Improving the mechanical and anti-wear properties of AlTiN coatings by the hybrid arc and sputtering deposition

Journal Pre-proof Improving the mechanical and anti-wear properties of AlTiN coatings by the hybrid arc and sputtering deposition Quan Zhang, Zhengtao Wu, Yu X. Xu, Qimin Wang, Li Chen, Kwang Ho Kim PII:

S0257-8972(19)31013-8

DOI:

https://doi.org/10.1016/j.surfcoat.2019.125022

Reference:

SCT 125022

To appear in:

Surface & Coatings Technology

Received Date: 26 June 2019 Revised Date:

23 September 2019

Accepted Date: 26 September 2019

Please cite this article as: Q. Zhang, Z. Wu, Y.X. Xu, Q. Wang, L. Chen, K.H. Kim, Improving the mechanical and anti-wear properties of AlTiN coatings by the hybrid arc and sputtering deposition, Surface & Coatings Technology (2019), doi: https://doi.org/10.1016/j.surfcoat.2019.125022. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier B.V.

1

Improving the mechanical and anti-wear properties of AlTiN coatings by the

2

hybrid arc and sputtering deposition

3 4

Quan Zhang a, Zhengtao Wu a,b, Yu X. Xu a,*, Qimin Wang a,*, Li Chen c,

5

and Kwang Ho Kim d a

6 7 8 9 10 11 12

School of Electromechanical Engineering, Guangdong University of Technology, Guangzhou 510006, China b Department of physics (IFM), Linköping University, Linköping SE-581 83, Sweden c State Key Laboratory of Powder Metallurgy, Central South University, Changsha 410083, China d Global Frontier R&D Center for Hybrid Interface Materials, Pusan National University, Busan 46241, South Korea

13 14

Abstract: Arc-evaporated AlTiN coatings with remarkable adhesion and thermal

15

stability are widely used in manufacturing. However, the large residual stress and

16

surface roughness limit their high-level applications for advanced machining. In this

17

study, a hybrid technique combining cathodic arc evaporation and magnetron

18

sputtering was utilized to enhance the mechanical properties and wear resistance of

19

AlxTi1−xN with x = ~0.65. The hybrid AlTiN coatings exhibited a nano-multilayer

20

architecture with alternating arc-evaporated and magnetron sputtered constituents,

21

which decreased their residual stress from −7.1 GPa to about −5.0 GPa. In addition, the

22

hybrid AlTiN showed higher indentation toughness than the monolithic coatings. The

23

hybrid AlTiN coating with a target sputtering power of 7.0 kW showed excellent wear

24

resistance at room and high temperature. At 800 °C, the hybrid coatings mainly showed

25

abrasive and oxidation wear, while the monolithic AlTiN coating showed

26

abrasive/oxidation wear and brittle failure.

27

Keywords: AlTiN; Hybrid deposition; Mechanical properties; Residual stress; Wear

28

resistance

29 30

*

31

E-mail: [email protected] (Y. Xu); [email protected] (Q. Wang)

Corresponding author.

1

1

Introduction

2

Owing to their excellent thermal stability, superior mechanical properties and high

3

oxidation resistance in comparison with traditional TiN, AlTiN coatings prepared by

4

physical vapor deposition (PVD) have been extensively used as protective materials for

5

various industrial applications such as cutting tools, casting moulds, and wear parts

6

[1-5]. When exposed to high-temperature environments, AlTiN coatings with a

7

metastable cubic structure decompose into cubic AlN- and TiN-rich nano-sized

8

domains with interrelated coherent interfaces, giving rise to age-hardening [6-8].

9

AlTiN has gained immense attention as a well-establish hard coating system. Various

10

efforts have been made to improve its properties in order to realize its practical

11

applications. The use of cutting-edge deposition techniques is an efficient approach to

12

develop high-performance AlTiN coatings [9, 10]. Greczynski et al. reported that

13

AlTiN coatings with high hardness and low stress can be developed by using a

14

combination of the high-power impulse and direct current magnetron sputtering

15

(DCMS) methods [11, 12].

16

Cathodic arc evaporation (CAE) with a high ionization rate is the most commonly

17

used technique to fabricate hard nitride coatings, which exhibit strong adhesion to

18

various substrates and dense microstructure with stoichiometric composition [13-15].

19

However, macro-particle growth defects, which increase the surface roughness [16]

20

and contribute to the fast outward diffusion of substrate materials [17], are inevitable

21

and limit the high precision applications of arc-evaporated AlTiN coatings. The use of 2

1

magnetically-filtered cathodic arc evaporation can suppress the macro droplet defect, to

2

a certain extent, at the expense of the deposition rate [18]. Nevertheless, the relatively

3

high residual stress generated during CAE, primarily resulting from the bombardment

4

of energetic particles during the deposition, limits the deposition of thick AlTiN

5

coatings [12]. Excessive residual compressive stress causes coating delamination and

6

chipping failure during machining applications. Magnetron sputtering is another vapor

7

deposition technique used to prepare excellent surface quality coatings with the

8

controllable residual stress state [19, 20]. On the other hand, sputtered AlTiN coatings

9

show low deposition rate and compactness because of their low ionization rates.

10

Coatings prepared using the combination of the CAE and MS techniques show high

11

deposition rate and compactness (arising from CAE) and low surface roughness and

12

residual stress level (arising from MS), and hence are ideal for industrial applications.

13

Hence, various hybrid techniques combining the CAE and MS methods have been used

14

for the deposition of TiAlSiN coatings [21-23]. Yu et al. [24] used a hybrid hollow

15

cathode discharge ion plating and medium frequency magnetron sputtering to improve

16

the milling performance of AlTiN coated cutting tools. Moreover, compositionally

17

modulated multilayers with alternating nitride layers can achieve superior properties,

18

involving synergistic improvement in hardness [25], toughness [26], and excellent

19

thermal stability [27, 28]. Hybrid deposition techniques that introduce a multilayer

20

architecture, which periodically stacks arc-evaporated sublayers with exceptional

21

adhesive properties and droplet-free magnetron sputtered constituents with low 3

1

residual stress, produce efficient coatings for high-speed heavy load cutting

2

applications.

3

To achieve highly efficient hybrid depositions, the optimization of process

4

parameters, such as the atmospheric pressure, is a prerequisite. For instance, the partial

5

pressure of nitrogen can significantly affect the crystallographic structure, crystallinity,

6

and composition of sputtered coatings in an Ar and N2 mixing atmosphere. The

7

hardness of reactively sputtered AlTiN films significantly depends on the partial

8

pressure of N2. With an increase in the N2 partial pressure, the coating hardness first

9

increases and then decreases [29, 30]. Cai and co-workers demonstrated that the

10

increase in the N2 partial pressure is beneficial for the reduction of macro-particles

11

during arc evaporation, and hence improves the mechanical properties of AlTiN

12

coatings [31]. Various studies have been carried out to fabricate AlTiN monolithic and

13

multilayer coatings and to investigate their structures and properties [2, 32, 33].

14

However, only a few studies are available on the fabrication of co-deposited AlTiN

15

coatings by CAE and MS. In this study, we fabricated AlTiN coatings using a

16

combination of the CAE and DCMS techniques with a gas ion source at various

17

sputtering target powers. Also, the microstructures, mechanical properties, and

18

tribological behaviours of the hybrid AlTiN coatings were investigated.

19

2. Experimental details

20

2.1 Coating deposition

21

The AlTiN coatings were synthesized on mirror polished cemented carbide 4

1

(WC-6 wt.% Co, 16 × 16 × 4 mm3 in dimension) and AISI 304 stainless steel sheets (50

2

× 10 × 0.8 mm3 in dimension for stress measurements with elastic modulus of 198 GPa

3

and Poisson’s ratio of 0.3) by the hybrid method combining CAE, DCMS, and gas ion

4

source (GIS) utilizing a self-build multifunctional PVD deposition system

5

(GDUT-HAS500). Before mounting to the sample holder, all substrates were

6

ultrasonically cleaned in deionized water dissolved with powdery metal cleaner and

7

anhydrous alcohol for 20 min at 65 °C, respectively. Two opposite Al67Ti33 (at.%)

8

circular targets (Φ100 × 20 mm2 for arc evaporation) and one Al67Ti33 (at.%) rotatable

9

cylindrical target (Φ70 × 597 mm2 for sputtering) were used for the deposition of

10

hybrid AlTiN coatings. The layout of targets is depicted in the form of top-view of Fig.

11

1. After reaching a base pressure less than 5 × 10−3 Pa, the GIS (leveled at 3.0 kW)

12

assisted etching procedure was performed under an Ar pressure of 1.0 Pa and a pulsed

13

bias voltage of −600 V (80% duty cycle and 80 kHz frequency) for 30 min to remove

14

contaminants from the substrate surface. During deposition, an arc-evaporated AlTiN

15

transition layer with a thickness of ~1.0 µm was prepared initially to work as the

16

support layer under an N2 pressure of 3.0 Pa, a target current of 80 A, and a pulsed bias

17

of −150 V. Subsequently, the hybrid AlTiN (H-AlTiN) coatings were co-deposited

18

simultaneously using the cathodic arc evaporation and magnetron sputtering with the

19

assistance of GIS (leveled at 0.5 kW) for 180 min. The deposition temperature was set

20

at 350 °C, the partial pressure of nitrogen was 0.6 Pa with constant total working

21

pressure (Ar + N2) of 0.8 Pa, and the bias voltage was −150 V (80% duty cycle and 80 5

1

kHz frequency). The target current for CAE, ICAE, was 80 A, and target power for

2

DCMS PMS was altered from 6.0 to 9.0 kW in the step of 1.0 kW. For comparison,

3

monolithic arc-evaporated (A-AlTiN) and sputtered AlTiN (S-AlTiN) coatings were

4

also synthesized utilizing the same deposition equipment under the same substrate

5

voltage and atmosphere in this work. The main deposition parameters for monolithic

6

and multilayered coatings are listed in Table 1.

7

2.2 Coating characterizations

8

The morphologies and elemental compositions of the AlTiN coatings were

9

investigated using a scanning electron microscope (SEM, FEI Nova NanoSEM 430)

10

equipped with an energy-dispersive X-ray spectrometer (EDX, Oxford instruments

11

X-MaxN). The phase structures of the co-deposited AlTiN coatings were investigated

12

using a grazing X-ray diffractometer (GIXRD, Bruker D8 Advance) operating at 40 kV

13

and 40 mA with an incidence angle of 1.0° and Cu Kα radiation as the X-ray source. A

14

transmission electron microscope (TEM, FEI Titan G2 60–300) operating at 300 kV

15

was used to examine the microstructures of the coatings. Scanning transmission

16

electron microscopy (STEM) investigations were also carried out on the same

17

instrument. The preparation of the TEM sample was conducted in a dual-beam focused

18

ion beam system (FEI Helios Nanolab 600i) following the lift-out procedure. Final

19

surface cleaning was conducted at 5 kV and 26 pA to minimize the damage and

20

artefacts.

21

The hardness and elastic modulus of the coatings were determined by carrying out 6

1

their nanoindentation tests following the Oliver and Pharr method [34] using an

2

instrumented indenter (Anton Paar TriTec, TTX-NHT2) equipped with a Berkovich

3

diamond tip. A normal load of 10 mN was used for all the coatings. A penetration depth

4

less than 10% of the coating thickness was used to minimize the effect of the substrate.

5

Fifteen measurements were conducted for each sample. The residual stress of the

6

coatings was calculated using the Stoney equation [35]. The curvature of the coated

7

sample was determined using a laser profiler (SuPro Instruments FST-1000). The

8

toughness of the coatings was qualitatively analysed by monitoring the contact-induced

9

damage using a micro-indentation instrument with a Vickers’s tip. An indentation load

10

of 4.905 N was used and the holding time at maximum load was 10 s. Five repeated

11

measurements were carried out for each sample. The adhesion strength between the

12

coatings and the substrate was evaluated using an Anton Paar Revetest scratch tester

13

with a scratch length of 5 mm at a scratch speed of 10 mm/min over the load range of

14

1–120 N.

15

The tribological properties of the coatings were evaluated using a ball-on-disc

16

tribometer (Anton Paar TriTec, HT-THT) under ambient atmosphere (50 ± 5 % of

17

relative humidity) using an Al2O3 ball (6 mm in diameter) as the counterpart material.

18

The normal load was 5 N, the wear track diameter was set at 4 mm, and the rotation

19

speed was maintained at 954.9 rpm (corresponding to a linear rate of 0.2 m/s) for all the

20

tests. The total number of laps was 14000 (175.9 m) for tests at room temperature (RT)

21

(23−25 °C) and 9000 (113.1 m) for tests at 800 °C. The friction coefficient of the 7

1

coatings was determined as a function of the number of sliding cycle. The wear rates, k,

2

of the coatings were calculated according to the Archard equation, k = V / (F × L),

3

where V is the wear volume loss, F is the applied load, and L is the total sliding distance.

4

The wear volume loss was equal to the cross-sectional area times the perimeter of the

5

wear track. The cross-sectional areas of the wear tracks were determined using a laser

6

scanning confocal microscope (Olympus LEXT OLS4000).

7

3. Results and discussion

8

3.1. Composition, morphology, and structure

9

The EDX results of the coatings are listed in Table 1. The H-AlTiN coatings

10

showed an Al/(Al + Ti) ratio of ~0.65. Figure 2 shows the SEM images of the fracture

11

cross-sections of the A-, S- and H-AlTiN coatings deposited at different PMS. All the

12

coatings exhibited good adhesion with the cemented carbide substrate without crack

13

initiation (scratch tests are discussed in the next section). The A-AlTiN and S-AlTiN

14

monolithic coatings showed columnar growth morphology (Figs. 2a and 2f). The

15

H-AlTiN coatings (Figs. 2b–2e) showed relatively finer crystalline grains as compared

16

to the A-AlTiN coatings. With an increase in the PMS, the coating thickness of the

17

H-AlTiN coatings increased gradually (Table 1), indicating the thickening of the

18

sputtered constituent sublayers.

19

The GIXRD patterns of the A-, S- and H-AlTiN coatings deposited onto the

20

cemented carbides at various PMS are shown in Fig. 3. Like the A-AlTiN coatings, all

21

the H-AlTiN coatings showed a face-centred cubic structure. The diffraction peaks of 8

1

A- and H-AlTiN coatings locate between those of standard cubic TiN and cubic AlN.

2

As compared to TiN (ICDD 00-38-1420), the diffraction peaks of the H-AlTiN coatings

3

shifted to higher angles mainly because of the formation of a solid solution of Al with a

4

smaller atomic size in the Ti−N lattice. The PMS showed no significant effect on the

5

crystalline structure of the H-AlTiN coatings. The phase structure of PVD AlTiN is

6

related to its Al content. The S-AlTiN coating with the Al/(Al + Ti) ratio of 0.71 mainly

7

exhibited a wurtzite structure because its Al content exceeded the solid solubility in

8

Ti−N under normal PVD deposition conditions [3, 36, 37]. In the S-AlTiN coatings, the

9

substitution of Al by Ti atoms with larger atomic radius resulted in a negative shift of

10

the diffraction peaks corresponding to the wurtzite phase as compared to w-AlN (ICDD

11

00-025-1133).

12

TEM and STEM investigations were carried out to examine the microstructure of

13

the H-AlTiN coating obtained at the PMS of 7.0 kW. The bright-field image shown in

14

Fig. 4a reveals that the coating consisted of fine columnar grains extending along the

15

growth direction with a column width less than 50 nm. The electron diffraction pattern

16

(inset Fig. 4a) revealed that the coating exhibited a single-phase face-centred cubic

17

structure. High-resolution TEM (HRTEM) observations are carried out to examine the

18

lattice structure of coatings. Individual columnar grains projected from the [011] zone

19

axis could be observed from the HRTEM image of the H-AlTiN coating (Fig. 4b). The

20

Z-contrast STEM high-angle annular dark field image (Fig. 4c) confirmed the

21

nano-scale multilayered structure of the H-AlTiN coating. The sublayers generated by 9

1

sputtering showed a thickness of ~2.4 nm, while the arc-evaporated AlTiN sublayers

2

showed a larger thickness of ~6.6 nm. The arc-evaporated sublayers showed lower Al

3

content and higher average atomic number, and hence a brighter contrast than the

4

sputtered layers. Moreover, the STEM image of the coating not only showed atomic

5

number contrast, but also diffraction contrast. The single columnar grains in the coating

6

were assembled by the alternating CAE- and DCMS-sublayers. The continuous lattice

7

fringes across the sublayers (Fig. 4c) indicate the occurrence of epitaxial growth via

8

coherent strain. The formation of coherent interfaces can be attributed to the

9

minimization of total energy, in which the decrease of interface overcomes the increase

10

of strain energy.

11

3.2. Mechanical properties

12

Figure 5 shows the top-view micro-indentation morphologies of the A-, S-, and

13

H-AlTiN coatings deposited on the cemented carbide substrates at various PMS. When a

14

load of 4.905 N and a load maintenance time of 10 s were used, A- (Fig. 5a) and S- (Fig.

15

5f) AlTiN coatings exhibited cracking and delamination at the edges of the indented

16

impression. However, the H-AlTiN coatings showed only some indentation-induced

17

rhombus cracks. With an increase in the PMS, both the number and width of the

18

rhombus cracks decreased significantly, thus improving the toughness of the coatings,

19

as revealed by the qualitative analysis results. This is consistent with the results

20

reported in Ref. [38]. The multilayer H-AlTiN coatings with finer crystalline grains

21

lead to greater critical failure stress for the onset of rhombus cracks and inclined shear 10

1

cracks [38]. Once the cracks formed, the crack propagation was hindered by crack

2

deflection and crack tip blunting at the interface between sublayers [39]. In addition, as

3

proven by STEM, H-AlTiN coating exhibited a compositionally modulated superlattice

4

structure, the enhanced fracture toughness of superlattice effect has been reported in

5

coherently grown TiN/CrN multilayer coatings [26]. As for the A- and S-AlTiN

6

monolithic coatings, once the rhombus and inclined shear cracks initiated, the

7

propagation along grain boundaries was hard to hinder due to the monogenous

8

composition and coarse columnar structure within the monolithic coatings [39], which

9

decreased the resistance to fracture.

10

The hardness (H) and elastic moduli (E) of A-, S- and H-AlTiN coatings deposited

11

at various PMS are shown in Fig. 6a. The H-AlTiN coating deposited at the PMS of 7.0

12

kW showed higher hardness (33.2 ± 0.4 GPa) and elastic modulus (457.9 ± 13.2 GPa)

13

than the A- and S-AlTiN coatings. For the modulation period of ten or a few

14

nanometres, the multilayer architecture of H-AlTiN coatings improves their hardness

15

[33]. Chu and Barnett [40] proposed a theoretical model based on the dislocation glide

16

across layers limited by the shear modulus difference and dislocation glide within

17

individual layers to explain superlattice strengthening/hardening. Besides, the

18

introduction of the multilayer structure hindered both the initiation and propagation of

19

microcracks when the H-AlTiN coated surface was exposed to a mechanical load

20

during the nanoindentation test, thus improving its hardness [41]. In addition, the TEM

21

results (Fig. 4) confirmed that the formation of the coherent/semi-coherent interfaces 11

1

generated an alternative strain field in the multilayer structure of the coatings, thus

2

blocking the dislocation movement [42, 43].

3

Figure 6b shows the adhesion force derived from the scratch test of the cemented

4

carbide and the residual stress on the AISI 304 stainless steel sheets for the A-, S-, and

5

H-AlTiN coatings. The adhesion force of the H-AlTiN coatings exceeded 100 N, which

6

is higher than those of the A- and S-AlTiN coatings. For hard coatings, failure in the

7

scratch test was mainly caused by interfacial defects between coating and substrate,

8

plastic deformation and fracture within the coatings [44]. Deposited on the same

9

substrate with the same transition layer, this higher adhesion force of the H-AlTiN

10

coatings can be attributed to the enhanced toughness (Fig. 5), and hence higher

11

resistance to plastic deformation and cracking fracture [45, 46]. On the other hand, as

12

compared to the S- and A-AlTiN coatings, the adhesion force of H-AlTiN coatings also

13

benefited from the high energy metallic ions produced by CAE and the modest residual

14

stress [47]. The H-AlTiN coatings exhibited a compressive stress level of −4.3 – −5.4

15

GPa, which is intermediate to those of the A-AlTiN (−7.1 GPa) and S-AlTiN (−2.1 GPa)

16

coatings. In general, at high residual compressive stress, PVD coatings show high

17

indentation hardness. However, the driving forces for the increase in the hardness and

18

residual stress in multilayers are different. Ion bombardment during the deposition is

19

the primary reason for the accumulation of compressive stress, Ion bombardment is

20

controlled by the deposition conditions such as the bias, pressure and target power. In

21

addition, Seidl et al. [48] demonstrated the strong dependence of the residual stress of 12

1

TiAlTaN/AlCrN multilayers with substrate materials. This dependence can be

2

attributed to the difference in the thermal expansion coefficients of the coating and the

3

substrate. The hardness of multilayers increases gradually with the diminution of the

4

modulation period from hundreds to several nanometres owing to interfacial

5

strengthening. However, the interfaces of multilayer architecture offer additional stress

6

relief by layer sliding, leading to a reduction in the residual stress with the optimization

7

of the modulation period. The multilayers of TiN/NbN [49] and AlTiN/MoSiB [50]

8

coatings show even lower residual stress than their monolithic constituents.

9

3.3. Tribological properties

10

Figure 7 shows the friction coefficient curves of the A-, S-, and H-AlTiN coatings

11

deposited various PMS as a function of the sliding distance at (a) RT and (b) 800 °C. The

12

friction coefficients of all the coatings showed a steep increase in the slope at the initial

13

running-in stage because of the transition in the contact conditions from two-body

14

abrasion to interfacial sliding [51]. A relatively steady state was achieved after the

15

transition. Among all the coatings, the S-AlTiN coating exhibits the lowest friction

16

coefficient at RT and 800 °C. Compared to the A-AlTiN coatings, the H-AlTiN coatings

17

showed low friction coefficients at both RT and 800 °C. Figure 8 shows the wear rates

18

of the A-, S-, and H-AlTiN coatings at RT and elevated temperatures. At RT, the wear

19

rates of the H-AlTiN coatings were less than 2.0 × 10-6 mm3/N·m. The H-AlTiN coating

20

with the PMS of 7.0 kW showed a minimum wear rate of 1.2 × 10-6 mm3/N·m, while the

21

A-AlTiN coating exhibited a higher wear rate of 5.7 × 10-6 mm3/N·m. However, the 13

1

wear rates of the coatings increased at 800 °C. The minimum wear rate of 3.2 × 10-6

2

mm3/N·m was observed for the H-AlTiN (PMS = 7.0 kW) coating. The wear rate of the

3

A-AlTiN coating (7.4 × 10-6 mm3/N·m) was almost twice that of the H-AlTiN (PMS =

4

7.0 kW) coating. The S-AlTiN coating showed wearing at both RT and 800 °C.

5

The SEM images of the wear morphologies of the A-AlTiN, S-AlTiN, and

6

H-AlTiN (PMS = 7.0 kW) coatings are shown in Figs. 9 and 10. The H-AlTiN (PMS = 7.0

7

kW) coating exhibited higher debris removal efficiency than the A-AlTiN coating (Figs.

8

9a and 9b). In the case of the A-AlTiN coating, the debris moved from the contact

9

surface and accumulated at the edge of the wear track. Furrows were observed in the

10

central zone of the wear tracks of all the coatings. These furrows were generated by the

11

abrasive wear caused by the breakage of the coating and the Al2O3 counterpart ball

12

during the wear process. The formation of furrows indicates that the abrasive grains

13

contributed to the tribological behaviour of the coating by causing abrasive wear. The

14

H-AlTiN (PMS = 7.0 kW) coating showed a smooth worn surface as compared to the

15

A-AlTiN coating with fewer and smaller furrows in the worn zone. Both the edges of

16

the wear tracks were O-rich (framed regions 2 and 5), indicating that the wear debris

17

were mainly generated by the counterpart Al2O3 ball, while the broken coating particles

18

contributed to the rest of the wear debris. W and Co were not observed in the central

19

area of the wear tracks of the A-AlTiN and H-AlTiN coatings, indicating that the

20

coatings did not wear out.

21

At the test temperatures of 800 °C, the H-AlTiN (PMS = 7.0 kW) exhibited higher 14

1

debris removal efficiency than the A-AlTiN coating (Figs. 10a and 10b). Besides, the

2

EDX data (Table 3) revealed that the debris accumulated at the edge of the wear tracks

3

consisted of titanium oxide, aluminium oxide, and wear particles of the coatings.

4

Moreover, at 800 °C, the oxidation of the AlTiN coatings occurred under ambient

5

atmosphere. Furthermore, the EDX results for the interior of the wear track showed that

6

the oxidative products of the A-AlTiN and H-AlTiN (PMS = 7.0 kW) coatings were a

7

mixture of Al2O3 and TiO2. As can be observed from the insets of Figs. 10a and 10b, the

8

dashed line framed wear zone of the A-AlTiN coating showed a much more coarse

9

morphology than the H-AlTiN (PMS = 7.0 kW) coating because of its brittle failure

10

mechanism [52]. The macro-cutting effect of the Al2O3 ball caused the detachment of

11

abrasive particles from the coatings through brittle failure mechanism. The

12

macro-cutting effect and low debris removal efficiency of the A-AlTiN coating [53],

13

resulted in the severe irregularity of the furrows in the wear track. This is consistent

14

with the results reported in a previous study [54]. The H-AlTiN (PMS = 7.0 kW) coating

15

showed fewer and smaller furrows than the A-AlTiN coating (which exhibited a

16

smoother surface) [55], because of its higher hardness, toughness, adhesive strength,

17

and debris removal efficiency. On the basis of these results, it can be stated that the

18

primary wear mechanism of the H-AlTiN (PMS = 7.0 kW) coating at high-temperatures

19

was a combination of abrasive and oxidation wear. On the other hand, the A-AlTiN

20

coating showed a combination of abrasive wear, oxidation wear, and brittle failure.

21

Consequently, the high hardness and adhesive strength of the H-AlTiN coating at both 15

1

RT and elevated temperatures improved its toughness, resistance, and debris removal

2

efficiency while reducing its friction coefficient.

3

4. Conclusions

4

In this study, AlTiN coatings were deposited using a hybrid

arc

5

evaporation/magnetron sputtering technique. The effect of the target sputtering power

6

on the structure, mechanical and tribological properties of the AlTiN coatings were

7

investigated. The main findings of this study are as follows:

8

1) The hybrid AlTiN coating exhibited a nano-scale multilayer structure,

9

consisting of arc-evaporated and sputtered constituents. The hybrid AlTiN coatings

10

incorporated with sputtered components showed residual compressive stress of 4.3 –

11

5.4 GPa and a thickness of ~6.0 µm. Moreover, as compared to the monolithic AlTiN

12

coating, the hybrid coatings showed improved toughness and adhesion strength, which

13

are beneficial for heavy-load cutting applications.

14

2) The sputtering target power affected the proportion of the sputtered components

15

in the multilayered hybrid AlTiN coatings. At the target power of 7.0 kW, the hybrid

16

AlTiN coatings showed a high hardness of 33.2 ± 0.4 GPa and excellent wear resistance

17

at RT and 800 °C. At 800 °C, the primary wear mechanism of the hybrid coatings was a

18

combination of abrasive and oxidation wear. On the other hand, the arc-evaporated

19

AlTiN coating showed a combination of abrasive wear, oxidation wear, and brittle

20

failure.

21 16

1

Acknowledgments

2

This work was supported by the National Natural Science Foundation of China

3

(Grant Nos. 51875109 and 51801032), the National Key Research and Development

4

Program of China (Grant No. SQ2017YFGH001672) and the Natural Science

5

Foundation of Guangdong Province (Grant No. 2018A030310546).

6

17

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41

References [1] [2]

[3]

[4]

[5]

[6]

[7]

[8]

[9]

[10]

[11]

L. Hultman, Thermal stability of nitride thin films, Vacuum 57 (2000) 1-30. https://doi.org/10.1016/S0042-207X(00)00143-3 S. PalDey, S.C. Deevi, Single layer and multilayer wear resistant coatings of (Ti, Al)N: a review, Materials Science and Engineering: A 342 (2003) 58-79. https://doi.org/10.1016/S0921-5093(02)00259-9 K. Kutschej, P.H. Mayrhofer, M. Kathrein, P. Polcik, R. Tessadri, C. Mitterer, Structure, mechanical and tribological properties of sputtered Ti1–xAlxN coatings with 0.5≤x≤0.75, Surface and Coatings Technology 200 (2005) 2358-2365. https://doi.org/10.1016/j.surfcoat.2004.12.008 L. Chen, M. Moser, Y. Du, P.H. Mayrhofer, Compositional and structural evolution of sputtered Ti-Al-N, Thin Solid Films 517 (2009) 6635-6641. https://doi.org/10.1016/j.tsf.2009.04.056 Y.X. Xu, L. Chen, F. Pei, Y. Du, Y. Liu, J.L. Yue, Influence of Hf on the structure, thermal stability and oxidation resistance of Ti-Al-N coatings, Thin Solid Films 565 (2014) 25-31. https://doi.org/10.1016/j.tsf.2014.06.043 P.H. Mayrhofer, A. Hörling, L. Karlsson, J. Sjölén, T. Larsson, C. Mitterer, L. Hultman, Self-organized nanostructures in the Ti–Al–N system, Applied Physics Letters 83 (2003) 2049-2051. https://doi.org/10.1063/1.1608464 L. Chen, Y. Du, P.H. Mayrhofer, S.Q. Wang, J. Li, The influence of age-hardening on turning and milling performance of Ti–Al–N coated inserts, Surface and Coatings Technology 202 (2008) 5158-5161. https://doi.org/10.1016/j.surfcoat.2008.05.036 D. McIntyre, J.E. Greene, G. Håkansson, J.E. Sundgren, W.D. Münz, Oxidation of metastable single‐phase polycrystalline Ti0.5Al0.5N films: Kinetics and mechanisms, Journal of Applied Physics 67 (1990) 1542-1553. https://doi.org/10.1063/1.345664 H. Elmkhah, T.F. Zhang, A. Abdollah-zadeh, K.H. Kim, F. Mahboubi, Surface characteristics for the TiAlN coatings deposited by high power impulse magnetron sputtering technique at the different bias voltages, Journal of Alloys and Compounds 688 (2016) 820-827. https://doi.org/10.1016/j.jallcom.2016.07.013 T. Shimizu, Y. Teranishi, K. Morikawa, H. Komiya, T. Watanabe, H. Nagasaka, M. Yang, Impact of pulse duration in high power impulse magnetron sputtering on the low-temperature growth of wurtzite phase (Ti, Al)N films with high hardness, Thin Solid Films 581 (2015) 39-47. https://doi.org/10.1016/j.tsf.2014.11.076 G. Greczynski, J. Lu, M. Johansson, J. Jensen, I. Petrov, J.E. Greene, L. Hultman, Selection of metal ion irradiation for controlling Ti1−xAlxN alloy growth via hybrid HIPIMS/magnetron co-sputtering, Vacuum 86 (2012) 1036-1040. https://doi.org/10.1016/j.vacuum.2011.10.027 18

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42

[12]

[13]

[14]

[15]

[16]

[17]

[18]

[19]

[20]

[21]

[22]

[23]

G. Greczynski, J. Lu, J. Jensen, S. Bolz, W. Kölker, C. Schiffers, O. Lemmer, J.E. Greene, L. Hultman, A review of metal-ion-flux-controlled growth of metastable TiAlN by HIPIMS/DCMS co-sputtering, Surface and Coatings Technology 257 (2014) 15-25. https://doi.org/10.1016/j.surfcoat.2014.01.055 G.H. Kang, H. Uchida, E.S. Koh, A study on the surface structure of Ti cathode and the macroparticle of TiN films prepared by the arc ion plating process, Surface and Coatings Technology 86-87 (1996) 421-424. https://doi.org/10.1016/S0257-8972(96)02980-5 H. Ljungcrantz, L. Hultman, J.E. Sundgren, G. Håkansson, L. Karlsson, Microstructural investigation of droplets in arc-evaporated TiN films, Surface and Coatings Technology 63 (1994) 123-128. https://doi.org/10.1016/S0257-8972(05)80016-7 A.J. Perry, J.R. Treglio, A.F. Tian, Low-temperature deposition of titanium nitride, Surface and Coatings Technology 76-77 (1995) 815-820. https://doi.org/10.1016/0257-8972(95)02628-2 C. Sateesh Kumar, S.K. Patel, Experimental and numerical investigations on the effect of varying AlTiN coating thickness on hard machining performance of Al2O3-TiCN mixed ceramic inserts, Surface and Coatings Technology 309 (2017) 266-281. https://doi.org/10.1016/j.surfcoat.2016.11.080 K. Dejun, G. Haoyuan, Friction-wear behaviors of cathodic arc ion plating AlTiN coatings at high temperatures, Tribology International 88 (2015) 31-39. https://doi.org/10.1016/j.triboint.2015.03.009 B.K. Tay, Z.W. Zhao, D.H.C. Chua, Review of metal oxide films deposited by filtered cathodic vacuum arc technique, Materials Science and Engineering: R: Reports 52 (2006) 1-48. https://doi.org/10.1016/j.mser.2006.04.003 U. Helmersson, M. Lattemann, J. Bohlmark, A.P. Ehiasarian, J.T. Gudmundsson, Ionized physical vapor deposition (IPVD): A review of technology and applications, Thin Solid Films 513 (2006) 1-24. https://doi.org/10.1016/j.tsf.2006.03.033 P.J. Kelly, R.D. Arnell, Magnetron sputtering: a review of recent developments and applications, Vacuum 56 (2000) 159-172. https://doi.org/10.1016/S0042-207X(99)00189-X I.-W. Park, S.R. Choi, J.H. Suh, C.-G. Park, K.H. Kim, Deposition and mechanical evaluation of superhard Ti–Al–Si–N nanocomposite films by a hybrid coating system, Thin Solid Films 447-448 (2004) 443-448. https://doi.org/10.1016/S0040-6090(03)01122-2 J. Wu, N. He, H. Li, X. Liu, L. Ji, X. Huang, J. Chen, Deposition and characterization of TiAlSiN coatings prepared by hybrid PVD coating system, Surface and Interface Analysis 47 (2015) 184-191. https://doi.org/10.1002/sia.5680 N. He, H. Li, L. Ji, X. Liu, H. Zhou, J. Chen, High temperature tribological properties of TiAlSiN coatings produced by hybrid PVD technology, Tribology 19

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42

[24]

[25]

[26]

[27]

[28]

[29]

[30]

[31]

[32]

[33]

[34]

[35]

International 98 (2016) 133-143. https://doi.org/10.1016/j.triboint.2016.02.034 D. Yu, C. Wang, X. Cheng, F. Zhang, Optimization of hybrid PVD process of TiAlN coatings by Taguchi method, Applied Surface Science 255 (2008) 1865-1869. https://doi.org/10.1016/j.apsusc.2008.06.204 F. Pei, H.J. Liu, L. Chen, Y.X. Xu, Y. Du, Improved properties of TiAlN coating by combined Si-addition and multilayer architecture, Journal of Alloys and Compounds 790 (2019) 909-916. https://doi.org/10.1016/j.jallcom.2019.03.248 R. Hahn, M. Bartosik, R. Soler, C. Kirchlechner, G. Dehm, P.H. Mayrhofer, Superlattice effect for enhanced fracture toughness of hard coatings, Scr. Mater. 124 (2016) 67-70. https://doi.org/10.1016/j.scriptamat.2016.06.030 Y.X. Xu, L. Chen, Z.Q. Liu, F. Pei, Y. Du, Improving thermal stability of TiSiN nanocomposite coatings by multilayered epitaxial growth, Surface and Coatings Technology 321 (2017) 180-185. https://doi.org/10.1016/j.surfcoat.2017.04.057 C. Hu, H.T. Wang, L. Chen, Y.X. Xu, H.J. Liu, Y. Du, Structural, mechanical and thermal properties of Ti1-xSixN/CrAlN (x = 0, 0.13 and 0.22) multilayers, Journal of Alloys and Compounds 800 (2019) 355-362. https://doi.org/10.1016/j.jallcom.2019.06.012 J. Musil, H. Hrubý, Superhard nanocomposite Ti1−xAlxN films prepared by magnetron sputtering, Thin Solid Films 365 (2000) 104-109. https://doi.org/10.1016/S0040-6090(00)00653-2 T. Zhou, P. Nie, X. Cai, P.K. Chu, Influence of N2 partial pressure on mechanical properties of (Ti, Al)N films deposited by reactive magnetron sputtering, Vacuum 83 (2009) 1057-1059. https://doi.org/10.1016/j.vacuum.2009.01.001 F. Cai, S. Zhang, J. Li, Z. Chen, M. Li, L. Wang, Effect of nitrogen partial pressure on Al–Ti–N films deposited by arc ion plating, Applied Surface Science 258 (2011) 1819-1825. https://doi.org/10.1016/j.apsusc.2011.10.053 A. Knutsson, M.P. Johansson, L. Karlsson, M. Odén, Thermally enhanced mechanical properties of arc evaporated Ti0.34Al0.66N/TiN multilayer coatings, Journal of Applied Physics 108 (2010) 044312. https://doi.org/10.1063/1.3463422 L. Chen, Y. Du, X. Xiong, K.K. Chang, M.J. Wu, Improved properties of Ti-Al-N coating by multilayer structure, International Journal of Refractory Metals and Hard Materials 29 (2011) 681-685. https://doi.org/10.1016/j.ijrmhm.2011.05.001 W.C. Oliver, G.M. Pharr, An improved technique for determining hardness and elastic modulus using load and displacement sensing indentation experiments, Journal of Materials Research 7 (2011) 1564-1583. https://doi.org/10.1557/JMR.1992.1564 G. Stoney George, A. Parsons Charles, The tension of metallic films deposited by electrolysis, Proceedings of the Royal Society of London. Series A, 20

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42

[36]

[37]

[38]

[39]

[40]

[41]

[42]

[43]

[44]

[45]

[46]

Containing Papers of a Mathematical and Physical Character 82 (1909) 172-175. https://doi.org/10.1098/rspa.1909.0021 J.Y. Rauch, C. Rousselot, N. Martin, Structure and composition of TixAl1−xN thin films sputter deposited using a composite metallic target, Surface and Coatings Technology 157 (2002) 138-143. https://doi.org/10.1016/S0257-8972(02)00146-9 P.H. Mayrhofer, D. Music, J.M. Schneider, Influence of the Al distribution on the structure, elastic properties, and phase stability of supersaturated Ti1−xAlxN, Journal of Applied Physics 100 (2006) 094906. https://doi.org/10.1063/1.2360778 S.J. Suresha, S. Math, V. Jayaram, S.K. Biswas, Toughening through multilayering in TiN–AlTiN films, Philosophical Magazine 87 (2007) 2521-2539. https://doi.org/10.1080/14786430701230245 S. Zhang, D. Sun, Y. Fu, H. Du, Toughening of hard nanostructural thin films: a critical review, Surface and Coatings Technology 198 (2005) 2-8. https://doi.org/10.1016/j.surfcoat.2004.10.020 X. Chu, S.A. Barnett, Model of superlattice yield stress and hardness enhancements, Journal of Applied Physics 77 (1995) 4403-4411. https://doi.org/10.1063/1.359467 M. Stueber, H. Holleck, H. Leiste, K. Seemann, S. Ulrich, C. Ziebert, Concepts for the design of advanced nanoscale PVD multilayer protective thin films, Journal of Alloys and Compounds 483 (2009) 321-333. https://doi.org/10.1016/j.jallcom.2008.08.133 Y.X. Xu, L. Chen, F. Pei, K.K. Chang, Y. Du, Effect of the modulation ratio on the interface structure of TiAlN/TiN and TiAlN/ZrN multilayers: First-principles and experimental investigations, Acta Materialia 130 (2017) 281-288. https://doi.org/10.1016/j.actamat.2017.03.053 L. Chen, Y.X. Xu, Influence of interfacial structure on the mechanical and thermal properties of CrAlN/ZrN multilayer coatings, Materials & Design 106 (2016) 1-5. https://doi.org/10.1016/j.matdes.2016.05.082 X. Zhang, X.-B. Tian, Z.-W. Zhao, J.-B. Gao, Y.-W. Zhou, P. Gao, Y.-Y. Guo, Z. Lv, Evaluation of the adhesion and failure mechanism of the hard CrN coatings on different substrates, Surface and Coatings Technology 364 (2019) 135-143. https://doi.org/10.1016/j.surfcoat.2019.01.059 W. Li, P. Liu, X. Zhu, D. Pan, K. Zhang, F. Ma, X. Liu, Effect of Si content on microstructural evolution and superhardness effect of TiN/CrAlSiN nanomultilayered films, Journal of Alloys and Compounds 650 (2015) 592-597. https://doi.org/10.1016/j.jallcom.2015.08.012 W. Li, P. Liu, X. Zhu, K. Zhang, F. Ma, X. Liu, X. Chen, D. He, Si content dependent microstructure and mechanical properties of CrN/TiSiN nanomultilayered films, Materials Science and Engineering: A 610 (2014) 28-32. https://doi.org/10.1016/j.msea.2014.05.030 21

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35

[47]

[48]

[49]

[50]

[51]

[52]

[53]

[54]

[55]

Q. Wan, B. Yang, H.D. Liu, J. Chen, J. Zhang, Microstructure and adhesion of MeN/TiSiN (Me = Ti, Cr, Zr, Mo, NbxAl1−x) multilayered coatings deposited by cathodic arc ion plating, Applied Surface Science (2019) 143602. https://doi.org/10.1016/j.apsusc.2019.143602 W.M. Seidl, M. Bartosik, S. Kolozsvári, H. Bolvardi, P.H. Mayrhofer, Influence of coating thickness and substrate on stresses and mechanical properties of (Ti, Al, Ta)N/(Al, Cr)N multilayers, Surface and Coatings Technology 347 (2018) 92-98. https://doi.org/10.1016/j.surfcoat.2018.04.060 H. Ljungcrantz, C. Engström, L. Hultman, M. Olsson, X. Chu, M.S. Wong, W.D. Sproul, Nanoindentation hardness, abrasive wear, and microstructure of TiN/NbN polycrystalline nanostructured multilayer films grown by reactive magnetron sputtering, Journal of Vacuum Science & Technology A 16 (1998) 3104-3113. https://doi.org/10.1116/1.581466 H. Riedl, E. Aschauer, C.M. Koller, P. Polcik, M. Arndt, P.H. Mayrhofer, Ti-Al-N/Mo-Si-B multilayers: An architectural arrangement for high temperature oxidation resistant hard coatings, Surface and Coatings Technology 328 (2017) 80-88. https://doi.org/10.1016/j.surfcoat.2017.08.032 Q. Zhang, Y. Xu, T. Zhang, Z. Wu, Q. Wang, Tribological properties, oxidation resistance and turning performance of AlTiN/AlCrSiN multilayer coatings by arc ion plating, Surface and Coatings Technology 356 (2018) 1-10. https://doi.org/10.1016/j.surfcoat.2018.09.027 A. Liu, J. Deng, H. Cui, Y. Chen, J. Zhao, Friction and wear properties of TiN, TiAlN, AlTiN and CrAlN PVD nitride coatings, International Journal of Refractory Metals and Hard Materials 31 (2012) 82-88. https://doi.org/10.1016/j.ijrmhm.2011.09.010 J.L. Mo, M.H. Zhu, Tribological oxidation behaviour of PVD hard coatings, Tribology International 42 (2009) 1758-1764. https://doi.org/10.1016/j.triboint.2009.04.026 J.H. Hsieh, A.L.K. Tan, X.T. Zeng, Oxidation and wear behaviors of Ti-based thin films, Surface and Coatings Technology 201 (2006) 4094-4098. https://doi.org/10.1016/j.surfcoat.2006.08.026 Y.-Y. Chang, C.-P. Chang, D.-Y. Wang, S.-M. Yang, W. Wu, High temperature oxidation resistance of CrAlSiN coatings synthesized by a cathodic arc deposition process, Journal of Alloys and Compounds 461 (2008) 336-341. https://doi.org/10.1016/j.jallcom.2007.06.084

36

22

1

Figure captions

2

Fig. 1 The deposition schematic of the setup for hybrid AlTiN coatings including the

3

relative position of targets. Substrate specimens kept a two-fold rotation both with a

4

rotation rate of 3 rpm during deposition. For the deposition of single coatings, the AlTi

5

target in position B or C was used for CAE or MS with GIS, respectively. At this

6

situation, the substrates were maintained individually secondary rotation in the front of

7

the corresponding target.

8 9 10

Fig. 2 SEM fracture cross-sectional morphologies of (a) A-AlTiN, (b – e) H-AlTiN with different PMS, and (f) S-AlTiN.

11 12

Fig. 3 GIXRD patterns of A-AlTiN coating, H-AlTiN coatings under various PMS and

13

S-AlTiN coating on cemented carbide substrates.

14 15

Fig. 4 (a) The panorama TEM bright-field image and selective area electron diffraction

16

pattern and (b) lattice resolved HRTEM micrograph of the H-AlTiN coating with PMS =

17

7.0 kW. (c) STEM HAADF image of this coating indicating multilayered structure with

18

a modulation period of ~9.0 nm.

19 20

Fig. 5 SEM top view image of the surface from (a) A-AlTiN, (b) H-AlTiN (PMS = 6.0

21

kW), (c) H-AlTiN (PMS = 7.0 kW), (d) H-AlTiN (PMS = 8.0 kW), (e) H-AlTiN (PMS = 23

1

9.0 kW), and (f) S-AlTiN coatings on the cemented carbide blocks after tested by

2

micro-indentation instrument with a Vickers tip.

3 4

Fig. 6 (a) Indentation hardness and elastic modulus, (b) adhesion force and residual

5

stress of A-AlTiN coating, H-AlTiN coatings under various PMS and S-AlTiN coating.

6 7

Fig. 7 Friction coefficient curves of A-AlTiN coating, H-AlTiN coatings under various

8

PMS and S-AlTiN coating as a function of sliding distance at (a) RT and (b) 800 °C.

9 10

Fig. 8 Wear rates of A-AlTiN, H-AlTiN, and S-AlTiN coatings after ball-on-disc tests at

11

RT and 800 °C.

12 13

Fig. 9 SEM secondary electron images of wear tracks for (a) A-AlTiN, (b) H-AlTiN

14

(PMS = 7.0 kW) and (c) S-AlTiN coatings after tribological tests at RT.

15 16

Fig. 10 SEM secondary electron images of wear tracks for (a) A-AlTiN, (b) H-AlTiN

17

(PMS = 7.0 kW) and (c) S-AlTiN coatings after tribological tests at 800 °C.

24

Tables and Table captions Table 1 Primary deposition conditions and chemical compositions detected by EDX for AlTiN coatings investigated. Coating

PMS /kW

ICAE /A

A-AlTiN

N/A 6.0 7.0 8.0 9.0 7.0

80

H-AlTiN

S-AlTiN

Elemental contents /at.% Al Ti N

Total Thickness /µm ~5.3 ~5.6 ~5.8 ~6.1 ~6.5 ~5.7

80

N/A

31.93 32.09 30.35 30.38 32.17 33.47

18.80 17.36 16.79 15.35 18.23 13.83

Al/(Al+Ti)

49.27 50.64 53.98 54.27 53.99 52.70

0.63 0.65 0.65 0.66 0.64 0.71

Table 2 EDX elemental contents of framed regions in Fig. 9. Coating

A-AlTiN

H-AlTiN (PMS = 7 kW)

S-AlTiN

Elemental contents /at.%

Framed regions

O

Al

Ti

N

C

W

Co

1

3.51

30.31

16.34

49.84

--

--

--

2 3 4 5 6 7

65.67 18.01 2.56 67.53 9.21 5.06

21.61 26.66 29.95 20.41 28.03 32.53

9.84 13.73 15.98 9.86 14.07 13.50

2.88 41.60 51.51 2.20 48.69 48.91

-------

-------

-------

8

9.41

2.45

17.37

--

45.00

21.69

4.09

9

40.91

24.66

8.40

20.57

5.21

0.26

--

Table 3 EDX elemental contents of framed regions in Fig. 10. Coating

A-AlTiN

H-AlTiN (PMS = 7 kW)

S-AlTiN

Framed regions

Elemental contents /at.% O

Al

Ti

N

C --

W --

Co ------8.10 --

1

14.84

28.57

14.73

41.86

2 3 4 5 6 7

40.71 46.66 18.35 29.70 11.91 49.78

25.88 25.91 27.42 25.43 28.70 27.56

13.02 11.99 15.29 14.48 15.16 7.34

20.39 15.44 38.93 30.39 44.24 15.31

-------

-------

8 9

60.66 58.41

3.05 26.76

8.06 8.85

---

8.54 4.98

11.59 1.01

i AlT

sputtered sublayer

GIS

arc-evaporated sublayer

3 rp

m

(p

A-AlTiN interlayer

rim

ar

Substrate

B_CAE

AlTi

AlTi

(s ec

A_CAE

y)

3 rpm

PMS

ond ary)

two-fold rotation A-AlTiN interlayer Substrate

Enlarged

(a)

2 1

3

50 µm

Enlarged

(b)

50 µm

4

7 5

200 µm

(c)

200 µm

6 200 µm

8

9

(a)

(b)

2 µm

(c)

2 µm

(d)

2 µm

(e)

2 µm

(f)

2 µm

2 µm

c -T iN ;

c -A lN ;

w -A lN

3 0

4 0

P

M S

= 9 k W

P

M S

= 8 k W

P

M S

= 7 k W

P

M S

= 6 k W

5 0

T w o -th e ta a n g le (d e g .)

H -A lT iN

In te n s ity (a r b . u n it)

S -A lT iN

6 0

7 0

A -A lT iN

(a)

(c)

(b)

18 nm

SAED

[200

220

Enlarged

]

2.4 6.6 nm

[011]

c-AlTiN

30 nm

200 nm

[_ [1 1_ 11 ] 111

25 n

m

20 nm

n gr ow th di re ct io

200

5 1/nm

311

c-AlTiN column

10 nm

200 nm

column 20 nm

(a)

(c)

(b)

5 µm

5 µm (d)

(f)

(e)

5 µm

5 µm

5 µm

5 µm

8 5 0

(a ) H

7 5 0

E

3 5

3 0

5 5 0 2 5

4 5 0

2 0

E la s tic m o d u lu s (G P a )

6 5 0

3 5 0

1 5

A -A lT iN

-1 4

6

7

8

9

T a r g e t p o w e r o f m a g n e tr o n s p u tte r in g (k W )

2 5 0

S -A lT iN

1 2 0

-1 2

(b )

1 0 0

-1 0 -8

8 0

-6 -4

s

-2 0

6 0

2

L c 2 A -A lT iN

6

7

8

9

T a r g e t p o w e r o f m a g n e tr o n s p u tte r in g (k W )

S -A lT iN

1 3

4 0

A d h e s io n s tr e n g th (N )

H a rd n e s s (G P a )

R e s id u a l s tr e s s (G P a )

4 0

1 .0

(a ) a t R T

F r ic tio n c o e ffic ie n t

0 .8

0 .6

0 .4 A -A H -A H -A H -A H -A S -A

0 .2

0 .0 0

1 .0

2 0 0 0

4 0 0 0

6 0 0 0

8 0 0 0

1 0 0 0 0

S lid in g la p s

lT lT lT lT lT lT

iN

iN 6 iN 7 iN 8 iN 9 iN

.0 k .0 k .0 k .0 k

1 2 0 0 0

W W W W 1 4 0 0 0

(b ) a t 8 0 0 ° C

F r ic tio n c o e ffic ie n t

0 .8

0 .6

0 .4

0 .2

0 .0 0

1 5 0 0

3 0 0 0

4 5 0 0

S lid in g la p s

6 0 0 0

A -A H -A H -A H -A H -A S -A

lT lT lT lT lT lT

7 5 0 0

iN iN iN iN iN iN

6 .0 7 .0 8 .0 9 .0

k W k W k W k W 9 0 0 0

Enlarged

(a)

1

2

3

50 µm

50 µm

4 200 µm

Enlarged

(b)

5

200 µm

6

(c)

7

200 µm

8

9




AlTiN nano-multilayers were deposited combining arc evaporation and magnetron sputtering.




Hybrid AlTiN coating with optimized sputtering power exhibited the highest hardness.




Moderate residual stress and excellent adhesion were achieved via hybrid deposition.




Hybrid AlTiN coatings obtained improved wear-resistant at 800 °C.