- Email: [email protected]
Hard and superhard nanocomposite coatings
Surface and Coatings Technology 125 (2000) 322–330 www.elsevier.nl/locate/surfcoat
Hard and superhard nanocomposite coatings J. Musil * Department of Physics, University of West Bohemia, P.O. Box 314, 30614 Plzenˇ, Czech Republic
Abstract This article reviews the development of hard coatings from a titanium nitride film through superlattice coatings to nanocomposite coatings. Significant attention is devoted to hard and superhard single layer nanocomposite coatings. A strong correlation between the hardness and structure of nanocomposite coatings is discussed in detail. Trends in development of hard nanocomposite coatings are also outlined. © 2000 Elsevier Science S.A. All rights reserved. Keywords: Hardness; Magnetron sputtering; Mechanical properties; Nanocomposite coatings; Structure
1. Introduction Hard coatings have been successfully used for protection of materials and particularly to enhance the life of cutting tools since the 1970s. Both the technological process of their production and their properties, i.e. hardness, wear and oxidation resistance, however, are continuously being improved. Important milestones in the development of hard coatings are briefly summarized in Table 1. This table shows a clear effort (i) to decrease the temperature T at which hard coatings are formed and (ii) to improve the properties of hard coatings, particularly to increase the hardness and oxidation resistance. The oxidation resistance should be increased up to approximately 1000°C because during high-speed machining the temperature of the tool tip can reach 1000°C and the coating should be stable at such high temperatures. As to the hardness, the coatings are usually divided into two groups: (1) hard coatings having a hardness <40 GPa, and (2) superhard coatings having a hardness >40 GPa. Compared to a large number of hard materials, there are only a few superhard materials, i.e. cubic boron nitride (c-BN ), amorphous diamond-like carbon (DLC ), amorphous carbon nitride (a-CN ) and polyx crystalline diamond. Moreover, these superhard materials are thermodynamically unstable. This is a serious disadvantage which strongly limits their utilization in some applications. For instance, the high chemical * Tel.: +420-19-279072; fax: +420-19-279071. E-mail address: [email protected] (J. Musil )
affinity of carbon to iron limits the applicability of diamond coated cutting tools to machining of aluminum, their alloys and wood only. Similar problems can be expected when the c-BN coating is used in cutting of steels due to the chemical dissolution of boron in iron. These problems stimulated intensive research in this field, and recently new superhard materials based on superlattices and nanocomposites were developed.
2. Hard superlattice coatings Superlattice coatings are nanometre-scale multilayers composed of two different alternating layers with a superlattice period, i.e. the bilayer thickness of two materials, ranging from 5 to 10 nm. The bilayers of these superlattices can be metal layers, nitrides, carbides or oxides of different materials or a combination of one layer made of nitride, carbide or oxide of one metal and the second layer made of another metal. According to the composition of the bilayer, superlattice coatings can be divided into five groups: (1) metal superlattices, (2) nitride superlattices, (3) carbide superlattices, (4) oxide superlattices and (5) nitride, carbides or oxides/metal superlattices. Experiments show that metal superlattices exhibit a relatively low hardness. On the contrary, single-crystal nitride superlattice coatings are superhard materials with a hardness ranging from 45 to 55 GPa, e.g. TiN/VN, 56 GPa [15]; TiN/( V Nb )N, 41 GPa [16 ]; 0.6 0.4 TiN/NbN, 51 GPa [17]; TiN/Nb, 52 GPa [18,19]; TiN/CN , 45–55 GPa [20]; ZrN/CN , 40–45 GPa [21]; x x
0257-8972/00/$ - see front matter © 2000 Elsevier Science S.A. All rights reserved. PII: S0 2 5 7- 8 9 7 2 ( 9 9 ) 0 0 58 6 - 1
J. Musil / Surface and Coatings Technology 125 (2000) 322–330
323
Table 1 Important steps in development of hard coatings Coating
Material
H (GPa)
Main characteristics
Single layer Single layer Multilayer Single layer Single layer Single layer Single layer Single layer Superlattices Single layer Single layer Single layer
TiN, TiC, Al O 2 3 TiN, TiC TiC/TiB 2 c-BN diamond TiAlN DLC CN x TiN/VN, TiN/NbN, etc. nc-MeN/a-nitride nc-MeN/metal Ti Al N 0.4 0.6
21, 28, 21 21, 28
CVD at T around 1000°C on cemented carbides PVD at T≤550°C on steel substrates About 103 phase boundaries TiC/TiB [2] 2 High chemical affinity of C to iron Chemical dissolution of B in iron [5] Oxidation resistance up to 800°C [6 ] Amorphous phase [7] Substoichiometric (x=0.2–0.35) turbostratic structure [8,9] Superlattice period 5–10 nm [10,11] Superlattice period 5–10 nm [10,11] Nanocomposite [13] Nanocomposite, oxidation resistance up to 950°C [14]
50 [3] 90 [4] 65 50–60 ~50 ~50 ~50 ~32
TiN /C–N, 20–50 GPa [22,23]. The overall hardness of x the superlattice coating is therefore greater than that of the materials of the individual components of the bilayer, e.g. H #52 GPa, H =21 GPa, H =14 GPa. TiN/NbN TiN NbN This hardness enhancement is a very complex phenomenon. In spite of this, several models which explain multilayers strengthening have already been developed [24–27]. The model of Shinn et al. [26 ] shows that (i) a difference in elastic modulus between the two layer materials is required to increase the hardness of the superlattice film, and (ii) the coherency strain at the interface between the two layers has only a minor effect. The model of Chu and Barnett [27] is based on restricted dislocation movement within and between layers in the superlattice coating. It predicts a peak in hardness when there is a difference in shear modulus between two layer materials and sharp interfaces between layers. So far, less attention has been devoted to carbide multilayer coatings. These coatings can also be superhard, up to 55 GPa, e.g. TiC/VC, 52 GPa [28]; TiC/NbC, 45–55 GPa [29]. Considerable attention has been devoted to nitride/metal superlattice coatings. This combination of bilayer materials, i.e. the hard nitride with a more ductile metal, makes it possible to improve the toughness of the coating while retaining its relatively high hardness (≥30 GPa), e.g. TiN/Ti, 36.8 GPa [30]; WN/W, 34 GPa [30]; HfN/Hf, 50 GPa [30]; TiN/Ni, 35 GPa [31]; TiN/Ni Cr , 32 GPa [31]; NbN/Mo, 0.4 0.1 33 GPa [32,33]; NbN/W, 30 GPa [32,33]. An improvement in the toughness of the coating increases its adhesion to the substrate, which is of fundamental importance for coating applications. At present, practically no data are available on oxide superlattice coatings. This is probably because the deposition rate of oxide films using both d.c. and r.f. reactive magnetron sputtering is, compared with that of metals, too low. Recently, this situation strongly changed. Considerable progress has been made in magnetron deposition of oxide films, such as Al O , ZrO , TiO , 2 3 2 2
etc. using pulsed d.c. sputtering [34,35]. For instance, the deposition rate for clear Al O films formed using 2 3 this technique can reach 78% of the metal deposition rate, which is about 25 times the rate obtainable with r.f. power [1]. Therefore, oxide films can be produced at economical rates and so we can expect that research will be intensified for oxide superlattice coatings. For instance, Sproul already reported on the deposition of multilayer nanometre-scale oxide films composed of alternating layers of Al O and ZrO at a high rate onto 2 3 2 glass, silicon, and high-speed steel substrates [1]. Results are very encouraging because the energy delivered to the growing film during the pulse operation can stimulate film crystallization at low deposition temperatures. In conclusion it is worthwhile to note that superhard coatings in the form of superlattices represent a very important milestone in the development of superhard materials and the understanding of the origin of the superhardness. However, the maximum hardness of the superlattice coating is very strongly dependent on the superlattice period l, see for instance Fig. 3 presented in Ref. [15]. The strong dependence of H on l may cause large variations in the coating hardness, H, when deposited in industrial machines because it is difficult to ensure the same thickness of all superlattice layers on all coated objects, particularly when they have a complex shape. Similar variations in H can also be caused by the interdiffusion of elements in neighbouring layers at high service temperatures. These problems can be avoided if the superlattice coating is replaced with a single-layer nanocomposite coating.
3. Nanocomposite materials Materials are composed of grains separated by grain boundaries. The size of grains in currently produced materials, which can be called the conventional materials, varies over a wide range from about 100 nm to
324
J. Musil / Surface and Coatings Technology 125 (2000) 322–330
several hundred millimetres, corresponding to monocrystals. This means that the number of atoms in grains is always considerably greater than that in boundary regions. The behaviour of such materials is determined mainly by the bulk of grains in which dislocations play a decisive role. Properties of these materials are continuously improved by optimizing their composition, structure and the technological processes used for their formation. No fundamental qualitative changes in properties of the conventional materials can, however, be expected. Completely new properties are exhibited by nanocrystalline materials with a grain size of about 10 nm or less. The behaviour of these materials is determined mainly by processes in boundary regions because the number of atoms in the grains is comparable to or smaller than that in the boundary regions. Under these conditions dislocations do not exist [36 ], because grain boundaries prevent their formation, and the boundary regions play a decisive role in the material deformation. A new deformation mechanism, called grain boundary sliding, replaces the dislocation activity which is the dominant deformation process in conventional materials [37]. All these facts result in new unique properties of nanocrystalline materials. In the case where the size of grains decreases below 5 nm, the participation of atomic forces in material formation has to be considered and the formation of nanocrystalline subatomic structures can be expected [38]. The new unique properties of nanocrystalline materials are the main driving force stimulating their development. These materials can be prepared only by a method which simultaneously ensures a high rate of nucleation and a low growth rate of grains. This can be achieved relatively easily when the nanocrystalline materials are produced in the form of films. The most suitable method for the production of nanocrystalline films is magnetron sputtering.
4. Methods to control the size and orientation of grains in sputtered films The main task in the development of nanocomposite materials is to master control of their growth mechanism. Therefore, further basic processes used for controlling the size and crystallographic orientation of grains in growing films are given. There are two fundamental processes: (1) low-energy ion bombardment and (2) mixing process [39,40]. 4.1. Low-energy ion bombardment The process of low-energy ion bombardment controls the growth mechanism of the film by the energy delivered to the growing film by bombarding ions. The ion
bombardment is a strongly non-equilibrium process which heats the growing film at an atomic level. Therefore, it is called atomic scale heating (ASH ). The ion bombardment significantly differs from conventional heating because the kinetic energy of bombarding ions is transferred into very small areas of atomic dimensions and then very quickly conveyed into their close vicinity, i.e. the ASH is accompanied by extremely fast cooling rates of about 1014 K/s [41]. ASH can replace conventional heating and so produce dense films corresponding to zone T in the Thornton structural zone model when sputtering is carried out at low pressures of about 0.1 Pa and lower [42]. Ion bombardment of the growing film can restrict the grain growth and permit the formation of nanocrystalline films. The size and crystallographic orientation of grains can be controlled by the energy and flux of bombarding ions. This control of the film structure is, however, accompanied by film heating and is not convenient for all applications. 4.2. Mixing process The mixing process is based on the addition of one or several elements to a base, one element material. As at least two elements are present in the film, alloy films are formed by this process. The mixing process is an efficient method convenient for production of nanocrystalline films. Compared with ion bombardment, no substrate bias and heating are necessary to form the films with nanocrystalline structure. Also, metastable hightemperature phases can be formed on unheated substrates [43]. This is connected with efficient ASH, caused by condensing sputtered atoms, and subsequent extremely fast cooling at an atomic level.
5. Nanocrystalline alloy films The structure of alloy films depends on the amount and type of elements added to a base, one element material. Recent experiments carried out in our laboratory show that there are two groups of binary metal alloy films. The films of the first group are characterized by relatively narrow X-ray reflection lines ( FWHM≤1°). The films of the second group are very fine grained (nanocrystalline) or X-ray amorphous films characterized by very broad low-intensity reflections ( FWHM>1°). Several examples are given in Table 2. Here, together with the grain size, d, two characteristic parameters of the alloyed materials, i.e (i) difference in atomic radii of the alloy elements and (ii) enthalpy of the alloy formation DH , are also given. These quantif ties are often used to predict the creation of an amorphous state [51–53]. This state is expected to be formed in the case when (i) the atomic size difference is greater
325
J. Musil / Surface and Coatings Technology 125 (2000) 322–330
Table 2 Typical features of X-ray reflection line from selected alloy films sputtered on unheated substrates (T=RT ) at I =1 A, U =U and p =0.5–0.7 Pa d s fl Ar Alloy
2H (°)
I (cps) hkl
FWHM (°)
d (nm)
DH (kJ/mol ) f
r (nm) A
r (nm) B
Dr/ ra
Reference
Narrow reflections NiCr (80/20 wt.%) CrNi (60/40 wt.%) ZrY (70/30 at.%) TiSi (90/10 at.%) CuCr (60/40 at.%)
51.74 46.50 46.56 52.21 51.57
22642 7627 89231 756 1500
0.2995 0.2788 0.2910 0.3770 0.9360
51.97 59.94 49.58 36.16 12.38
−4 −9 14 −65 20
0.1246
0.125
0.003
[44]
0.160 0.144 0.117
0.181 0.1176 0.125
0.123 0.202 0.066
[45] [46 ] [47]
Broad reflections TiCu (50/50 at.%) TiAl (60/40 at.%) ZrCu (70/30 at.%)
46.70 45.80 43.55
15 45 135
4.7020 5.1018 9.8495
1.79 1.49 0.93
−9.6 −57 −12
0.144 0.144 0.160
0.117 0.143 0.117
0.207 0.007 0.255
[48] [49] [50]
a Dr/ r is the difference in atomic radii, r=(r +r ). A B
than 0.15 [52] and (ii) the enthalpy of alloy formation is negative and large. As can be seen from Table 2, these two materials parameters, Dr/ r and DH , are not, f however, sufficient to predict the structure of binary metal alloy films formed by sputtering. At present, it is not known which combinations of these two elements will form films with narrow and broad X-ray reflection lines, respectively. Despite these problems there exists a solution which allows the preparation of nanocrystalline alloy films. Alloy films with narrow X-ray reflection lines can be converted into films with broad X-ray reflection lines if nitrogen is added, i.e. when the nitride of the alloy film is formed, see e.g. Ref. [44]. This means that either binary metal alloys or their nitrides form nanocrystalline films. The structure of the nanocrystalline film can be controlled by the substrate bias, substrate temperature and by the amount of nitrogen incorporated into the film, i.e. by the energy delivered to the growing film.
6. Nanocomposite films based on nitrides of binary metal alloys Nanocomposite films consist of at least two phases. Experiments show that the incorporation of N into the growing film is not sufficient to produce the nanocomposite film with fully separated phases. To achieve this, the substrate has to be heated, see for instance the formation of ZrCu–N [50], NiCr–N [44] and TiNi–N [54] nanocomposite films. The formation of nanocomposite structures is connected with a segregation of the one-phase to grain boundaries of the second phase, and this effect is responsible for stopping of the grain growth. There are, however, two open questions. (1) What is a minimum temperature T necessary to start the segregation proseg cess. and are there some factors which could be used to decrease T ? (2) What is a thermal stability of nanoseg composite films? Recent experiments indicate that rare-
earth elements, such as yttrium, drastically reduce the grain size in metals, see for instance Ref. [55].
7. Hard and superhard nanocomposite coatings At present, significant effort is devoted to master the formation of nanocomposite coatings using magnetron sputtering because this technology can easily be scaled up for industrial use. The nanocomposite coatings are produced using so-called selective reactive magnetron sputtering [56 ]. In this deposition process, one element of the alloy is converted into nitride and the second element of the alloy is transported into the growing film unreacted. Recently, new hard (<40 GPa) supertough material of the type nc-TiC/a-C [57] with a remarkable plasticity (40% during nanoindentation deformation) and new superhard (≥40 GPa) materials of the type nc-MeN/a-Si N [12] with a high elastic recovery (up to 3 4 80%) in the form of nanocomposite coatings were developed; here nc- and a- denote the nanocrystalline and amorphous phases, respectively, and Me=Ti, W, V, Zr, etc. are transition metals. Even though both types of material are composed of nanocrystalline grains embedded in an amorphous matrix, they exhibit completely different physical properties. This is due to the different structures of the supertough and superhard nanocomposite coatings. A systematic investigation carried out on systems ZrCu–N [13] and TiAl–N [49] showed that superhard nanocomposite coatings can be composed not only of two hard phases as proposed by Veprˇek et al. [58] but also in the case when only one phase is hard, e.g. a nc-ZrN/Cu nanocomposite [13]. This means that there are two groups of superhard nanocomposite coatings: 1. nc-MeN/nitride (e.g. a-Si N , a-TiB , etc.); 3 4 2 2. nc-MeN/metal (e.g. Cu, Ni, Y, Ag, Co, etc.). The hardness H of films of both groups can be continuously varied from low values of about 10 GPa to very
326
J. Musil / Surface and Coatings Technology 125 (2000) 322–330
high values, achieving up to 50–70 GPa. Moreover, it is worthwhile to note that the hardness of nanocomposite films is strongly correlated with their structure. 7.1. Structure of superhard nanocomposite coatings Superhard films with H≥40 GPa are formed only when their structure is close to X-ray amorphous. This structure corresponds to a transition from the crystalline structure to an amorphous one. The main factors which govern the formation of films with an X-ray amorphous structure are: the energy delivered to the growing film, the substrate temperature T , the type of elements forms ing the film [their (i) mutual solubility or immiscibility, (ii) ability to form intermetallic compounds, (iii) chemical affinity, and (iv) binding energy], gap of elements immiscibility, enthalpy of the alloy formation and the content of individual phases in the nanocomposite film, e.g. the content of soft phase in a composite of the type nc-Me/metal. The structure of hard (<40 GPa) and superhard (≥40 GPa) nanocomposite coatings is significantly different. A systematic investigation of the correlation between the hardness and structure in the ZrCu–N and TiAl–N nanocomposite films showed that (a) the hard films are characterized by many reflections from polyoriented grains of both phases, (b) the superhard films are two-phase nanocomposites one phase of which has a nanocrystalline structure and the second is X-ray amorphous, and (c) the maximum hardness is achieved only when all grains are oriented in the same direction and the size of grains has an optimum value of approximately several tens of nanometres. 7.2. Classification of hard nanocomposite coatings Till now, various different types of hard nanocomposite coating have been prepared: 1. nc-MeN/a-nitride, e.g. nc-MeN/a-Si N (Me=Ti, W, 3 4 V ) [58,61], nc-TiN/a-Si N [62]; 3 4 2. nc-MeN/nc-nitride, e.g. nc-TiN/nc-BN [58]; 3. nc-MeC/a-C, e.g. nc-TiC/DLC [7]; 4. nc-MeN/ metal, e.g. nc-ZrN/Cu [13], nc-( Ti,Al )/AlN [49,60], nc-CrN/Cu [47]; 5. nc-MeN or MeC/a-boron compounds, e.g. nc-Ti(B,O)/quasi-a-(TiB , TiB and B O ) [63], Ti– 2 2 3 B–C [64]; 6. nc-WC+nc-WS /DLC [65]; 2 7. nc-MeC/a-C+a-nitride, e.g. nc-Mo C/a-C+a2 Mo N [59]. 2 This survey shows that all hard nanocomposite coatings contain one or two hard crystalline phases. The second phase is more complicated. It is either amorphous (e.g. a-Si N ) or crystalline (e.g. nc-BN [58]). Sometimes, the 3 4 content of the second phase in the nanocomposite coating is very low at approximately 1–2 wt.% (e.g. Cu
in nc-ZrN/Cu films [13]). In such a case, it is very difficult, without a HRTEM investigation, to determine if the second phase is crystalline or amorphous because the X-ray reflections from a small quantity of grains are below the detection limit. This means that the second phase can be incorrectly interpreted as amorphous when determined from X-ray diffraction analysis only. On the basis of these facts hard nanocomposite coatings can be divided into two main groups: (1) crystalline/amorphous nanocomposites and (2) crystalline/crystalline nanocomposites. At present, no definite methodology exists on how to select the combination of elements to produce films with nanocrystalline and/or X-ray amorphous structure. Such films can be formed from nitrides of alloys composed of elements which exhibit a wide miscibility gap and contain one element forming hard nitride, for instance nitrides of binary alloys formed of immiscible elements such as Cu–Cr, Zr–Y, etc. or transition metal nitride/boride and nitride/carbide systems [64]. A miscibility gap between intermetallics is also sufficient to produce hard nanocomposite coating, e.g. ZrCu–N film. There is also a possibility to produce the nanocomposite coating from nitrides of alloys whose elements form a solid solution, e.g. Ti Al N film. This is enabled by 1−x x the existence of a gap in the alloy composition x where the structure of the film is quasi-X-ray amorphous [66 ]. 7.3. Mechanical properties of hard nanocomposite coatings A short survey of basic mechanical characteristics of some recently prepared nanocomposite films is given in Table 3. For comparison, the characteristics of some selected hard bulk materials and hard amorphous carbon films are also given in this table. The hard nanocomposite coating is characterized not only by its hardness H but also by its Young’s modulus E and elastic recovery W . The determination of these e quantities for thin films is, however, difficult because they strongly vary with the load L used in their measurement. Despite these problems, Figs. 1 and 2 display H as a function E1=E/(1−n2) and the elastic recovery W as a function H, because these dependencies were e measured under the same conditions in one laboratory and on different nanocomposite systems. The high hardness of the material is only one parameter which ensures scratch and abrasion resistance. Protective overcoat films must be highly resistant also to plastic deformation during contact events. This requires a low Young’s modulus E since, according to Johnson analysis, the load P needed to initiate plastic deformation when a y rigid sphere of radius r is pressed into the coating is proportional to H3/E2 [74]. The ratio H3/E2 is a parameter which controls the resistance of materials to plastic deformation. The likelihood of plastic deformation is
327
J. Musil / Surface and Coatings Technology 125 (2000) 322–330 Table 3 Comparison of hard bulk materials, hard single layer films and selected hard and superhard nanocomposite coatings Material Bulk materials Diamond Boron Sapphire Amorphous films DLC a-C (cathodic arc) Nanocomposite single layer films nc-TiN/Si N 3 4 nc-TiN/BN nc-W N/a-Si N 2 3 4 Ti–B–C Ti–B–N Zr Cu N 98 2 W Ni N 86.7 8.3 5 W Si N 68 14 18 nc-Mo C/a-(C+Mo N ) 2 2 Ti Al N 45 55 Ti Al N 60 40 ZrY–N CrNi–N Ti Si N 75 25 Ti C ( TiC/a-C ) 0.32 0.68
H (GPa)
E1=E/(1−n2) (GPa)
100 35 30
1050 470 441
65 >59
550 >395
48 69 51 71 54 54 55 45 49 47 40 41 32 29 32
~565 585 560 486 ~500 394 510 – 440 409 650 319 253 256 370
W (%) e
H3/E12
d (nm)
Reference
0.91 0.19 0.14
[67] [68] [69]
80–90
0.91 ~1.3
[7] [74]
–a – – 80.5 – 81 – – 67 74 – 77 74 67 60
~0.34 0.96 0.42 1.52 0.63 1.03 0.64 – 0.61 0.62 0.15 0.66 0.50 0.36 0.239
4.5 9 3.5 ~1 ~1 35 – – 27 30 – – – – 10–50
[70] [58] [58,70,71] [64] [64] [13] [75] [61] [59] [49] [14] [45] [72] [46 ] [73]
a Denotes data not given in the references or not determined.
Fig. 1. H as a function of E/(1−n2) for Zr–Cu–N, Zr–Y–N, Cr–Ni–N, Ti–Si–N and Ti–Al–N nanocomposite films sputtered at different deposition conditions, i.e. T , U , i and p . s s s N2
reduced in materials with high hardness and low modulus [74]. In general, a low modulus is also desirable as it allows the given load to be distributed over a wider area. The ratio H3/E2 of nanocomposite films spreads over a very wide range from about 0.15 to 1.52, see Fig. 3 and Table 3. Data given in Table 3 and in Figs. 1– 3 show that the elastic recovery W and the resistance e
of materials to plastic deformation can be controlled by the film hardness H and its elastic modulus E. A spread of experimental points around the straight lines in Figs. 1–3 is connected with the variation in the film structure induced particularly by different (i) deposition conditions and (ii) chemical composition of the film. Therefore, H and E of nanocomposite coatings
328
J. Musil / Surface and Coatings Technology 125 (2000) 322–330
Fig. 2. Elastic recovery upon indentation W as a function of H for Zr–Cu–N, Zr–Y–N, Cr–Ni–N, Ti–Si–N and Ti–Al–N nanocomposite films e sputtered at different deposition conditions, i.e. T , U , i and p . s s s N2
Fig. 3. The ratio H3/E2 characterizing the resistance of the material to plastic deformation as a function of H for Zr–Cu–N, Zr–Y–N, Cr–Ni–N, Ti–Si–N and Ti–Al–N nanocomposite films sputtered at different deposition conditions, i.e. T , U , i and p . s s s N2
prepared by plasma CVD and magnetron-assisted laser deposition processes differ from those prepared by magnetron sputtering. A further systematic investigation is necessary to master the preparation of materials with prescribed properties.
8. Conclusions The main results obtained in the development of hard and superhard nanocomposite coatings can be summarized as follows. 1. Hard (<40 GPa) coatings are characterized by a high
plastic deformation increasing with decreasing H up to about 70% for H#10 GPa. 2. Superhard (≥40 GPa) coatings are characterized by a high elastic recovery increasing with increasing H up to about 85% for H#70 GPa. 3. There are two types of superhard nanocomposite film: (1) nc-MeN/a-nitride and (2) nc-MeN/metal. This means that the superhard film can be composed either of two hard phases or of one hard phase and the second soft phase. 4. The hardness of nanocomposite coatings correlates well with their structure. A decrease of the film crystallinity, characterized by a decrease of the inten-
J. Musil / Surface and Coatings Technology 125 (2000) 322–330
sity of X-ray reflection lines and by an increase of their broadening, results in an increase of hardness. 5. The structure of superhard films is close to X-ray amorphous. The nanocomposite films are fascinating materials since novel structures and new physical properties are expected. The development of nanocomposite films is, however, only at the beginning. Many problems are unsolved and many questions remain open. Therefore, further systematic research in this field is required. Further research is focusing on the following problems: (i) understanding of the origin of superhardness; (ii) correlation between the mechanical parameters of materials and process parameters; (iii) dramatic changes in crystallographic orientation of grains in alloy films, their hardness and elastic recovery induced by incorporation of nitrogen; (iv) formation of nanocomposite films with controlled hardness, elastic modulus and elastic recovery and new functional properties; and (v) investigation of materials having very small grains of about 1 nm in size.
Acknowledgements The author would like to thank his Ph.D. students H. Hruby´, I. Leipner, H. Pola´kova´, F. Regent, Z. Soukup and P. Zeman for the preparation of many nanocomposite coatings of different chemical composiˇ erstvy´ and Dr. M. Kolega for X-ray tion and Dr. R. C characterization of these coatings. This work was supported in part by the Grant Agency of the Czech Republic under Project No. 106/96/K245 and by the Ministry of Education of the Czech Republic under Project Nos. VS 96/059 and ME 173/1999.
References [1] W.D. Sproul, Science 273 (1996) 889. [2] H. Holleck, Ch. Ku¨hl, H. Schultz, J. Vac. Sci. Technol. A 3 (6) (1985) 234. [3] H. Holleck, J. Vac. Sci. Technol. A 4 (6) (1986) 2661. [4] C.A. Brookes, in: J.E. Field (Ed.), The Properties of Diamond, Academic Press, New York, 1979, p. 383. [5] B.M. Kramer, P.K. Judd, J. Vac. Sci. Technol. A 3 (1985) 2439. [6 ] W.D. Mu¨nz, J. Vac. Sci. Technol. A 4 (6) (1986) 2717. [7] A.A. Voevodin, J.S. Zabinski, Diamond Relat. Mater. 7 (1998) 463. [8] H. Sjo¨stro¨m, S. Stafstro¨m, M. Boman, J.-E. Sundgren, Phys. Rev. Lett. 76 (1996) 56. [9] H. Sjo¨stro¨m, L. Hultman, J.-E. Sundgren, S.V. Hainsworth, T.F. Page, G.S.A.M. Teunissen, J. Vac. Sci. Technol. A 14 (1996) 56. [10] U. Helmersson, S. Todorova, S.A. Barnett, J.-E. Sundgren, L.C. Markert, J.E. Greene, J. Appl. Phys. 62 (1987) 481. [11] X. Chu, S.A. Barnett, M.S. Wong, W.D. Sproul, Surf. Coat. Technol. 57 (1993) 13. [12] S. Veprˇek, S. Reiprich, S. Li, Appl. Phys. Lett. 66 (20) (1995) 2640. [13] J. Musil, P. Zeman, H. Hruby´, P. Mayrhofer, ZrN/Cu nanocom-
329
posite film — novel superhard material, Proc. ICMCTF’99, April 12–18, San Diego, CA (1999), paper no. B1-2-9. [14] Y. Min, Y. Makino, M. Nose, K. Nogi, Thin Solid Films 228 (1999) 204. [15] U. Helmersson, S. Todorova, S.A. Barnett, J.-E. Sundgren, L.C. Markert, J.E. Greene, J. Appl. Phys. 62 (1987) 481. [16 ] P.B. Mirakami, L. Hultman, S.A. Barnett, Appl. Phys. Lett. 57 (25) (1990) 2654. [17] M. Shinn, L. Hultman, S.A. Barnett, J. Mater. Res. 7 (4) (1992) 901. [18] X. Chu, M.S. Wong, W.D. Sproul, S.L. Rohde, S.A. Barnett, J. Vac. Sci. Technol. A 10 (4) (1992) 1604. [19] X. Chu, S.A. Barnett, M.S. Wong, W.D. Sproul, Surf. Coat. Technol. 57 (1993) 13. [20] D. Li, X.W. Lin, S.C. Chen, V.P. Dravid, Y.W. Chung, M.S. Wong, W.D. Sproul, Appl. Phys. Lett. 68 (9) (1996) 1211. [21] M.L. Wu, W.D. Qian, Y.W. Chung, Y.Y. Wang, M.S. Wong, W.D. Sproul, Thin Solid Films 308/309 (1997) 113. [22] H. Jensen, J. Sobota, G. Sorensen, Surf. Coat. Technol. 94/95 (1997) 174. [23] H. Jensen, J. Sobota, G. Sorensen, J. Vac. Sci. Technol. A 16 (3) (1998) 1180. [24] J.E. Kryanowski, Mater. Res. Soc. Symp. Proc. 239 (1992) 509. [25] S.A. Barnett, in: M. Fracombe, J.A. Vossen ( Eds.), Physics of Thin Films, Academic Press, New York, 1993, p. 1. [26 ] M. Shinn, L. Hultman, S.A. Barnett, J. Mater. Res. 7 (1992) 901. [27] X. Chu, S.A. Barnett, J. Appl. Phys. 77 (1995) 4403. [28] B.G. Wendler, P. Kula, K. Jakubowski, S. Fauvry, Ph. Kapsa, L. Vincent, D. Heper, Proc. 36th Tagung der Deutschen Gesellschaft fu¨r Obrfla¨chen- und Galvanotechnik, October 7–9, Schwabisch Gmu¨nd, Germany (1998) 6–11. [29] B.G. Wendler, M. Molinaro, Proc. 10th Int. Summer School on Modern Plasma Surface Technology, May 10–14, Mielno, Poland (1998) 423–436. [30] K.K. Shih, D.B. Dove, Appl. Phys. Lett. 61 (6) (1992) 654. [31] X. Chu, M.S. Wong, W.D. Sproul, A.S. Barnett, Mechanical properties and microstructures of polycrystalline metal/ceramic superlattices, TiN/Ni and TiN/Ni Cr , paper presented at 0.9 0.1 ICMCTF-20, San Diego, CA, 1993. [32] A. Madan, X. Chu, S.A. Barnett, Appl. Phys. Lett. 68 (16) (1996) 2198. [33] A. Madan, Y.-Y. Wang, S.A. Barnett, C. Engstro¨m, H. Ljungcrantz, L. Hultman, M. Grimsditch, J. Appl. Phys. 84 (2) (1998) 776. [34] J.M. Schneider, W.D. Sproul, A.A. Voevodin, A. Matthews, J. Vac. Sci. Technol. A 15 (3) (1997) 1084. [35] J.M. Schneider, W.D. Sproul, Reactive pulsed dc magnetron sputtering and control, Handbook of Thin Film Process Technology, IOP, Bristol, 1998, pp. A5.1:1–A5.1:12. [36 ] S. Veprˇek, S. Reiprich, Thin Solid Films 268 (1995) 64. [37] R.W. Siegel, Nanophase materials: Structure, defects and properties, Proc. ATC Int. Symp. (ACTA), October 29–30, Tokyo, Japan (1996) 1. [38] P.Yu. Butyagin, Colloid J. 59 (1997) 112. [39] J. Musil, J. Vlcˇek, Czech. J. Phys. 48 (10) (1998) 1209. [40] J. Musil, J. Vlcˇek, Magnetron sputtering of alloy and alloy-based films, Thin Solid Films 343–344 (1999) 47. [41] K.D. Leedy, J.M. Rigsbee, J. Vac. Sci. Technol. A 14 (1996) 2202. [42] J. Musil, Vacuum 50 (3/4) (1998) 363. [43] J. Musil, A.J. Bell, J. Vlcˇek, T. Hurkmans, J. Vac. Sci. Technol. A 14 (4) (1996) 2247. [44] J. Musil, F. Regent, J. Vac. Sci. Technol. A 16 (6) (1998) 3301. [45] J. Musil, H. Pola´kova´, Hard nanocomposite Zr–Y–N coatings. Correlation between hardness and structure, 2nd Asian–European Int. Conf. on Plasma Surface Engineering, September 15–19, Beijing, China (1999), Paper No. Sat-0A3-2.
330
J. Musil / Surface and Coatings Technology 125 (2000) 322–330
[46 ] J. Musil, H. Pola´kova´, V. Cibulka, Czech. J. Phys. 49 (3) (1999) 359. [47] J. Musil, I. Leipner, M. Kolega, Surf. Coat. Technol. 115 (1) (1999) 32. [48] J. Musil, Z. Soukup, unpublished results. [49] J. Musil, H. Hruby´, Superhard nanocomposite Ti Al N films 1−x x prepared by magnetron sputtering, Proc. 14th Int. Symp. on Plasma Chemistry, August 2–6, Praha, Czech Republic, Vol. III (1999) 1605–1610. [50] J. Musil, P. Zeman, Vacuum 52 (1999) 269. [51] W.L. Johnson, Progr. Mater. Sci. 30 (1986) 81. [52] G. Weigang, H. Hecht, G. von Minnigerode, Z. Phys. B 96 (1995) 349. [53] J.R. Ding, D.Y. Che, H.B. Yhang, K. Tao, B.X. Liu, Appl. Phys. Lett. 60 (8) (1992) 994. [54] M. Misˇina, J. Musil, S. Kadlec, Surf. Coat. Technol. 110 (1998) 168. [55] Z. Liu, R. Singh, K. Poole, R.J. Diefendorf, J. Harris, K. Cannon, J. Vac. Sci. Technol. B 15 (1997) 1990. [56 ] L. May, W.R. Allen, A.L. Glower, J.C. Mabon, J. Vac. Sci. Technol. B 13 (1995) 361. [57] A.A. Voevodin, S.V. Prasad, J.S. Zabinski, J. Appl. Phys. Lett. 82 (2) (1997) 855. [58] S. Veprˇek, P. Nesla´dek, A. Niederhofer, F. Glatz, M. Jı´lek, M. Sˇ´ıma, Surf. Coat. Technol. 108/109 (1998) 138. [59] M. Benda, J. Musil, Vacuum 55 (1999) 171. [60] Y. Min, Y. Makino, M. Nose, K. Nogi, Thin Solid Films 339 (1/2) (1999) 238. [61] C. Louro, A. Cavaliero, Hardness versus structure in W–Si–N sputtered coatings, 6th Int. Conf. on Plasma Surface Engineering, PSE-98, September 14–18, Garmisch-Partenkirchen, Germany (1998), paper no. ThWB5.
[62] M. Diserens, J. Patscheider, F. Levy, Surf. Coat. Technol. 108/ 109 (1998) 241. [63] C. Mitterer, P. Losbichler, F. Hofer, P. Warbichler, W. Gissler, Vacuum 50 (3/4) (1998) 313. [64] C. Mitterer, P.H. Mayrhofer, M. Beschliesser, P. Losbichler, P. Warbichler, F. Hofer, P.N. Gibson, W. Gissler, H. Hruby´, J. Musil, J. Vlcˇek, Proc. ICMCTF’99, April 12–16, San Diego, CA (1999), paper no. BP-16. [65] A.A. Voevodin, J.P. O’Neill, J.S. Zabinski, Surf. Coat. Technol. (1999) submitted. [66 ] U. Wahlstro¨m, L. Hultman, J.-E. Sundgren, F. Abidi, I. Petrov, J.E. Greene, Thin Solid Films 235 (1993) 62. [67] J.E. Field, The Properties of Diamond, Academic Press, London, 1979. [68] S.M. Gorbatkin, R.L. Rhoades, T.Y. Tsui, W.C. Oliver, Appl. Phys. Lett. 65 (1994) 2672. [69] W.C. Oliver, G.M. Pharr, J. Mater. Res. 7 (1991) 1564. [70] S. Veprˇek, Thin Solid Films 317 (1998) 449. [71] S. Veprˇek, Thin Solid Films 297 (1997) 145. [72] J. Musil, F. Regent, Structure and microhardness of magnetron sputtered nanocrystalline CrNi–N films, Proc. 14th Int. Symp. on Plasma Chemistry, August 2–6, Praha, Czech Republic, Vol. III (1999) 1617–1622. [73] A.A. Voevodin, J.S. Zabinski, J. Mater. Sci. 33 (1998) 319. [74] T.Y. Tsui, G.M. Pharr, W.C. Oliver, C.S. Bhatia, R.L. White, S. Anders, A. Anders, I.G. Brown, Mater. Res. Soc. Symp. Proc. 383 (1995) 447. [75] A. Cavaleiro, B. Trindale, M.T. Viera, Deposition and characterization of fine-grained W–Ni–C/N ternary films, 6th Int. Conf. on Plasma Surface Engineering, PSE-98, Garmisch-Partenkirchen, Germany, Book of Abstracts (1998) 265.
Hard and superhard nanocomposite coatings J. Musil * Department of Physics, University of West Bohemia, P.O. Box 314, 30614 Plzenˇ, Czech Republic
Abstract This article reviews the development of hard coatings from a titanium nitride film through superlattice coatings to nanocomposite coatings. Significant attention is devoted to hard and superhard single layer nanocomposite coatings. A strong correlation between the hardness and structure of nanocomposite coatings is discussed in detail. Trends in development of hard nanocomposite coatings are also outlined. © 2000 Elsevier Science S.A. All rights reserved. Keywords: Hardness; Magnetron sputtering; Mechanical properties; Nanocomposite coatings; Structure
1. Introduction Hard coatings have been successfully used for protection of materials and particularly to enhance the life of cutting tools since the 1970s. Both the technological process of their production and their properties, i.e. hardness, wear and oxidation resistance, however, are continuously being improved. Important milestones in the development of hard coatings are briefly summarized in Table 1. This table shows a clear effort (i) to decrease the temperature T at which hard coatings are formed and (ii) to improve the properties of hard coatings, particularly to increase the hardness and oxidation resistance. The oxidation resistance should be increased up to approximately 1000°C because during high-speed machining the temperature of the tool tip can reach 1000°C and the coating should be stable at such high temperatures. As to the hardness, the coatings are usually divided into two groups: (1) hard coatings having a hardness <40 GPa, and (2) superhard coatings having a hardness >40 GPa. Compared to a large number of hard materials, there are only a few superhard materials, i.e. cubic boron nitride (c-BN ), amorphous diamond-like carbon (DLC ), amorphous carbon nitride (a-CN ) and polyx crystalline diamond. Moreover, these superhard materials are thermodynamically unstable. This is a serious disadvantage which strongly limits their utilization in some applications. For instance, the high chemical * Tel.: +420-19-279072; fax: +420-19-279071. E-mail address: [email protected] (J. Musil )
affinity of carbon to iron limits the applicability of diamond coated cutting tools to machining of aluminum, their alloys and wood only. Similar problems can be expected when the c-BN coating is used in cutting of steels due to the chemical dissolution of boron in iron. These problems stimulated intensive research in this field, and recently new superhard materials based on superlattices and nanocomposites were developed.
2. Hard superlattice coatings Superlattice coatings are nanometre-scale multilayers composed of two different alternating layers with a superlattice period, i.e. the bilayer thickness of two materials, ranging from 5 to 10 nm. The bilayers of these superlattices can be metal layers, nitrides, carbides or oxides of different materials or a combination of one layer made of nitride, carbide or oxide of one metal and the second layer made of another metal. According to the composition of the bilayer, superlattice coatings can be divided into five groups: (1) metal superlattices, (2) nitride superlattices, (3) carbide superlattices, (4) oxide superlattices and (5) nitride, carbides or oxides/metal superlattices. Experiments show that metal superlattices exhibit a relatively low hardness. On the contrary, single-crystal nitride superlattice coatings are superhard materials with a hardness ranging from 45 to 55 GPa, e.g. TiN/VN, 56 GPa [15]; TiN/( V Nb )N, 41 GPa [16 ]; 0.6 0.4 TiN/NbN, 51 GPa [17]; TiN/Nb, 52 GPa [18,19]; TiN/CN , 45–55 GPa [20]; ZrN/CN , 40–45 GPa [21]; x x
0257-8972/00/$ - see front matter © 2000 Elsevier Science S.A. All rights reserved. PII: S0 2 5 7- 8 9 7 2 ( 9 9 ) 0 0 58 6 - 1
J. Musil / Surface and Coatings Technology 125 (2000) 322–330
323
Table 1 Important steps in development of hard coatings Coating
Material
H (GPa)
Main characteristics
Single layer Single layer Multilayer Single layer Single layer Single layer Single layer Single layer Superlattices Single layer Single layer Single layer
TiN, TiC, Al O 2 3 TiN, TiC TiC/TiB 2 c-BN diamond TiAlN DLC CN x TiN/VN, TiN/NbN, etc. nc-MeN/a-nitride nc-MeN/metal Ti Al N 0.4 0.6
21, 28, 21 21, 28
CVD at T around 1000°C on cemented carbides PVD at T≤550°C on steel substrates About 103 phase boundaries TiC/TiB [2] 2 High chemical affinity of C to iron Chemical dissolution of B in iron [5] Oxidation resistance up to 800°C [6 ] Amorphous phase [7] Substoichiometric (x=0.2–0.35) turbostratic structure [8,9] Superlattice period 5–10 nm [10,11] Superlattice period 5–10 nm [10,11] Nanocomposite [13] Nanocomposite, oxidation resistance up to 950°C [14]
50 [3] 90 [4] 65 50–60 ~50 ~50 ~50 ~32
TiN /C–N, 20–50 GPa [22,23]. The overall hardness of x the superlattice coating is therefore greater than that of the materials of the individual components of the bilayer, e.g. H #52 GPa, H =21 GPa, H =14 GPa. TiN/NbN TiN NbN This hardness enhancement is a very complex phenomenon. In spite of this, several models which explain multilayers strengthening have already been developed [24–27]. The model of Shinn et al. [26 ] shows that (i) a difference in elastic modulus between the two layer materials is required to increase the hardness of the superlattice film, and (ii) the coherency strain at the interface between the two layers has only a minor effect. The model of Chu and Barnett [27] is based on restricted dislocation movement within and between layers in the superlattice coating. It predicts a peak in hardness when there is a difference in shear modulus between two layer materials and sharp interfaces between layers. So far, less attention has been devoted to carbide multilayer coatings. These coatings can also be superhard, up to 55 GPa, e.g. TiC/VC, 52 GPa [28]; TiC/NbC, 45–55 GPa [29]. Considerable attention has been devoted to nitride/metal superlattice coatings. This combination of bilayer materials, i.e. the hard nitride with a more ductile metal, makes it possible to improve the toughness of the coating while retaining its relatively high hardness (≥30 GPa), e.g. TiN/Ti, 36.8 GPa [30]; WN/W, 34 GPa [30]; HfN/Hf, 50 GPa [30]; TiN/Ni, 35 GPa [31]; TiN/Ni Cr , 32 GPa [31]; NbN/Mo, 0.4 0.1 33 GPa [32,33]; NbN/W, 30 GPa [32,33]. An improvement in the toughness of the coating increases its adhesion to the substrate, which is of fundamental importance for coating applications. At present, practically no data are available on oxide superlattice coatings. This is probably because the deposition rate of oxide films using both d.c. and r.f. reactive magnetron sputtering is, compared with that of metals, too low. Recently, this situation strongly changed. Considerable progress has been made in magnetron deposition of oxide films, such as Al O , ZrO , TiO , 2 3 2 2
etc. using pulsed d.c. sputtering [34,35]. For instance, the deposition rate for clear Al O films formed using 2 3 this technique can reach 78% of the metal deposition rate, which is about 25 times the rate obtainable with r.f. power [1]. Therefore, oxide films can be produced at economical rates and so we can expect that research will be intensified for oxide superlattice coatings. For instance, Sproul already reported on the deposition of multilayer nanometre-scale oxide films composed of alternating layers of Al O and ZrO at a high rate onto 2 3 2 glass, silicon, and high-speed steel substrates [1]. Results are very encouraging because the energy delivered to the growing film during the pulse operation can stimulate film crystallization at low deposition temperatures. In conclusion it is worthwhile to note that superhard coatings in the form of superlattices represent a very important milestone in the development of superhard materials and the understanding of the origin of the superhardness. However, the maximum hardness of the superlattice coating is very strongly dependent on the superlattice period l, see for instance Fig. 3 presented in Ref. [15]. The strong dependence of H on l may cause large variations in the coating hardness, H, when deposited in industrial machines because it is difficult to ensure the same thickness of all superlattice layers on all coated objects, particularly when they have a complex shape. Similar variations in H can also be caused by the interdiffusion of elements in neighbouring layers at high service temperatures. These problems can be avoided if the superlattice coating is replaced with a single-layer nanocomposite coating.
3. Nanocomposite materials Materials are composed of grains separated by grain boundaries. The size of grains in currently produced materials, which can be called the conventional materials, varies over a wide range from about 100 nm to
324
J. Musil / Surface and Coatings Technology 125 (2000) 322–330
several hundred millimetres, corresponding to monocrystals. This means that the number of atoms in grains is always considerably greater than that in boundary regions. The behaviour of such materials is determined mainly by the bulk of grains in which dislocations play a decisive role. Properties of these materials are continuously improved by optimizing their composition, structure and the technological processes used for their formation. No fundamental qualitative changes in properties of the conventional materials can, however, be expected. Completely new properties are exhibited by nanocrystalline materials with a grain size of about 10 nm or less. The behaviour of these materials is determined mainly by processes in boundary regions because the number of atoms in the grains is comparable to or smaller than that in the boundary regions. Under these conditions dislocations do not exist [36 ], because grain boundaries prevent their formation, and the boundary regions play a decisive role in the material deformation. A new deformation mechanism, called grain boundary sliding, replaces the dislocation activity which is the dominant deformation process in conventional materials [37]. All these facts result in new unique properties of nanocrystalline materials. In the case where the size of grains decreases below 5 nm, the participation of atomic forces in material formation has to be considered and the formation of nanocrystalline subatomic structures can be expected [38]. The new unique properties of nanocrystalline materials are the main driving force stimulating their development. These materials can be prepared only by a method which simultaneously ensures a high rate of nucleation and a low growth rate of grains. This can be achieved relatively easily when the nanocrystalline materials are produced in the form of films. The most suitable method for the production of nanocrystalline films is magnetron sputtering.
4. Methods to control the size and orientation of grains in sputtered films The main task in the development of nanocomposite materials is to master control of their growth mechanism. Therefore, further basic processes used for controlling the size and crystallographic orientation of grains in growing films are given. There are two fundamental processes: (1) low-energy ion bombardment and (2) mixing process [39,40]. 4.1. Low-energy ion bombardment The process of low-energy ion bombardment controls the growth mechanism of the film by the energy delivered to the growing film by bombarding ions. The ion
bombardment is a strongly non-equilibrium process which heats the growing film at an atomic level. Therefore, it is called atomic scale heating (ASH ). The ion bombardment significantly differs from conventional heating because the kinetic energy of bombarding ions is transferred into very small areas of atomic dimensions and then very quickly conveyed into their close vicinity, i.e. the ASH is accompanied by extremely fast cooling rates of about 1014 K/s [41]. ASH can replace conventional heating and so produce dense films corresponding to zone T in the Thornton structural zone model when sputtering is carried out at low pressures of about 0.1 Pa and lower [42]. Ion bombardment of the growing film can restrict the grain growth and permit the formation of nanocrystalline films. The size and crystallographic orientation of grains can be controlled by the energy and flux of bombarding ions. This control of the film structure is, however, accompanied by film heating and is not convenient for all applications. 4.2. Mixing process The mixing process is based on the addition of one or several elements to a base, one element material. As at least two elements are present in the film, alloy films are formed by this process. The mixing process is an efficient method convenient for production of nanocrystalline films. Compared with ion bombardment, no substrate bias and heating are necessary to form the films with nanocrystalline structure. Also, metastable hightemperature phases can be formed on unheated substrates [43]. This is connected with efficient ASH, caused by condensing sputtered atoms, and subsequent extremely fast cooling at an atomic level.
5. Nanocrystalline alloy films The structure of alloy films depends on the amount and type of elements added to a base, one element material. Recent experiments carried out in our laboratory show that there are two groups of binary metal alloy films. The films of the first group are characterized by relatively narrow X-ray reflection lines ( FWHM≤1°). The films of the second group are very fine grained (nanocrystalline) or X-ray amorphous films characterized by very broad low-intensity reflections ( FWHM>1°). Several examples are given in Table 2. Here, together with the grain size, d, two characteristic parameters of the alloyed materials, i.e (i) difference in atomic radii of the alloy elements and (ii) enthalpy of the alloy formation DH , are also given. These quantif ties are often used to predict the creation of an amorphous state [51–53]. This state is expected to be formed in the case when (i) the atomic size difference is greater
325
J. Musil / Surface and Coatings Technology 125 (2000) 322–330
Table 2 Typical features of X-ray reflection line from selected alloy films sputtered on unheated substrates (T=RT ) at I =1 A, U =U and p =0.5–0.7 Pa d s fl Ar Alloy
2H (°)
I (cps) hkl
FWHM (°)
d (nm)
DH (kJ/mol ) f
r (nm) A
r (nm) B
Dr/ ra
Reference
Narrow reflections NiCr (80/20 wt.%) CrNi (60/40 wt.%) ZrY (70/30 at.%) TiSi (90/10 at.%) CuCr (60/40 at.%)
51.74 46.50 46.56 52.21 51.57
22642 7627 89231 756 1500
0.2995 0.2788 0.2910 0.3770 0.9360
51.97 59.94 49.58 36.16 12.38
−4 −9 14 −65 20
0.1246
0.125
0.003
[44]
0.160 0.144 0.117
0.181 0.1176 0.125
0.123 0.202 0.066
[45] [46 ] [47]
Broad reflections TiCu (50/50 at.%) TiAl (60/40 at.%) ZrCu (70/30 at.%)
46.70 45.80 43.55
15 45 135
4.7020 5.1018 9.8495
1.79 1.49 0.93
−9.6 −57 −12
0.144 0.144 0.160
0.117 0.143 0.117
0.207 0.007 0.255
[48] [49] [50]
a Dr/ r is the difference in atomic radii, r=(r +r ). A B
than 0.15 [52] and (ii) the enthalpy of alloy formation is negative and large. As can be seen from Table 2, these two materials parameters, Dr/ r and DH , are not, f however, sufficient to predict the structure of binary metal alloy films formed by sputtering. At present, it is not known which combinations of these two elements will form films with narrow and broad X-ray reflection lines, respectively. Despite these problems there exists a solution which allows the preparation of nanocrystalline alloy films. Alloy films with narrow X-ray reflection lines can be converted into films with broad X-ray reflection lines if nitrogen is added, i.e. when the nitride of the alloy film is formed, see e.g. Ref. [44]. This means that either binary metal alloys or their nitrides form nanocrystalline films. The structure of the nanocrystalline film can be controlled by the substrate bias, substrate temperature and by the amount of nitrogen incorporated into the film, i.e. by the energy delivered to the growing film.
6. Nanocomposite films based on nitrides of binary metal alloys Nanocomposite films consist of at least two phases. Experiments show that the incorporation of N into the growing film is not sufficient to produce the nanocomposite film with fully separated phases. To achieve this, the substrate has to be heated, see for instance the formation of ZrCu–N [50], NiCr–N [44] and TiNi–N [54] nanocomposite films. The formation of nanocomposite structures is connected with a segregation of the one-phase to grain boundaries of the second phase, and this effect is responsible for stopping of the grain growth. There are, however, two open questions. (1) What is a minimum temperature T necessary to start the segregation proseg cess. and are there some factors which could be used to decrease T ? (2) What is a thermal stability of nanoseg composite films? Recent experiments indicate that rare-
earth elements, such as yttrium, drastically reduce the grain size in metals, see for instance Ref. [55].
7. Hard and superhard nanocomposite coatings At present, significant effort is devoted to master the formation of nanocomposite coatings using magnetron sputtering because this technology can easily be scaled up for industrial use. The nanocomposite coatings are produced using so-called selective reactive magnetron sputtering [56 ]. In this deposition process, one element of the alloy is converted into nitride and the second element of the alloy is transported into the growing film unreacted. Recently, new hard (<40 GPa) supertough material of the type nc-TiC/a-C [57] with a remarkable plasticity (40% during nanoindentation deformation) and new superhard (≥40 GPa) materials of the type nc-MeN/a-Si N [12] with a high elastic recovery (up to 3 4 80%) in the form of nanocomposite coatings were developed; here nc- and a- denote the nanocrystalline and amorphous phases, respectively, and Me=Ti, W, V, Zr, etc. are transition metals. Even though both types of material are composed of nanocrystalline grains embedded in an amorphous matrix, they exhibit completely different physical properties. This is due to the different structures of the supertough and superhard nanocomposite coatings. A systematic investigation carried out on systems ZrCu–N [13] and TiAl–N [49] showed that superhard nanocomposite coatings can be composed not only of two hard phases as proposed by Veprˇek et al. [58] but also in the case when only one phase is hard, e.g. a nc-ZrN/Cu nanocomposite [13]. This means that there are two groups of superhard nanocomposite coatings: 1. nc-MeN/nitride (e.g. a-Si N , a-TiB , etc.); 3 4 2 2. nc-MeN/metal (e.g. Cu, Ni, Y, Ag, Co, etc.). The hardness H of films of both groups can be continuously varied from low values of about 10 GPa to very
326
J. Musil / Surface and Coatings Technology 125 (2000) 322–330
high values, achieving up to 50–70 GPa. Moreover, it is worthwhile to note that the hardness of nanocomposite films is strongly correlated with their structure. 7.1. Structure of superhard nanocomposite coatings Superhard films with H≥40 GPa are formed only when their structure is close to X-ray amorphous. This structure corresponds to a transition from the crystalline structure to an amorphous one. The main factors which govern the formation of films with an X-ray amorphous structure are: the energy delivered to the growing film, the substrate temperature T , the type of elements forms ing the film [their (i) mutual solubility or immiscibility, (ii) ability to form intermetallic compounds, (iii) chemical affinity, and (iv) binding energy], gap of elements immiscibility, enthalpy of the alloy formation and the content of individual phases in the nanocomposite film, e.g. the content of soft phase in a composite of the type nc-Me/metal. The structure of hard (<40 GPa) and superhard (≥40 GPa) nanocomposite coatings is significantly different. A systematic investigation of the correlation between the hardness and structure in the ZrCu–N and TiAl–N nanocomposite films showed that (a) the hard films are characterized by many reflections from polyoriented grains of both phases, (b) the superhard films are two-phase nanocomposites one phase of which has a nanocrystalline structure and the second is X-ray amorphous, and (c) the maximum hardness is achieved only when all grains are oriented in the same direction and the size of grains has an optimum value of approximately several tens of nanometres. 7.2. Classification of hard nanocomposite coatings Till now, various different types of hard nanocomposite coating have been prepared: 1. nc-MeN/a-nitride, e.g. nc-MeN/a-Si N (Me=Ti, W, 3 4 V ) [58,61], nc-TiN/a-Si N [62]; 3 4 2. nc-MeN/nc-nitride, e.g. nc-TiN/nc-BN [58]; 3. nc-MeC/a-C, e.g. nc-TiC/DLC [7]; 4. nc-MeN/ metal, e.g. nc-ZrN/Cu [13], nc-( Ti,Al )/AlN [49,60], nc-CrN/Cu [47]; 5. nc-MeN or MeC/a-boron compounds, e.g. nc-Ti(B,O)/quasi-a-(TiB , TiB and B O ) [63], Ti– 2 2 3 B–C [64]; 6. nc-WC+nc-WS /DLC [65]; 2 7. nc-MeC/a-C+a-nitride, e.g. nc-Mo C/a-C+a2 Mo N [59]. 2 This survey shows that all hard nanocomposite coatings contain one or two hard crystalline phases. The second phase is more complicated. It is either amorphous (e.g. a-Si N ) or crystalline (e.g. nc-BN [58]). Sometimes, the 3 4 content of the second phase in the nanocomposite coating is very low at approximately 1–2 wt.% (e.g. Cu
in nc-ZrN/Cu films [13]). In such a case, it is very difficult, without a HRTEM investigation, to determine if the second phase is crystalline or amorphous because the X-ray reflections from a small quantity of grains are below the detection limit. This means that the second phase can be incorrectly interpreted as amorphous when determined from X-ray diffraction analysis only. On the basis of these facts hard nanocomposite coatings can be divided into two main groups: (1) crystalline/amorphous nanocomposites and (2) crystalline/crystalline nanocomposites. At present, no definite methodology exists on how to select the combination of elements to produce films with nanocrystalline and/or X-ray amorphous structure. Such films can be formed from nitrides of alloys composed of elements which exhibit a wide miscibility gap and contain one element forming hard nitride, for instance nitrides of binary alloys formed of immiscible elements such as Cu–Cr, Zr–Y, etc. or transition metal nitride/boride and nitride/carbide systems [64]. A miscibility gap between intermetallics is also sufficient to produce hard nanocomposite coating, e.g. ZrCu–N film. There is also a possibility to produce the nanocomposite coating from nitrides of alloys whose elements form a solid solution, e.g. Ti Al N film. This is enabled by 1−x x the existence of a gap in the alloy composition x where the structure of the film is quasi-X-ray amorphous [66 ]. 7.3. Mechanical properties of hard nanocomposite coatings A short survey of basic mechanical characteristics of some recently prepared nanocomposite films is given in Table 3. For comparison, the characteristics of some selected hard bulk materials and hard amorphous carbon films are also given in this table. The hard nanocomposite coating is characterized not only by its hardness H but also by its Young’s modulus E and elastic recovery W . The determination of these e quantities for thin films is, however, difficult because they strongly vary with the load L used in their measurement. Despite these problems, Figs. 1 and 2 display H as a function E1=E/(1−n2) and the elastic recovery W as a function H, because these dependencies were e measured under the same conditions in one laboratory and on different nanocomposite systems. The high hardness of the material is only one parameter which ensures scratch and abrasion resistance. Protective overcoat films must be highly resistant also to plastic deformation during contact events. This requires a low Young’s modulus E since, according to Johnson analysis, the load P needed to initiate plastic deformation when a y rigid sphere of radius r is pressed into the coating is proportional to H3/E2 [74]. The ratio H3/E2 is a parameter which controls the resistance of materials to plastic deformation. The likelihood of plastic deformation is
327
J. Musil / Surface and Coatings Technology 125 (2000) 322–330 Table 3 Comparison of hard bulk materials, hard single layer films and selected hard and superhard nanocomposite coatings Material Bulk materials Diamond Boron Sapphire Amorphous films DLC a-C (cathodic arc) Nanocomposite single layer films nc-TiN/Si N 3 4 nc-TiN/BN nc-W N/a-Si N 2 3 4 Ti–B–C Ti–B–N Zr Cu N 98 2 W Ni N 86.7 8.3 5 W Si N 68 14 18 nc-Mo C/a-(C+Mo N ) 2 2 Ti Al N 45 55 Ti Al N 60 40 ZrY–N CrNi–N Ti Si N 75 25 Ti C ( TiC/a-C ) 0.32 0.68
H (GPa)
E1=E/(1−n2) (GPa)
100 35 30
1050 470 441
65 >59
550 >395
48 69 51 71 54 54 55 45 49 47 40 41 32 29 32
~565 585 560 486 ~500 394 510 – 440 409 650 319 253 256 370
W (%) e
H3/E12
d (nm)
Reference
0.91 0.19 0.14
[67] [68] [69]
80–90
0.91 ~1.3
[7] [74]
–a – – 80.5 – 81 – – 67 74 – 77 74 67 60
~0.34 0.96 0.42 1.52 0.63 1.03 0.64 – 0.61 0.62 0.15 0.66 0.50 0.36 0.239
4.5 9 3.5 ~1 ~1 35 – – 27 30 – – – – 10–50
[70] [58] [58,70,71] [64] [64] [13] [75] [61] [59] [49] [14] [45] [72] [46 ] [73]
a Denotes data not given in the references or not determined.
Fig. 1. H as a function of E/(1−n2) for Zr–Cu–N, Zr–Y–N, Cr–Ni–N, Ti–Si–N and Ti–Al–N nanocomposite films sputtered at different deposition conditions, i.e. T , U , i and p . s s s N2
reduced in materials with high hardness and low modulus [74]. In general, a low modulus is also desirable as it allows the given load to be distributed over a wider area. The ratio H3/E2 of nanocomposite films spreads over a very wide range from about 0.15 to 1.52, see Fig. 3 and Table 3. Data given in Table 3 and in Figs. 1– 3 show that the elastic recovery W and the resistance e
of materials to plastic deformation can be controlled by the film hardness H and its elastic modulus E. A spread of experimental points around the straight lines in Figs. 1–3 is connected with the variation in the film structure induced particularly by different (i) deposition conditions and (ii) chemical composition of the film. Therefore, H and E of nanocomposite coatings
328
J. Musil / Surface and Coatings Technology 125 (2000) 322–330
Fig. 2. Elastic recovery upon indentation W as a function of H for Zr–Cu–N, Zr–Y–N, Cr–Ni–N, Ti–Si–N and Ti–Al–N nanocomposite films e sputtered at different deposition conditions, i.e. T , U , i and p . s s s N2
Fig. 3. The ratio H3/E2 characterizing the resistance of the material to plastic deformation as a function of H for Zr–Cu–N, Zr–Y–N, Cr–Ni–N, Ti–Si–N and Ti–Al–N nanocomposite films sputtered at different deposition conditions, i.e. T , U , i and p . s s s N2
prepared by plasma CVD and magnetron-assisted laser deposition processes differ from those prepared by magnetron sputtering. A further systematic investigation is necessary to master the preparation of materials with prescribed properties.
8. Conclusions The main results obtained in the development of hard and superhard nanocomposite coatings can be summarized as follows. 1. Hard (<40 GPa) coatings are characterized by a high
plastic deformation increasing with decreasing H up to about 70% for H#10 GPa. 2. Superhard (≥40 GPa) coatings are characterized by a high elastic recovery increasing with increasing H up to about 85% for H#70 GPa. 3. There are two types of superhard nanocomposite film: (1) nc-MeN/a-nitride and (2) nc-MeN/metal. This means that the superhard film can be composed either of two hard phases or of one hard phase and the second soft phase. 4. The hardness of nanocomposite coatings correlates well with their structure. A decrease of the film crystallinity, characterized by a decrease of the inten-
J. Musil / Surface and Coatings Technology 125 (2000) 322–330
sity of X-ray reflection lines and by an increase of their broadening, results in an increase of hardness. 5. The structure of superhard films is close to X-ray amorphous. The nanocomposite films are fascinating materials since novel structures and new physical properties are expected. The development of nanocomposite films is, however, only at the beginning. Many problems are unsolved and many questions remain open. Therefore, further systematic research in this field is required. Further research is focusing on the following problems: (i) understanding of the origin of superhardness; (ii) correlation between the mechanical parameters of materials and process parameters; (iii) dramatic changes in crystallographic orientation of grains in alloy films, their hardness and elastic recovery induced by incorporation of nitrogen; (iv) formation of nanocomposite films with controlled hardness, elastic modulus and elastic recovery and new functional properties; and (v) investigation of materials having very small grains of about 1 nm in size.
Acknowledgements The author would like to thank his Ph.D. students H. Hruby´, I. Leipner, H. Pola´kova´, F. Regent, Z. Soukup and P. Zeman for the preparation of many nanocomposite coatings of different chemical composiˇ erstvy´ and Dr. M. Kolega for X-ray tion and Dr. R. C characterization of these coatings. This work was supported in part by the Grant Agency of the Czech Republic under Project No. 106/96/K245 and by the Ministry of Education of the Czech Republic under Project Nos. VS 96/059 and ME 173/1999.
References [1] W.D. Sproul, Science 273 (1996) 889. [2] H. Holleck, Ch. Ku¨hl, H. Schultz, J. Vac. Sci. Technol. A 3 (6) (1985) 234. [3] H. Holleck, J. Vac. Sci. Technol. A 4 (6) (1986) 2661. [4] C.A. Brookes, in: J.E. Field (Ed.), The Properties of Diamond, Academic Press, New York, 1979, p. 383. [5] B.M. Kramer, P.K. Judd, J. Vac. Sci. Technol. A 3 (1985) 2439. [6 ] W.D. Mu¨nz, J. Vac. Sci. Technol. A 4 (6) (1986) 2717. [7] A.A. Voevodin, J.S. Zabinski, Diamond Relat. Mater. 7 (1998) 463. [8] H. Sjo¨stro¨m, S. Stafstro¨m, M. Boman, J.-E. Sundgren, Phys. Rev. Lett. 76 (1996) 56. [9] H. Sjo¨stro¨m, L. Hultman, J.-E. Sundgren, S.V. Hainsworth, T.F. Page, G.S.A.M. Teunissen, J. Vac. Sci. Technol. A 14 (1996) 56. [10] U. Helmersson, S. Todorova, S.A. Barnett, J.-E. Sundgren, L.C. Markert, J.E. Greene, J. Appl. Phys. 62 (1987) 481. [11] X. Chu, S.A. Barnett, M.S. Wong, W.D. Sproul, Surf. Coat. Technol. 57 (1993) 13. [12] S. Veprˇek, S. Reiprich, S. Li, Appl. Phys. Lett. 66 (20) (1995) 2640. [13] J. Musil, P. Zeman, H. Hruby´, P. Mayrhofer, ZrN/Cu nanocom-
329
posite film — novel superhard material, Proc. ICMCTF’99, April 12–18, San Diego, CA (1999), paper no. B1-2-9. [14] Y. Min, Y. Makino, M. Nose, K. Nogi, Thin Solid Films 228 (1999) 204. [15] U. Helmersson, S. Todorova, S.A. Barnett, J.-E. Sundgren, L.C. Markert, J.E. Greene, J. Appl. Phys. 62 (1987) 481. [16 ] P.B. Mirakami, L. Hultman, S.A. Barnett, Appl. Phys. Lett. 57 (25) (1990) 2654. [17] M. Shinn, L. Hultman, S.A. Barnett, J. Mater. Res. 7 (4) (1992) 901. [18] X. Chu, M.S. Wong, W.D. Sproul, S.L. Rohde, S.A. Barnett, J. Vac. Sci. Technol. A 10 (4) (1992) 1604. [19] X. Chu, S.A. Barnett, M.S. Wong, W.D. Sproul, Surf. Coat. Technol. 57 (1993) 13. [20] D. Li, X.W. Lin, S.C. Chen, V.P. Dravid, Y.W. Chung, M.S. Wong, W.D. Sproul, Appl. Phys. Lett. 68 (9) (1996) 1211. [21] M.L. Wu, W.D. Qian, Y.W. Chung, Y.Y. Wang, M.S. Wong, W.D. Sproul, Thin Solid Films 308/309 (1997) 113. [22] H. Jensen, J. Sobota, G. Sorensen, Surf. Coat. Technol. 94/95 (1997) 174. [23] H. Jensen, J. Sobota, G. Sorensen, J. Vac. Sci. Technol. A 16 (3) (1998) 1180. [24] J.E. Kryanowski, Mater. Res. Soc. Symp. Proc. 239 (1992) 509. [25] S.A. Barnett, in: M. Fracombe, J.A. Vossen ( Eds.), Physics of Thin Films, Academic Press, New York, 1993, p. 1. [26 ] M. Shinn, L. Hultman, S.A. Barnett, J. Mater. Res. 7 (1992) 901. [27] X. Chu, S.A. Barnett, J. Appl. Phys. 77 (1995) 4403. [28] B.G. Wendler, P. Kula, K. Jakubowski, S. Fauvry, Ph. Kapsa, L. Vincent, D. Heper, Proc. 36th Tagung der Deutschen Gesellschaft fu¨r Obrfla¨chen- und Galvanotechnik, October 7–9, Schwabisch Gmu¨nd, Germany (1998) 6–11. [29] B.G. Wendler, M. Molinaro, Proc. 10th Int. Summer School on Modern Plasma Surface Technology, May 10–14, Mielno, Poland (1998) 423–436. [30] K.K. Shih, D.B. Dove, Appl. Phys. Lett. 61 (6) (1992) 654. [31] X. Chu, M.S. Wong, W.D. Sproul, A.S. Barnett, Mechanical properties and microstructures of polycrystalline metal/ceramic superlattices, TiN/Ni and TiN/Ni Cr , paper presented at 0.9 0.1 ICMCTF-20, San Diego, CA, 1993. [32] A. Madan, X. Chu, S.A. Barnett, Appl. Phys. Lett. 68 (16) (1996) 2198. [33] A. Madan, Y.-Y. Wang, S.A. Barnett, C. Engstro¨m, H. Ljungcrantz, L. Hultman, M. Grimsditch, J. Appl. Phys. 84 (2) (1998) 776. [34] J.M. Schneider, W.D. Sproul, A.A. Voevodin, A. Matthews, J. Vac. Sci. Technol. A 15 (3) (1997) 1084. [35] J.M. Schneider, W.D. Sproul, Reactive pulsed dc magnetron sputtering and control, Handbook of Thin Film Process Technology, IOP, Bristol, 1998, pp. A5.1:1–A5.1:12. [36 ] S. Veprˇek, S. Reiprich, Thin Solid Films 268 (1995) 64. [37] R.W. Siegel, Nanophase materials: Structure, defects and properties, Proc. ATC Int. Symp. (ACTA), October 29–30, Tokyo, Japan (1996) 1. [38] P.Yu. Butyagin, Colloid J. 59 (1997) 112. [39] J. Musil, J. Vlcˇek, Czech. J. Phys. 48 (10) (1998) 1209. [40] J. Musil, J. Vlcˇek, Magnetron sputtering of alloy and alloy-based films, Thin Solid Films 343–344 (1999) 47. [41] K.D. Leedy, J.M. Rigsbee, J. Vac. Sci. Technol. A 14 (1996) 2202. [42] J. Musil, Vacuum 50 (3/4) (1998) 363. [43] J. Musil, A.J. Bell, J. Vlcˇek, T. Hurkmans, J. Vac. Sci. Technol. A 14 (4) (1996) 2247. [44] J. Musil, F. Regent, J. Vac. Sci. Technol. A 16 (6) (1998) 3301. [45] J. Musil, H. Pola´kova´, Hard nanocomposite Zr–Y–N coatings. Correlation between hardness and structure, 2nd Asian–European Int. Conf. on Plasma Surface Engineering, September 15–19, Beijing, China (1999), Paper No. Sat-0A3-2.
330
J. Musil / Surface and Coatings Technology 125 (2000) 322–330
[46 ] J. Musil, H. Pola´kova´, V. Cibulka, Czech. J. Phys. 49 (3) (1999) 359. [47] J. Musil, I. Leipner, M. Kolega, Surf. Coat. Technol. 115 (1) (1999) 32. [48] J. Musil, Z. Soukup, unpublished results. [49] J. Musil, H. Hruby´, Superhard nanocomposite Ti Al N films 1−x x prepared by magnetron sputtering, Proc. 14th Int. Symp. on Plasma Chemistry, August 2–6, Praha, Czech Republic, Vol. III (1999) 1605–1610. [50] J. Musil, P. Zeman, Vacuum 52 (1999) 269. [51] W.L. Johnson, Progr. Mater. Sci. 30 (1986) 81. [52] G. Weigang, H. Hecht, G. von Minnigerode, Z. Phys. B 96 (1995) 349. [53] J.R. Ding, D.Y. Che, H.B. Yhang, K. Tao, B.X. Liu, Appl. Phys. Lett. 60 (8) (1992) 994. [54] M. Misˇina, J. Musil, S. Kadlec, Surf. Coat. Technol. 110 (1998) 168. [55] Z. Liu, R. Singh, K. Poole, R.J. Diefendorf, J. Harris, K. Cannon, J. Vac. Sci. Technol. B 15 (1997) 1990. [56 ] L. May, W.R. Allen, A.L. Glower, J.C. Mabon, J. Vac. Sci. Technol. B 13 (1995) 361. [57] A.A. Voevodin, S.V. Prasad, J.S. Zabinski, J. Appl. Phys. Lett. 82 (2) (1997) 855. [58] S. Veprˇek, P. Nesla´dek, A. Niederhofer, F. Glatz, M. Jı´lek, M. Sˇ´ıma, Surf. Coat. Technol. 108/109 (1998) 138. [59] M. Benda, J. Musil, Vacuum 55 (1999) 171. [60] Y. Min, Y. Makino, M. Nose, K. Nogi, Thin Solid Films 339 (1/2) (1999) 238. [61] C. Louro, A. Cavaliero, Hardness versus structure in W–Si–N sputtered coatings, 6th Int. Conf. on Plasma Surface Engineering, PSE-98, September 14–18, Garmisch-Partenkirchen, Germany (1998), paper no. ThWB5.
[62] M. Diserens, J. Patscheider, F. Levy, Surf. Coat. Technol. 108/ 109 (1998) 241. [63] C. Mitterer, P. Losbichler, F. Hofer, P. Warbichler, W. Gissler, Vacuum 50 (3/4) (1998) 313. [64] C. Mitterer, P.H. Mayrhofer, M. Beschliesser, P. Losbichler, P. Warbichler, F. Hofer, P.N. Gibson, W. Gissler, H. Hruby´, J. Musil, J. Vlcˇek, Proc. ICMCTF’99, April 12–16, San Diego, CA (1999), paper no. BP-16. [65] A.A. Voevodin, J.P. O’Neill, J.S. Zabinski, Surf. Coat. Technol. (1999) submitted. [66 ] U. Wahlstro¨m, L. Hultman, J.-E. Sundgren, F. Abidi, I. Petrov, J.E. Greene, Thin Solid Films 235 (1993) 62. [67] J.E. Field, The Properties of Diamond, Academic Press, London, 1979. [68] S.M. Gorbatkin, R.L. Rhoades, T.Y. Tsui, W.C. Oliver, Appl. Phys. Lett. 65 (1994) 2672. [69] W.C. Oliver, G.M. Pharr, J. Mater. Res. 7 (1991) 1564. [70] S. Veprˇek, Thin Solid Films 317 (1998) 449. [71] S. Veprˇek, Thin Solid Films 297 (1997) 145. [72] J. Musil, F. Regent, Structure and microhardness of magnetron sputtered nanocrystalline CrNi–N films, Proc. 14th Int. Symp. on Plasma Chemistry, August 2–6, Praha, Czech Republic, Vol. III (1999) 1617–1622. [73] A.A. Voevodin, J.S. Zabinski, J. Mater. Sci. 33 (1998) 319. [74] T.Y. Tsui, G.M. Pharr, W.C. Oliver, C.S. Bhatia, R.L. White, S. Anders, A. Anders, I.G. Brown, Mater. Res. Soc. Symp. Proc. 383 (1995) 447. [75] A. Cavaleiro, B. Trindale, M.T. Viera, Deposition and characterization of fine-grained W–Ni–C/N ternary films, 6th Int. Conf. on Plasma Surface Engineering, PSE-98, Garmisch-Partenkirchen, Germany, Book of Abstracts (1998) 265.