Microstructure and mechanical behavior of arc-evaporated Cr–N coatings

Surface and Coatings Technology 114 (1999) 39–51

Microstructure and mechanical behavior of arc-evaporated Cr–N coatings Magnus Ode´n a, *, Claes Ericsson a,1, Greger Ha˚kansson b, Henrik Ljungcrantz b,2 a Division of Engineering Materials, Department of Mechanical Engineering, Linko¨ping University, 581 83 Linko¨ping, Sweden b Tixon Brukens Sverige AB, Svedengatan 2, 582 73 Linko¨ping, Sweden Received 25 November 1998; accepted 12 January 1999

Abstract Cr–N coatings were grown by arc evaporation onto high speed steel substrates. The coatings were each grown using a different negative substrate bias voltage, V , between 20 and 400 V. X-ray diffraction showed a microstructure containing primarily the S NaCl structured CrN phase along with the BCC-Cr and HCP-Cr N. Auger electron spectroscopy and transmission electron 2 microscopy indicate a substoichiometric (N/Cr=0.85±0.08) composition and dense columnar microstructure, respectively. At V =400 V, equiaxed grains and microcracks parallel the substrate–coating interface were observed. The fiber textured coatings S are in a compressive residual stress state that increases from 2.9 (V =20 V ) to 8.8 GPa (V =100 V ). At higher bias voltages, a S S decrease of the compressive residual stress is seen, which is discussed in terms of lattice defect annihilation in the collision cascade and lattice defect diffusion during deposition. Nanoindentation showed a maximum hardness at V =100 V of 29 GPa. The critical S loads for cohesive failure in a scratch test decreased monotonically with increasing negative substrate bias. The scratch results suggest a transition in deformation mechanism from plastic deformation to cracking, which occurs at lower applied loads when V is increased. Similar behavior was also seen in a crater grinding wear test where a shift in wear mechanism from plastic S deformation to chipping occurred at V =200 V. The influence of the microstructure on the deformation transition is discussed in S terms of lattice defect density and the presence of equiaxed grains. © 1999 Elsevier Science S.A. All rights reserved. Keywords: Arc evaporation; Cr–N coatings; Mechanical behaviour; Microstructure

1. Introduction The use of hard coatings as wear protection in various engineering applications is widespread today. CrN can, for example, be found as a protective coating on tools in applications such as metal sheet forming. Many engineering applications where a coating would be useful employ substrate material with a low tempering temperature such that it cannot withstand high deposition temperatures. Consequently, there is a demand for CrN coatings with good mechanical properties that can be deposited at low temperatures. Optimization of such coatings includes understanding and control of the complete chain from choice of deposition parameters to obtained microstructure, resulting mechanical and tribo* Corresponding author. Fax: +46-13-282505. E-mail address: [email protected] (M. Ode´n) 1 Present address: Avesta-Sheffield AB, 774 80 Avesta, Sweden. 2 Present address: Impact Coatings AB, Box 1533, 581 15 Linko¨ping, Sweden.

logical properties, and final function for a specific application. The complete development chain for CrN is almost never reported in the literature, while the number of papers concerned with only a single step of the chain is rapidly growing. However, comparisons between studies are often difficult to make since mechanical property investigations lack information on microstructure and vice versa. Depending on the choice of deposition parameters, arc-evaporated Cr–N coatings can have a wide range of different microstructures and phase compositions (i.e. Cr, Cr N and CrN ). Normally, a columnar growth is 2 seen with reported grain sizes varying from a few to several hundreds of nm [1–4]. The preferred orientation and, consequently, the elastic behavior is also influenced by the deposition parameters. A 220 preferred orientation is often reported for arc-evaporated CrN [1,2,5–8] but 200 and 111 have also been observed [1]. Due to lack of single crystal data of the elastic stiffness of CrN, the influence of this texture on elastic anisotropy remains unknown.

0257-8972/99/$ – see front matter © 1999 Elsevier Science S.A. All rights reserved. PII: S0 2 5 7- 8 9 7 2 ( 9 9 ) 0 0 01 9 - 5

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In addition to the influence from the microstructure on the mechanical performance of the coating, the residual stress state must be controlled [9]. The success in fabrication and performance of coatings is often related to the residual stress state. A compressive residual stress state can be beneficial during service due to increased resistance to crack initiation and growth. However, a too high compressive residual stress may cause cohesive failure during fabrication or when mechanically loaded, and thereby limit the practical achievable thickness of the coating. A normal commercial CrN coating is today about 5 mm thick. A thicker coating would increase the mechanical properties if the residual stress state and surface roughness could be controlled. In this paper, we report on the influence of microstructural and residual stress changes on mechanical properties of CrN coatings deposited with increasing energy of the energetic particles incident on the growing coating. The as-deposited microstructures were studied by X-ray diffraction ( XRD) and cross-sectional transmission electron microscopy ( TEM ). Chemical composition was studied by Auger electron spectroscopy (AES), and preferred orientations and residual stress states were determined by XRD. The mechanical performance was assessed by indentation, scratch, and cratergrinding tests [10].

2. Experimental details 2.1. Materials and deposition conditions A powder metallurgical high speed steel [ASP 23, Erasteel kloster AB, Sweden; chemical composition: (wt%) 1.28 C, 4.2 Cr, 5.0 Mo, 6.4 W, and 3.1 V ] hardened to 62-64 HRC was used as substrate material (20×20×10 mm3). The substrate surfaces were ground, with the final step performed with a 500 grit SiC paper, and then degreased with alkaline and deionized water in an ultrasonic cleaning line. All samples were deposited in a commercially available 20 in Multi-Arc system containing three 50 mm diameter Cr sources. A single rotating fixturing was used, designed such that the samples continually faced the sources. The shortest cathode to substrate distance was 200 mm. The system was pumped down below 1×10−3 Pa, after which the substrates were sputter cleaned. This step was performed by applying a 45 A current to each of the three Cr sources and a 1000 V bias voltage to the substrates. The etching and cleaning step was interrupted when a substrate temperature of 300°C was reached, as measured with a calibrated pyrometer. After cleaning, the CrN coatings were deposited, using a nitrogen gas pressure of 8 Pa and a cathode

Table 1 Deposition parameters: negative substrate bias voltage; nitrogen partial pressure; cathode current; deposition time; final deposition temperature V (V ) S

P (Pa) N2

I

20 50 100 200 400

8 8 8 8 8

45 45 45 45 45

Cathode

(A)

t (min)

T (°C ) Final

220 220 220 220 220

200 200 270 310 420

current of 45 A for all the deposition runs. Five different coatings were produced by varying the negative substrate bias voltage, V =20, 50, 100, 200, and 400 V. The final S temperatures, measured with a pyrometer, together with other deposition data, are presented in Table 1. 2.2. Microscopy and spectroscopy N:Cr stoichiometry for all coatings was determined with a VG Scientific Micro Lab 310F Auger electron spectrometer. The relative intensities of the N KL1 and Cr KM2 peaks, obtained after removing surface contamination by Ar+ ion etching, were used. The measured relative intensities were then compared with values obtained from a stoichiometric single-crystal sample (N/Cr=0.99±0.02, as measured with Rutherford backscattering spectroscopy). The surface roughness was measured using a surface profilometer (Mitutoyo Surftest Analyser), and the coating thickness was determined with the ball cratering technique. A Jeol 6400 scanning electron microscope (SEM ) operated at 10–20 kV was used to study the as-coated as well as worn surface appearances. The growth defect number densities reported in Table 2 have been point counted on the surface and only growth defects larger than 0.5 mm were recorded. The microstructure of the as-deposited coatings was studied using a Philips EM400T transmission electron microscope ( TEM ) operated at 120 kV. Cross-sectional sample preparation for TEM consisted of mechanical grinding and polishing followed by ion-beam sputter etching [11]. 2.3. X-ray diffractometry The X-ray diffractometry experiments were all conducted on a XRD 3000 PTS diffractometer ( Rich. Seifert) using Cu Ka radiation. Phase identification and peak widths were determined from diffraction patterns obtained in conventional Bragg–Bretano geometry with a bent graphite secondary monochromator and scintillation detector. Three background corrected pole figures ({111}, {200}, and {220}) were measured using the Schultz reflection method. With this data, the orientation distri-

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Table 2 Coating thickness, roughness, growth defect number density, and chemical composition. The error indicates one standard deviation V (V) S

Thickness (mm)

Coating roughness, R (mm) a

Coating roughness, R (mm) z

Growth defect number density×105 (cm−2)

20 50 100 200 400

9.6±1.4 10.4±0.5 8.6±0.5 9.2±0.5 9.6±0.6

1.06±0.07 0.78±0.23 0.94±0.14 0.42±0.18 0.10±0.05

8.0±0.4 5.4±1.3 6.4±0.7 3.0±0.8 1.0±0.5

70 20 9 10 6

Substrate

–

0.05

0.4

–

bution functions (ODF ) were calculated according to Bunge’s series expansion method [12], which was truncated at the order l=22. The residual stress measurements were conducted in V configuration with line focus and a bent graphite secondary monochromator. The CrN 331 peak position (at ~112° 2h) was determined at a minimum of seven y angles ranging from ±45° by fitting a pseudo-Voigt function to the diffraction data. The stress was then evaluated using the sin2 y method [13]. The d versus sin2 y plots show weak tendencies to oscillatory behavior which could indicate influence of texture and a stress gradient. The elastic modulus of CrN reported in the literature varies between 220 and 400 GPa [3,14], which makes any analysis that involves the elastic properties of CrN uncertain. We have chosen to use isotropic values of E =280 GPa and n =0.2, which are consistent c c with the choice made by Ehrlich et al. [2] and Grant et al. [15], when evaluating the residual stress in similar coatings in order to facilitate comparisons.

L and first exposure of the substrate material C1 L .L and L were determined by viewing the C2 C1 C2 scratches in an optical microscope and SEM. The inherent coating abrasive wear resistance was evaluated using the crater grinding test [10]. A commercial dimple grinder ( E.A. Fischione Instruments 2000) was used to grind spherical cap-shaped wear scars into the coated surface. The grinding was interrupted at regular intervals in order to measure the crater diameter and calculate the worn volume. The obtained wear constant (k ) of the coating is a measure of its wear C rate, that is, a high wear rate corresponds to a high wear constant. A grinding procedure with an abrasive medium of 1 mm diamond slurry and a grinding load of 20 g was used. The wear scars were studied by SEM.

2.4. Mechanical and wear testing

3.1. Coating chemistry and morphology

The film microhardness was measured with a Leco M-400 hardness tester on polished and tapered crosssections using a load of 100 gf. In order to minimize the influence of the substrate, indents were made across a mechanically polished (1 mm diamond paste) tapered section, starting at the interface and moving across the coating until a plateau in hardness was reached. The nanoindentation response was recorded with a Nano Indentor@ II, using a maximum load of 10 mN, and the hardness was calculated from the first 30% of the unloading curve [16 ]. The shape of the Bercovich diamond indentor tip was calibrated for contact depth versus projected area using fused silica. The calibration was checked after completing the measurements and no change was observed for all investigated samples. A commercial scratch tester (CSEM Revetest@), equipped with a Rockwell-C diamond stylus, was used to assess the critical normal load of the coating–substrate composites. Scratches were made using linearly increasing load ranging from 40 to 140 N at a load rate of 10 N/min and a horizontal stylus speed of 10 mm/min. The critical loads were defined as first cohesive failure

The Cr–N coatings were found to be substoichiometric, with an average N/Cr ratio of 0.85±0.08. The coating thickness, roughness, and surface growth defect (>0.5 mm) number density are presented in Table 2. The measured coating thicknesses are all in the range 8 to 11 mm. Any influence from resputtering on the deposition rate as the bias voltage was increased could not be established nor excluded due to the limited resolution of the measurement technique and variation in deposition conditions associated with using an industrial system. The measured coating surface roughness values (R and R ) are 8 to 20 times higher compared with the a z uncoated polished substrates, and decrease monotonically with increasing substrate bias voltages. A similar trend is observed for the growth defect number density which decreases from 70×105 cm−2 at V =20 V to S between 6 and 20×105 cm−2 at larger bias voltages. These trends can be visually seen in Fig. 1, where SEM micrographs of the as-deposited surfaces are shown. In addition to the decrease in growth defect number density, a decrease in size of growth defects with increasing bias voltage can also be observed.

3. Results

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Fig. 1. Scanning electron micrographs of the surface appearance of arc-evaporated CrN coatings.

3.2. Microstructure The XRD patterns of the coatings, Fig. 2, show the presence of cubic CrN and some minor peaks, which indicate the presence of BCC-Cr. The diffraction line positions given by Ref. [17] have also been included as lines in Fig. 2. The Cr phase is attributed to incorporated droplets, which have previously been found to have a core consisting of pure metal phase [18]. The hexagonal Cr N phase is only detected in the coating with 2 V =400 V. However, the presence of Cr N cannot be S 2 excluded in the other samples since Cr N peaks may be 2 hidden in the broad CrN peaks. The relative intensities change substantially between samples, in particular, the intensity of the 220 peak increases with increasing bias. The peak full-width at half-maximum, FWHM, of the 111, 200, and 220 peaks was determined by fitting each peak to a pseudo-Voigt function. The peak broadening was affected by the substrate bias, and Fig. 3

shows FWHM, without contributions from Ka , as a 2 function of V . The FWHM increases with increasing S V from 20 to 200 V. This increase in peak width suggests S a decrease in diffracting crystallite size and/or an increase in inhomogeneous strain. At V =400 V, a substantial S decrease in peak width is observed. The textures of the coatings, as determined by XRD, are presented in Fig. 4 as inverse pole figures. All coatings exhibit fiber textures and the degree of the texture increases with increasing bias voltage except for the sample with V =400 V. For V =20 V, a maximum S S texture intensity of 1.8 multiplier of a random distribution (mrd ) was observed while the sample with V =200 V showed a maximum texture intensity of S 5 mrd. At V =100 and 200 V the 110 preferred direcS tion is dominant. However, at lower bias voltages (i.e. V =20 and 50 V ) the texture is weaker and several S different components are present. It can also be noted that changing V from 50 to 200 V results in a rotation S

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Fig. 2. X-ray diffraction patterns of the investigated arc-evaporated CrN coatings. The lines indicate diffraction line positions given by Ref. [17].

Fig. 3. Full-width at half-maximum, FWHM, of the 111, 200, and 220 diffraction peaks as a function of negative substrate bias voltage in the investigated arc-evaporated CrN coatings. The error bars indicate one standard deviation.

of the texture around the 100 zone axis towards a 110 preferred orientation. The samples with V =20 and S 400 V show textures diverting from this trend with

211 and 122 preferred orientations, respectively, as the strongest component. Cross-sectional TEM was used to investigate the microstructure of coatings deposited with V =20, 50 S and 400 V. All coatings exhibited a columnar microstructure with dense grain boundaries and a high defect density. The defect density was too high to resolve individual dislocations or defect clusters. Fig. 5(a) shows a micrograph of the V =20 V sample with a typical S

columnar structure. The interface between substrate and coating is distinct: no voided or structural divergence is observed. A nucleation and growth region is visible (~150 nm thick) above the substrate coating interface with narrower columns (5–15 nm) compared with columns further away from the interface. Above the nucleation and growth region the column width increases to 50–100 nm after approximately 400 nm of growth and then remains constant, see Fig. 5(a). Most columns extend through the total thickness of the coating starting from the nucleation and growth region. Fig. 5(b) shows the microstructure of the sample with V =50 V in the S

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(c) (a)

(b)

(d)

Fig. 4. Inverse pole figures of arc-evaporated CrN coatings for (a) V =20 V, (b) V =50 V, (c) V =100 V, (d ) V =200 V, (e) V =400 V, and (f ) S S S S S indicates where preferred orientations would appear.

center of the coating. This coating shows a microstructure similar to the 20 V sample except for a narrower column width (25–50 nm) above the nucleation and growth region next to the substrate coating interface. All samples show incorporation of droplets with diameters in the range 0.1–0.3 mm [see for example Fig. 5(c)]. On top of the droplets, CrN grows in a columnar manner with a column width close to the size of the droplet diameter, hence much larger than the overall coating microstructure. Such columns extend all the way to the coating top surface and contribute to the surface roughness of the coating. The voids beneath the droplets were toroidal shaped, and the projection in Fig. 5(c) shows two bright contrast ‘wings’ below the droplet. This type of void shape has previously been reported for arc-evaporated TiN coatings [18] and can be explained by shadowing effects.

The 400 V sample had a larger columnar width (100– 200 nm). It also contains regions with more equiaxed grains in the columnar microstructure. An example of such region is shown in Fig. 5(d ). The defect density in the equiaxed grains was also too high for individual defects such as dislocations or defect clusters to be resolved. Fig. 5(e) shows cracks perpendicular to the growth direction in the 400 V sample. The cracks are always found near a droplet and 1–3 mm below the surface. The void beneath the droplet may have acted as the crack initiation site. These types of crack were not found in the other samples. The biaxial residual stress state in the CrN coatings is presented as a function of substrate bias voltage in Fig. 6. All coatings are in a state of compressive residual stress. The compressive residual stress state increases with increasing substrate bias from about 2.9 GPa at

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3.3. Mechanical and wear properties Table 3 summarizes the hardness, scratch, and weartest data obtained for the different coatings. Both the microhardness and nanoindentation results indicate an increased hardness with bias voltage to a maximum at V =100 V of about 2500 HV and 29 GPa, respectively. S Coatings deposited with higher bias voltages show a decreased hardness, which was more pronounced for the microhardness than the nanoindentation measurements. The scratch test results are schematically shown in Fig. 7. The drawings are deduced from SEM micrographs taken along the whole scratch distance. All coatings exhibit a similar scratch response starting with side cracks at low loads and the formation of internal transverse cracking at approximately the critical load for cohesive fracture, L . Then follows a region of C1 increasing plastic deformation and internal cohesive fractures until the critical load for substrate exposure, L , is reached. As can be seen, both L and L C2 C1 C2 decrease with increasing substrate bias voltage. Despite these decreases, the relative difference between L and C1 L remains almost constant at approximately 50 N. C2 The SEM investigations of the scratches also show that a stick–slip behavior is visible at applied loads close to L for all samples. When comparing the zone above C2 L between the coatings, a clear difference in the C2 amount of cohesive flaking at the rim of the scratch track is visible. The coatings produced with lower bias voltages are less affected in this respect than the samples deposited at higher bias values where relatively large areas surrounding the rim have been affected. The wear constant k of the substrate material was C about five times higher than k values measured for the C CrN coatings. The wear constants are seen to be similar for V =20 to 100 V, while the higher biased coating, S V =200 V, showed an increase in wear constant comS pared with the other three samples. Micrographs of wear scars from samples prepared with V =100 and S 200 V are shown in Fig. 8(a) and (b), respectively. The

Fig. 4 (continued)

V =20 V to a maximum of 8.8 GPa at V =100 V. At S S the highest bias voltage levels, V =200 and 400 V, the S trend is reversed and a decrease in compressive residual stress is observed.

Table 3 Result from hardness, scratch, and wear tests of arc-evaporated CrN coatings. The error indicates one standard deviation V (V) S

20 50 100 200 400 Substrate

Hardness

Scratch test: critical load (N )

Vickers, 100 g load (HV )

Berkovich, 10 mN load (GPa)

First cohesive failure, L C1

First exposure of substrate, L C2

2101±46 2213±28 2474±97 2422±36 2027±57 925±31

17.5±1.4 24.7±0.4 29.0±0.7 27.3±0.6 27.6±1.9 –

88 77 62 58 56 –

129 127 120 110 107 –

Coating abrasive wear constant (mm3/mmN )

24.5±2.2 21.6±2.3 21.8±1.2 31.1±3.8 – 111±8.1

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M. Ode´n et al. / Surface and Coatings Technology 114 (1999) 39–51

Fig. 5. Cross-sectional TEM micrographs of arc-evaporated CrN coatings. The growth direction is from bottom to top in all micrographs. (a) Micrograph showing coating with V =20 V including the steel substrate (substrate coating interface is marked by an arrow). (b) Micrograph of a S sample fabricated with V =50 V, 4 mm above the substrate. (c) An almost circular shaped droplet incorporated in the V =50 V coating is marked S S by ‘A’, marker ‘B’ points out the narrow column width below the droplet, and ‘C’ the coarser column grown on top of the droplet. A ‘wing-shaped’ void beneath the droplet has bright contrast. (d) Micrograph of coating with V =400 V showing an equiaxed grains region, and (e) shows a crack S (marked with arrows) located approximately 6 mm above the substrate–coating interface. The crack is parallel to the substrate–coating interface and located near a droplet.

worn surface appearance is different for the two coatings, indicating different wear mechanisms. The V =200 V S coating in Fig. 8(b) shows areas where cohesive failures have occurred. This is not observed in the V =100 V S coating, where the wear scar is very smooth, indicating a pure abrasive wear situation.

4. Discussion The complete chain from choice of deposition parameters, to generated microstructure, resulting mechanical properties, and final function of a coating is seldom completely reported in the literature, mainly due to the

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Fig. 6. Residual stress in arc-evaporated CrN coatings as a function of negative substrate bias voltage. The error bars indicate one standard deviation.

complexity and breadth of such studies. However, the importance of understanding all steps in the chain is obvious when tailoring a coating for a specific application. This paper contains microstructural and mechanical behavior results from CrN coatings grown on high speed steel substrates with arc evaporation at various negative substrate bias voltages (V =20 to 400 V ). The S results are discussed below, based on the above mentioned chain starting with the influence of increasing V on the coating microstructure, preferred orientation, S morphology, and residual stress state. The influence of these microstructural differences on mechanical and abrasive wear behavior is then addressed. 4.1. Substrate bias vs. microstructure Although the overall grain structure in all coatings was columnar, appreciable effects of bias voltage in the coating microstructures was found. The column width decreases significantly as V is increased from 20 to S 50 V, which is in agreement with previously reported results [19]. This result is attributed to an increased renucleation probability as the ion energy, or equivalently the bias voltage, is increased [20]. However, in the coating grown at V =400 V and a deposition temperS ature of 420°C a larger column width is seen along with regions with equiaxed grains. A similar increase in grain size has been reported for arc-evaporated CrN coatings when the substrate temperature was increased from 350 to 500°C [19]. The increase in column width can then be interpreted as a result of higher adatom mobility induced by higher ion energy and increased temperature [21]. Perry and Chollet [22] have reported on defect diffusion in TiN during tempering experiments for tem-

peratures larger than 400°C. The chemical stability of CrN is lower compared with TiN due to a larger amount of filled antibonding electron states [23]. It is therefore reasonable to assume that lattice defect diffusion occurs in CrN during deposition at a deposition temperature of 420°C. The substoichiometry also indicates a driving force for the formation of Cr N. We note that nucleation 2 and growth of Cr N grains, as an effect of diffusion 2 during deposition, can generate regions with equiaxed grains. The lattice defect density in all coatings is too high for individual dislocations or defect clusters to be resolved in TEM, consistent with observations of arcevaporated TiN and Ti(CN ) coatings [4,9]. The high lattice defect density will cause significant inhomogeneous strain, which contributes to the observed XRD line broadening (see Fig. 3) [24]. XRD line broadening due to small diffracting domain size is not expected in a columnar microstructure; therefore, changes in line broadening primarily reflect changes in lattice defect density. The defect density initially increases with bias voltage, which is due to increased numbers of implanted atoms at higher ion energies. The subsequent decrease at V =400 V is attributed to an increase in defect S annihilation rate with V [25]. S The preferred crystallographic orientation of the columnar microstructure was strongly influenced by V , S see Fig. 4. Common for most studies of texture in CrN is that the preferred orientation is determined from the relative peak intensities in an XRD (h–2h) pattern. When instead pole figures, measured for several diffraction lines, are used the complete orientation distribution function (ODF ) can be obtained. The ODF contains additional information that cannot be directly observed

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Fig. 7. Schematic presentation of scratch test results deduced from optical and scanning electron microscopy of arc-evaporated CrN coatings.

by using relative peak intensities in an XRD (h–2h) pattern. When viewing the obtained ODFs as inverse pole figures, both an increase of the intensity of the texture and a rotation of the fiber texture around the 100 zone axis towards 110 orientation as the bias voltage increases is seen. The relationships between the mechanisms governing texture evolution and deposition conditions for CrN are presently unknown. However, some comments can be made. A 220 preferred orientation has been reported several times for arc-evaporated CrN coatings when the nitrogen partial pressure is high enough to form single phase CrN [1,2,5–8]. Gautier and Machet [1] found 200 and 111 preferred orientations at low substrate bias voltage (V <70 V ) that S shifted to 220 when the bias was increased (70
4.2. Substrate bias vs. residual stress The residual stress state in the coatings was found to rapidly increase in compression from ~2.8 GPa at −20 V to ~8.8 GPa at −100 V, see Fig. 6. Further increase of the substrate bias voltage resulted in a decrease of the compressive residual stresses. Since the measured residual stress can be generated by both thermal and intrinsic sources, it is instructive to examine their relative contributions. Assuming a biaxial stress state, the thermal contribution was estimated as s =DaDTE /(1–n ), where Da is the difference in th c c thermal expansion coefficients between coating and substrate, and DT is the temperature change after deposition. Taking a =2.3×10−6 K−1 [23] and CrN a =12.1×10−6 K−1, s is calculated to be 0.7– Sub th 1.4 GPa. The measured residual stress is thus dominated by the intrinsic stress since the contribution from s is th small. The value of the coefficient of thermal expansion

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Fig. 8. SEM micrographs of the crater wear pattern in arc-evaporated CrN coatings: (a) V =100 V, (b) V =200 V. Cohesive failure is S S marked with an arrow in (b).

for CrN is uncertain. Reported values of other transition metal nitrides are higher, for example TiN and VN have values of 9.3 and 8.1×10−6 K−1, respectively [23]. A larger value of a would result in an even smaller s CrN th and will consequently not alter the following discussion. A similar influence of V on the residual stress state S has been observed for arc-evaporated TiN and Ti(CN ) coatings [9,26 ] where it also was shown that such behavior could be well predicted by a model proposed by Davis [25] superimposed by thermally induced stresses. Davis’ model is based on the steady state of defect creation and annihilation rates within the collision cascade from the particle bombardment. The model includes parameters such as the necessary excitation energy for an atom in a metastable position in the coating to move to the surface, and the energy of the bombarding ions. It is not possible today to perform calculations for CrN with Davis’ model with a high degree of confidence since several of the included parameters are unknown. However, upon attempting to fit the unknown parameters to the experimental data it become clear that Davis’ model overestimates the compressive residual stress at high V . S

49

Additional stress relaxation mechanisms must therefore be active to account for the observed relaxation. Plastic yielding in sputtered Ti films [27] has, for example, been suggested as one possible mechanism for residual stress saturation observed at high V . Plastic S yielding in CrN, however, would require a decrease in yield strength by a factor of ~3 when the temperature is raised from 310 to 420°C in order to account for the complete compressive residual stress decrease. Instead, we propose annihilation of defects by diffusion [28] in addition to annihilation of defects in the collision cascade [25] as responsible for stress relaxation at high V S [29]. The formation of both b-Cr N and equiaxed grains 2 at the highest deposition temperature suggests that the temperature is sufficiently high for diffusion to take place during growth. The formation of microcracks is an additional mechanism for stress relaxation. Microcracks, perpendicular to the growth direction, were only observed for the V =400 V sample [see Fig. 5(e)]. A possible initiation S site for such cracks is the area beneath the observed droplets, where the droplet–coating interface is weak due to the presence of voids. Assuming that the grain boundaries are weaker in terms of crack resistance than the grain interior, the propagation in this direction will be enhanced by the presence of equiaxed grains, as observed in the 400 V sample. However, it should be pointed out that the density of observed cracks is not high enough to account for the total amount of stress relaxation observed. We also cannot exclude the possibility that the cracks seen in Fig. 5(e) are introduced during TEM sample preparation. 4.3. Microstructure vs. mechanical behavior The hardness in the coatings increases with increasing substrate bias voltage up to –100 V, after which a decrease is observed. Since a similar chemical composition was found in all coatings, the bonding structure is not expected to contribute to the observed differences in coating hardness. Possible microstructural features that can explain the hardness increase are the influence of column width by a Hall–Petch relation and an increase in lattice defects [30]. Observations from TEM and XRD peak broadening analysis show that both mechanisms can be active and are promoted by increasing bias voltage. The decrease in hardness seen at high substrate bias voltages is in agreement with a microstructure with a larger column width, a lower lattice defect density and a lower resistance for microcrack formation. The critical load level obtained from the scratch test was influenced by V . The critical load for cohesive S failure is 1.57 times higher in the sample prepared with V =20 V than the one prepared with V =400 V. Scratch S S results are influenced by parameters such as substrate hardness, surface roughness, coating thickness, etc. [31].

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However, in this study the variation of these parameters cannot explain the observed differences seen between the samples since they remain almost constant between the samples. Instead, microstructure and residual stress differences are the primary sources for scratch test differences. Several deformation mechanisms are active simultaneously, such as plastic deformation and crack propagation. It is apparent that cracks form and propagate more easily when a higher substrate bias is used. The amount of cohesive flaking outside the rim of the scratch track at high contact loads also suggest that cracks propagate longer in samples deposited with a high V compared with low biased samples. As discussed S in the previous section, the presence of more equiaxed grains observed at V =400 V introduces a larger fraction S of grain boundaries, which may enhance crack propagation perpendicular to the growth direction. The CrN coatings were found to offer a five times improvement in wear rate over the bare substrate. Comparing the coatings, only the V =200 V coating S can be distinguished as having a higher wear rate. This sample also shows a different wear scar appearance. While the other samples show grinding grooves [Fig. 8(a)] without chipping the V =200 V sample shows S less grooves and pronounced chipping [Fig. 8(b)]. The difference in wear mechanisms is, in the same way as the scratch test, related to the load level for different deformation mechanisms to become active, i.e. a higher scratch load is required for cohesive failure to occur at V =100 V than at V =200 V. For samples deposited at S S V =20 to 100 V, plastic deformation is the dominating S deformation mechanism, while fracturing dominates for V =200 V under the load used in the wear test. The S critical load for transition from plastic deformation to fracturing is controlled by the combination of microstructural features including lattice defect density, columnar size, grain boundary properties, residual stress state, and elastic and plastic anisotropy.

5. Conclusions Arc evaporation of Cr in nitrogen gas with a high nitrogen partial pressure (8 Pa) can be used to deposit thick single phase, dense, and well-adherent CrN coatings on high speed steels. An increase in lattice defect density, a decrease in column width, and an increase in compressive residual stress occurs as the substrate bias voltage is increased from 20 to 100 V, and these changes are considered responsible for an increase in hardness and decrease in critical load for cohesive flaking. At higher substrate bias voltages (V =200–400 V ) the comS pressive residual stress decreases, which is proposed to be a result of defect annihilation in the collision cascade and by diffusion. The observed presence of equiaxed grains at V =400 V is suggested to promote a further S

decrease in both resistance to cohesive flaking and hardness.

Acknowledgements The authors thank Jonathan Almer and Lars Hultman for useful discussions and Ivan Petrov for the supply of a reference sample for chemical analysis. The financial support of Uddeholm Tooling AB’s research foundation is gratefully acknowledged.

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