Microstructure and mechanical properties of cathodic arc ion-plated (Al,Ti)N coatings

Surface and Coatings Technology 163 – 164 (2003) 135–143

Microstructure and mechanical properties of cathodic arc ion-plated (Al,Ti)N coatings K. Satoa,*, N. Ichimiyab, A. Kondoa, Y. Tanakaa a

MMC Kobelco Tool Co., Ltd., 179-1, Nishioike, Kanagasaki, Uozumicho, Akashi, Hyogo 674-0971, Japan b Mitsubushi Materials Corporation, 1511, Furumagi, Ishige-machi, Yuuki-gun, Ibaraki 300-2795, Japan

Abstract Cathodic arc ion-plated (Al,Ti)N coatings were deposited on WC–Co substrates by several bias voltages. The residual stress of the films, measured by the X-ray diffraction (XRD) 2u–sin2 c method, was found to vary significantly with the negative bias voltages. The composition and microstructure of the films were investigated by electron probe microanalysis and transmission electron microscopy, respectively, and these were also changed by the bias voltages. The changes in composition and microstructure were related to the film strains induced by ion bombardment, and were evaluated by the peak shifts and half-value widths of XRD profiles. At higher bias voltages, the films were mainly composed of granular structures with a size of 100 nm. With the decrease of negative substrate bias voltage, the microstructure changed to a columnar structure with a column size of 120 nm. Carbide drills deposited with (Al,Ti)N coating with lower bias voltages showed better cutting performance in the machining of carbon steel compared with those deposited with coating under high bias voltages. The wear resistance of drill margins is critical for the drilling performance in this test, and the dense columnar (Al,Ti)N films with low strains showed greatly improved wearresistant characteristics of the drills. 䊚 2002 Elsevier Science B.V. All rights reserved. Keywords: Microstructure; (Al,Ti)N; X-ray diffraction; Residual stress; Cutting performance

1. Introduction Over the last decade, (Ti,Al)N coatings with high hardness and high oxidation resistance have been developed for use in cutting tools that meet the demands for high efficiency machining fabrication w1–6x. (Ti,Al)N has a B1 (NaCl-type) structure in which Ti atoms of the TiN lattice are partially substituted by Al atoms. The coating structure becomes deformed and strengthened with the incorporation of Al atoms and this structural distortion is considered to cause the high hardness. One reason is that the size of Al atoms are smaller than that of Ti atoms w7,8x. The high oxidation resistance is caused by the formation of dense Al2O3 films at the coating surface during exposure to high temperatures w4x. *Corresponding author. Tel.: q81-78-936-7405; fax: q81-78-9355673; http:yywww.kobelcotool.co.jp. E-mail address: [email protected] (K. Sato).

Many investigations have been performed to improve the cutting performances of (Ti,Al)N coatings. Optimization of the Al contents is one of the promising approaches. The cathodic arc ion-plated (Al,Ti)N coating, which has a higher Al content of 60 at.% than that of (Ti,Al)N coatings, exhibits higher hardness of 29 GPa and higher oxidation temperature of 800 8C w9,10x. (Al,Ti)N-coated carbide endmills demonstrate improved cutting performance in the high speed machining of hard materials w11,12x. To further improve the cutting performance of (Al,Ti)N-coated tools, it is important to consider the relationship between the microstructure and properties of the film. Generally, cathodic arc ion-plated coatings have compressive residual stress that is generated by the energy of ions accelerated by bias voltage. And although this residual stress has a large impact on the film properties w13x, there have been few investigations into the microstructural properties of (Al,Ti)N films.

0257-8972/03/$ - see front matter 䊚 2002 Elsevier Science B.V. All rights reserved. PII: S 0 2 5 7 - 8 9 7 2 Ž 0 2 . 0 0 6 1 0 - 2

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Fig. 1. Schematic illustration of the cathodic arc ion plating system.

The aim of this paper is to investigate the properties of (Al,Ti)N films deposited by various negative bias voltages, and to discuss the relationships between the film characteristics and cutting performance. 2. Experimental procedures 2.1. Deposition conditions (Al,Ti)N coatings were prepared using the cathodic arc ion-plating system schematically illustrated in Fig. 1. The arc discharge was carried out on a powder metallurgy target of Al–40 at.% Ti composition. The substrate was negatively biased and the ionized vapor generated from target materials was accelerated and deposited to the substrate. The properties of the deposited films could be adjusted by changing the negative bias voltage. M30 WC–Co carbide inserts and JIS Z10 WC–Co carbide drills were ultrasonically cleaned in alkali and organic solvents and used as substrates. Prior to the deposition, the chamber was evacuated to a pressure of 4=10y3 Pa, and the substrate was heated to 450 8C by an electron radiation heater. The substrate was then sputtered and cleaned by ion bombardment to attain sufficient adhesion of the films. The deposition of (Al,Ti)N was carried out in a nitrogen atmosphere with a pressure of 4 Pa, and different negative bias voltages (VB) between 0 and y200 V were applied. Sample films of 4 mm thickness were achieved by controlling the deposition time under a constant arc current of 100 A.

area of ⭋ 30 mm using a JEOL JXA-8800RL electronprobe X-ray analyzer with an acceleration voltage of 10 kV. The strain of the films introduced during the deposition was evaluated by the peak shifts and half-value widths of X-ray diffraction (XRD) profiles using a MAC SCIENCE MXP-18HF system with Cu Ka radiation (ls0.1541 nm). The microstructure observation of films was carried out by transmission electron microscopy (TEM) using a JEOL JEM-2010F TEM with an acceleration voltage of 200 kV. Sample preparations for TEM were performed by the focused ion beam technique using a JEOL JFIB-2100 system. By this technique, samples were thinned to approximately 0.1 mm by Ga ions with the acceleration voltage and current of the 30 kV and 2.0 mA, respectively. The hardness and Young’s modulus of the films were measured using a CSEM nano-hardness tester. An increasing load was applied to the films to a maximum value of 10 mN. The load was then reduced until relaxation of the films was achieved. Load–displacement curves were monitored and the Young’s modulus and hardness were calculated. The residual stress of the films was measured by the 2u–sin2 c method of XRD. This method was performed by measuring the diffraction angle 2u in the c tilt axis against the direction of stress measurement and generating 2u–sin2 c diagrams. The stress s value was calculated as ssy

E p ≠(2u) cotu0 2(1qn) 180 ≠Žsin2c.

where ≠(2u)y≠(sin2c) is the gradient of the 2u–sin2 c diagrams that is approached to linear relationship. u0, E and n are the standard Bragg diffraction angle, Young’s modulus and Poisson ratio of the films, respectively. In this study, the Young’s modulus was measured and the Poisson ratio was taken as V?.20, which is the value used for TiN w14x. 2.3. Cutting performances Cutting performance was investigated by a drilling test using carbide drills with a diameter of 6.8 mm using carbon steel, AISI 1049 (HB200), as a work material. Drills were coated with (Al,Ti)N at a thickness of 4 mm with negative bias voltages of 20, 50 and 100 V. Table 1 gives the details of the cutting conditions. The tool life and wear conditions of (Al,Ti)N-coated carbide drills, examined by optical microscopy, were evaluated. 3. Results

2.2. Characterization of the films

3.1. Compositions of the films

The compositions of the films were measured by electron probe microanalysis (EPMA) with an analysis

Fig. 2 shows the Aly(AlqTi) ratio in the (Al,Ti)N films as a function of negative bias voltage. Cemented

K. Sato et al. / Surface and Coatings Technology 163 – 164 (2003) 135–143 Table 1 Cutting conditions for coated drills Drill size Work material Cutting speed (mymin) Feed rate (mmyrevolution) Depth of cut (mm) Overhang (mm) Coolant

f6.8 AISI 1049 (HB200) 80 0.22 16 (through) 50 Emulsion

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carbide inserts, ISO SNGN120408, were chosen as the substrate, and the measurement was carried out with the inserts in two different positions to investigate the influence of geometry on the composition. The measuring points were the center and corner of the rake face, respectively, as shown schematically in Fig. 2. With smaller VB of less than 20 V, the Aly(AlqTi) ratios were similar to the target composition and almost identical between the center and corner of the rake face. With increasing VB, the Aly(AlqTi) ratios decreased at both positions. Compared with that at the center, the Aly(AlqTi) ratio at the corner edge decreased significantly with an increase of VB to more than 20 V. As VB increased, the appearance of the inserts was also changed, and the color of the films was rendered brighter due to the decrease in Al content. 3.2. XRD analysis

Fig. 2. Aly(AlqTi) ratio in the (Al,Ti)N films as a function of negative bias voltage.

Fig. 3 shows the XRD profiles of the (Al,Ti)N films. All films were indexed as B1 (NaCl-type) structures, and the peaks presenting hexagonal (Al,Ti)N or AlN were not observed. (1 1 1), (2 0 0) and (2 2 0) peaks were apparent in all films, and (3 1 1) and (2 2 2) peaks were confirmed only in the films with VB of 0 and 10 V, as shown in Fig. 3. Fig. 4 shows the details of the (1 1 1), (2 0 0) and (2 2 0) peaks, where dot marks indicate the corresponding peaks and the scale of intensity is the same in each peak. The results show that the peaks are broadened and shifted to a low diffraction angle with the increase of VB. Fig. 5 shows the half-value width of each peak as a

Fig. 3. XRD profiles of the (Al,Ti)N films at various negative bias voltage.

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Fig. 4. (1 1 1), (2 0 0) and (2 2 0) peaks in XRD profiles at various negative bias voltage.

function of negative bias voltage. The peaks, which were overlapped with other peaks such as WC and Co, were isolated and calculated as independent (Al,Ti)N film peaks. Peak broadening is represented by the change of half-value widths, and the results clearly indicate that the peaks broadened with increasing VB. Fig. 6 shows the change of lattice constants as a function of negative bias voltage. The lattice constants were calculated by the position of each peak, and varied from 0.414 to 0.420 with increasing VB.

3.3. Residual stress and hardness

Fig. 5. Half-value width of each peak as a function of negative bias voltage.

Fig. 6. Lattice constant of the (Al,Ti)N films as a function of negative bias voltage.

In general, the 2u–sin2 c method of XRD is suitable for residual stress measurement in isotropic materials. When applied to non-isotropic materials, however, the relationship between 2u and sin2 c deviates from a linear function, causing large measurement errors. To

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increasing VB of up to 100 V. Further increase of VB caused a decrease in hardness. The behaviors of film hardness and Young’s modulus changes were symmetrical to that of residual stress. 3.4. Microstructure Cross-sectional TEM micrographs of the (Al,Ti)N films deposited with VB of 20, 50 and 100 V are shown in Fig. 9. All films had few defects, which is typical for cathodic arc ion plating, since this method generates relatively high ion energy bombardment during deposition, even with low substrate bias. At the higher VB of 100 V, the film was composed of granular structures with an average grain size of 100 nm. On the other hand, at the lower VB values of 20 and 50 V, the films had columnar structures with an average column size of 120 nm. These results indicate that the microstructural evolution of the films changed drastically with the substrate bias voltage, while there were no remarkable differences between grain size and column sizes. Fig. 7. Residual stress of the (Al,Ti)N films as a function of negative bias voltage.

minimize such error, it is necessary to choose peaks in the XRD profiles which: 1. exhibit few orientations compared with the standard diffraction pattern; and 2. are isolated and show high intensity against the background, and exist at a high diffraction angle, preferably 2u of more than 1008. In this study, the XRD profiles of the films (Fig. 2) were strongly oriented to the (1 1 1) or (2 0 0) direction compared with the standard diffraction pattern, and few peaks at higher diffraction angle were observed. For this reason, (2 2 0) peaks were selected for the calculation of the residual stress. In order to detect the small orientation of (2 2 0) compared with (1 1 1) and (2 0 0), the rocking curve technique was employed. Cr Ka (ls0.2291), which has a longer wavelength than Cu Ka was used to locate the (2 2 0) peaks at a high diffraction angle. Fig. 7 shows the relationship between residual stress and negative bias voltage. Each value contained approximately "10% measurement error, and at VB of 0 or 10 V the films showed tensile stress. In contrast, with VB of 20, 50, 100 and 200 V, compressive stress was observed. The values of compressive residual stress increased to 5.6 GPa with VB of 100 V, then decreased to 3 GPa with VB of 200 V. A nano-hard tester was used to measure film hardness and Young’s modulus. The results are shown in Fig. 8. Hardness increased to a maximum value of 31 GPa with

3.5. Cutting performance Fig. 10 shows the number of holes drilled by the (Al,Ti)N-coated carbide drills deposited with VB of 20, 50 and 100 V. The drills deposited with smaller VB showed the better cutting performance, with those coated at 20 V having an improved tool life of 7392 drill holes, corresponding to a cutting length of 118 m. In this cutting test, the wear at the drill margin has a pronounced and critical effect on the tool life. Fig. 11

Fig. 8. Hardness and Young’s modulus of the (Al,Ti)N films as a function of negative bias voltage.

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4. Discussion In the present study, (Al,Ti)N films deposited using negative bias voltages ranging from 0 to 200 V exhibited a variety of characteristics, even though they were prepared by discharging the same target material (Al– 40 at.% Ti). The Al content of the films decreased with the increase of negative bias voltage. With VB of 0 V, the ions generated from the target were not accelerated, and they accumulated on the substrate only with the ion energy caused by the discharge. With increasing VB, the ionized materials were accelerated and deposited with higher energy. Ti ions and Al ions are ionized and deposited separately in the cathodic arc discharge, and Al ions are considered to be resputtered more than Ti ions during the deposition, since the sputtering ratio of Al is about twice as large as that of Ti. With the higher VB, the difference between the number of resputtered Al and Ti ions was therefore increased and the Al content of the films was more decreased. From a geometric point of view, this phenomenon was significant at the corner edge than the center of the insert (Fig. 2), probably because the resputtering effect was more distinct at the corner edge, where more ions were focused during the deposition. All of the (Al,Ti)N films deposited in this study showed B1 (NaCl-type) structures by XRD analysis. The shifting of peaks to low angle (Fig. 4) indicated an increase in the interplaner spacing within crystals. This increase is quantitatively illustrated in Fig. 6, where the lattice constant, calculated by the peak position, was

Fig. 9. Cross-sectional TEM micrographs of the (Al,Ti)N films at the negative bias voltage of (a) 20 V; (b) 50 V and (c) 100 V.

shows the optical micrographs of wear conditions at the drill margin after 3696 holes. The bright area of the drill margins corresponds to the absence of (Al,Ti)N coating. The film deposited with VB of 100 V was wellworn, and a large area of substrate was exposed at the drill margin. The film coated with VB of 50 V, on the other hand, exhibited small wear and very little exposure of substrate. And the film coated at 20 V showed even greater wear resistance. As these results clearly demonstrate, the drilling performance was much improved by the significant wear resistance of (Al,Ti)N coatings at the drill margin, and this property was strongly affected by the application of negative bias voltage during the deposition.

Fig. 10. Number of holes at the tool lives of the (Al,Ti)N coated drills with negative bias voltages of 20, 50 and 100 V.

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the intrinsic value of the XRD apparatus. The half-value width caused by crystal size Wc can be calculated by Scherrer’s equation: Wcs

l Dcosu

where l is the X-ray wavelength (Cu Ka: 0.1541 nm) and D is the crystal size which can be measured by TEM. Assumption of the half-value width by inhomogeneous strain Ws follows Gauss distribution according to the following equation w16x: Wss

ŽWmyW.qŽWm2 yWmW.1y2 2

where Wm is the half-value width measured in XRD profiles, and W is a half-value width of crystal size in addition to intrinsic value of XRD apparatus defined by WsŽWc2qWi2.1y2 Here, Wi is the intrinsic half-value width of the XRD apparatus. After calculating Wc using the equation above, the inhomogeneous strain (Ddyd) was estimated using the differentiated Bragg’s equation: B Dd E

Wcsy2C D

d

Ftanu G

Fig. 12 shows the inhomogeneous strain (Ddyd) of each peak as a function of negative bias voltage. In each peak, the inhomogeneous strain increased with increasing VB up to 100 V, then decreased with further

Fig. 11. Optical micrographs of wear condition at the drill margin of (a) 20 V; (b) 50 V and (c) 100 V at the drilling of 3696.

increased with increasing VB. Similar to the compositional change, this change in lattice constant was caused by preferential resputtering of Al from the lattice site during deposition. The ion radius of Ti is larger than that of Al, and by replacing the Al site by Ti, the lattice was expanded and the lattice constant increased. In addition, at the high negative bias voltage, high implanted ion energy in the plane of the film will increase the lattice spacing parallel to the surface according the Poisson’s effect w15x. The strains can be predicted by the broadening of individual peaks, which is determined by measuring the half-value widths of the XRD profiles. The peak broadening is caused by the implantation of highly accelerated ions through the application of negative bias voltage. The half-value widths of the XRD profiles are influenced by inhomogeneous strain in the films, crystal size, and

Fig. 12. Inhomogeneous strain (Ddyd) of each peak in XRD profiles as a function of negative bias voltage.

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increase in VB. This means that the implanted ions make strains inhomogeneously with the increase of VB up to 100 V, and with the further increase of VB, the replacement of the Al site by Ti in the crystals yielded a relaxation of inhomogeneous strain. Moreover, high VB with the high ion energy raises temperature of the substrate and increase the mobility of the atoms w17x. So, especially at VB of 200 V, the inhomogeneous strains decrease significantly. The residual stress, measured by the 2u–sin2 c method of XRD, includes the thermal stress caused by a mismatch in the coefficient of thermal expansion (CTE) between the films and substrates. The CTE of the WC– Co substrate is 5.5=10y6 Ky1, and that of (Al,Ti)N is estimated as 7.2=10y6 Ky1 using a linear relationship between the CTE of TiN, aTiNs9.4=10y6 Ky1, and AlN, aAlNs5.7=10y6 Ky1, respectively w18x. The thermal stress sth can be calculated as sths

Ef Žafyas.ŽTsyTa. 1yn

where Ef, nf and af are the Young’s modulus, Poisson’s ratio and CTE of the film, respectively w19x, and as and Ts are the CTE and deposition temperature of the substrate, respectively. The thermal stress calculated with this equation is 0.38 GPa in tensile mode, and this value was in good correspondence with the measured residual stress (0.32 GPa) of the (Al,Ti)N film deposited with VB of 0 V (Fig. 7). With VB of 0 V, ions generated from the target materials were accumulated on the substrate without being accelerated by the bias voltage. Therefore, at 0 V there was little intrinsic and thermal stress caused by the mismatches of CTE between the film and the substrate. With further increase of negative bias voltage, the compressive stress caused by accelerated ions was increased. At VB of 100 V, the compressive stress reached a maximum value of 5.6 GPa, mainly due to the high energy ion bombardment accelerated by the negative bias voltage. It is also considered that the difference of ion radius between Al and Ti may have added structural strain to the films. At 200 V, the compressive residual stress was lower than that at 100 V, and this result is considered that the replacement of Al sites by Ti in the crystals and the increased mobility of atoms at raised temperature accumulated strain relaxation, whose amount is larger than the multiplication of stresses by ion bombardment. In this study, the residual stresses in the (Al,Ti)N films deposited with various negative bias voltages were affected by three main factors: 1. Thermal stresses caused by CTE mismatches between the films and the substrates; 2. Multiplication of strains by high ion bombardment energies; and

3. Stain relaxation caused by the replacement of Al sites by Ti, and the increased mobility of atoms in the crystals. With VB of 0–10 V, factor (1) was dominant and tensile stress was generated. With VB of 20–100 V, factor (2) created high compressive stress, which increased almost linearly as a function of negative bias voltage. And at VB of 200 V, factor (3) became dominant and the compressive residual stress was reduced. The hardness of the films was also changed by the negative bias voltages, and this trend was quite symmetrical to that of the residual stresses. In this study, film hardness was defined as the deformation resistance of film materials against indentation. Since the residual stress was the main factor preventing deformation, higher residual stresses resulted in greater hardness. The microstructures of the films were also changed by negative bias voltages. Although the grain and column sizes were nearly identical, the microstructure evolution changed drastically at different bias voltages, as shown in the TEM micrographs in Fig. 10. With VB of 100 V, a granular structure was dominant and there were many bright and dark fields, indicating the random orientation of the grains. It is considered that the high bias voltage introduced a large strain that fragmented the grain growth as well as the enough effect of resputtering. On the other hand, with VB of 20 or 50 V, the film showed a mainly columnar structure, and more individually developed grains were observed than in the granular structure at VB 100 V. With lower VB of 20 or 50 V, grain growth in the surface direction became dominant, since the grains were subjected to small strains in addition to the smaller amount of ion resputtering during the deposition, and there were few anisotropic crystal growths. The film deposited at 20 V showed the most dense and obvious columnar structure, running the full depth of the film from substrate to film surface. It should be noted that there were very few defects between columns in these films, since cathodic arc ion plating can generate ion energy sufficient to develop a dense microstructure even under condition of a small bias voltage to the substrate. The (Al,Ti)N-coated carbide drills deposited with lower VB showed better drilling performance. In the cutting test, there were no clear abnormal damages, such as chipping or peeling, and the differences in performance were mainly attributed to the wear-resistance characteristics of the coating. In this cutting test, wear at the drill margins is critical to the tool life, and the (Al,Ti)N films deposited at lower VB showed better wear resistance at the drill margins and therefore better cutting performance. In the case of coated tools for milling or other types of interrupt cutting, it has generally been considered that coatings with higher hardness and compressive stress will perform better. In the present study,

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however, the best drilling performance was demonstrated by the coating with lower hardness and lower compressive stress, i.e. the coating deposited at 20 V. In this context, it is important to note that drilling is performed with the margins tied down and kept in contact with the drilling holes, this sometimes causes greater wear at the drill margins rather than the flank faces. The dense columnar structure, which has less strain, is considered to be more suitable for resisting sliding wear of the drill margins, since the dense column-structured (Al,Ti)N films might wear stably and gradually from the top of the column surface. Compressive residual stresses are necessary in order for the coatings to resist cracking or chipping, but from a microstructural point of view, a dense columnar structure is very important for the wearresistant characteristics of drilling applications. 5. Conclusions Cathodic arc ion-plated (Al,Ti)N films were deposited with various negative bias voltages, and the relationship between the microstructure and mechanical properties was investigated. The results can be summarized as follows: 1. The Aly(AlqTi) ratio in (Al,Ti)N films decreased with increasing negative bias voltage, because Al ions are resputtered more than Ti ions during the deposition. 2. Inhomogeneous strains and residual stress were multiplied with increasing negative bias voltage of up to 100 V. A further increase of VB resulted in relaxation of strains due to the replacement of Al sites by Ti in the crystals and the increased mobility of atoms at the raised temperature by high VB. 3. The dense columnar-structured (Al,Ti)N film deposited with low negative bias voltage of 20 V exhibited

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