films deposited from the target without cobalt, where this element is not present in the film, the failure mode is always the same regardless of the substrate bias applied during the deposition. The relationship between the type of failure and the cobalt content is not so evident in films deposited by d.c. sputtering. However, different failure modes have also been observed related to the percentage of cobalt in the film. The influence of the substrate bias on the critical load of W-C-(Co) films is represented in Fig. 5. In the films that show two types of failure mode (type I and II) some controversy exists about the failure criterion which determines the critical load. Using the procedure recognized by other authors [6, 8-10], we have defined two critical loads, one related to the first cracking observed (Lcl), represented in Fig. 5 by full symbols, and the other corresponding to the first sign of flaking in lateral cracking (Lc2), represented b y open symbols. Some authors consider that the initial cracking corresponds to the failure of the coating

634

Fig. 2. Scratch track micrographs of W - C - ( C o ) films (SEM). (a) r.f.; WC + 6 wt.% Co; p = 1 Pa; P = 6.25 W c m - 2; V~ = 0 V: (al) tensile cracking; (bl) lateral cracking; (c I ) end of the scratch. (b) r.f.; WC + 15 wt.% Co; p = 1 Pa; P = 6.25 W cm - 2; V, = 300 V: (a I ) tensile cracking; (b~) lateral cracking; (c~) end of the scratch.

%

Fig. 3. Scratch track micrographs of a W - C - ( C o ) film (r.f.; W C + 1 5 wt.% Co; p = l Pa; P = 6 . 2 5 W c m - 2 ; V,=0V) showing the lateral cracking: (a)optical microscope; (b) SEM.

and this must be the event which determines the critical load [7, 11], while others propose the first flaking as the event which defines the critical load [12]. However, the evolution of critical load as a function of the substrate bias is analogous for Lcl o r tc2. Microscopy was the technique used for the evaluation of L c. The acoustic emission signals obtained during scratch testing were very difficult to interpret, different acoustical diagrams being observed even for the same sample. Sometimes the initial cracking was associated with a sudden increase in the acoustic sign and no significant variation was observed for the load corresponding to the flaking (Fig. 6). In other cases, the initial cracking did not cause an increase in the acoustic sign, this being observed only for the flaking (Fig. 6). Both of these behaviours were observed by other authors for different coating/substrate systems [7-9, 13-17]. However, as can be seen in Fig. 7, there are some scratches where the acoustic signal is consistent with microscopical observations. In this figure, the signal increases moderately when the diamond tip provokes the initial track cracking and increases suddenly if lateral cracking or spalling occurs. Similar behaviour was found by Perry

635

[1 1] for a CVD TiC coating deposited on a cold working steel. Considering the films deposited from targets with different cobalt contents without substrate bias, the critical load (L cI ) increases linearly with the cobalt content of the film. This behaviour reveals the important role of cobalt on the toughness of W-C-(Co) films deposited without substrate bias, as is observed in bulk cemented carbide where the increase in cobalt content contributes to a significant increase in toughness [18]. If we compare the results presented in Fig. 5 with those shown in Fig. 8 it is possible to deduce that the critical load is interrelated with the hardness of the coatings. For amorphous films resulting from the sputtering of a high cobalt content target ( 15 wt.% Co) and of a low cobalt content target (6 wt.% Co) without substrate bias, the percentage of cobalt in the film has no influence on the film hardness [2], nor on the critical loads Lc, and Lc2. However, it is important to remark that in the films with high cobalt content, Lc2 was defined as the first value which leads to lateral cracking without flaking. The flaking has been detected only for one film

deposited from the target WC + 15 wt.% Co with the highest substrate bias (Fig. 5, curve a) and for the films obtained from the targets with lower cobalt content. In the films with higher cobalt content it is very difficult to detect the flaking for the range of loads applied in the scratch test. This behaviour shows the role of cobalt in the Lee value. It is interesting to note that amorphous films with the same cobalt content, where flaking is present, obtained from different targets and different substrate bias (target 6 wt.% Co, V~= 0 V and target 15 wt.% Co, V, = 400 V) present the same Lcl and Lc2 (Fig. 5, curve a). Concerning the crystalline films, which have a cobalt content less than or equal to 6 wt.% [4, 5], the influence of the substrate bias on the hardness or critical load results not only from the variation in cobalt content and structure of the films with this deposition parameter, but also from the internal stress developed during the sputtering. In fact,

50

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C3

~3c ..j

I

I

I

100

200

300

I

4 O0 "

5UB5TRATE BIA5 ( V )

Fig. 5. Evolution of the critical load of W-C-(Co) films as a function of substrate bias: o, o, WC + 15 wt.% Co (curve a) flaking in the lateral cracking; A, / x WC + 6 wt.% Co; m []. WC. Full symbols, Let; open symbols, Lc: (see text).

Fig. 4. Scratch track micrographs of W-C-(Co) films. (a) d.c.; W C + 6 wt.% Co; p = 11 Pa; P= 3.5 Wcm 2; V,= 100 V. The spalling was observed by optical microscopy. (b) d.c.; W C + 15 wt.% Co; p = 11 Pa; P=3.5 Wcm :; I/,=200 V (SEM): (a t) tensile cracking; (b,) spalling; (c,) end of the scratch.

636

50 "~ ~0 ::zz

"4 '0

U3

~20 t

o

20

40

60 LOAD (N)

I

Ii

dO

700

Fig. 6. Accoustic emission vs. load graphs obtained on the same W-C-(Co) film which shows two different failure modes (r.f.; WC + 6 wt.% Co; V~= 100 V; t = 2/~m).

0

I

0

I

I

100 200 .300 5UBSTRATE BIAS ( V )

I

L

400

Fig. 8. Evolution of the hardness of W-C-(Co) films as a function of substrate bias: o WC + 15 wt.% Co; zx, WC + 6 wt.% Co; % WC.

5O L..

{,

.-..1

-

[340 o .4 30

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w 10 I

20

40

(a)

60 LOAD (N)

i

Ij

1o0

8O

0

2

4 6 ~ I'HICKNES5 (,4~m)

10

Fig. 9. Evolution of the critical load of W-C-(Co) films as a function of their thickness, x, d.c.; W C + 6 wt,% Co; p = l l Pa; P=3.5 Wcm--~; V,=100V; HV=2500. T, v , d.c.; W C + I 5 wt.% Co; p = l l Pa; P = 3 . 5 W c m 2; V~=200V; HV =2300. m, t3, r.f.; WC; p = l Pa; P=6.25 Wcm-2; V~=50V; HV=4600. e, o, r.f.; W C + 6 wt.% Co; p = 0 . 3 Pa; P=6.25 Wcm-2; V~= 100 V; HV=4200. zx, r.f.; W C + 6 wt.% Co; p = 0 . 3 P a ; P=6.25 Wcm-2; V~= 100 V; HV= 2300, Full symbols, L~q; open symbols, Lc2 (see text). I

o

(b)

20

40

60 L OA D (N)

~0

t

Ii

I00

Fig. 7. Accoustic emission vs. load graphs obtained on W-C-(Co) films which shows two different failure modes: (a) r.f.; W C + 1 5 wt.% Co; V~=I00V; t=2/~m; (b) d.c.; WC + 15 wt.% Co; V, = 200 V; t = 3.7/tm.

a detailed analysis of the X-ray spectra of W-C-(Co) films [2, 4, 5] allows us to conclude that the hardness of the films is interrelated with the degree of lattice parameter distortion due to the internal stress of the coating. This result is in accordance with those obtained by Bumett and Rickerby [19, 20] for TiN crystalline films deposited by sputter ion plating; conversely, it is in disagreement with the conclusions of other authors [21, 12] who, in one case, have

not detected in TiN crystalline films any change of L c with the hardness, in spite of L c increasing with the substrate bias, and in the other case have obtained the highest critical loads for the highest hardness values. The analysis of the evolution of critical load values with substrate bias allows us to conclude that, apart from the influence of cobalt content, other factors which determine the film hardness such as structure and internal stress have an important role on the Lc values.

3.2. Influence offilm thickness In Fig. 9, where the critical loads are presented as a function of film thickness, the critical loads refer to initial cracking (full symbols) and to flak-

637

ing or spalling (open symbols). Three different behaviours should be considered as the thickness of the film increases: films obtained by r.f. sputtering from the targets with low cobalt content, critical load decrease; films obtained from the target W C + 15 wt.% Co, critical load increase; and films obtained from the target WC + 6 wt.% Co by d.c. sputtering, there is a critical thickness for which the critical load is a maximum. The increase in the critical load with coating thickness is a common observation made by several authors [6-11, 13, 15, 17, 22] and is attributed to the shear stress distribution below the stylus. Thicker layers may require a greater surface stress to achieve the same shear stress at or near the interface, thus resulting in higher critical loads for thicker coatings. However, when the coating has high internal stress, the coating failure is dominated by the stress levels already present within the films and the critical load can decrease with the increase in coating thickness [6]. This seems to be the case for some W-C-(Co) coatings. In fact, for the films deposited by r.f. sputtering from the target WC + 6 wt.% Co, it was observed that for thicknesses higher than 6 /~m the film cracked during the deposition, without any additional mechanical energy input, showing a cohesive failure but without spalling. This means that a high stress level must be present in the coating which is responsible for the coating failure. 3.3. Influence of substrate hardness In Fig. 10 are shown the critical loads obtained for two films deposited on the same HSS sub-

5O

"~~o (b

~ 3o

0

I

0

2

i

[

i

I

J

I

~

I~

4 6 8 2 10 SUBS TRA TE HA RDNESS (HVx 10 )

Fig. ll). Evolution of the critical load of W - C - ( C o ) films as a function of substrate hardness, a, D, d.c.; W C + 15 wt.% Co; p=ll Pa;P=3.5Wcm 2; V , = 2 0 0 V ; H V = 2300. e, o, r.f.; W C + 6 wt.% C o ; p = 0 . 3 Pa; P = 6 . 2 5 W cm 2; V , = I 0 0 V ; HV = 420(I. Pull symbols, L ,. ~; open symbols, L ~2 (see text ).

strate heat treated differently so as to obtain several hardnesses. The symbols mean the same as previously. For both cases, the critical load increases with the substrate hardness, which is commonly observed for this kind of film [6, 8, 12, 15, 17, 22, 23]. The deformation of the coating/ substrate assembly caused by scratch testing depends mainly on the substrate deformation. The critical load of a coating must be related to the degree of deformation of the assembly. Therefore, when the substrate hardness increases, the deformation of the coating/substrate assembly is lower and there is need for a higher load to obtain the same level of deformation as that observed for lower substrate hardness. Then the critical load should increase with the substrate hardness.

4. Conclusions

We have presented an analysis of the dependence on failure mode and critical load films of the sputtering technique and of the substrate bias, film thickness and substrate hardness. From our results obtained by the scratch tester, the sputtering technique has an important role in the failure mode and critical load. The W-C-(Co) films deposited by r.f. show a cohesive failure for all deposition conditions studied, while those obtained by d.c. show an adhesive failure. The critical load is generally lower for the films deposited by d.c. For the same technique the substrate bias indirectly determines the failure mode, owing to its influence on film cobalt content, which depends also on the cobalt percentage in the target. The influence of V, on L c is less evident than in the failure mode. In fact, other factors which are directly related to the hardness must be taken in account. In W-C-(Co) films the change in L c with the substrate bias is the inverse of that obtained for their hardness. The variation in L c with film thickness is different depending on the cobalt content of the films. Films with lower cobalt content (r.f.) present a decrease in L c with increasing film thickness. Conversely for the higher cobalt content films (r.f. and d.c.) L c increases with the film thickness. When the film has a low cobalt percentage and the deposition technique is d.c. sputtering, L c has a maximum for a critical value of the thickness.

638

The increase in substrate hardness gives rise to an increase in critical load, particularly if we consider the Lc2 values.

References 1 A. Cavaleiro, M. T. Vieira, G. Lemperiere, C. Sfi Furtado and J. M. Poitevin, in E. Broszeit, W. D. Munz, H. Oechsner, K.-T. Rie and G. K. Wolf (eds.), Production and optimization of hard coatings of tungsten carbide with cobalt, in E. Broszeit et al. (eds.), Plasma Surface Engineering, Vol. l, DGM lnformationsgesellschaftVerlag, Oberursel, 1989, p. 595. 2 A. Cavaleiro, Tese de Doutoramento, Coimbra, Portugal, 1990. 3 A. Ramalho, M. T. Vieier and A. S. Miranda, in S. K. Ghosh (ed.), Proc. Int. Conf. on Developments in Forming Technology, Lisbon, Portugal, 12-14 September 1990. 4 A. Cavaleiro, M. T. Vieira and G. Lemperiere, Thin Solid Films, 185(1990) 199. 5 A. Cavaleiro, M. T. Vieira and G. Lemperiere, Thin Solid Films, 197(1991) 237. 6 R J. Burnett and D. S. Rickerby, Thin Solid Films, 154 (1987) 403. 7 J. H. Je, E. Gyarmati and Naoumidis, Thin Solid Films, 136 (1986) 57.

8 A.J. Perry, Proc. 8th Int. Conf. on CVD, Electrochemical Society, Pemmington, NJ, U.S.A., 1981, p. 737. 9 A.J. Perry, Thin Solid Films, 78(1981)77. 10 A.J. Perry, Thin Sofid Films, 81 ( 1981 ) 357. 11 A.J. Perry, Thin Solid Films, 107(1983)167. 12 H. Hintermann, Wear, 100(1984) 381. 13 J. Valli, U. Makela, A. Matthews and V. Murawa, J. Vac. Sci. Technol. A, 3(1985) 2411. 14 A.J. Perry, Su~ Eng., 2(1986) 183. 15 E. Hummer and A. J. Perry, Thin Solid Films, 101 (1983) 143. 16 J. P. Bucher, K. P. Ackermann, E W. Buschor and A. J. Perry, Thin Solid Films, 122 (1984) 63. 17 M.J. AI-Jaroudi, H. G. T. Hentzell and J. Valli, Thin Solid Films, 154(1987) 425. 18 D.J. Bettle, Int. Conf. on New Frontiers in Tools Materials Cutting Techniques, Metal Forming, London, 1979, session 4, com. 2. 19 D. S. Rickerby, J. Vac. Sci. Technol. A, 4 (1986) 2809. 20 D. S. Rickerby and P. J. Burnett, Thin Solid Films, 154 (1987) 125. 21 Th. Roth, E. Broszeit and K. H. Kloos, in E. Broszeit, W. D. Munz, H. Oechsher, K.-T. Rie and G. K. Wolf (eds.), Plasma Surface Engineering, Vol. 2, DGM Informationsgesellschaft-Verlag, 1989, p. 837. 22 P.A. Steinmann and H. E. Hintermann, J. Vac. Sci. Technol. A, 3(1985) 2394. 23 P. Billgren, Speedsteel Technical Report SD22/84.