Machining high-performance engineering materials, faces challenges including low machining efficiency, poor workpiece surface integrity, and rapid tool wear, which restrict high quality and efficient machining. Ultra-high-speed machining (UHSM) has been expected to address these issues. However, the material removal mechanisms involved in UHSM remain unclear and need further exploration. This paper reviews the criteria for evaluating the ductile and brittle behaviors of high-performance materials subjected to machining, as well as the developmental history of the material’s ductile-brittle transition induced by machining, proposing the concept of relativization of ductile-brittle property. Additionally, it further summarizes three typical material removal mechanisms: ductile-mode removal based on shear stress, brittle-mode removal based on tensile stress, and extrusion removal based on compressive stress, clarifying the universality of the brittle-mode removal. On this basis, this paper focuses on the discussion of the material removal mechanisms in UHSM, including high strain-rate-induced material embrittlement, UHSM-induced skin effect of damage, and the thermal effect in UHSM. Furthermore, it provides a detailed description of the typical characteristics of chip morphology in the ductile-brittle transition region (DBTR) under the high strain rate condition and, for the first time, elucidates the material removal mechanisms in the DBTR from a microstructural dislocation perspective, enriching the basic theory of UHSM. In the discussion section, it standardizes the definition for the UHSM, and explores the dislocation movement at high strain rates and the crack propagation in the UHSM. Finally, based on the current status of the UHSM technology, it summarizes the relevant research hotspots. For the first time, this paper brings up the brittle-mode removal mechanism under ultra-high-speed conditions, which is helpful to promote the UHSM for industrial applications.
Since the birth of machining technology, in response to the various needs of the societal development in the steam engine, electrical, and information eras, various traditional machining methods such as turning, milling, grinding, and drilling have been developed, and thus evolved into various non-traditional machining methods, such as ultrasonically-assisted machining (UAM), electrochemical machining (ECM), and laser machining. Nowadays, in the face of challenges in extreme environments, high-performance materials have emerged at the historic moment, posing increasing demands for machining. Currently, the primary methods for machining high-performance materials include laser-assisted machining (LAM) [1], [2], UAM [3], [4], cryogenic machining (CGM) [5], [6], [7], ECM [8], and ultra-high-speed machining (UHSM) [9], [10]. Among them, LAM can reduce material hardness and improve machinability by introducing energy to soften the material in the machining zone, while the generated thermal effect can lead to material ablation, melting, and re-solidification, potentially causing thermal stress, and even producing cracks in the material. The high-frequency vibration of UAM can generate local high stress and high strain rate on the workpiece surface, thereby accelerating the material deformation and removal, reducing the cutting force, and extending the tool life. Nevertheless, with an increase in the contact area between the tool and workpiece, the damping increases, and the ultrasonically auxiliary effect gradually weakens, which limits the machining depth to some extent and reduces machining efficiency. CGM and ECM can increase machining costs, and the electrolyte in ECM may cause environmental pollution. Therefore, it is important to reduce or even avoid additional machining damage while ensuring machining efficiency. UHSM has become the choice for meeting the above requirement based on the considerations of material embrittlement [11] and the skin effect of machining damage [12].
In this paper, machining that exceeds the critical speed for the ductile-brittle transition of a workpiece material is defined as UHSM [13], [14], while machining at a speed between that of UHSM and conventional machining (CM) is high-speed machining (HSM). UHSM is a key technology in the advanced manufacturing field. In recent decades, the application of motorized spindles has significantly improved the machining speed of machine tools, bringing UHSM into reality. UHSM has been used in the manufacturing of aircraft, automobiles, industrial molds and dies, and so forth, and has gradually become a hot topic in the manufacturing community. UHSM is an upgraded new generation machining technology based on HSM. It is a highly nonlinear thermo-mechanical coupling process, with a high strain rate on material deformation, notably short contact time between the tool and workpiece, and rapid stress release during the material removal, resulting in a completely different mechanism from that of CM.
The concept of UHSM originated from the conjecture curve of cutting temperature variation with cutting speed proposed by the German scientist Carl Salomon, who suspected that when the cutting speed exceeded a certain critical value, the cutting temperature would decrease, leading to a lower cutting temperature in the HSM/UHSM range [15]. Following this, numerous scholars conducted research on the material removal mechanisms in UHSM, including forces [16], heat [17], chip morphology [18], specific cutting energy [19], tool wear [20], and surface integrity [21]. However, the aforementioned research content mainly focused on macroscopic phenomena. A more in-depth understanding of the material removal mechanisms in the UHSM requires a comprehensive analysis from both a microscopic perspective in terms of physics and a macroscopic perspective in terms of thermodynamics.
The logical framework of this paper is based on above considerations and is illustrated in Fig. 1. The first part provides an overview of the criteria for evaluating the ductility-brittleness property of material, and introduces the ductility-brittleness transition induced by machining, proposing the concept that there are no absolutely ductile or absolutely brittle materials in the world. Instead, the concept of relativization of ductility-brittleness property is introduced. In the second part, the ductile-mode removal mechanism is discussed from the perspective of the adiabatic shear bands (ASBs), the brittle-mode removal mechanism is discussed from the perspective of the generation and propagation of cracks, and the extrusion removal mechanism is discussed from the atomic or close-to-atomic scale. The effects of size, strain rate, and temperature are considered in the analysis of the applicable range of material removal mechanisms, and the universality of the brittle-mode removal is emphasized. On the basis of the first two parts, the third part delves into the state-of-the-art material removal mechanisms in the UHSM, including high strain rate induced material embrittlement, “skin effect” of machining damage, and the thermal effect analysis. Multiple perspectives from the fields of physics and mechanics are involved. Moreover, based on the UHSM experiment conducted by the author, a new mechanism is herein proposed, namely ductile-brittle mixture removal, that is, the material removal mechanisms in the ductile-brittle transition region (DBTR) at a high strain rate, which supplements and enhances the UHSM theoretical framework. Furthermore, the discussion section further standardizes the definition for UHSM, mainly exploring the transition of dislocation movement state at high strain rates, as well as the impact of the UHSM on crack propagation. Finally, the research hotspots and future perspectives of UHSM are prospected.
2. Transition of material properties
Ductility or brittleness is an inherent property of materials, and is closely related to their engineering applications. For example, low-temperature brittle fracture leads to shipwrecks [22], and metal embrittlement during welding causes premature weld failure [23], all of which illustrate the universal existence of the ductile-brittle transition. Before discussing the machining-induced property transition, it is necessary to understand the criteria that determine whether a material exhibits ductile or brittle behavior.
2.1. Criteria for ductility-brittleness property
To date, it is widely accepted that the formation of shear bands is a typical feature of plastic deformation, while when a material is in a brittle state, the formation and propagation of cracks are mainly manifested. The ductile or brittle behavior was initially attributed to the relative value of energy associated with crystal cleavage and energy required for dislocation nucleation [24]. Materials with $G_{\text {disl }}<G_{\text {cleav }}$ exhibit ductility; otherwise, they show brittleness [25]. However, as and are difficult to obtain, this criterion is challenging to apply in practical applications. In recent years, Christensen [26], [27] has proposed a new theory for assessing the material state based on theoretical derivations to enable a simpler and more direct differentiation between ductile and brittle behaviors. He pointed out that the material’s ductility-brittleness property has an inherent connection with its elastic constant, with the macroscopic criteria represented by Eqs. (1), (2) and the microscopic criteria by Eq. (3).
where v is Poisson’s ratio, μ denotes the shear modulus, k represents the bulk modulus, and $k_{\mathrm{A}}$ and $k_{\mathrm{B}}$ represent the effective stiffness coefficients for bond stretching and bond bending/shear between atoms, respectively. These criteria have been widely confirmed in both pure metals [28] and metallic glasses [29]. Analyzing the ductility and brittleness of materials in terms of atomic bonds provides a novel viewpoint; however, a single parameter is insufficient to comprehensively describe the material’s ductility and brittleness. The above criteria are limited to homogeneous or isotropic materials, and only uniaxial tensile states are considered, making the conditions overly idealized. In comparison, Lawn and Marshall [30] proposed a brittleness index, comprehensively considering the dual impact of micro-hardness and fracture toughness. The dimension of the brittleness index is in the power of length. However, this index lacks relative rationality to evaluate the ductility or brittleness property by simply using the length dimension. In summary, there is a need for further improvement in the evaluation system for ductility and brittleness that is applicable to various materials under complex stress states during the machining.
2.2. Ductility-brittleness transition induced by machining conditions
There are no absolutely ductile or brittle materials in the world; the so-called ductility or brittleness of a material is relative. In the field of material preparation, researchers aim to lower the ductile-brittle transition temperature to ensure that the material maintain strength and toughness over a wide temperature range without easily failing. They also strive to enhance the material’s high-temperature resistance to maintain excellent performance in high temperature working environments. The development of materials for nuclear energy application also faces significant challenges, including the embrittlement failure of components owing to neutron irradiation [31], [32] and high-temperature helium embrittlement [33]. These problems require urgent resolution.
The ductility-brittleness transition mechanism significantly influences the tool wear, machined surface integrity [34], [35], material removal mechanisms, and energy consumption [36], which in turn affect parts’ quality, manufacturing costs, and production efficiency. In the machining field, whether the machined material behaves as ductile or brittle is closely related to the machining environment and parameters. It is possible to induce ductility-brittleness transition by changing the machining environment. Rao et al. [37] conducted experiments on the reaction-bonded silicon carbide ceramic using laser-assisted variable-depth machining. They found that the scratch length of plastic deformation increased with temperature. Korte et al. [38] revealed that for the micro-pillar compression experiment of single-crystal Si, the failure mode was mainly brittle fracture at room temperature, but completely transformed into plastic deformation at 500 °C. Ding and Hong [39] reduced the temperature of chip to the embrittlement temperature by cooling the cutting zone, thereby improving the chip-breaking ability of machining low carbon steel AISI1008. Kang et al. [40] improved the material property of a SiC workpiece surface through ion implantation, extending the ductile-mode removal range.
Meanwhile, the ductility-brittleness transition can be achieved by altering the machining parameters. Bifano et al. [41] conducted small-depth grinding experiments on engineering ceramics and optical glasses. They obtained a ductile-regime machining surface without cracks, and proposed the concept of “ductile-regime grinding” for hard and brittle materials. However, this concept is based on machining surface and does not reveal the essence of ductile-regime removal in hard and brittle materials. Liu et al. [42] and Zhang et al. [43] found that there are still a large number of micro-cracks below the smooth ground surface, which challenges the concept of “ductile-regime grinding.” Subsequently, Maas et al. [44], Li et al. [45], and Kumar and Melkote [46], found that under ductile regime removal, the subsurface deformation of hard and brittle materials consists of three parts: plastic flow zone, micro-crack zone and median cracks, in which no cracks occur in the plastic flow zone. Ductile-mode removal is comprehensively determined from the two aspects of surface and subsurface topography.
In the field of brittle-mode removal of ductile materials, Eda et al. [47] first proposed that when the cutting speed exceeds the static plastic wave propagation speed, the plastic strain is greatly restrained. Zhou et al. [48] found no apparent plastic flow in the metal’s subsurface after the cutting speed exceeded the static stress wave propagation speed. Liu and Su [49] discovered that, as the cutting speed increased, the workpiece material became brittle, with extremely small non-ductile particulate chips at a cutting speed of 116.67 m·s−1. In addition, the development of bulletproof armor equipment, and the application of ultrasonically-assisted condition [50] are related to the ductility-brittleness transition.
3. Typical material removal mechanisms and application range
Based on dominant stresses in the material removal, the existing material removal mechanisms can be classified into ductile (shear stress dominant), brittle (tensile stress dominant), and extrusion (compressive stress dominant) modes.
3.1. Ductile-mode removal mechanism
In the ductile-mode removal, the abrasive particle/tool cuts into the workpiece to create a plastic deformation zone. The internal stress accumulates continuously and eventually reaches the material’s yield strength, forming continuous chips. Currently, the theory of ductile-mode removal focuses on shear removal. In this process, shear stress dominates, and the material is removed along a specific shear plane. The adiabatic shear phenomenon was first discovered in the plastic flow of steel by Zener and Hollomon [51]. Subsequent research on ASBs was mainly based on the microscale dislocation theory and mesoscale void theory. The essential reason for a material undergoing plastic deformation is the dislocation evolution. As the machining speed increases within the CM, the dislocation slip speed increases. To adapt to the high slip speed, the random vibration of the atoms intensifies, and the energy cannot be transmitted to distant regions in a short time, resulting in significant local heating. An adiabatic effect comes into play, and the adiabatic shear frequency increases with an increase in machining speed [52], [53]. In addition to dislocation movement, twins [54], [55] and phase transformation [56] also play significant roles in the ductile-mode removal.
During the plastic deformation, ductile fracture occurs when the shear stress exceeds the material’s yield strength [57], [58]. Currently, micro-voids are regarded as the main cause of ASBs failure, with Gurson [59] being the first to consider the influence of micro-void fraction in describing material behavior. Subsequently, many scholars conducted studies on the nucleation and growth of micro-voids [60], [61], [62]. Nucleation primarily occurs in two aspects: for single crystals, micro-voids form at dislocation networks and vacancy clusters; for alloys and polycrystals, impurities, second phases, and grain boundaries are the main locations [63]. The evolution of micro-voids is shown in Fig. 2 [64], [65]. Under the influence of stress, micro-voids nucleate within the shear band and grow with increasing stress (segment AB). When the stress reaches point B—the material yield strength, necking occurs between adjacent micro-voids. The growth of micro-voids continues to increase the micro-void fraction, ultimately leading to coalescence (segment BC). When the fraction reaches a threshold, the ductile material releases energy in the form of cracks (segment CD), resulting in ductile fracture. Dimples is a typical feature of the ductile fracture.
3.2. Brittle-mode removal mechanism
The premise of brittle-mode removal is that micro-cracks can nucleate and propagate. Transgranular or intergranular cracks are typical characteristics. Therefore, the nucleation and propagation mechanisms of brittle cracks form the foundation for research on brittle-mode removal. During the machining, the material is subjected to the impact of an abrasive point, resulting in a sharp increase in stress near its surface. Compared with ductile materials, hard and brittle materials mostly dissipate stress in the form of cracks, with only a minimal portion of the stress being borne by plastic deformation. As a result, the efficiency of the spatial redistribution of stress is relatively low. For polycrystalline materials, stress concentration often occurs at grain boundaries’ intersections. For single crystals, the stress concentration at dislocation piles up owing to the significant micro-deformation. When the stress reaches the fracture strength, the initial atomic bond breaks, and the stress released after bond breakage is transmitted to the adjacent atoms. Under the domino effect, a series of atomic bonds breakage occurs, forming micro-cracks [66], [67]. These micro-cracks further penetrate to form macro-cracks, and the crack tip undergoes further propagation through a dislocation kink mechanism [68], [69].
The brittle-mode removal mechanism originates from the indentation-fracture mechanics theory of hard and brittle materials. The distribution of ideal cracks is shown in Fig. 3 [11]. Under the influence of high-pressure stress, an elastic-plastic deformation zone initially forms near the abrasive. When the tensile stress reaches the fracture strength of the workpiece material, median cracks are generated and extend deep into the subsurface. The strain mismatch between the inelastic deformation zone and the surrounding elastic deformation zone results in a complex residual stress field, determining the final length of crack propagation [70]. During the unloading stage, the median crack recovers, and the surrounding region prevents this process, generating shear stress. The shear stress is integrated along the edge of the inelastic deformation zone and reaches its maximum value near the middle of the arc, where the crack nucleates and further expands to form a lateral crack [71]. Based on the energy balance principle and the significant hindrance of grain boundaries to crack propagation, lateral cracks extend toward the lower-energy free surface and eventually form chips [72]. Compared with the ductile-mode removal, brittle-mode removal requires lower cutting forces. This is partly because the impact resistance of hard and brittle materials is much lower than that of ductile materials; moreover, the friction coefficient between brittle materials and tool is lower [73]. Therefore, brittle-mode removal often consumes less energy, aligning with the development concept of low-energy industrial production [74].
3.3. Extrusion removal mechanism
At a macroscopic scale, material removal is dominated by either shear or tensile stress. While at a microscopic scale, compressive stress plays a critical role [75], [76]. In the nano-cutting, the cutting depth $t_{\mathrm{uc}}$ is not significantly bigger or even smaller than the tool edge radius $R_{\mathrm{t}}$, resulting in a negative effective rake angle $\gamma_{\mathrm{E}}$. Chip formation is mainly caused by extrusion [77], [78]. As shown in Fig. 4(a) [79], extrusion removal has two typical characteristics. First, high-pressure stress concentration region exists at the front end of the tool edge radius, where hydrostatic pressure can effectively inhibit the void nucleation, and exceeding the threshold can induce phase transformation [80]. The second characteristic is a stagnation region relative to the cutting tool, which separates the atoms forming the chip from those being pressed into the workpiece [79]. Single atomic layer removal can be achieved as shown in Fig. 4(b) [81].
The effective rake angle $\gamma_{\mathrm{E}}$ formed between the tool and workpiece is shown in Fig. 4(c) [77], and its calculation using Eq. (4) can be deduced, where $\gamma_{\mathrm{w}}$ represents the atomic radius of the workpiece. However, in the atomic-scale cutting (ASC), owing to its smaller cutting depth, another effective rake angle $\gamma_{\mathrm{e}}$ formed between the atom at the very bottom of the tool and the workpiece atoms should also be considered, as shown in Figs. 4(d)–(f) [77]. When the $t_{\mathrm{uc}}$ is less than the sum of the atomic radii of the tool and workpiece, $\gamma_{\mathrm{e}}$ is a negative rake angle; otherwise, it is positive [77]. According to Eq. (5), the impact of $\gamma_{\mathrm{e}}$ on the ASC is primarily achieved through $t_{\mathrm{uc}}$. When the $t_{\mathrm{uc}}$ is small, the workpiece undergoes only elastic deformation, making it impossible to achieve effective atomic removal [82]. Alternatively, it may lead to partial atomic removal from the workpiece, where residual atoms can form flaky defects, significantly deteriorating the surface morphology [81], [83]. When the $t_{\mathrm{uc}}$ is large, severe plastic deformation occurs, introducing subsurface damage (SSD).
where the $C, \bar{O}_{\mathrm{t}}$, and $O_{\mathrm{w}}$ is the point in Fig. 4.
Chips formation of ASC is based on dislocation slip, which is fundamentally different from that primarily based on shear stress of the CM [84], [85]. Additionally, shear stress is present in the extrusion removal but does not participate in chip formation [86]. In ASC, where the de Broglie wavelength is greater than the system scale and particle spacing, the machining becomes discrete. Therefore, the mechanisms involved cannot be explained solely from the classical Newtonian mechanics, and achieving single atomic layer removal requires considering the influence of quantum theory, which is also the challenge faced in Manufacturing Paradigm III [87], [88].
3.4. Application range of material removal mechanisms
The material removal is influenced by the combined effects of strain rate, size, and temperature. Fig. 5 [89] shows the relationship among these three effects on the deformation behavior. First, the size effect mainly manifests in the chip load, where a smaller cutting depth leads to higher strain rates on the workpiece; simultaneously, a smaller cutting depth can intensify the shear deformation of the workpiece, thereby enhancing the temperature gradient effect. Moreover, an increased strain rate can alter material properties in the machining zone, subsequently enhancing or weakening the temperature effect. These three effects interact and constrain each other, jointly affecting the dislocation movement. Among them, the smaller the scale is, the fewer defects occur, making it easier for dislocations to overcome obstacles; the higher the temperature is, the more chance that dislocations climb over energy barriers. These findings indicate that both size and temperature effects can promote dislocation movement, leading to more plastic deformation, shown as ductile-mode removal. However, higher strain rates make dislocations more prone to jamming, leading to the nucleation and propagation of cracks, resulting in brittle-mode removal.
It is feasible to discuss the applicability of material removal mechanisms from these three effects [89]. Fig. 6 illustrates the impact of the strain rate, size, and temperature on the material removal mechanisms. The centroid of the triangle formed by connecting any state of the three effects corresponds to the material removal mechanisms. During the CM, ductile materials undergo ductile-mode removal, while hard and brittle materials correspond to brittle-mode removal. Therefore, applicability of these two removal mechanisms needs to be compared under different machining conditions. The strain rate effect on the material removal mechanisms is mainly reflected in the machining speed. Once the machining speed exceeds a critical value, brittle-mode removal can be achieved on ductile materials. In terms of the size effect, hard and brittle materials are usually machined at a very small scale to achieve ductile-mode removal. In addition, the temperature effect is reflected in auxiliary machining. Ductile-mode removal of hard and brittle materials typically involves heat assistance, with preheating the workpiece below the phase transformation temperature. Brittle-mode removal of the ductile material typically involves low-temperature cooling. Based on the analysis of the strain rate effect, the brittle-mode removal has a broader range of applicability. Furthermore, the division of brittle-mode and ductile-mode removal is based on the mechanical deformation of materials. In contrast, only recoverable deformation occurs in the extrusion removal, which is fundamentally different from either brittle or ductile-mode removal. Therefore, current research on extrusion removal only considers the size effect, and extrusion removal occurs when the cutting depth is at the atomic or close-to-atomic scale.
Compared with the ductile-mode removal, the brittle-mode removal features high-efficiency and high-quality machining and is more versatile. The high-precision machining capabilities of ASC can achieve non-damage removal, making it crucial in high-end precision manufacturing fields. However, the edge radius of the current tool is often at the micron level. Achieving a transition from micrometers to nanometers or even angstroms level still requires an amount of time, resulting in a considerable gap in the practical industrial applications.
4. Material removal mechanisms in UHSM
4.1. Material embrittlement induced by high strain rate
In the field of impact engineering, Dummer et al. [90] discovered that as the strain rate increased, deformation mechanisms of polycrystalline tungsten transitioned from dislocation slip, twinning, to intergranular fracture. Klepaczko [91] found that when the loading strain rate exceeded 104 s−1, the maximum shear strain of C-Cr-Mo hot-rolled steel rapidly decreased. Meanwhile, when the impact velocity exceeded 100 m·s−1, the fracture energy of VAR 4340 steel dramatically decreased. These abrupt changes in mechanical properties indicate that the material undergoes a ductile-brittle transition. Escobedo and Gupta [92] and Tang et al. [93] (Fig. S1 in Appendix A) initially observed that, at a lower impact speed, metallic glass exhibited ductile fracture with relatively rough fracture surfaces and clear ductile dimple features. At a higher impact speed, the fracture surface became smoother and displayed a cup-and-cone structure. In the field of machining, the phenomenon of chip formation, transitioning from continuous to serrated, and then to fragmented chips with increasing machining speed, indicates the embrittlement of ductile metallic materials [49], [94], [95] (Fig. S2 in Appendix A).
4.1.1. Mechanical analysis
According to the spontaneous emission of dislocations at a cleavage crack, Rice and Thomson [96] established a criterion for brittle fracture of crystalline materials:
$\frac{G b}{\gamma_{\mathrm{s}}} \geq 7.5-10$
where G represents the shear modulus, b denotes the Burgers vector, and $\gamma_{\mathrm{s}}$ means the material’s surface energy. Eq. (6) quantitatively separates brittle fracture from ductile fracture. However, owing to the difficulty in accurately obtaining the material’s surface energy $\gamma_{\mathrm{s}}$, the above equation is challenging to apply in practical scenarios.
For ductile materials, stress-strain curves at different strain rates are shown in Fig. 7. With an increase in strain rate, the tensile and yield strengths rise sharply, while the plastic failure strain decreases noticeably. Furthermore, as the strain rate increases, the material transitions from the plastic region to the elastic region, indicating that a ductile-brittle transition occurs as long as the strain rate is sufficiently high [9].
Whether a material exhibits ductility or brittleness depends on the ratio of the theoretical shear strength to the theoretical tensile strength. This can be further summarized as the difference between ultimate tensile strength $\sigma_{\mathrm{b}}$ and yield strength $\sigma_{\mathrm{s}}$ at different strain rates [97] (Fig. S3 in Appendix A). Yang and Zhang [11] (Fig. S4 in Appendix A) summarized the difference $\Delta \sigma$ for different materials at various strain rates, with their corresponding relationship expressed as Eq. (7).
where $k_{0}$ is a constant determined by a specific material, $k_{s}$ represents the strain rate sensitivity coefficient of material embrittlement, and $\dot{\varepsilon}$ is the strain rate.
The influence of strain rate on hard and brittle materials follows the same trend. Lawn and Marshall [30] proposed using the ratio of hardness to fracture toughness to assess the degree of material brittleness (Bm), as shown in Eq. (8).
$B_{\mathrm{m}}=\frac{H}{K_{\mathrm{C}}}$
where H represents the microhardness, and $K_{\mathrm{C}}$ indicates the fracture toughness.
The relationship between the microhardness and strain rate is described by Eq. (9) [98]. With the increase in strain rate, the dislocation density relatively increases, resulting in an increase in microhardness owing to the strain rate strengthening effect.
Fracture toughness represents the material’s ability to resist crack propagation. It is often quantified as the energy absorbed by the material prior to fracture. Therefore, as shown in Eq. (10), the fracture toughness ($K_{\mathrm{C}}$) can be expressed as the integral of stress with respect to strain, where stress is influenced by strain (${\varepsilon}$), strain rate ($\dot{\varepsilon}$), and temperature (T). It can be expressed as $\sigma=f(\varepsilon, \dot{\varepsilon}, T)$.
where $\varepsilon_{f}$ represents the fracture strain. Generally, with an increase in strain rate, the material’s strength rises; however, the material’s fracture strain significantly decreases, leading to a decrease in fracture toughness, as shown in Fig. 7. In most materials, hardness and fracture toughness generally exhibit an inverse relationship, the brittleness degree increases with an increase in the strain rate.
4.1.2. Stress wave analysis
Stress wave theory is effective for a comprehensive understanding of the UHSM. Von Karman and Duwez [99] first proposed, considering the effect of stress waves, that under a tensile impact exceeding a critical speed, ductile materials would experience brittle fracture, with plastic strain being negligible. To investigate the critical conditions for ductile-brittle transition, Eda et al. [47], Zhou et al. [48], and Wang et al. [97] conducted UHSM experiments on metallic materials. They concluded that, by increasing the machining speed, the propagation speed of the stress wave could eventually exceed that of the plastic wave (Eq. (11)), causing a lag in plastic deformation and a reduction in the tendency for hardening, leading to the material exhibiting brittleness.
where V represents the machining speed, $v_{\mathrm{p}}$ indicates the propagation speed of the plastic wave, and $\rho_{0}$ is the material’s density.
Currently, two main experimental methods exist for studying stress wave propagation. One involves direct observation through photoelastic imaging techniques, which requires the tested components to have high transparency. The other indirectly infers the stress wave propagation by analyzing crack distribution. In addition, some scholars have used simulation methods to analyze stress wave propagation. The most significant feature of simulation is its ability to characterize the instantaneous dynamic propagation of the stress wave. Liu et al. [100] (Fig. S5 in Appendix A) used the dislocation density model to update the material’s constitutive behavior and introduced a strain rate effect term to simulate HSM. By comparing the CM and HSM conditions, they found that the primary deformation direction shifted from shear along the 45° direction to almost perpendicular to the tool movement direction. The finding indicated that shear deformation is suppressed with an increase in machining speed. Jiang et al. [101] used time-resolved photoelastic imaging techniques to observe the propagation of elastic stress waves in a polycarbonate workpiece and conducted simulations by establishing the lattice model. Figs. 8(a) and (b) show the experimental and simulated stress propagation, demonstrating good consistency. By comparing the stress variations at positions S1-S3 at different speeds (Fig. 8(c)), it is evident that, as the cutting speed increases, significant fluctuation in stress occurs at all points at a speed of 75 m·s−1. It is worth noting that, at a cutting speed of 75 m·s−1, owing to the superposition of different stress wave, the stress field exhibits a strong distortion. The stress states at S1 and S2 undergo a turning point from compressive to tensile stress before and after the stress wave, contributing to crack generation. Additionally, the stress wave distribution is primarily concentrated at the tool-workpiece contact surface and the workpiece’s free surface, making it essential to explore the influence mechanism of the workpiece’s free surface on the stress wave propagation.
According to Hopkinson [102], compressive stress waves can convert into tensile stress waves through reflection at the free surface. Wang [103] elaborated on the stress distribution after the reflection of compressive stress waves at the end of the rod, and found that the surface reflection can alter the stress wave’s propagation path and subsequently change the internal stress state of the workpiece (Fig. S6 in Appendix A). Zhang et al. [104] corroborated this viewpoint by simulating crack propagation paths inside cast iron HT250 at different cutting speeds using the discrete element method (Fig. S7 in Appendix A). With increasing cutting speed, the location of crack initiation changes from the tool edge region to an area near the workpiece’s free surface, accompanied by a significant increase in the number of cracks. Similar phenomena were observed in Hopkinson bar tests on granite, where cracks were concentrated near the free surface [105]. The effect of surface reflection on changing stress states similarly applies to scenarios involving high-energy impacts, such as explosions and laser shock [63]. Moreover, during the UHSM, the loading and unloading become more intense and concentrated, warranting further in-depth research into their influence on stress wave propagation.
4.2. UHSM induced skin effect of damage
There are many practical examples of the skin effect of damage in mechanical machining. Zhang and Yin [12] summarized the SSD depth for various hard and brittle materials at different speeds. They proposed a linear relationship between the SSD depth and the power of the strain rate, as seen in Eq. (12), indicating that with an increase in the strain rate, SSD depth (δ) exhibits an overall linear decreasing trend.
This rule is also applicable to ductile materials. With an increase in grinding speed, the depth of dynamic recrystallization zone (DRXZ) in the Al6061T6 alloy significantly decreases [106] (Fig. S8 in Appendix A). The same phenomenon also occurs in composite materials [107] (Fig. S9 in Appendix A).
In UHSM, a noticeable characteristic of damage is “small range, large gradient,” where state variables are confined to a small area but have a large span. The propagation of stress waves follows the same pattern. Fig. 9 shows the internal stress distribution of the workpiece at different cutting speeds [101]. Under CM (25 m·s−1), high-intensity stress is generated ahead of the cutting tool and propagates along the primary deformation zone (PDZ). With an increase in speed, the stress wave undergoes a nonlinear change. The energy of the stress wave becomes more concentrated, leading to more severe deformation, but with a reduced influence range. Moreover, the propagation path changes to align with the cutting direction, and the normal stress gradually replaces the shear stress. However, the above analysis mainly focuses on the stress wave behavior induced by the strain rate effect in ductile-mode removal. At higher strain rates, the workpiece surface becomes brittle and ultimately fractures. Countless tiny cracks are distributed near the workpiece surface, forming a distinct boundary with the continuous medium inside the workpiece, further impeding the propagation of stress waves into the workpiece’s interior.
4.3. Analysis of thermal effect in UHSM
The energy consumption during material removal includes several forms: external friction, plastic deformation, new surface formation, and transfer of momentum. Among these, external friction and plastic deformation dominate, and are eventually converted into thermal energy. This section discusses the generation and transfer of heat from a fresh perspective on brittle-mode removal, providing a clearer understanding of the thermal effect in UHSM.
4.3.1. Generation and transfer of cutting heat
To distinguish the similarities and differences in heat generation between brittle and ductile-mode removals, it is necessary to first understand the generation and transfer of heat in machining systems under ductile-mode removal mechanism. When a material undergoes elastic deformation, the energy required for deformation is stored in the form of strain energy. In contrast, when the material undergoes plastic deformation, the energy required for deformation is dissipated in the form of heat. As shown in Fig. 10(a) [108], in the ductile-mode removal, the sources of cutting heat mainly include plastic deformation in the PDZ, friction between the tool and chip in the secondary deformation zone (SDZ), and friction between the tool and workpiece in the tertiary deformation zone (TDZ). Among these, plastic deformation in the PDZ generates internal frictional heat $Q_{\mathrm{ps}}$, which is closely related to the dislocation movement and accounts for approximately 75% of the total heat generation. Assuming that all mechanical work ($W_{\mathrm{c}}$) is transformed into heat, the equation for calculating $Q_{\mathrm{ps}}$ is given by Eq. (13). External frictional heat $Q_{\mathrm{sd}}$ in the SDZ constitutes about 20% of the total, with its calculation equation provided in Eq. (14). Heat generation $Q_{\mathrm{td}}$ in the TDZ is extremely minimal and is typically neglected. Heat is carried away by chip $Q_{\mathrm{c}}$, with a portion being transferred to the workpiece $Q_{\mathrm{w}}$ and tool $Q_{\mathrm{t}}$, satisfying Eq. (15).
where $F_{\mathrm{V}}$ represents the tangential cutting force along the cutting speed direction, denotes the chip thickness ratio, and $F_{\mathrm{fr}}$ is the total shear force acting on the tool’s rake face.
In the early stage of ductile-mode removal, with an increase in cutting speed, material plastic deformation intensifies, leading to increased heat generation in the PDZ and SDZ. This results in an overall rise in heat. However, the heat cannot infinitely increase; thus, after it reaches the maximum, further increasing the cutting speed has minimal impact on heat generation. Nonetheless, the contact time between the chips and tool continues to shorten. During the machining process, approximately 70%-90% of the heat is absorbed by the chips [109], [110]. From an energy balance perspective, heat is transferred from the high-temperature chip area to the workpiece and tool through thermal conduction and convection. Consequently, as the cutting speed increases, the majority of the heat is trapped inside the chips and is carried away, resulting in an overall reduction in the heat in the cutting region [111]. Therefore, increasing the cutting speed is beneficial for reducing the heat entering the tool and workpiece. The high-speed hobbing experiment conducted by Ueda et al. [112] (Fig. S10 in Appendix A) effectively illustrated above points.
Figs. 10(b) and (c) illustrate the heat distribution under CM and UHSM [108]. The material embrittlement is severe under UHSM, which inhibits the shear slip of dislocations. This leads to energy consumption in the form of atomic/molecular bond breakage rather than the heat generated through internal friction. Therefore, under the brittle-mode removal of ductile material caused by UHSM, the heat generated by plastic deformation in the PDZ can be neglected. For the heat generated by external friction in the SDZ, more heat is generated due to the friction between the tool and chip during the ductile-mode removal stage. As the cutting speed increases, the material enters the brittle-mode removal, forming fragmented or powder chips, rather than sliding out along the tool’s rake face. This significantly reduces the contact area between the tool and chip, effectively avoiding the friction between them, and the heat generation in this region is reduced.
In the brittle-mode removal stage, with an increase in machining speed, material fragmentation intensifies, more kinetic energy is consumed for the fracture of chips rather than being transformed into heat through plastic deformation, and plastic flow is further suppressed, resulting in smaller and smaller chip deformation and size [113]. Moreover, the higher machining speeds, the higher the dispersion speed of fragmented chips; and more heat is taken away. Additionally, the substantial separation and disintegration of chips lead to a reduction in external frictional heat, and the increased machining speed accelerates airflow circulation in the tool-chip contact area, further enhancing heat dissipation.
4.3.2. Localization of thermal effect
When the machining speed exceeds the critical speed of the ductile-brittle transition, the removal mechanism changes from ductile-mode to brittle-mode removal. It is important to emphasize that brittle-mode removal only refers to the way that chips leave the workpiece in a brittle manner, while the newly formed machining surface is still influenced by the ductile removal characteristics owing to external frictional heat generated by TDZ [108]. The brittle-mode removal shares some similarities with preheating treatments or LAM, all of which can effectively reduce cutting forces through the thermal softening effect [114]. The difference lies in the extremely short contact time between the tool and workpiece in UHSM, and the temperature lag effect of strain rate response limits the penetration of heat deep into the workpiece. Moreover, the influence of phonon resistance at high strain rates cannot be ignored. The phonon resistance affects the material’s thermal conductivity, hindering the transfer of heat. These two factors together result in the localization of the thermal effect layer (TEL) [115].
The localized thermal effect can be supplemented by relevant experimental observations. By detecting the oxygen element, Guo et al. [116] (Fig. S11 in Appendix A) found that, as the machining speed increases, the distribution of oxygen elements tends to be near the surface. Additionally, the higher the grinding speed, the more pronounced the trend of temperature attenuation [106] (Fig. S12 in Appendix A). As shown in Fig. 11, Liu and Su [49] found that a significant amount of melting chips is observed at 50 m·s−1. With increasing machining speed, the number of melting chips gradually decreases. Furthermore, this law can be observed through changes in the color of the chips, tool oxidation [49], [115] and variations in the oxygen content in the PDZ and SDZ [117].
4.4. DBTR at high strain rate
As the inherent property, ductility and brittleness have duality in themselves, but under the influence of external factors, there inevitably exists a DBTR. In the field of geology, the temperature variation leads to strain localization, and the reduction in grain size induced by recrystallization is the primary reason for the ductile-brittle transition of feldspar [118]. In the field of mechanics, the existence of DBTR can be confirmed from stress-strain curves [119] (Fig. S13 in Appendix A). At axial strain rates lower than 650 s−1, the curve exhibits the typical characteristics of ductile material, which is divided into three stages: a linear increase in axial stress, a nonlinear increase to the maximum, and a nonlinear decrease. At an axial strain rate of 950 s−1, the compression process is simplified into two stages: a linear increase and a nonlinear decrease in axial stress, with distinct brittle characteristics. While at strain rates of 750 and 850 s−1, influenced by the DBTR, the curves exhibit significant fluctuations and can be divided into four stages: a linear increase to a local maximum, a reduction to a local minimum, a subsequent increase to a local maximum of axial stress, and a nonlinear decrease leading to material failure. In the process of cutting, the occurrence of DBTR can be indirectly confirmed by recording acoustic emission signals [36], [40]. Currently, ductile-brittle transition behavior has been widely observed, but research focuses primarily on the critical condition for the DBTR, such as the energy required for the formation of dislocation loops at the sharp crack tip [96], [120], and whether dislocations formed around the crack can blunt the crack to shield further propagation [24], [121], [122]. There is almost no trace of the material removal mechanisms in the DBTR.
4.4.1. Chip morphology of DBTR
Fig. 12 illustrates the evolution of chip morphology for Inconel 718 alloy (Shanghai Pingpu Industrial Co., Ltd., China) from the ductile removal (130 m·s−1) to the DBTR (160 and 190 m·s−1), with a scratching depth of 5 μm. Figs. 12(a)-(c) show that at the scratching speed of 130 m·s−1, the chips exhibit the typical serrated chip morphology with a distinct adiabatic shear instability feature. However, during the DBTR, two distinct characteristics are evident: a significant reduction in chip size, and the coexistence of shear bands and cracks. As shown in Figs. 12(d)-(f), at the scratching speed of 160 m·s−1, the chips exhibit a regular fan shape with the local generation of tiny (ductile) cracks, but shear bands remain the main feature. Therefore, this speed can be considered as a slight extension of ductile-mode removal. At the scratching speed of 190 m·s−1, chip morphology becomes diverse, with claw-shaped, fan-shaped, and strip-shaped chips coexisting. The chips show clear (brittle) cracks penetrating the interior, without the presence of dimples, as shown in Figs. 12(g)-(i). Therefore, this speed can be regarded as the very beginning of brittle-mode removal. The diversity of chip morphology is essentially caused by the serious instability of the adiabatic shear, which is consistent with the significant stress fluctuations in Ref. [119]. According to the macro-scale chip evolution patterns and the results of molecular dynamics simulations [49], [65], [94], the overall chip size will further decrease, and irregular fragmentary chips are typical features of brittle-mode removal. Fig. 12(j) illustrates the schematic of chip morphology with changing machining speed, successively showing the continuous chip under CM, the serrated chip under HSM, transition stage chips in the DBTR, irregular fragmented chips under UHSM, and powdered chips at even higher speeds.
4.4.2. Material removal mechanisms in DBTR
Discussion about the material failure mechanism has mainly been divided into two categories: the models based on dislocation nucleation [96] and dislocation mobility [123], [124], [125]. The dislocation nucleation model posits that brittle fracture occurs when dislocations cannot effectively nucleate at the crack tip before it propagates. Some evidence negates this interpretation, suggesting that at a temperature slightly below the brittle-ductile transition temperature, dislocations can still nucleate at the crack tip, but the stress required to initiate dislocation nucleation is too small to support rapid dislocation movement. Therefore, it fails to blunt the crack propagation, resulting in brittle fracture [122], [124], [126]. However, the above analysis is based on the existence of pre-cracks. The existence of pre-cracks can significantly reduce the stress needed for material failure. Without pre-cracks, the critical stress required for failure is significantly higher than the Peierls stress. Once dislocations nucleate, they move rapidly, leading to ductile fracture [127]. Therefore, the dislocation nucleation model offers a more effective explanation for the transition from brittleness to ductility. Correspondingly, the criteria for determining whether ductile or brittle fracture occurs shift to the relative size of the nucleation energy barriers of dislocations and cracks [128]. During the DBTR induced by the high strain rate, the material has accumulated a large number of dislocations in the ductile stage. The model based on dislocation mobility plays a predominant role in explaining the mechanism.
According to the Orowan’s dislocation bypass mechanism [129], the probability of successfully bypassing an obstacle during dislocation slip is directly related to the size of the obstacle. With an increase in strain rate, dislocation density increases sharply. Moreover, shear deformation is suppressed, and extrusion deformation plays a dominant role. As a result, the dislocation movement changes from parallel shear slip bands to a disordered state, facilitating the entanglement of different dislocations and the formation of dislocation clusters or cells.
As shown in Fig. 13, when the size of dislocation clusters or cells is small, there is a high probability that dislocations will bypass them, and ductile-mode removal occurs. However, when the size is large or the quantity is substantial, the spacing between dislocation clusters/cells decrease significantly. Few of the dislocation line can successfully bypass the obstacles. Instead, they contribute to further nucleation and growth of dislocation clusters/cells. At this point, the energy expended by dislocation movement is insufficient to offset the energy. Some energy is transferred to promote the generation of small cracks [121], and the overall size of chips decreases significantly, marking the beginning of the DBTR. Subsequently, with further increases in the strain rate, dislocation nucleation and movement are further suppressed, leading to the formation of penetrating cracks, signaling the imminent end of the DBTR. The suppression of dislocation movement leads to a significant increase in the proportion of immobile dislocations, and eventually results in dislocation avalanches, producing a large number of crisscrossing cracks. Chips leave the workpiece in irregular fragments, thus entering the brittle-mode removal.
5. Discussion
5.1. Definition of UHSM
The definition of UHSM was initially divided into two types: one considering a spindle speed exceeding 60 000 r·min−1 as UHSM, and the other defining it as the machining linear speed surpassing 150 or 166.7 m·s−1 [130], [131]. The former is a definition based on spindle speed and is applicable for evaluating spindle equipment but cannot be used in the machining field. The latter definition ignores the influence of the workpiece material’s intrinsic properties during machining, and is not applicable to all workpiece materials. Therefore, defining the machining method that exceeds the critical speed of the ductile-brittle transition of the machined material as UHSM is more universally applicable and standardized, necessitating further refinement. In a machining system, the phenomenon of ductile-brittle transition is primarily manifested in the chip and workpiece. As illustrated in Fig. 11, Fig. 12, fragmented chips are a typical manifestation of material embrittlement and can serve as one of the characteristics of UHSM. On the other hand, according to the theoretical analysis of UHSM induced skin effect of damage, with increasing machining speed, the SSD depth of the machined workpiece gradually decreases but does not become infinitely small. After the occurrence of ductile-brittle transition, damage caused by plastic deformation will be significantly restrained, leading to fundamental changes in energy absorption, transitioning from dislocations to twinning [132]. Consequently, the SSD depth tends to stabilize, representing the second characteristic of UHSM.
5.2. Dislocation movement at high strain rates
Dislocations are crucial for coordinating material deformation. Their movement state under the high strain rate condition is the essence of material embrittlement and the skin effect of machining damage.
5.2.1. Strain rate effect on movement state of dislocations
According to Orowan’s theory, the relationship between dislocation movement and strain rate can be deduced as follows [133]:
$\frac{\mathrm{d} \varepsilon}{\mathrm{~d} t}=\rho b \nu$
where $\rho $ represents dislocation density, and $\nu $ is the velocity of dislocations [134], [135].
By substituting the strain in Eq. (16) with Eq. (17), the above equation can be transformed from a transient equation into a continuous one as Eq. (18), thereby representing the dynamic behavior of dislocations.
$\varepsilon=\rho b L $
$\frac{\mathrm{d} \varepsilon}{\mathrm{~d} t}=\frac{\mathrm{d}(\rho b L)}{\mathrm{d} t}=\frac{\mathrm{d} \rho}{\mathrm{~d} t} b L+\rho b \frac{\mathrm{~d} L}{\mathrm{~d} t}=\frac{\mathrm{d} \rho}{\mathrm{~d} t} b L+\rho b \nu^{\prime} $
where L represents the average displacement of dislocations, and $\nu^{\prime} $ is the dynamic velocity of dislocations. After transformation, the influence of strain rate on dislocations can be decomposed into two parts: the first term on the right side represents the static behavior of dislocations, and the second term represents the dynamic behavior. The velocity of dislocations is primarily driven by the applied shear stress $\tau $ and can be expressed by Eq. (19) [136]:
$C_{\mathrm{r}} \nu=b \tau $
where $ C_{\mathrm{r}}$ is the resistance coefficient related to lattice viscosity.
Fig. 14 shows the results of the molecular dynamics simulation for single-crystal SiC subjected to nano-scratching [137]. With an increase in scratching speed, the length of perfect dislocation lines reaches its maximum at 100 m·s−1, and before which, there is a certain regularity in the dislocation distribution. At a scratching speed of 200 m·s−1, the length of dislocation lines decreases, and their slip movement transforms into a disordered state. The same evolution also occurs in the SiC particles of the composite material at different grinding speeds [116] (Fig. S14 in Appendix A). This change is related to the previous mention that shear deformation is suppressed as scratching speed increases. It is worth noting that the depth of the damage layer at a scratching speed of 200 m·s−1 is significantly smaller compared to that at 150 m·s−1 (Fig. 14(e)). According to Eq. (19), the reduction in shear stress leads to a decrease in dislocation speed in the workpiece. As a result, dislocation movement is suppressed and confined to the shallow subsurface, exhibiting the skin effect of machining damage.
The influence of UHSM on dislocation movement is not only reflected in the shear stress, but also in terms of contact time. Increasing the strain rate can reduce the contact time between the tool and workpiece, thereby inhibiting dislocation nucleation and migration, to some extent reducing the SSD depth [138].
5.2.2. Immobile dislocations
Based on the movement state, dislocations can be categorized into two types: mobile and immobile dislocations. The density of mobile dislocations determines the material’s ductility, while the density of immobile dislocations dictates the brittleness. Currently, one of the typical theories describing dislocation movement is the thermal activation theory [139] (Fig. S15 in Appendix A). Building upon this theory, Follansbee and Kocks [140] introduced a generalized Eq. (20) to fit any representation of barrier potential shapes by altering constants p and q.
where $\Delta G_{0}$ represents the total energy required for dislocations to overcome obstacles, $ \Delta G_{\mathrm{a}}$ is the added thermal energy, $\tau_{0}$ represents the flow stress at 0 K temperature, and $\tau$ is the flow stress at different temperatures.
At the critical temperature T, the energy of dislocations exceeds the barrier energy $\left(\Delta G_{0}-\Delta G_{\mathrm{a}}\right)$, and the frequency of dislocation is $v_{1}$, which can be expressed by Eq. (21):
$ v_{1}=v_{0} A \exp \left(-\frac{\Delta G_{\mathrm{a}}}{k_{\mathrm{d}} T}\right)$
where A represents a value related to entropy change, $v_{0}$ denotes the vibration frequency of dislocations, $v_{1}$ is the frequency of dislocations which successfully overcome the barriers, and $k_{d}$ denotes the Boltzmann constant. Therefore, the waiting time for thermal activation of dislocations at the barrier can be represented as.
Material internal defects mainly include hard phases, voids, and grain boundaries. Regardless of these uncertain factors, the speed $v_{f}$ of dislocation movement between adjacent barriers satisfies Eq. (22) [141]:
$b \sigma_{\mathrm{f}}=D v_{\mathrm{f}}$
where D represents the drag coefficient, and $\sigma_{\mathrm{f}}$ is the driving force for the dislocation slip. Assuming l is the average distance between adjacent barriers, dislocation slip time from one barrier to an adjacent one $ V_{\mathrm{dis}}$ can be expressed as Eq. (23):
$t_{\mathrm{f}}=\frac{l}{v_{\mathrm{f}}}$
Dislocation migration between adjacent barriers experiences two states: thermal activation and dislocation slip states. Therefore, the average speed can be expressed by Eq. (24) [142]:
At low strain rates (< 104 s−1), dislocation thermal activation plays a major role $ t_{\mathrm{w}} \gg t_{\mathrm{f}}$ [143],, and dislocation resistance can be neglected; at high strain rates (104–107 s−1), the phonon drag dominates [65], [140], [144], [145], [146], $ t_{\mathrm{w}} \approx t_{\mathrm{f}}$ and dislocation emission requires overcoming a significant energy barrier. At this point,, and dislocation resistance cannot be ignored. Dislocation resistance increases with an increase in strain rate, restricting further dislocation slip. Moreover, owing to interactions between dislocations (entanglement, attraction, hindrance, etc), mobile dislocations are confined and gradually transformed into immobile dislocations. The slowdown of dislocation movement hinders the further nucleation of subsequent dislocations, making it impossible to effectively blunt microcracks formed during deformation. This leads to stress concentration at the crack tip and promoting crack propagation, resulting in the material exhibiting brittle characteristics [22], [147]. When the strain rate is high to a certain extent, dislocation avalanche may occur, leading to a sharp increase in the immobile dislocation density, and eventual embrittlement.
5.3. Crack propagation in UHSM
Compared with the CM, the material removal mechanisms in UHSM undergo a fundamental change, where larger impacts provide sufficient energy for crack nucleation, leading to a significant increase in crack nucleation density. In addition, tensile stress plays a dominant role, and the energy consumed by tensile-stress-dominated fracture is significantly smaller than that of shear-stress-dominated fracture, which is more conducive to crack propagation [148]. Although the crack propagation speed increases with the strain rate, it ultimately converges to the Rayleigh wave speed of 0.5-0.7. Relying solely on a single crack propagation cannot timely dissipate the enormous energy, thus forming a network of numerous branched microcracks ahead of the crack tip [149].
The crack propagation path is mainly influenced by the internal stress state and material resistance. The stress state is greatly affected by the strain rate. With an increase in strain rate, the contact time between the tool and workpiece greatly shortens, leading to a larger stress gradient and faster decay [101]. Meanwhile, the formation of a microcrack network results in rapid dissipation of the loading energy, causing the overall crack propagation to concentrate near the workpiece surface. The resistance magnitude is closely related to the internal microstructure. The elastic energy W consumed by workpiece deformation can be calculated using Eq. (25) [150].
$ W=\frac{1}{2} E(1-N)(\varepsilon r)^{2}$
where E is the elastic modulus, N denotes the number of activated microcracks, andγrepresents the radius of the elastic deformation region.
Therefore, the accumulation of elastic energy $ \Delta W$ and the activation of microcracks $ \Delta N$ satisfy Eq. (26):
$ \Delta W=\frac{1}{2} E \Delta N(\varepsilon r)^{2}$
Meanwhile, according to the Griffith’s criterion, the propagation of cracks needs to meet Eq. (27):
$ \Delta W \geq \Delta I$
where $\Delta I $ represents the surface energy required for crack propagation.
According to Eqs. (25), (26), (27), Fig. 15(a) [151] illustrates the mechanism of crack propagation, influenced by factors such as the orientation relationship between grains [152] and internal defects. The workpiece material can be classified into two regions: weak and strong. In the weak region, the surface energy required for crack propagation is relatively small, while in the strong region, it is larger. Under low strain rate conditions, the accumulation of elastic energy is relatively slow, making it difficult to reach the surface energy required for crack propagation in the strong area; however, under high strain rate conditions, the accumulation of elastic energy can simultaneously meet the requirements for crack propagation in both the weak and strong regions.
Figs. 15(b) and (c) further summarize the crack propagation paths of different materials under CM and UHSM conditions. For hard and brittle materials such as ceramic matrix composites, as illustrated in Fig. 15(b), taking Al2O3-SiO2 ceramics as an example. Some interfaces between SiO2 particles and Al2O3 matrix become weak regions owing to inadequate sintering, while the Al2O3 matrix, especially the SiO2 particles, serves as strong regions. Under CM conditions, initial cracks primarily form along cleavage planes, then mainly propagate along the interface between SiO2 particles and the Al2O3 matrix as this path consumes the least energy. The machining energy is concentrated on a limited number of cracks, thus leading to cracks propagating deep into the workpiece. Under UHSM conditions, the path of crack propagation undergoes fundamental changes owing to the influence of high strain rates, with more cracks extending through the Al2O3 matrix or even within the SiO2 particles [151]. Under high strain rate conditions, the composition difference between the matrix and the reinforcing phase decreases [116], reducing the gap in surface energy required for crack propagation, leading to a more uniform propagation path. As shown in Fig. 15(c), for ductile materials such as metals and alloys, a ductile-brittle transition occurs under high strain rate conditions, lacking sufficient ductility to suppress the propagation of microcracks. Simultaneously, the nucleation rate of twins in materials with low stacking fault energy increases significantly. The interaction between twins, as well as between twins and dislocations, accelerates the initiation and propagation of cracks at twin boundaries [153]. New twin interfaces become the preferred paths for crack propagation, favoring transgranular fracture [54], [154]. The abundance of twin boundaries significantly increases the propagation path of cracks [155]. Moreover, under low strain rate conditions, cracks are prone to deflection; however, under high strain rate conditions, the influence of inertia on the redistribution of fracture energy increases the probability of crack penetration through interfaces, leading cracks to propagate along straight lines [156], [157].
In general, the cracks distribution is few and deep after CM; while after UHSM, it is characterized by numerous but shallow cracks. Compared with CM, UHSM can achieve indiscriminate treatment of different materials, which also illustrates the universality of the brittle-mode removal.
6. Conclusions and outlook
This paper deals with the material property transition and three typical material removal mechanisms, and focuses on the material removal mechanisms in the UHSM, which is based on the relevant knowledge in the fields of thermodynamics and physics. The following conclusions are drawn:
(1) There are no absolute ductile or brittle materials in the world, and the so-called ductility or brittleness is relative, depending on loading or surrounding conditions.
(2) ASC adds extrusion removal to the typical material removal mechanism array, playing an important role in ultra-precision machining. However, the edge radius of a cutting tool is in the micron range at present, and achieving a transition from micrometer to nanometer or even angstrom level still requires a considerable amount of time for technological accumulation and breakthroughs.
(3) In the field of extreme manufacturing, the brittle-mode removal is more universal than the ductile-mode removal, and will play a vital role in the UHSM. Compared with ductile-mode removal, brittle-mode removal can effectively inhibit heat production and alleviate work hardening for ductile materials.
(4) At high strain rate, the probability of dislocation entanglement and pinning increases, causing mobile dislocations to transform into immobile dislocations, which induces the embrittlement of ductile materials and aggravates the embrittlement degree of hard and brittle materials. This is of great significance to reduce the cutting force and prevent built-up edges or grinding blockage.
(5) In the UHSM, the stress transfer changes significantly, and the dislocation movement is no longer the ordered movement dominated by shear stress (dislocation slip band form) but the disordered movement dominated by normal stress, making it easier for dislocations to entangle and pin to form clusters, limiting them to the near-surface layer and resulting in the skin effect of machining damage.
(6) Owing to the effect of material embrittlement in UHSM, most of the heat generated by internal friction and partial external friction would be suppressed. At the same time, the contact time between the tool and the workpiece is greatly reduced, resulting in a significant reduction in machining problems caused by temperature.
(7) In the DBTR, the mechanical properties of materials fluctuate obviously, and the chip morphology becomes diverse, which is not conducive to the stable machining. Therefore, it is advisable to avoid material removal during the DBTR.
(8) The model based on dislocation mobility, modified by the dislocation bypass mechanism, can provide a reasonable explanation for the material removal mechanisms in the DBTR induced by the high strain rate.
In addition to the mechanisms of material embrittlement, skin effect of machining damage and thermal effect localization discussed above, there are many other aspects that need to be further explored in UHSM. For example,➀ based on the concept of relativization of ductile-brittle property mentioned earlier, UHSM can narrow the difference between ductile and brittle components for composite materials, thereby enhancing the post-machining surface integrity [116]. ➁ Under the influence of the high strain rate, UHSM can alter the microstructure of the workpiece subsurface, which fundamentally differs from CM. ➂ For ASC, under the condition of ultra-high-speed, the extrusion effect also transforms into hydrostatic stress-dominated isobaric effect, where atomic layers are ultimately removed in a jetting manner [65], and no pile-up occurs on both sides, thereby evolving into the UHSM, further elucidating the universality of brittle-mode removal. ➃ Current material constitutive models used in machining simulations assign significant weight to the temperature-softening term. However, the temperature effect becomes increasingly less pronounced during brittle-mode removal. Therefore, a material constitutive model suitable for the mechanism of brittle-mode removal in the context of UHSM is urgently needed.
Furthermore, there are many pending hot issues in other fields related to UHSM, such as real-time monitoring of dynamic balance [158], aerodynamic study of the air barrier under UHSM conditions [159], [160], development of high-pressure cooling systems [161], [162], ultra-high-speed grinding wheels [163], [164], [165], and temperature measurement devices suitable for UHSM [166].
CRediT authorship contribution statement
Hao Liu: Writing - review & editing, Writing - original draft, Validation, Methodology, Investigation, Formal analysis, Data curation, Conceptualization. Jianqiu Zhang: Formal analysis, Data curation, Conceptualization. Qinghong Jiang: Formal analysis, Data curation, Conceptualization. Bi Zhang: Conceptualization, Supervision, Project administration, Funding acquisition.
Declaration of competing interest
The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.
Acknowledgments
This work was supported by the Shenzhen Science and Technology Innovation Commission (KQTD20190929172505711, JSGG20210420091802007, GJHZ20210705141807023, JSGG20220831110605009, and JCYJ20210324115413036), the Guangdong Basic and Applied Basic Research Foundation (2021B1515120009), and the Department of Guangdong Science and Technology (2019JC01Z031).
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