Modeling and Simulation for Material Selection and Mechanical Design Part 12 docx

and the workpiece under test) to the short-time tensile yield strength of the softer contact body at the test temperature The value of t is simply a measure of the resistance of the joint to shear The friction condition at the surface of a cutting tool will be similar to that for which the value of t was measured The results of the wear resistance tests are given in Fig 6 As can be seen, the wear resistance of HSS tools is 2.0–3.5 times lower than that of DCPM tools This reduction was associated with a significantly lower friction parameter of DCPM compared to HSS (Fig 7), and a broadening of the range of normal friction (Fig 6) Within the normal friction range, the rate of wear for DCPM is much lower than that for HSS (Fig 6, curves 1–3) While the hardness and heat resistance values of the HSS T15 and DCPM are similar, the wear resistance of the latter is significantly higher (Table 3) In our opinion the lower wear intensity of the DCPM-tool material is related to the presence of titanium carbides in the structure and their subsequent transformation to oxygen-rich compounds during cutting When studied by secondary ion mass spectroscopy (SIMS), the analysis of typical wear craters revealed the formation of oxygen-containing phases The data in Fig 8 demonstrate that the transformation of titanium carbide into an oxygen-containing phase starts in the initial stage of wear (during the running-in process, Fig 8a) With further operation, there is increased surface oxide formation at the bottom of the wear crater This process is accompanied by stabilization of the wear processes (Fig 6 and 8b,c) and an expansion of the normal friction range Evidently, this is determined by the phenomenon of self-organization that is connected with the emergence of secondary structures (titanium–oxygen compounds), which play the role of stable solid lubricants [25] Figure 6 The dependence of the flank wear value of cutting tools on the cutting time: (1) HSS M2; (2) HSS T15; (3) DCPM Turning test data acquired with 1040 steel Parameters of cutting: speed (m=min): 55; depth (mm): 0.5; feed (mm=rev): 0.28 Copyright 2004 by Marcel Dekker, Inc All Rights Reserved Figure 9 Microphotograph of tool friction surface with films of secondary structures: (a) general view of the surface using secondary electrons; (b) distribution of oxygen close to the ‘‘built-up-crater’’ contact surface (SI, intensity of signals, arbitrary units) Copyright 2004 by Marcel Dekker, Inc All Rights Reserved electron spectroscopy, are given in Fig 9(b) In the left part of the micrograph (a), a build-up of 1040 C steel can be seen The right part of the micrograph shows the distribution of dispersed hardening phases in the HSS-based DCPM Angular (dark) particles of titanium carbide (less than 8 mm in cross-section) as well as dispersed tungsten and molybdenum carbides (less than 0.2–1.5 mm in diameter) are uniformly distributed in the HSS matrix In the surface layers of the tool material, we can observe a zone of intense plastic deformation less than 5 mm in depth There, dispersed particles of a titanium-containing phase have been drawn out parallel to the wear surface, forming a discontinuous film The titanium carbides in the wear zone have been transformed into oxides (Fig 8 and [23]) Titanium oxides are known to be much more plastic than titanium carbides, accounting for the plastic deformation of the particles in the surface layers of the HSS-DCPM on cutting These results are confirmed by Auger-spectroscopy Fig 9(b) represents the distribution of the intensity of the characteristic Auger KLL lines for O, C, and the LMM (418 eV) line of Ti along the I–I direction in Fig 9(a) The analysis volume includes the built-up layer (of 1040 steel), the built-up layer=wear crater boundary, and the DCPM volumes beneath the wear crater At the interface, the titanium compounds show an increased concentration of oxygen and a decreased carbon content The observed change in chemical composition is related to the instability of titanium carbide Due to the high cutting temperatures (in excess of 4508C) and pronounced affinity for oxygen, titanium adsorbs the latter from the environment and forms thin films of oxygen-containing compounds, in agreement with the SIMS data presented in Fig 8 The total plastic deformation of these particles at the wear surface is greater than 600% The crystal structure of these compounds is believed to differ from the titanium oxides that would be obtained under equilibrium conditions (see below) An understanding of the self-organizing phenomenon is critically important for the development of advanced tool materials A major interest in these studies (from the point of view of materials science) is the nature of the secondary structures forming under severe cutting conditions According to the principles of current tribology, one of the main methods to control friction is the creation of stable secondary structures at the tool surface The more stable are the secondary structures, the greater will be the tool life The development of protective secondary structures can be manipulated by alloying or by surface treatment technologies The type of secondary structures formed during cutting depends strongly on the conditions of cutting and the type of the material under analysis A detailed study of the physico-chemical parameters of the SSs formed during cutting using a tool made of DCPM was done using AES, ELS, and Copyright 2004 by Marcel Dekker, Inc All Rights Reserved EELFS methods To interpret the atomic structure of the tool wear surface, data obtained in this work by the EELFAS method were compared to TiC and TiO2 standards Fig 10 presents the Fourier transform of data obtained on analyzing the extended electron energy loss fine structure (EXELFS) for titanium carbide (TiC) with a cubic (B1) structure The positions of the main peaks (Fig 10a) are consistent with the interatomic distances for a (1 0 0) plane in the cubic lattice of titanium carbide (see Fig 10b) Figure 10 (a) Fourier transform of EELFS close to the line of back-scattered electrons for TiC specimen, Ep¼1500 eV; (b) cubic lattice of titanium carbide (¼Ti; ¼ C)  Copyright 2004 by Marcel Dekker, Inc All Rights Reserved Fig 11(a) shows data for TiO2 with the rutile (C4) structure We can identify the type of bonds by using partial functions F(R) obtained from the analysis of the fine structure of spectra close to the characteristic Auger lines of oxygen and titanium By comparing these data with those given in Fig 9(a), we can see that TiO2 has a more complex crystalline structure than TiC This explains the greater number of F(R)-function peaks The positions of the main peaks are again in good agreement with the Figure 11 (a) Fourier transform of EELFS close to the line of back-scattered electrons for TiO2 specimen with rutile structure, Ep¼1500 eV; (b) cubic lattice of titanium oxide (¼Ti; ¼O)  Copyright 2004 by Marcel Dekker, Inc All Rights Reserved interatomic distances for a (1 0 0) plane in the TiO2 lattice The complete analysis of all the peak positions by the Fourier transform method shows that the interatomic distances O–O and Ti–O in the secondary structures are different from those discussed in the literature (see Fig 11b) This may be related to a deviation from stoichiometry, or the interatomic distances, measured from an analyzed volume that is only several angstroms thick near the surface, are different from the equilibrium values The evolution of the atomic structure in the surface films on the wear crater of the cutting tool is well illustrated by the data given in Fig 12 The oxygen-containing films in the wear crater are significantly enriched with titanium and oxygen after only 5 min of cutting (see Fig 12a–d) As this takes place, a periodicity in the arrangement of atoms of various types is observed both in the nearest coordination sphere and at greater interatomic ˚ distances, up to approximately 7 A (see Fig 12b) As noted above, the interatomic distances in these oxygen-containing films differ from those observed in equilibrium titanium oxides, including rutile (compare with Fig 11) The very thinnest films may be 2D (twodimensional) phases whose atomic structure is close to the supersaturated a-solid solutions of oxygen in titanium After 15 min of cutting, the degree of long-range order is reduced, while the intensities of peaks from higher order coordination spheres are less pronounced (see Fig 12c) After 30 min of cutting (Fig 12d), the translational symmetry at large interatomic ˚ distances disappears, and peaks at R > 4 A are lost The adaptability of the surface layer to external thermo-mechanical effects is the physical basis of such evolution The surface is gradually converted to an amorphous state during the wear process (after cutting times of about 15 min) When a steady-state condition is reached, i.e., after the development of the SS is completed, the surface generates amorphous-like films having an effective protective function The lattice instability of the solid solution of oxygen in titanium finally leads to complete amorphization of the water surface A similar effect was observed earlier from EELFS data for TiN-coatings on worn cutting tools [26] Typical EELS spectra of TiC, TiN, TiO2 obtained with a 30.0 eV primary electron beam are shown in Fig 13 The elastic peak has a 30 meV FWHM (full width at half maximum) The high-resolution structures of the spectra are represented at 1000 Â magnification after normalizing The experimental curves are approximated by Gaussian peaks in each spectrum in the range of 1–9 eV energy loss In the series of titanium compounds TiC–TiN–TiO, the number of 4s electrons in the atomic sphere of the metal decreases from the carbide to the oxide These electrons are transferred to the 2p orbital of the non-metal atom, and the band energy is lowered due to the increasing Ti–X attraction Copyright 2004 by Marcel Dekker, Inc All Rights Reserved this orbital is more completely filled in the case of oxygen than in either nitrogen or carbon For this reason, when the 3d Ti- and 2p X orbitals are hybridized, the contribution of the ds-electrons of Ti is less pronounced in the oxide and more expressed in the nitride and carbide Consequently, the oxide has considerably less strength and hardness than the carbide or nitride [27] The interaction of Ti–Ti atoms is realized at the expense of dp-electrons The metallic nature of this compound is related to the high density of dp-electrons As seen in Fig 13, the intensity of t1g ! t2g transitions in the p-band is relatively insignificant in TiC, but it is much higher in TiO2 This implies that the density of conduction electrons is low in TiO2 but it is higher for TiC and TiN This is consistent with the electrical conductivity data of these compounds, which is extremely low for the dielectric TiO2, but 16,400 (Ohm m)À1 for TiC and TiN, respectively [27] The replacement of carbon with oxygen in titanium compounds was shown to change their properties significantly Thus, the oxidation of TiC at 823 K for 30 min influences the electronic structure of the material, the electron spectrum acquiring some features specific to TiO2 (see Figs 14a Figure 14 Representative ELS spectra: (a) after oxidation of TiC by heating up to 823 K for 30 min in air; (b) wear surface of DCPM cutting tool after 5 min of operation; (c) wear surface of DCPM cutting tool after 30 min of operation The spectra were obtained using a 30.0 eV primary electron beam Copyright 2004 by Marcel Dekker, Inc All Rights Reserved and 13c) After oxidation of TiC, the intensity of the lines at 6.7 and 1.6 eV is substantially enhanced On oxidation, the titanium carbide loses its metallic properties and acquires those of a dielectric In this case, we observe a reduced concentration of conduction electrons and a localization of the electron density both in the metal and non-metal atoms This is shown by the increased intensity of the peak at 1.6 eV corresponding to p-states in the 3dd-band of titanium (see Fig 13a) It was noted earlier that an intense oxidation of TiC could be observed during the operation of a DCPM tool [20,22] In this case, the nature of the phase transformation differs significantly from that found on heating a TiC standard up to 823 K for 30 min Figure 14(b) and (c) presents electron energy loss spectra from the wear crater after 5 and 30 min of DCPM-tool operation As the wear time increases, the spectra display a somewhat increased intensity of peaks at 6.8 and 3.1 eV Peaks corresponding to plasmon losses (p1 and p2) appear, while the peak at 1.6 eV is significantly attenuated The thin SS films in the wear crater of the cutter are associated with the formation of supersaturated solid solution of oxygen in titanium due to the oxidation of titanium carbide In this case, we observe an increase in the electron density in the 2p orbital of the non-metal (peak 6.8 eV) as well as an enhanced filling of the ds-electron band of titanium atoms (peak 3.1 eV) These effects are similar to those encountered in the model oxidation of TiC (Fig 14) There are, however, substantial differences As the cutting time increases, the effects brought about by the crystalline structure of phases become significantly weakened in the electron spectrum The splitting of the 3d orbital into p and s-states degenerates, the intensity of t1g ! t2g transitions is reduced as well as the density of p-electrons which are related to the long-range Ti–Ti bonds in the lattice (along the diagonals in the (1 0 0) planes) These distinctive features of the electron structure are related to amorphization and to the increasing role of short-range interatomic bonds Of considerable interest is the appearance of plasmon loss peaks p1 and p2 in the spectra of Figs 14(b,c) due to the growing concentration of conduction electrons The delocalization of p-electrons close to the titanium atoms enhances the metallic nature of bonds in the amorphous films developed on the friction surface These specific traits of electronic and atomic structural change might help to explain the unique mechanical properties in the secondary structures of the first type The high wear resistance and good frictional properties of DCPM tools are associated with complex structural and phase transformations on the surface, among them TiC oxidation and the development of thin protective amorphous films The SSs are saturated or supersaturated (amorphous) solid solutions of oxygen in titanium, whose electron structure is characterized by a high density of conduction electrons giving metallic characteristics Copyright 2004 by Marcel Dekker, Inc All Rights Reserved These results show that the SSs formed during cutting not only increase the DCPM-tool life but also change friction characteristics as well The amorphous-like secondary structures of the first type behave like a solid lubricant with enhanced tribological properties [4] Additional alloying of the DCPM might be beneficial For example, the partial substitution of titanium carbide by aluminum oxide, which is stable under cutting leads to a decrease in the friction coefficient (Fig 15) Figure 15 Impact of the test temperature on the wear and friction characteristics as determined from wear contact tests for the DCPM with 15% TiC þ5% Al2O3; 20% TiCþ2% BN and 20% TiCN additions Copyright 2004 by Marcel Dekker, Inc All Rights Reserved and in an increase in the wear resistance of the tool (Fig 16) The decrease of the friction coefficient when Al2O3 is added is important not only as it increases the wear resistance but also because it lowers the cutting temperature at the tool surface [28,29] Alloying often cannot be implemented by the traditional metallurgical methods since this may induce an undesirable change in the properties of the cutting material We took a different approach by making small additions of low-density compounds, which are relatively unstable at the operational temperatures This allowed us to use this compound in relatively small quantities (up to 2 w%) with minimal possible impact on the bulk properties The solid lubricant (hexagonal BN) was chosen as the additional alloying compound [28] The high probability of oxygen-containing secondary phases formed from BN during cutting was also taken into account The possibility that TiC and BN might oxidize and generate thin surface oxide films for exploitation in cutting tools can be assessed by a thermodynamic approach [27] Secondary ion mass spectroscopy investigations have shown that on cutting DCPM with a boron nitride addition, oxygen-containing compounds develop at the wear-crater surface, associated with a set of parallel disassociation reactions of BC, BN, TiC leading to the formation of BO, TiO and TiB,N Figure 17a–c presents spectra of the positive and negative Figure 16 Wear curves of friction contact materials: (1) DCPM with 20% TiC; (2) DCPM with 20% TiC and 2% BN; (3) DCPM with 15% TiC and 5% Al2O3; (4) DCPM with 20% TiCN Turning test data acquired with 1040 steel Parameters of cutting: speed (m=min): 90; depth (mm): 0.5; feed (mm=rev): 028 Copyright 2004 by Marcel Dekker, Inc All Rights Reserved Figure 18 Change in absolute and relative values of the peak intensity of the secondary ion mass spectra for the BN-doped DCPM tool at different depths beneath the surface of the wear crater developed after 4 min cutting spectroscopy was done to demonstrate this phenomenon for the phases that form on the wear crater surfaces of M2 high-speed steel, DCPM and BNdoped DCPM tools Figure 19a–c show Fourier transforms obtained from data collected from the surface of wear craters in HSS and DCPM alloyed with either TiC or TiC with BN The F(R) functions feature pronounced ˚ peaks in the range 1–2, 4–5 and 7–8 A for all cases Using partial F(R) functions obtained by analysis of the fine structures close to the Auger lines of C, B, and Ti, it was possible to interpret the nearest-neighbor interatomic bonds It was found that Fe–O bonds are typical of the HSS sample, while B–O and Ti–B bonds are observed at the wear crater of the BN-doped DCPM sample The results of these investigations have shown that the composition of the tool material not only determines the composition of the phases developed at the friction surface in the cutting zone, but it also exerts an influence on the perfection of the crystalline structure of the new phases The thinnest films of iron oxide formed on the HSS-tool friction surface are crystalline, Copyright 2004 by Marcel Dekker, Inc All Rights Reserved the TiO family The amorphization of secondary structures probably depends on the DCPM composition and on increases in the level of BNalloying One can see that the wear resistance of this material is increased by 80% compared to the carbide steel having a base composition with 20% TiC (Fig 16) This suggests that alloying enhances the stability of the secondary structures developed during friction This is promoted by the presence of BO-type compounds that act as liquid lubricants [at elevated cutting temperatures [28]] and promote the stability and the self-organization of the complex compounds This is of paramount importance for tool wear resistance The thickness of the stable secondary structures layer does not exceed 0.1–0.15 mm (Fig 18) Finally, it is possible that such alloying of DCPM provides both a reduction in the friction coefficient and a broadening of the normal wear stage In our opinion, the same goal can also be achieved in HSS-based DCPM by the substitution of TiC with TiCN Figures 15 and 16 demonstrate that this substitution is extremely effective, decreasing the friction coefficient to abnormally low values (in the range of 0.03–0.05) at a service temperature of 500–550 8C, and significantly increasing the tool life (Fig 16) In this case, the self-organization mechanism differs somewhat from the process found for materials alloyed with BN With TiCN the diffusional transfer of nitrogen into the chip arises from dissociation of TiCN during cutting [26] The increased nitrogen concentration on the contact surface of the chip is a direct consequence of the mass transfer of nitrogen, formed by dissociation of nitrides and carbo-nitrides Such mass transfer takes place under the extreme temperature and stress conditions encountered in the friction zone Therefore, the selection of titanium carbo-nitride for cermets and titanium nitride as the hard phase in Sandvik Coronites seems to be completely reasonable from the standpoint of tribological compatibility The behavior of titanium nitride should be comparable to a carbo-nitride [24,30] and this has been confirmed experimentally by the analysis of the self-organizing phenomenon of PVD TiN coatings (see below) Another important factor is the thermal stability of DCPM compared to cemented carbides The importance of the thermal stability of this type of material is evident at elevated cutting speeds One way to improve this property is to increase the volume fracture of the hard phase in the DCPM; e.g., Sandvik Coronite has 50% of the hard phase (TiN) [30] and its wear resistance is higher than either HSS or cemented carbides (Fig 20) So the benefits of an adaptive material are obvious from the standpoint of both wear resistance and tribological compatibility However, the application of these materials is currently limited, although the problems encountered might be partially solved by improved surfaceengineering techniques For example, state-of-the-art coating technologies Copyright 2004 by Marcel Dekker, Inc All Rights Reserved with high strength and toughness of the core The intermediate layer has graded properties that lie between those of the core and the surface Promising developments in this field have been reported by the Laboratory of Materials Processing and Powder Metallurgy of the Helsinki University of Technology They developed a novel functionally graded material having a surface ceramic layer, a graded WC-cermet composition with high crack resistance and a cemented carbide core with excellent toughness A high crack resistance parameter value (K1C¼25 MPa m1=2) at a hardness of 1500 HV (typical for tool steels with half the hardness) was found (http:==www.hut.fi=Units=LMP=) There are two methods used for FGM processing The first one, noted above, is a surface-engineering method This method has unique possibilities and versatility But other methods, principally those based on powder metallurgy, are also widely used Combinations of these two methods have recently been put into practice Thus, tools manufactured by ordinary sintering processes, having high-toughness cemented carbide substrates with high wear resistant ceramic coatings and functionally graded interlayers show excellent wear resistance [31] Functionally graded materials can have superior wear resistance, resistance to fracture, and good thermal shock resistance in the comparison to conventional cermets, with a beneficial compressive residual stress distribution [32] Ceramics have also been recently developed with functionally graded structures In order to combine high hardness and high toughness, graded ceramics of Al2O3 þ TiC (surface)=Al2O3þTi (inner core) and sialon (inner core) have been successfully developed [33] Functionally graded ceramic tools can exhibit better cutting performance than regular ceramic tools [34] Functionally graded materials have also been successfully used for milling applications [35] Functionally graded powder materials are normally used for high-speed cutting, but they can also be successfully employed in the domain of HSS tools application particularly with functionally graded cement carbide and hard PVD coatings [36] III TRENDS IN THE DEVELOPMENT OF SURFACE-ENGINEERED TOOL MATERIALS Surface engineering has recently become one of the most effective ways of improving the wear resistance of tool materials The principal beneficial effects associated with surface engineering for this application are shown in Table 4 Copyright 2004 by Marcel Dekker, Inc All Rights Reserved to vary the parameters of deposition, using a regular PVD unit, to optimize the coating properties and coating design A study designed to optimize the deposition parameters for TiN coatings [38] determined that the nitrogen pressure is the most critical process parameter responsible for changes in the coating structure and properties of the films For the deposition conditions described in Ref [38], an increase in the nitrogen pressure up to 0.4–0.6 Pa leads to stoichiometric TiN (53 at.% N2) Further increases in the nitrogen pressure lead to a decrease in the nitrogen concentration of the film (to 43 at.% N2, Fig 21a) This is probably caused by a decreased intensity of the plasma-chemical reaction as a result of a reduction in the flux and energy of the impinging ions The phase composition changes from a-Ti þ Ti2N at very low nitrogen pressures to TiN at higher nitrogen pressures (0.4–0.6 Pa) The structural parameters of the film also depend on the nitrogen concentration in the coating The lattice parameter and subgrain size (or equivalently the dislocation density) are related to the nitrogen concentration in the coating (Fig 21b–d) and have maximum values at a composition corresponding to stoichiometric TiN An increase in the nitrogen pressure leads to a pronounced axial texture (Fig 21e), with (1 1 1) planes in the TiN being parallel to the substrate surface The (1 1 1) axial texture increases to 95% as the gas pressure is raised to 0.6 Pa and remains practically unchanged with subsequent pressure increases The residual compressive stress in the coating shows similar trends (Fig 21f), with the stress increasing from 200 to 1300 MPa as the gas pressure is raised to 0.6 Pa The rate of increase in the compressive stress slows with a further increase in the nitrogen pressure up to 2.6 Pa, reaching a maximum value of 1600 MPa It is important to realize that the trends shown by these structural-dependent parameters depend primarily on the deposition conditions An increase in the nitrogen pressure over 1.3 Pa decreases the ion energy, and the effective temperature at the TiN crystallization front decreases At a nitrogen pressure in excess of 1.3 Pa, the coatings form under conditions similar to those encountered in the balanced magnetron sputtering process (for the PVD method) These conditions can be indirectly characterized by the deposition rate, which should not exceed 5–6 mm hrÀ1 A decrease in the deposition rate enhances both the axial texture and the magnitude of the residual stresses The most probable cause of the high compressive residual stress found in thin condensed films deposited at nitrogen pressures greater than 1.3 Pa is a high density of point defects [39] An increase in the nitrogen pressure also decreases the crystallite dimensions (Fig 21f) The microhardness of the coatings depends on their phase composition The maximum microhardness (H0.5¼45 GPa) is achieved when the two-phase a-TiþTiN composition changes into the three-phase composition Copyright 2004 by Marcel Dekker, Inc All Rights Reserved Figure 22 pressure The dependence of the properties of TiN PVD coatings on the nitrogen from stoichiometry causes a decrease in the Palmquist toughness, particularly when a second phase is formed in the coating (e.g., Ti2N) This result seems unexpected at first, but we should remember that the coating should be regarded as a quasi-brittle material, whose toughness is determined primarily by its crack propagation resistance The optimum nitrogen concentration corresponds to the largest lattice parameter [40], although the Copyright 2004 by Marcel Dekker, Inc All Rights Reserved microhardness of the coating with this structure is not relatively high In addition to the intrinsic mechanical properties of the coating, the level of the residual compressive stress is important for crack initiation and propagation The value of the residual stress is about À800 MPa in stoichiometric TiN coatings (Fig 21f) The fracture resistance also appears to depend on the columnar grain size, which again can be controlled if balanced deposition conditions can be achieved The adhesion of the coating to the substrate and the shear load resistance (i.e the cohesion of the coating) both decrease as the nitrogen pressure is increased The nitrogen atoms (ions) in the plasma scatter the Ti ions, so that the net effect of Ti ion bombardment of the coating is reduced as the nitrogen gas pressure rises As noted earlier, both the axial texture and the residual stress gradient at the coating– substrate interface increase as the pressure rises, and consequently the adhesion of the coating falls (Fig 22b) The coating wear resistance during dry sliding friction under high loads (Fig 22c) reaches its maximum value in the three-phase field area a-TiþTiNþTi2N Cutters made of a high-speed steel with TiN coating are used under adhesive wear conditions [3] The tool life of TiN coated parts then depends on the friction coefficient under cutting conditions that are close to galling or seizure during wear testing (Fig 22, e) The overall conclusion of this study is that the optimum combination of properties of the coating for adhesion wear is obtained at low deposition rates (for the PVD method) of 5 mm hrÀ1 and a stoichiometric composition of TiN This can be achieved by optimizing the deposition parameters In this case, the hardness and toughness increase, while the shear resistance decreases A coating with the optimum structure will crack by shear failure at or near to the surface of the coating rather than forming deep cracks leading to a catastrophic failure of the whole tool However, at the same time, the shear stress resistance of the coating should be strong enough to allow for easy flow of the chip (the value of cohesion should be about 0.2) Since a monolithic TiN coating usually has a low adhesion to the substrate, adhesive sublayers are necessary to achieve high efficiency from this type of coating A number of principles guiding the selection of the processing of TiN multi-layer coatings for adhesive wear conditions can be proposed from this study The coating should have at least three sublayers: (1) An adhesion sublayer, deposited with substoichiometric nitrogen These deposition conditions provide the maximum kinetic energy of the ions and a low nitrogen concentration in the layer (up to 35%) At the same time the (1 1 1) axial texture should not exceed 50%, while the residual stress at the coating–substrate interface should be low (not Copyright 2004 by Marcel Dekker, Inc All Rights Reserved more than 200 MPa) When this combination is achieved, the adhesion of the sublayer is high (2) A transition layer deposited with a gradual increase in the nitrogen pressure to provide: (a) development of an axial (1 1 1) texture (from 48% up to 100%) from the substrate to the top layer; (b) a residual compressive stress increase from 200 MPa in the adhesion layer up to 1700 MPa; and (c) a gradual transition from a three-phase structure a-TiþTiNþTi2N to single phase TiN (3) A working (contact) layer, deposited under high nitrogen pressures and balanced conditions (the deposition rate for the CAPDP method is 5– 6.5 mm hrÀ1) In addition the deposition conditions at this stage should be chosen to yield stoichiometric TiN, having a nearly perfect axial texture, a high residual compressive stress (more than 2000 MPa), and a fine columnar grain size containing minimal ‘‘droplet’’ phases A multi-layer coating is required to meet these diverse requirements This coating offers many advantages (in comparison to monolithic coatings) in satisfying the broad range of mechanical properties needed in these applications It has high adhesion to the substrate but low adhesion to the workpiece (i.e a minimal friction coefficient), a high microhardness (H 0.5¼35 GPa) and a high toughness (more than 50 J mÀ2, see Table 5) This favorable blend of structural and mechanical properties has many advantages for wear resistance during cutting operations (Table 5) Table 5 Comparative Characteristics of TiN Coatings Deposited by the PVD Method on the Coating Deposition Process [38] Parameters PVD method Regular arc deposition Filtered arc deposition Coating design Monolithic coating Multi-layer coating Multi-layer coating Coefficient Palmquist of adhesion Microhardness toughness to the Wear resistance (N mÀ2) (GPa) substrate on cutting 25.0 26.0 0.5 1.0 30–35 50–60 0.8 1.5–2.0 35–37 150–200 0.8 2.0–2.5 Copyright 2004 by Marcel Dekker, Inc All Rights Reserved The same multi-layer coating could be used for filtered PVD coatings with the additional advantages offered by this technology [41] These systems not only eliminate the ‘‘droplet’’ phase from the coating, but they can also be used to control the deposition conditions so that an excellent microstructure and properties are obtained in the film An extremely fine-grained structure (the grain size is 10 nm in comparison to a grain size of several microns in regular coatings) can be achieved [42], with excellent mechanical properties The nano-grained structure of the film is due to: the higher ionization rates achievable in filtered arc-evaporated plasmas, which are thought to enhance the nucleation rate and depress the growth rate of coarse crystals; and (2) a lower overall deposition rate for the filtered arc PVD process that results in a drop in the temperature at the growth front of the film (1) Although lower deposition rates are found with this process and can lead to a loss of productivity, these disadvantages can be offset by improvements in the quality of the film The hardness and Palmquist toughness of a TiN coating deposited by this method can be increased up to 35–37 GPa (instead of 25 GPa) and 150–200 J mÀ2 (instead of 26 J mÀ2), respectively The adhesion of this coating is also very high (kadh¼0.8, see Table 5) In addition, the wear resistance of filtered coatings is usually much better than regular coatings (Table 5) The principles outlined above for a multi-layered TiN coating can also be successfully applied to filtered coatings 1 Frictional Wear Behavior and Self-Organization of TiN PVD Coatings Hard TiN coatings act as surface lubricants by inhibiting the adhesion of the tool surface to the workpiece [1] The friction parameter of this coating is also low at the operating temperature (Fig 23) The wear behavior changes when TiN coatings are applied to cutting tools (Fig 24a) [43]: the initial rate of wear (during the running-in stage) is significantly lower and the range of normal wear is expanded As a result, both reduced friction and wear control can be achieved (see Fig 2) To explain the enhanced wear characteristics imparted by TiN PVD coatings, a study of the self-organization behavior of the tool was performed [26] Protective secondary structures (SS-I) of the Ti–O type form at the surface of hard PVD TiN coatings during cutting The transition from the running-in stage to the normal wear stage is marked by the development of a supersaturated solid solution of oxygen in titanium (Fig 25a–c) At the same time, as the friction drops, Copyright 2004 by Marcel Dekker, Inc All Rights Reserved  ... phase transformation differs significantly from that found on heating a TiC standard up to 823 K for 30 Figure 14(b) and (c) presents electron energy loss spectra from the wear crater after and 30... the formation of BO, TiO and TiB,N Figure 17a–c presents spectra of the positive and negative Figure 16 Wear curves of friction contact materials: (1) DCPM with 20% TiC; (2) DCPM with 20% TiC and. .. nitrogen, formed by dissociation of nitrides and carbo-nitrides Such mass transfer takes place under the extreme temperature and stress conditions encountered in the friction zone Therefore, the selection