Abstract
Titanium alloys have been extensively applied due to their outstanding features, including high specific strength and corrosion resistance. However, the machining of these alloys presents challenges, leading to significant tool wear and reduced tool lifetime, ascribed to the difficult-to-machine features of Ti alloys. In the context of pursuing environmental sustainability and enhanced productivity, high-speed dry-cutting technology coupled with high-quality tool coatings are prospective. This paper comprehensively explores wear patterns and recent advancements in tool coatings for high-speed dry cutting of Ti-alloys. Key coating properties such as resistance to abrasion, adhesion, oxidation, fatigue, hot hardness, and self-adaptability are briefly introduced. Design strategies for enhancing micro-hardness, toughness, and adhesion, including nanocomposite, multilayer, and gradient architectures, are also discussed. Furthermore, the paper highlights several Boron (B) doped hard coatings as promising candidates for high-speed dry cutting against Ti-alloys, exploring the impact of B on microstructure, hardness, thermal stability, oxidation resistance, and lubrication. In summary, this work provides a systematic overview of novel B-doped tool coatings applied in high-speed dry cutting against Ti-alloys.
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1 Introduction
Ti-alloys have been applied in-depth in the fields of aerospace, defense equipment, automobile manufacturing, and medical components owing to their splendid specific strength, high corrosion/heat resistance, and antibacterial properties [1]. The mechanical properties of both pure Ti and commonly used Ti-alloys are listed in Table 1. Despite their advantageous properties, Ti-alloys are typical difficult-to-machine materials due to their unique machining features, such as a work-hardening effect, low thermal conductivity, short chip-tool contact length, and high chemical reactivity [2, 3]. Consequently, the cutting tools usually experience high temperatures and drastic wear loss due to the concentration of heat and stresses, resulting in short tool life, unacceptable surface quality, high energy consumption, and production costs [4, 5]. In practice, the cutting speed is generally maintained below 60 m/min during the machining of Ti-alloys, significantly lower than that for steel cutting, to ensure a prolonged tool life and better surface quality [6].
The morphologies of Ti-alloys chips serve as focal points for research concerning the machinability and cutting performances of tools. Numerous studies have demonstrated that the chips often exhibit serrated patterns during high-speed cutting of Ti-alloys, and the saw-tooth morphologies depend strongly on the cutting parameters. Wang et al. pointed out that the degree of chip serration (i.e., the ratio of serration height to total chip thickness) is closely correlated to the feed rate (f) during the dry turning of Ti-6Al-4 V alloy. At a feed rate (f) of 0.05 mm/rev, the chips displayed fuzzy and unevenly arranged saw-tooth shapes. As the feed rate increased, the degree of serration gradually intensified, presenting a more periodic arrangement (see Fig. 1) [9]. Barry et al. claimed that the periodicity of the serration was related to the cutting speed and depth during the orthogonal turning of Ti-6Al-4 V alloy [10]. Younas et al. further supported these findings, stating that the degree of chip serration increased with higher cutting speeds in the turning process of Ti-6Al-4 V alloy when the uncoated inserts and a feed rate of 0.16 m/rev were employed [7]. Nonetheless, the formation mechanism of serrated chips of Ti-alloys remains unclear to date, with the following two recognized theories attempting to explain this phenomenon. The first theory is the adiabatic shear mechanism, which insists that high-frequency strains cause rapid local heating-up during the cutting of Ti-alloys, inducing the local thermal softening and thermoplastic instability effects. These effects promote the occurrence of concentrated shear slip** deformation adjacent to the unaffected areas, resulting in uneven internal deformation and, ultimately, the serration of Ti-alloys chips [11]. The second theory is the cyclic fracture mechanism, suggesting that the Ti-alloys undergo internal ductile fracture generating cracks during cutting. Hence, the shear resistance of Ti-alloys is depressed, leading to the formation of concentrated shear deformation and shear bands, and further resulting in the slip of machined materials along the shear band and, ultimately, the formation of serrated chips [12]. Generally, the formation of serrated chips is accompanied by unsteady ultra-high frequency vibrations of cutting force and stress during the cutting of Ti-alloys, which leads to high-frequency impact, crack initiation and propagation, and eventual fatigue damage to the coating and tool.
The evolution of chips morphologies versus. feeding rate during dry cutting of Ti-6Al-4V alloy [9]
Metal-cutting processes account for a considerable part of the machining industry [13]. Utilization of cutting fluids is a common way to enhance surface quality and productive efficiency in metal cutting despite issues such as environmental damage and policy restrictions. The manufacture, usage, treatment, and discharge of cutting fluids are reported to pollute the environment and cause respiratory and neurological diseases due to the release of smoke, oil mist, and particles from the machines. Meanwhile, “secondary pollution” occurs when residual cutting fluids are removed or high-temperature evaporated from the surface of the finished parts [14]. Additionally, the usage of cutting fluid accounts for about 16% of the total manufacturing cost and increases the operating burdens of enterprises [15]. To meet the growing demands for environmental protection and productivity, the high-speed dry-cutting technique, which eliminates the usage of cutting fluid, has emerged not only to ensure high cutting speed, high surface quality, and long tool durability, but also to offer advantages over traditional cutting techniques, such as: (1) improved machining efficiency by 3 ~ 5 times arising from increased feed rate and cutting speed [16]; (2) limited thermal deformations and internal stresses in tools and workpieces due to the accelerated removals of chip and dissipation of cutting heat at the tool-chip contact area [17]; (3) reduced manufacturing energy consumption and wear losses of tools attributed to the decreased cutting forces when the cutting speed overpasses a certain range [18]; (4) improved machining accuracy owing to avoidance of resonance phenomenon between tools and machines [17].
However, for high-speed dry cutting of difficult-to-machine materials, such as the Ti-alloys, challenges such as high cutting force, difficulty in heat dissipation, and rapid tool wear need to be addressed. Non-metallic elements (such as B, Si, C, etc.) doped Al-based hard coatings, which combine the self-lubrication or oxidation resistance effect of the former and the excellent thermal stability of the latter, are prospective for application in high-speed dry cutting of Ti alloys [19, 20]. This paper aims to present a concise review of the actual performances of coated tools, the correlation with crucial properties of tool coatings, the evolution of the modern machining industry, the principal structure design of tool coatings, and recent investigations on the B-containing hard coatings.
2 Coated Tool Performance in Ti-Alloys Machining
2.1 Wear patterns
As shown in Fig. 2, in the course of Ti-alloys machining, each part of coated tool exhibits intricate wear patterns dependent on cutting conditions. Rake face of tool is where chips flow over, maximum cutting force and heat accumulation are usually detected at the tool-chip contact area of rake face [21]. Ti-alloys chips are firstly adhered on the rake face due to synergistic effects of cutting load, cutting heat and chemical reaction between chips and coating, then the accumulated cutting heat (stemming from low thermal conductivity of Ti-alloys) result in thermal expansion of coating and tool substrate, the mismatch of thermal expansion coefficient and shear stress caused the peeling of coating and exposure of tool substrate [7]; removal of sticked Ti-alloys during cutting also pluck the tool material away [22], and oxidation of coating and tool materials further degrade their strength [23]; repeat of these processes finally induce the formation of crater on the rake face under high-temperature adhesion, chemical diffusion, and oxidation [6]. Flank face of tool is usually in contact with machined surface of workpiece so that cutting heat accumulation is mild thereat, the flank faces of cemented carbide tools typically endure abrasive and adhesive wear, forming a broad wear band, especially at lower cutting speeds [24]. The cutting edges are susceptible to fatigue and generate cracks under the cyclic impact of heat/force and may undergo plastic deformation or chip** once the cutting pressure exceeds the endurance of the tool bodies [25,26,27]; and sticked Ti-alloy on cutting edge is called build-up edge (BUE), removal of BUEs probably contribute to severe chip**s or breakages of cutting edges and even notch wear on flank faces, as exhibited in Fig. 2 [15, 28].
2.2 Cutting performance of prevailing coated tools
The adoption of hard coatings significantly extends the lifespan of cutting tools during Ti-alloys machining. Priarone et al. reported a 200% increase in the life of a nanocomposite AlSiTiN-coated milling cutter compared to an uncoated one when dry milling against the Ti-48Al-2Cr-2Nb alloy. Additionally, the coated tool with post-polishing treatment mitigated the work-hardening effect on the surface [29]. An et al. conducted a comparative study on the wear patterns and lifespan of uncoated, CVD-Ti(C, N)/Al2O3/TiN-coated, and PVD-(Ti, Al)N/TiN-coated tools in face milling of Ti-6242S and Ti-555 alloys. The results revealed that the PVD-coated tools exhibited the smallest flank wear for both Ti-alloys. In contrast, severe chip** and coating peeling occurred on CVD and uncoated tools due to concurrent adhesive and diffusion wear, resulting in shorter lifetimes [5]. Nouari et al. observed that the service life of cutting tools was susceptible to brittle fracture (such as crack, breakage, and peeling) and local wear loss on the primary edge, as demonstrated by a study focusing on the failure modes of uncoated and CVD-TiC/TiCN/TiN-coated tools during high-speed dry cutting of Ti-6242S alloy [30].
As the dedicated hard coatings for cutting Ti-alloys, the coatings containing self-adaptive/lubricative elements such as Cr, V, etc., are reported to achieve superior cutting performance [31]. For instance, CrN-coated tools exhibited lower coefficient of friction (COF), flank wear rates, and milder crater wear than AlTiN-coated ones during high-speed turning of Ti-6Al-4 V alloy, owing to the formation of self-adaptive Cr2O3 oxide film acting as a thermal barrier layer to cool the cutting area, accelerate the flow of chips, and reduce the tool-chip contact length [32]. AlTiN/VN multi-layer coated tools, as reported by Biksa et al., showed a ~ 40% increase in cutting life and a ~ 20% decrease in COF, presented a smooth morphology of chips compared to the AlTiN-monolayer coated one, attributed to the self-lubricating effect of V2O5 oxides [33]. Kumar et al. investigated the wear patterns and chip removals of TiSiN and TiSiVN-coated ceramic inserts after a short-term cutting of Ti-6Al-4 V, the TiSiVN-coated insert exhibited smaller crater depth and chip curvature, promoting chip flow and inhibiting adhesion, oxidation, and diffusion on the tool-chip contact area [2].
Nevertheless, the high hardness and low friction of coatings alone cannot guarantee superior cutting tool performance in Ti-alloys machining. The stability of microstructure and mechanical properties, along with self-adaptations, plays a more crucial role, especially under the thermal conditions of dry cutting. For instance, diamond coatings typically exhibit extremely high hardness (50 ~ 80 GPa) and low COF [34,35,36]. However, a drop in hardness occurs due to oxidation and graphitization at around 600 °C, making them difficult to service at ~ 1000 °C, a typical temperature range under high-speed dry cutting conditions against Ti-alloys [37, 38]. Aslantas et al. reported that the AlCrN coating with lower hardness (3200 HV0.05) and COF (0.35) showed better-cutting forces, tool wear, and machining quality than nanocrystalline diamond coating with higher hardness (9000 HV0.05) and lower COF (0.15), according to a comparative study on the dry milling performance against Ti-6Al-4 V alloy [39].
Furthermore, the tool wear is associated with the impact fatigue resistance of hard coatings during high-speed cutting of Ti-alloys. Coated tools are subject to large and low-frequency cyclic impact loads when the feed rate is large in the intermittent cutting process (i.e., the rough milling). TiSiN/TiAlN multi-layer coated tools, for instance, exhibit low cutting force and wear loss during the milling of Ti-alloys due to their better low-frequency (16.6 Hz) impact fatigue resistance. However, cracks in coating and severe plastic deformation on the cutting edge were observed for TiSiN/TiAlN dual-layer coated tools, indicating inferior low-frequency impact fatigue resistance [40]. During continuous cutting (i.e., turning) of Ti-alloys, high-frequency cyclic impact (2 ~ 35 kHz) is imposed on coated tools [41, 42]. Literature [9, 43, 44] has presented the failure mode of TiSiN/TiAlN coatings after dry turning of Ti-alloys at different cutting speeds. The results revealed that the impact fatigue resistance depended intensively on their microstructure: the multi-layer structure possessed better fatigue resistance under a small impact load, whereas the dual-layer structure was more splendid in fatigue resistance when the impact load was large.
It is worth noting that under dry cutting conditions of titanium alloy, the tool coating is prone to delamination, peeling, and wear due to the synergistic effect of cutting heat and mechanical impact, resulting in the large cutting force, aggravated wear loss, and even chip** of cutting edge. Therefore, the cutting speed is generally limited to below 60 m/min in practice, and it is considered as high-speed cutting if it exceeds this limit [6]. Table 2 lists the lifetime of coated carbide inserts for high-speed turning of Ti-6Al-4 V (flank wear of 0.2 mm is adopted as the criterion of tool failure). It can be observed that the reported coated tools have a short lifespan during high-speed dry machining of Ti-alloys, posing a significant challenge for further prolongation of tool lifetimes [44, 45].
3 Crucial Properties of Tool Coatings
Since the 1970s, with the widespread commercialization of CVD and PVD technologies, technological innovations in the machining industry have been closely linked to the development of hard coatings [46]. Computer Numerical Controlling (CNC) machines became prevalent worldwide after the 1980s, and the demand for productivity drove the extensive application of high-hardness coatings, such as TiC, TiN, TiAlN, etc. After 2010, the effective combination of CNC and high-speed motors led to higher machining speeds. Concurrently, the compositions and structures of hard coatings were diversified to accommodate the evolving machining technologies [47,48,49,50,51]. To meet the increasing requirements of the machining industry for cutting speed, lubrication conditions, and machining quality, hard coatings for cutting tools are progressing in in-service properties, which can be mainly classified into the following aspects.
3.1 Resistance to abrasive and adhesive wear
In the early decades of their development, the primary objective of hard coatings was to improve the wear rate and resistance to adhesion of tools by virtue of the high hardness [46]. The chips flowing across the rake face, along with the friction occurring between the flank face and pieces, result in severe abrasion and even crater wear [21]. Additionally, serious plastic deformation or notch wear can occur on the cutting edge, particularly for the workpiece materials exhibiting a work-hardening effect. Moreover, when the atomic bonds in the tool-chip sliding contact system are stronger than those in the tool itself, the atomic bonds of the cutting tool are torn apart. From a microscopic perspective, the materials of the cutting tool are carried by the flowing chips, leading to adhesive wear [22, 30], which can be further aggravated by asperities as a result of the wear loss on the cutting tool [52, 53].
Table 3 shows the hardness of coatings applied in early applications, which are basically above 20 GPa, well meeting the demands for resistances to abrasive and adhesive wear. Among them, TiC and TiN coatings were the first to be applied on the cutting tools, followed by the development of TiCN coating, a hardening ternary compound (up to 30 GPa or more) achieved by incorporating carbon atoms into the TiN lattice [54]. Compared with TiN coating, TiCN typically exhibits better COF and abrasive wear resistance [54]. Ruppi et al. [55] demonstrated that the TiCN-coated inserts showed the least flank wear and limited crater wear on the rake face compared with TiC, TiN, and Al2O3-coated ones during dry cutting of 42CrMo4 steel. Subsequently, Al atoms were added into the TiN lattice to form a metastable fcc-(Ti, Al)N solid solution through substitution and/or occupation of interstitial sites. The incorporation of Al atoms induced both solid solution and fine-grain strengthening effects, arising from a smaller radius of Al atoms (Ti: 0.200 nm, Al: 0.182 nm [56]). As a result, the hardness, toughness, and cutting performance were enhanced for TiAlN coatings [57, 58]. Martinez et al. indicated that the flank wear rate of TiAlN-coated tools was 50% less than that of TiN-coated ones at the cutting speeds of 150 and 200 m/min during dry cutting of cast iron [59]. To date, TiAlN remains a prevalent coating for milling and turning steels [60].
3.2 Hot hardness
During high-speed dry machining, cutting heat is generated and accumulates inside the tool, with peak temperature typically observed at the tool-chip contact area on the rake face [21]. Figure 3 depicted the change of peak temperature and cutting forces as functions of cutting time during machining of Ti-6Al-V4 alloy at different cutting speed, the peak temperature increased from ~ 650 °C to above 900 °C, while the cutting forces remained nearly constant as the cutting speed raised from 46 m/min to 91 m/min during, it can be summarized that the temperature is rather susceptible to cutting speed [21]. Therefore, the hot hardness of the hard coating is a crucial and essential property for enhancing wear resistance at high service temperatures. The conventional hard coatings, such as TiN, TiC, and TiCN, were reported to undergo structural recovery, recrystallization, and softening at high temperatures, resulting in their lower service temperature range of 400 °C ~ 500 °C (Table 3) [61,62,63,64].
2D FEA simulations of cutting forces and temperature (as a function of time) while machining a Ti6AlV4 alloy at (a) 46 m/min and (b) 91 m/min [21]
In contrast, Al-containing coatings (Me, Al)N (where, Me: Ti, Cr, etc.) exhibited higher hot hardness at elevated temperatures, surpassing even their as-deposited state. Spinodal decomposition from the metastable fcc-(Me, Al)N solid solution into Al-rich (Al, Me)N phases and Me-rich (Me, Al)N phases resulted in a complex mixture of residual and new solid solutions, leading to a coherent strengthening effect and increased of hardness (Fig. 4(a)) [66,67,68,69,70]. With further temperature elevation, the Al- and Me-rich phases decomposed into fcc-MeN and soft hexagonal AlN (h-AlN) phases, causing a decline in coating hardness with the vanishing of the coherent structure (Fig. 4(b)) [71]. Veprek et al. highlighted that the onset temperature of spinodal decomposition is closely correlated with the atomic ratio of Al/(Al + Me) in hard coatings. When the ratio exceeded 0.4, decomposition initiated at 700 °C and accelerated at 900 °C; however, if the ratio was lower than 0.3, the onset temperature depended on the synthesis method, e.g., 450 °C and 250 °C for multi-arc ion plating and ion beam-assisted CVD deposition, respectively [71]. Following vacuum annealing at 1000 °C, the hardness of the AlTiN coating was 25.5 GPa, 31.0 GPa, and 33.0 GPa for Al/(Al + Ti) ratios for 0.4, 0.6, and 0.7, elucidating the positive effect of increased Al content on the hot hardness [72]. Another vacuum annealing experiment (600 °C ~ 1100 °C) showed that the hardness of Al0.34Ti0.64N (atomic ratio) coating increased with annealing temperature until reaching a peak value at 900 °C (38 GPa) and then decreased monotonically [67]. Generally, Al-rich AlTiN coatings can maintain superior cutting performance up to 700 °C ~ 850 °C (Table 3).
Illustration of (a) spinodal decomposition of Ti1-xAlxN into fcc-TiN and fcc-AlN rich domains and (b) dual-phase structure (right) after transformation of fcc-AlN into w-AlN at higher annealing temperatures [60]
3.3 Oxidation resistance
As cutting speed and temperature increase, the primary wear patterns of tools alter from abrasive and adhesive wear to the co-existence of oxidation, diffusion, abrasion, and adhesion [46]. The service temperature of the coating is limited by its oxidation resistance; for instance, TiN coatings transform into TiO2 with a subsequent rapid decrease in hardness above 500 °C [73]. Compared to Ti-based coatings, Al-based hard coatings exhibit significantly improved oxidation resistance at high temperatures [74, 75]. This improvement is attributed to the diffusion of Al atoms to the outermost layer of the coating, forming a dense Al2O3 oxide layer (similar to the passivation effect), which prevents the inward diffusion of O atoms. Meanwhile, Al2O3 has excellent chemical stability and hot hardness, which improves the cutting performance of hard coatings under dry cutting conditions [46]. The diffusion coefficient of Al atoms increases due to cutting heat, and the higher migration velocity of Al is beneficial for repairing the outermost worn-out Al2O3 oxide layer [76]. Higher Al content in TiAlN coatings is considered to enhance the oxidation resistance; however, Klaus et al. claimed that both Al content and temperature determine the formation of an Al concentration gradient and the generation of a protective oxide layer [77].
3.4 Fatigue resistance and self-adaptability
Under intermittent cutting conditions, such as milling, tool coatings are susceptible to fatigue and cracking during periodic cutting into and out of the workpiece, which ultimately results in coating flaking and tool wear as cracks expand [78]. In contrast, during continuous machining (e.g., turning) of high-hardness steels and difficult-to-machine materials (e.g., Ti alloys and high-temperature Ni-based alloys), high-frequency fatigue impacts on the coated tools also occurred due to the formation of serrated chips [42]. Proper design of composition and structure is considered essential to improve fatigue resistance in hard coatings.
Researchers have employed cyclic impact testing to evaluate the fatigue resistance of hard coatings. For instance, Sergejev et al. designed a rotating disc-impact hammer dynamic loading system [79], testing fatigue on TiCN and AlCrN coatings by accelerating the rotating disc to drive the impact of the hammer at impact frequencies ranging from 25 to 100 Hz. Bouzakis et al. [80] used an adjustable piezoelectric device to drive a hard metal indenter for fatigue testing of TiAlN coatings at frequencies up to 50 Hz However, the impact frequencies of these testing methods are much lower than the actual working conditions (2 ~ 35 kHz) during the cutting of titanium alloys [41, 42]. To better simulate real working conditions, Zha et al. utilized an ultrasonic generator to drive a diamond indenter to vertically impact TiAlN/TiSiN hard coatings at ultrahigh-frequency impacts of 20 kHz. This study delineated the damage process of the coatings into four states: (i) the initial impact, (ii) the steady state of the initial fatigue, (iii) severe damage, and (iv) the steady state of the re-fatigue [43].
Under the harsh cutting conditions of high-speed dry cutting, the concept of "adaptive" coatings has attracted attention in recent years. Adaptive coatings undergo reversible transformations in surface chemistry, structure, and mechanical behavior based on applied load and operating conditions, ultimately maintaining low friction and wear loss in frictional contact. These coatings produce lubricating oxide with easy slip shear surfaces and low melting points under severe conditions, such as the Magneli oxide phase of V, W, and Mo [81]. Sputtered VN coatings, for example, begin to oxidize and form the V2O5-Magneli phase at 520 °C under the thermal and frictional oxidation, transitioning into a liquid oxide lubricating film at 660 °C, ultimately reducing high temperature (700 °C) coefficient of friction to less than 0.2 when employed against steel and alumina counter-pairs [82]. The addition of V into TiAlN coatings has also been reported to exert a lubricating effect [83]. Rabinovich et al. [84, 85] analyzed the irreversible entropy-generating thermodynamics based on the frictional behaviors of TiAlCrSiYN hard coatings during dry cutting of AISI H13 steel. The results showed that the cyclic depletion and generation of lubricating oxides on the tool surface dissipated most of the energy in a cutting system, which reduced systematic entropy (more ordered) and also brought the "self-adaptive" feature of hard coatings under harsh cutting conditions and effectively mitigated wear loss of the tools.
4 Structural Design of Tool Coatings
In addition to compositional optimization, structural optimization of hard coatings is another effective strategy to enhance the cutting performance. Some structural design approaches are briefly introduced below.
4.1 Nanocomposite structure
The hardness of a material usually increases with the decrease of the grain size (Hall–Petch effect). However, when grain size further decreases to tens or even several nanometers (Anti-Hall–Petch effect), this rule no longer applies [86]. To explore the hardening mechanisms of coatings at the nanoscale, Veprek et al. proposed a nanocomposite structure hardening model. This model involved TiN/TiSi2 nanocrystals (diameter below 10 nm) (Fig. 5(a)) wrapped by one or two atomic layers of Si3N4 amorphous phase, based on the study of superhard coatings (> 40 GPa) of Si-doped TiN in 1995 [87, 88]. The nanocomposite structure originates from the spinodal decomposition of immiscible phases in a solid solution. The Gibbs free energy curve of A1-xBx solid solution as a function of B-phase content at different temperatures (T2 > T1) is shown in Fig. 5(b) [89]. When the second derivative of the curve is negative, any infinite and small changes in composition cause the decomposition of A1-xBx solid solution into A and B two-phase mixture without the process of nucleation and growth, a phenomenon known as "spinodal decomposition". This process forms a fine nanostructure, and its structure and size depend on the free energy change during decomposition, the composition gradient, and the elastic strain energy derived from the coherent interfaces between two-phase products. At T1, the second derivative of the curve is constantly negative; thus, a mixture of phase A and phase B would be formed rather than the solid solution of A1-xBx. As for T2, when phase B content is between x1 and x2, the second derivative of the curve is negative, and spinodal decomposition occurs; however, when phase B content is in ranges of x’< x < x1 and x2 < x < x”, the second derivative of the curve is positive, and the solid solution decomposition process requires nucleation. Moreover, the decomposition products rely on temperature strongly, which are pure phase A and phase B at T1, and A1- x’B x’ and A 1-x” B x” at T2 [81].
Researchers developed a series of nc-MeN/a-Si3N4 nanocomposite structure coating (nc: nanocrystalline, a: amorphous), where Me represents transition metals such as W, Ti, V, Zr, etc. It was found that the amorphous layer acted as a strong "glue" among MeN nanocrystalline, resulting in super-hardness by inhibiting grain boundary slip**, imparting high hot-hardness, and enhancing oxidation resistance [86]. Moreover, elements such as B and C can also promote the formation of amorphous phases. Consequently, numerous ternary or quaternary nanocomposite structural hard coatings have been thoroughly investigated and also applied in various industries. Examples of such nanocomposites include nc-TiN/a-BN [89], c-CrN/a-Si3N4 [90], nc-TiAlN/a-Si3N4 [91], nc-TiAlN/a-C [92, 93], etc.
4.2 Multilayer structure
In addition to the nanocomposite structure, fabricating alternating deposition of multilayer coatings is also a promising method to enhance the cutting performance of hard coatings, primarily due to the improved adhesive strength and toughness [74, 86]. When the coating surface is subjected to mechanical loads (such as cutting), the durability of the coating is typically determined by the generation, propagation, and energy dissipation of cracks. The construction of a multilayer structure serves to interrupt the columnar growths of coatings, allowing cracks to be split, deflected, and energy dissipated either at the grain/phase boundaries (grain boundary toughening or strengthening) or at the interfaces between sub-layers (interface toughening or strengthening) [94]. As a result, this impedes the crack propagation within the coatings. Moreover, local delamination may occur at nano-vacancies at the interfaces, leading to local stress relaxation and even nanoscale plastic enhancement [95]. The crack propagation also interacts with periodic strain–stress fields within multilayer structures. Figure 6 illustrates the hardening/toughening mechanism of the multilayer structure.
Schematic of toughening and strengthening mechanisms in multilayer coatings [95]
Moreover, achieving higher hardness is more attainable for multilayer coatings compared with monolayer ones, attributed to the hindered dislocation motion at interfaces [96]. Kao et al. reported the TiAlN/CrSiN multilayer coating, prepared by pulsed DC reaction magnetron sputtering, had a hardness of 35 GPa, higher than either TiAlN (28 GPa) or CrSiN (30 GPa) monolayer coating [94]. Additionally, an AlCrN/AlCrMoN multilayer coating (~ 3600 HK0.05), produced by multi-arc ion plating, also showed a higher hardness compared to the AlCrMoN monolayer coating (~ 3250 HK0.05) [97]. Furthermore, the multilayer structure has been reported to enhance the hot-hardness and oxidation resistance of the coatings [25, 98] and facilitate the diffusion of elements, such as Al, Cr, Si, B, V, etc., to the outmost surface of coatings, promoting the formation of protective oxide films during cutting [84, 85].
4.3 Gradient structure
High hardness and wear resistance in hard coatings are often accompanied by large residual stresses, originating from internal stresses due to lattice distortion, dislocation, and other defects during coating growth [92, 99], as well as thermal stress generated during the cooling process after coating deposition [100]. Therefore, there is a mismatch in the distribution of stresses between hard coatings and tool substrates, leading to the formation of cracks at substrate-coating interfaces and, in severe cases, spalling or breakage of hard coatings. Furthermore, variations in the thermal expansion coefficient (TEC) between coatings and tool substrates (for example, the TEC of AlTiBN coating and cemented carbide is 7.5 × 10−6 K−15.5 × 10−6 K−1 [101, 102], respectively), may result in a substantial mismatch in thermal expansion volume, potentially causing coating peeling during the cutting process.
Fortunately, the design of a gradient structure is considered a practical approach to avoid premature failure of hard coatings induced by sharp changes in structure and service properties. The implementation of a smooth transition in structure, chemical composition, and function proves instrumental in enhancing mechanical properties such as micro-hardness and adhesive strength. Cai et al. reported that an AlCrSiN coating with the smoothest Si content gradient (0% ~ 2% ~ 4% ~ 6%, mass ratio) exhibited the highest micro-hardness and adhesive strength, extending tool life by ~ 90% compared to an AlCrSiN coating without a gradient structure, as demonstrated in milling tests against 20CrMo gear steel [103]. Wang et al. also claimed that CrAlSiN coating with a smooth gradient structure (Fig. 7) exhibited a significant improvement in toughness, scratch adhesive strength and crack propagation resistance [86, 104].
Cross-sectional image of CrASiN coatings with gradient change of Si content [86]
5 Research on Boron-containing Coatings
The addition of B into hard coatings has attracted extensive attention from researchers, aiming not only to reinforce hardness, toughness, and high-temperature properties by constructing nanocomposite structure but also to enhance self-adaptability through the formation of oxides with an in-situ self-lubricating effect. Boron-containing hard coatings can be categorized based on their components and applications.
5.1 BN based hard coatings
Cubic boron nitride (c-BN) is known for its high hardness, second only to diamond, making c-BN hard coatings promising for cutting applications. However, synthesized c-BN coatings typically consist of nanocrystals, amorphous phases, and hexagonal BN phase (h-BN), impacting their hardness. The synthesis process is still under exploration [105]. Notably, h-BN, also known as white graphene, has a lamellar structure similar to graphite, allowing slippage along laminates due to weak Van der Waals forces [106]. Various synthesis methods of h-BN are well developed, including PVD, CVD, Vapor deposition infiltration, and dip-coating [106, 107], encouraging its application in industries such as metal processing (e.g., aluminum die-casting), automotive, cosmetics, and others requiring high lubricity [106]. Improved mechanical properties can be achieved by compounding h-BN with other hard coatings. For instance, Nose et al. synthesized CrAlN/h-BN dual-phase hard coatings by co-plating, i.e., pulsed DC and RF magnetron sputtering. The results showed enhanced hardness and elastic modulus, reaching peak values of 46 GPa and 440 GPa, respectively, by increasing h-BN phase content; however, the properties declined for h-BN phase content exceeding 7% (volume ratio) [108].
5.2 Ti-B based hard coatings
TiB2 coating, despite early commercialization, remains a focus of research due to its excellent properties, including high hardness, high wear, and corrosion resistance. In addition, the low affinity of TiB2 with aluminum makes it crucial for die-casting or dry-cutting of aluminum alloys [109]. Fabricated primarily through CVD, novel methods like plasma-assisted CVD have enabled TiB2 coating synthesis at lower temperatures (below 400 °C) with nanocrystals for exceptionally high hardness, reaching 50 GPa [110]. However, it is necessary to adjust B content within coatings since high B content results in a large proportion of amorphous soft phase that decreases the hardness of TiBx films. To further improve the performance of TiB2 coating, multilayer structures containing TiB2 and other hard coatings are usually adopted; for example, TiN/TiB2, TiCN/TiB2, TiC/TiB2, and TiN/BN/TiB2, to address issues of poor adhesion, high porosity, and low hardness [109, 111]. Introducing Boron into TiC coating improves thermal stability, hardness, friction and wear properties, corrosion, and oxidation resistance. Larhlimi et al. pointed out that the Ti-B-C ternary nanocomposite coating prepared by magnetron sputtering showed a complex phase structure, including TiC, TiB2, a-C, B4C and ternary TiBxCy, and a super hardness of 70 GPa was achieved, much higher than that of TiC or TiB2 binary coating [63].
Similarly, the addition of B into TiN coating enhances hardening and structural refinement. The structure of TiBN coating is closely dependent on B content. When the B content is higher than 18%, it’s the nanocomposite structure composed of TiN nanocrystals surrounded by amorphous TiB2 phase or the binary-phase structure composed of TiN/TiB2 nanocrystals, while at less than 18%, the nanocomposite structure including TiN nanocrystals and amorphous BN tissue phases (and possibly amorphous TiB phase) is formed [89, 109, 112]. Karvankova et al. fabricated nc-TiN/a-BN nanocomposite coatings with varying B content by adjusting the flow rates and the ratio of BCl3 to TiCl4 gas [113]. The study revealed that the grain size of TiN nanocrystals decreased as the B content increased. The hardness of the coating was found to be closely associated with the thickness of the amorphous BN layer. Specifically, the highest hardness, reaching approximately 50 GPa, was achieved when there was a continuous amorphous monolayer of BN. However, the introduction of two or more layers led to a rapid decrease in hardness, drop** below 30 GPa [113].
Due to the reinforcing effect of B, Ti-B-N coatings typically demonstrate superior wear resistance compared to TiN coatings. The wear rate of Ti-B-N coating was reported to be one order of magnitude lower than that of TiN coating when a GCr15 steel ball (HRC60 ~ 62) was employed as a counterpart in the pin-on-disk wearing tests [114]. Furthermore, the presence of element B contributes to the enhancement of the hot hardness of TiN coating. In annealing tests conducted in an Ar atmosphere, Ti-B-N coating with 4.1 at. % B content maintained a high hardness (> 40 GPa) at 1000 °C. In contrast, TiN coating experienced a drastic reduction in hardness above 500 °C [89, 113]. Moreover, the oxidation resistance of the Ti-B-N coating was significantly improved, especially up to 700 °C, which was attributed to the formation of a protective B2O3 top layer. In contrast, TiN coating was oxidized severely, ascribed to the formation of destructive oxides, e.g., r-TiO2, which accelerated the inward penetration of O atoms and brought overall oxidation of coating [115].
Although element B shows substantial superiority by constructing a Ti-B-based hard coating, it is necessary to compare it with another prevalent strengthening element, Si, especially for application in high-speed dry-cutting areas. Firstly, Si exhibits good compatibility with TiN nano-grains in the nc-TiN/a-Si3N4 coating system. Thus, Si enhances the oxidation resistance more prominently due to denser Si3N4 amorphous grain boundaries and encapsulated TiN nano-grains, which prevent O atoms from penetrating across the grain boundaries and diffusing inward at high temperatures [89]. In addition, the formation of dense and stable oxides-scale, such as SiO2, is reported to enhance the oxidation resistance up to 1600 °C [69]. In contrast, B shows a higher surficial mismatch with TiN nano-grains in the nc-TiN/a-BN system. This results in larger TiN grain sizes and a reduced number of grain boundaries, leading to severe inward diffusion of O atoms along the grain boundaries. Moreover, the solid-to-liquid phase transformation of B2O3 occurs at 450 °C, leading to fracture of the oxide layer and further deterioration of oxidation resistance of Ti-B-N coating [89]. Secondly, Si shows an inferior high-temperature lubricative effect compared to B. According to Polcar et al., the high-temperature COF of AlCrSiN is significantly higher than that of CrAlN due to a lack of a solid lubricative phase [116]. As for element B, it was reported to form a lubricative liquid product of B2O3 at high temperatures [112, 115]. As B2O3 reacts with water vapor in the atmosphere, a further product of H3BO3 oxide film is generated above 900 °C [117]. The H3BO3 is a material in which compact packed and firmly bonded B, O, and H atoms form a loose lamellar structure. The inter-lamellar spaces are large to facilitate the slip** and detachment of lamellae, thereby reducing the COF [117,118,119].
5.3 Al-Me-B based hard coatings
As discussed above, while the inclusion of B has exhibited outstanding high-temperature lubricative effects, it has shown shortcomings in terms of oxidation resistance. In order to further improve its oxidation resistance, thermal stability, and other mechanical properties, adding Al into the coating seems to be promising to meet the stringent requirements of high-speed dry cutting. Therefore, explorations and studies have been carried out on Al-Me-B-based (Me = Ti, Cr, etc.) hard coatings.
5.3.1 Effect of element B on microstructure and mechanical properties
The addition of B has a significant effect on the grain size of Al-Me-B-based hard coatings [120]. Tritremmel et al. fabricated AlCrBN coatings with different B content through multi-arc ion plating. They reported a sharp decrease in grain size from 35 to 20 nm with the addition of a small amount of B (2.3 at. %), and the grain size continued to decrease with further increase of B content [100]. The primary reason for these results is the incorporation of small B atoms into the (Al, Cr)N solid solution, promoting the formation of (Al,Cr,B) N via replacing the Al and Cr atoms or occupying the interstitial positions, which is beneficial to reduce the grain size. Another contributing factor is that the growth of Al-Cr-(B)-N microcrystals is interrupted by amorphous BN, and the increase in B content leads to a higher volume fraction of a-BN and smaller grain size [100]. Cai et al. claimed that cubic BN grew coherently at the grain boundaries of (Al,Cr,B)N solid solution when the B content was 1.5 at. % in AlCrBN coating. Increasing the B content to 2.9 at. % promoted the formation of a nanocomposite structure comprising nanograins with a size of ~ 5 nm and surrounding a-BNx layer with a thickness of ~ 1.3 nm. Further increasing it to 4.8 at. % resulted in the reduction of the sizes of (Al,Cr,B)N nanograins to 2 nm, while the thickness of a-BNx increased to 1.7 nm, as illustrated in Fig. 8 [121].
Schematic of microstructural evolution of AlCrBN coatings with increasing B content [121]
The mechanical properties of Al-Me-B-based hard coating, including hardness, toughness, and adhesion, can also be improved by the addition of B. For example, when compared to AlCrN coating, the hardness and adhesive strength of multi-arc ion plating AlCrBN was increased by ~ 3.7 GPa and 6.5 N, respectively [121]. Nevertheless, the strengthening mechanism of Al-Me-B-based hard coating is contingent on the B content as follows:
-
a)
Solid Solution Strengthening with Low B Content: With low B content, B atoms are mainly incorporated into the (Al, Me)N lattice as the solid solute, resulting in local/residual stress field, lattice strain, and finally hindering dislocation movement. This mechanism is known as solid solution strengthening [100]. In this scenario, added B atoms fill lattice vacancies and substitute Al, Me, or other metal atoms until the solubility limit is reached in the (Al, Me) N lattices [99, 122].
-
b)
Combination of Hall–Petch Effect and Solid Solution Strengthening with Moderate B Content: When B content is moderate, B atoms precipitate and aggregate at the grain boundary to form a-BN phase, but the amount of a-BN is insufficient to completely encapsulate (Al, Me) N grains (in the of case of B content equal to 1.5 at. %, as shown in Fig. 8). The strengthening mechanism is thus dominated by both Hall–Petch effect of grain refinement and solid solution strengthening [100, 123].
-
c)
Grain Boundary Strengthening with higher B Content: When the B content is further increased, there are enough a-BN phases to encapsulate (Al, Me) N grains. Besides the Hall–Petch effect, grain boundary strengthening occurs, enhancing the cohesive strength of coating (in the case of B content equal to 2.9 at. %, as shown in Fig. 8) [100, 124]. However, if the critical B content is exceeded (when the B content is 4.8 at. %, as shown in Fig. 8), the strengthening effect on mechanical properties is declined. This could be attributed to the annihilation of dislocations at the grain boundaries, vanishment of the Hall–Petch effect, an increased volume fraction of the amorphous phase, and extremely small grain sizes [125]. In addition, the slip** of a-BN grain boundaries is more likely to occur due to an increased number of atoms [100].
5.3.2 Effect of element B on thermal stability
The addition of boron (B) has a positive effect on the thermal structural stability and hot hardness of (Al, Me)N/a-BN coatings. The nanocomposite structure of these coatings helps to inhibit the coarsening of grains at elevated temperatures, ensuring the high hardness of the coating [126, 127]. For example, Mendez et al. detected an abrupt decrease in hardness for an Al0.6Ti0.4N coating at 850 °C, indicating the occurrence of severe thermal decomposition with the appearance of a soft h-AlN phase. In contrast, Al0.6Ti0.4BN coating delayed the thermal decomposition until 1000 °C and maintained high hardness (above 32GPa) until 900 °C [128]. Moraes et al. reported that the hardness of TiAlN coating peaked at 600 °C and then declined. However, TiAlBN coatings showed upward trends in hardness up to 800 °C, and the Ti–Al-10B-N coating with the highest B content continued this upward trend until 1000 °C (Fig. 9) [129]. Therefore, the addition of B plays a positive role in the thermal structural stability and hot hardness of AlTiN coating. A similar positive effect was found for CrAlBN coating as well. While the CrAlN coating decomposed into h-AlN and h-Cr2N at 1000 °C, the addition of B delayed the process by 100 °C ~ 200 °C [126].
Dependence of micro-hardness on B content in AlTiBN coatings after annealing tests [129]
5.3.3 Effect of element B on oxidation resistance
The addition of boron (B) to hard coatings, particularly in Al-based coatings such as AlTiN and AlCrN, has a significant impact on oxidation resistance at elevated temperatures.
The addition of Al into the hard coating forms a dense Al2O3 passivation layer at high temperatures to prevent the inward diffusion of O atoms, which enhances the oxidation resistance of Al-based coatings, such as AlTiN and AlCrN, to 800 °C ~ 900 °C [130, 131]. However, with the recent development of high-speed dry cutting, the temperature of cutting tools is reported to be around 1000 °C, and there is a further demand for facilitating the high-temperature oxidation resistance of Al-based coatings [132, 133]. As for TiAlN coatings, the protective passivation film generated on the surface is usually a mixed oxide composed of TiO2 and Al2O3 or a bilayer structure comprising an Al-rich top layer and a Ti–rich bottom layer. TiO2 has a hardly positive effect on oxidation resistance at high temperatures. The phase transformation from a-TiO2 (Anatase) to r-TiO2 (Rutile) is accompanied by a change in volume, resulting in the generation of pores and cracks and potential destruction of the protective Al2O3 oxide layer [134, 135]. This phase transformation can result in a rapid decrease in hardness as oxygen atoms diffuse into the coating through the compromised passivation layer [76]. Therefore, restraining the transformation of Ti oxides from the a-TiO2 phase to the r-TiO2 phase is particularly important.
According to the studies on oxides of TiAlBN coatings with different B content at 1000 °C, the r-TiO2 phase could only be detected in TiAlN coating. In contrast, the a-TiO2 phase was still detected in the TiAlBN coating. Furthermore, the ratio (r-TiO2/a-TiO2) of XRD peak intensity was inversely proportional to B content, indicating the inhibitory effect of B on the phase transformation of TiO2 [134]. The distribution of B2O3 along grain boundaries was identified as the key factor in delaying the phase transformation and hindering the diffusion of O atoms [134]. Zhang et al. synthesized AlTiBN coatings with varying B content by multi-arc ion plating and studied the oxidation mechanism at different temperatures [136]. The results showed that a dense Al2O3 passivation layer with a thickness of only 0.5 μm was formed on top of AlTiN coating at 800 °C, while an amorphous oxide layer with a thickness of 3 μm containing Al, Ti, and B was formed on top of AlTiBN coating. At 900 °C, AlTiN coating was completely oxidized, and a large amount of r-TiO2 was formed. Meanwhile, AlTiBN coating demonstrated a better oxidation resistance due to the formation of mixed oxide products of Al2O3, a-TiO2, and r-TiO2 by crystallization of the amorphous oxide layer at higher temperatures, preventing the penetration of O atoms (Fig. 10). The proportion of the rutile phase (r-TiO2) decreased with the increasing B content [136]. From these results, it can be summarized that increasing B content in AlTiBN coating inhibits the formation of the r-TiO2 phase and thus improves its oxidation resistance. Furthermore, the effect of incorporated B on oxidation resistance is also reflected in the AlCrBN coating system.
Schematic of oxidation process of AlTiBN coating [136]
The results suggest that the benefits of improved oxidation resistance are more pronounced at higher temperatures, particularly above 900 °C, making it a relevant consideration for applications in high-speed dry cutting.
5.4 Application for Ti-alloys machining
Some researchers enabled studies on application of Boron-containing coatings for Ti-alloys machining. During dry turning against Ti-6Al-4 V, TiB2 coating was reported to improve the tool life by 60% and 73% in comparison to uncoated and TiAlN-coated tools, respectively. This was attributed to its outstanding plasticity and adhesive strength, avoiding rapid brittleness fracture and peeling during cutting, along with the in-situ formation of B-O liquid tribo-film at high temperatures, playing a role similar to cutting fluids [137]. According to high-speed dry machining tests against Ti-6Al-4 V, the cutting forces of AlTiBN coated tools were reduced by 10 ~ 20 N compared with Boron-free AlTiN coated ones, and the reducing effect was more pronounced with increasing Boron content; nonetheless, the durability of AlTiBN coated tools was intensified when minor (3.9 at. %) or mediate (5.0 at. %) content of B was adopted owing to enhanced thermal stabilities and hot hardness, while a further increasing of B content to 7.9 at. % promoted thermal decomposition of coating and surge of soft h-AlN phase, ultimately deteriorating the cutting performance of coated tool [138]. Meanwhile, synthesizing Boron-containing coatings with multilayer and/or gradient structure was considered as superb strategies improving the cutting performance against Ti-alloys further. As evidenced by high-speed dry turning of Ti-6Al-4 V (cutting speed: 100 m/min), the AlTiN/AlTiBN multilayer coating prolonged the lifetime by 180% in comparison to AlTiN coating, which was more significant than monolithic AlTiBN coating (81%), the coherent interfaces between AlTiN/AlTiBN sublayers enhanced mechanical properties such as micro-hardness and impacting fatigue resistance, resulting in splendid resistance to crack generation and propagation on the rake/flank face of tools, and finally longer service life [139].
6 Conclusion
In conclusion, the novel high-speed dry-cutting technology has successfully addressed challenges associated with traditional cutting methods, especially for the machining of steel and composite materials. However, the application of this technology to Ti-alloys is confronted with various difficulties due to their unique characteristics, such as working hardening, limited thermal conductivity, and high-frequency impacting behaviors, rendering severe attritions of coatings and tool bodies, and even poor cutting performance of coated tools under synergism of high-temperature adhesion, chemical diffusion, oxidation, and fatigue.
On the other hand, with increasing academic and industrial interests in protective hard coatings, their service properties have progressed successively to accommodate the development of modern machining technologies. The revolution of tool coatings has shifted from focusing on monotonic functions to multifaceted ones involving resistance to abrasion, adhesion, oxidation, and fatigue, along with high-temperature hardness and self-adaptability. Strategies involving chemical composition and structural design are usually adopted to achieve these complex service properties. Boron, in particular, is conducive to forming a nanocomposite structure in which nanograins are encapsulated by amorphous B-containing phases, irrespective of Ti-B-based or Al-Me-B-based hard coatings. In general, this structure brings excellent properties such as ultra-high hardness, enhanced thermal stability, and oxidation resistance. However, these properties are susceptible to the B content.
B-doped tool coatings exhibit self-adaptability for high-speed dry cutting of Ti-alloys due to the formation of in-situ lubricative B-O oxides that significantly reduce friction. Al-Me-B-based coatings (where Me = Ti, Cr, etc.), combining the self-lubrication effect of element B and outstanding high-temperature properties of element Al, exhibited significant improvements in lifetime and surface quality when cutting against Ti-6Al-4 V without the use of cooling fluids, compared to conventional AlTiN coatings. The optimization of Al-Me-B-based tool coatings with structural enhancements, such as constructing multilayer or gradient architecture, is expected to play a crucial role in improving toughness, adhesive strength, resistances to crack generation, propagation, and high-frequency impacts induced fatigue in this field.
Availability of data and materials
Data and material will be made available on request.
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This work was supported by the Natural Science Foundation of China (No. 52271047) and the Natural Science Foundation of Anhui Province (No. 2208085ME106).
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Wei Fang contributes to the Sections 1 and 5, Langlang Tao contributes to the Section 2, Benchi Zhang contributes to the Section 3, Ju Huang contributes to the Section 4, Fei Cai contributes to the conceptualization of whole article and Section 6, Gobinda Gyawali contributes to the review & editing of the whole article, Shihong Zhang contributes to the conceptualization of whole article.
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Fang, W., Tao, L., Zhang, B. et al. Design and application of B-containing hard coatings for high-speed dry cutting against Ti-alloys- A review. Surf. Sci. Tech. 2, 19 (2024). https://doi.org/10.1007/s44251-024-00048-3
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DOI: https://doi.org/10.1007/s44251-024-00048-3