Progress in Mechanical Properties of Gradient Structured Metallic Materials Induced by Surface Mechanical Attrition Treatment

Abstract

It is well known that the bulk nanostructured metallic materials generally exhibit high strength but poor ductility, which greatly hinders their applications. In most cases, material failures usually start from the surface. Therefore, the surface modification is crucial to improving the mechanical properties of metallic materials. The surface mechanical attrition treatment (SMAT) is one of the most effective surface modification methods. It can be used to manufacture gradient nanostructured materials that there is no interfaces between the surface and the coarse-grained matrix. Additionally, the gradient nanostructured metallic materials fabricated by SMAT exhibit a preferable combination of strength and ductility compared with conventional homogeneous materials. It is generally believed that the high strength of the SMAT-ed metallic materials is owing to the surface fine-grain strengthening. Whereas the improved ductility can be attributed to the coarse-grained matrix and the superior work hardening ability of the gradient structured (GS) materials.

In this overview, the research progress of the GS metallic materials fabricated by SMAT is summarized. It mainly introduces the microstructure characteristics of the GS layer and the mechanical properties of GS metallic materials. Finally, in order to find the optimal match between the strength and ductility in the GS materials, the several factors affecting the mechanical properties of GS materials are summarized.

1. Introduction

In the past decades, bulk nanostructured materials have attracted widespread attention due to the high strength.^1,2) However, the research on bulk nanostructured materials revealed that there seemed to be an inherent contradiction between their strength and ductility, that is strength-ductility trade-off.^3–5) Generally, the materials with grain sizes below 100 nm are considered to be nanocrystalline (NC) materials. And the materials with grain sizes between 100 nm and 1000 nm are defined as ultrafine-grained (UFG) materials.⁶⁾ Therefore, the above two materials are collectively referred to as NC/UFG materials because they all have high strength but low tensile ductility.

After decades of development, synthesizing methods of the NC/UFG materials can be classified into two major approaches -“bottom-up” and “top-down”.⁷⁾ Typical examples of the “bottom-up” methods include inert gas condensation,⁸⁾ electrodeposition,⁹⁾ ball-milling.¹⁰⁾ However, some “bottom-up” methods are complicated and the materials prepared by these methods contain high impurity and/or artificial porosity content. This may deteriorate the mechanical properties of the materials.^11,12)

In contrast, some “top-down” approaches, (e.g. severe plastic deformation (SPD)), are free of artificial porosity and contamination. Therefore, the mechanical properties of the materials prepared by SPD methods are improved compared with those of the materials prepared by “bottom-up” methods. The most commonly used SPD techniques are equal channel angular pressing (ECAP),¹³⁾ high-pressure torsion (HPT)¹⁴⁾ and so on. Although the strength of the NC/UFG materials prepared by SPD has been further enhanced, the ductility remains low.

As we known, material failures usually occur on the surface. Some previous studies have shown that the surface modification methods by surface nanocrystallization (SNC) may effectively improve the overall properties of materials.¹⁵⁾ Conventional nanostructured surface layers can be prepared by means of various coating and deposition techniques such as physical vapor deposition (PVD), chemical vapor deposition (CVD), sputtering, and electrodeposition. But the interface between the coated layer and the matrix may have a non-negligible negative effect on the performance of materials.¹⁵⁾

As a form of the heterogeneous materials,^16,17) the gradient structured (GS) materials have become a new research field in recent years because of their superior global properties compared with conventional homogeneous materials. The structures of GS materials change gradually and continuously from the surface to interior and thus there is no significant interface.^18,19) They exhibit high strength and good ductility,^20,21) high diffusion rate,^22–24) corrosion resistance,²⁵⁾ improved fatigue, friction and wear properties.^26–28)

The gradient change of grain sizes can effectively avoid the performance mutation caused by the sudden change of the grain sizes. Therefore, the structures with various grain sizes can be coordinated with each other and outstanding performance can be obtained.²⁹⁾ As shown in Fig. 1, the gradient nanostructures, whose top surface is nanoscale, are generally divided into the following categories: grain size gradient, twin thickness gradient, lamellar thickness gradient, and columnar size gradient.²⁹⁾

Fig. 1

The classification of gradient nanostructures with grain size gradient (a), twin thickness gradient (b), lamellar thickness gradient (c) and columnar size gradient (d).²⁹⁾

Typical processing techniques used to fabricate gradient nanostructured materials include surface mechanical attrition treatment (SMAT),¹⁵⁾ surface mechanical grinding treatment (SMGT),³⁰⁾ surface mechanical rolling treatment (SMRT),³¹⁾ friction sliding,³²⁾ high-energy milling,³³⁾ air blast shot peening³⁴⁾ and laser-induced projectile impact testing.^35,36)

K. Lu and J. Lu proposed the SMAT process in 1999 and reported the mechanisms of SNC for this process.¹⁵⁾ In the past 20 years, the SMAT has been widely applied in various kinds of materials including pure metals and alloys.^18,37–46) The mechanical properties of the material can be improved to some extent under appropriate SMAT process parameters.

In this overview, the SMAT mechanism, the microstructure characteristics of the GS layer induced by SMAT and the investigations on mechanical properties of different SMAT-ed materials would be presented and discussed in the following sections.

2. The Surface Mechanical Attrition Treatment (SMAT)

2.1 The device and principle of SMAT

The SMAT device (Fig. 2) consists mainly of a vibration generator and a treatment chamber placing the balls as well as holding the sample.⁴⁷⁾ The materials and sizes of the balls can be adjusted according to the type of the processed materials. High-strength steel is usually used to make the balls with the diameter of 1–10 mm.⁴⁰⁾

Fig. 2

Schematic illustration of the SMAT device (a) and the plastic deformation in the surface layer caused by the repeated impact of the flying balls on the sample surface in multiple directions (b).⁴⁷⁾

The plate sample is fixed to the upper wall of the treatment chamber and the balls are placed at the bottom. In order to prevent the surface layer of the treated sample from being oxidized or contaminated by impurities in the air, the treatment chamber is usually evacuated or filled with an inert gas (such as argon or nitrogen gas).⁴⁰⁾ The vibration generator drives the treatment chamber at a frequency of 50 to 20000 Hz and resonates the balls. And then the surface of the sample is continuously struck by the balls at high velocity from different angles, which promotes the refinement process of surface grains.⁴⁸⁾ During this process, different slip systems or twins are driven, and a large amount of dislocation movement or twinning behavior ultimately continuously divides the coarse grains in the top surface layer into nanoscale.^40,48)

The cryogenic temperature can prevent dynamic recovery and promote the generation of mechanical twinning, which causes the grains to be further refined. Therefore, the SPD processing at cryogenic temperatures is receiving more and more attention.⁴⁴⁾ Figure 3 shows the SMAT setup at cryogenic temperature, the details of which are discussed in Ref. 49). The effect of SMAT at cryogenic temperature on the mechanical properties of materials will be discussed in Section 4.

Fig. 3

Physical map (a) and schematic illustration (b) of the SMAT setup at cryogenic temperature.

The SMAT device mentioned in this overview mainly refers to those shown in Fig. 2 and Fig. 3. In addition, the high-energy milling (Fig. 4 and Fig. 5) is also a kind of SMAT.

Fig. 4

Assembled SPEX 8000 mixer/mill (a) and tungsten carbide vial set (b).³³⁾

Fig. 5

Improved vial set and sample.

2.2 SMAT process parameters

The ball velocity is an important SMAT parameter because it is proportional to the strain rate in the surface layer of the sample. Higher strain rates can facilitate more ideal nanostructures in the top surface layer of materials. The ball velocity is generally 1–20 m/s, which mainly affected by the vibration frequency, the height of treatment chamber and the size, density, number of balls.⁴⁸⁾ For example, the smaller mass and fewer balls result in the higher ball velocity in the vertical direction, due to the reduced collision among the balls in a limited space. It is reported that the high density of nano-twins was observed in the GS layer of the SMAT-ed 304 stainless steel by using 100 balls of 3 mm diameter, which leads to the higher tensile strength.⁵⁰⁾

On the other hand, the SMAT time can be selected according to the kinds of materials, the plate thickness, the ball velocity and the required thickness of the GS layer. Appropriate SMAT time is necessary (generally controlled within 60 min), however excessive treating time makes it difficult to further refine the grains due to balance between the dislocations generation and the dislocations annihilation. A recent study⁵¹⁾ points out that the short SMAT time (within 1 min) can increase the strength of the material with less sacrifice of ductility. This can be attributed to a unique bimodal structure consisting of the “hard phase” and the “soft matrix”. The “hard phase” with nano-twins increase the strength of the material and the coarse-grained “soft matrix” can prevent the propagation of the surface micro-cracks after necking.

In addition, A.M. Gatey et al.⁵²⁾ proposed an important term “100% peening time” (i.e. the time required for the indentation of the ball to cover 100% of the sample surface). This may provide a direction for the manufacture of GS materials that have high strength while maintaining superior ductility. Rajib Kalsar et al.⁴⁶⁾ reported a composite gradient structure consisting of twins, dislocation density and grain size gradients in twinning induced plasticity (TWIP) steel treated by SMAT. The results show that the yield strength of SMAT-ed TWIP steel is improved while having little ductility loss (i.e. uniform elongation is only reduced by 8%).

A non-negligible detail is that the thickness of GS layer in the sample is only about 50 µm. This may be a key reason for obtaining the superior combination of strength and ductility. Although the author did not provide SMAT time in the paper, from the perspective of “100% peening time”, we can speculate that this very thin GS layer thickness is likely to be produced near “100% peening time”. So theoretically, with other factors being the same, if we can find the “100% peening time” in other materials, the excellent combination of strength and ductility may be realized. For the determination of “100% peening time”, the way of numerical simulation combined with experiment can be considered.

3. Microstructure Characteristics of GS Layer under Different GS Materials and SMAT Processes

Firstly, we can define the region of GS layer and the volume fraction of GS layer here. As shown in Fig. 6, the GS layer consists of three parts: a nanostructured layer, an ultrafine-grained structure layer, and a deformed coarse-grained layer. The volume fraction of GS layer generally refers to the ratio of GS layer volume to total sample volume.

Fig. 6

Schematic illustration of microstructure characteristics and strain and strain rate distribution at different surface layer depths.⁴⁸⁾

Reference 48) shows that the gradual change of strain and strain rate from the top surface to the almost deformation-free matrix in the GS material results in a corresponding change of grain size gradient distribution from a few nanometers to several micrometers (Fig. 6).

As we all know that the observation of microstructure using transmission electron microscopy (TEM) is very local and strongly depends on the exact sampling position.⁵³⁾ Samih et al.⁵⁴⁾ determine the thickness of three regions in the GS layer by combining the distribution map of geometrically necessary dislocations (GNDs) densities with statistics. It is believed that the method based on average grain size statistics is useful for the accurate partition of various regions in the GS layer. However, electron backscatter diffraction (EBSD) technique is not sufficient to obtain the detailed features of the UFG region. Therefore, Gwénaëlle Proust et al.⁵⁵⁾ used the newly developed high-resolution transmission Kikuchi diffraction (TKD) technique to better characterize the UFG region. On the whole, the microstructure in the GS layer can be better characterized by the combined use of the above three methods.

Generally, the main plastic deformation mechanisms for metallic materials include dislocation slip and twinning. Different types of materials have different grain refinement mechanisms during SMAT, which mainly depends on the crystal structure types and stacking fault energy (SFE).⁴⁸⁾

The related studies have shown that the grain refinement processes of materials with high SFEs are mainly dominated by dislocation movement. However, for the metals with low SFEs, the plastic deformation mechanism may change from dislocation slip to mechanical twins.⁵⁶⁾ On the other hand, the effects of different SMAT parameters and SMAT at cryogenic temperature on the microstructures of the GS layer are also two interesting aspects. For some materials, the phase transition process is even included.

In the following sections, as the description of the microstructures and grain refinement mechanisms has been discussed in detail in Ref. 57), two classical grain refinement mechanisms will be briefly introduced. In addition, the effects of SMAT at cryogenic temperature and different SMAT parameters on the microstructures of GS layer will be discussed.

3.1 Dislocations-induced grain refinement mechanism

Figure 7 shows a schematic of SMAT-ed pure iron with a high SFE during the grain refinement. The following steps are involved in the grain refinement process:⁴⁰⁾ (i) the development of dense dislocation walls (DDWs) and dislocation tangles (DTs); (ii) the transition from DDWs and DTs to sub-boundaries with small misorientation; (iii) the evolution of the sub-boundaries to highly misoriented grain boundaries.

Fig. 7

Schematic illustration of grain refinement induced by plastic deformation in SMAT-ed pure iron: (i) development of dense dislocation walls (DDWs) and dislocation tangles (DTs) in the original grains and the further refined cells; (ii) transformation of DDWs and DTs into sub-boundaries with small misorientations; (iii) evolution of sub-boundaries to highly misoriented grain boundaries; and (iv) formation of nanostructures in the top surface layer.⁵⁷⁾

3.2 Twins-induced grain refinement mechanism

The grain refinement mechanism in the GS layer of the SMAT-ed AISI 304 stainless steel with a low SFE is different from that in the GS layer of the materials with high SFEs (such as pure Fe⁴⁰⁾ and Al-alloy⁴¹⁾). The grain refinement process can be listed as follows:⁵⁸⁾ (i) the forming of planar dislocation arrays and twins, (ii) the intersections of multidirectional twins, and (iii) the formation of randomly oriented nanocrystallites.

3.3 Effect of SMAT process parameters on microstructures of AISI 304L steel

The previous research⁵⁸⁾ has shown that deformation-induced martensite (α′) is formed near the surface of the treated sample and the high-density shear bands (mechanical twins) are observed in this region. Recently, A.M. Gatey et al.⁵²⁾ studied the effects of different SMAT process parameters on the microstructures of AISI 304L steel. The results show that most of the α′ is formed inside the grains at the intersection of the shear bands and along the shear bands, some α′ is also present near the grain boundaries. With the SMAT time and the number of balls increase, the proportion of α′ increases.

The effect of the ball diameter on the volume fraction of α′ can be divided into two cases⁵²⁾ in Fig. 8. When the percentage of the ball is only 25%, the strain rate of the ball with 3 mm diameter will be higher than that of the ball with 8 mm diameter, and the adiabatic heating associated with high strain rate deformation inhibits the formation of α′. In the case of a large percentage of the ball (50% and 75%), because the number of balls with 3 mm diameter is more than that of balls with 8 mm diameter and the strain rate is lower, the former corresponds to a higher volume fraction of α′.

Fig. 8

The volume percent of α′-martensite phase for various percentages of balls of 3 and 8 mm diameter respectively; a constant 15 mm gap is used.⁵²⁾

3.4 Effect of SMAT at cryogenic temperature on microstructures at different surface depths in GS 304L steel

It is reported that the SMAT-ed 304L steel undergoes a γ → ε → α′ transition at room temperature, where ε martensite is generally stress-assisted and forms in the early stages of deformation.⁵⁹⁾ The strain-induced γ → α′ transition seems to occur more easily when γ-twins are preferentially formed.⁶⁰⁾ A recent paper investigated the characteristics of martensite (ε, α′) content at different surface depths in the 304L steel treated by SMAT at cryogenic temperature.⁴³⁾

The results show that although the stress-assisted ε martensite transformation occurred at 150–450 µm below the surface treated by SMAT at cryogenic temperature, it is not found at any depth below the surface after SMAT treatment at room temperature. The difficulty of deformation induced transformation can be offset by promoting the formation of stress-assisted ε martensite at cryogenic temperature. Other than that, the maximum amount of α′ martensite always appears on the subsurface, regardless of the treatment at room temperature or cryogenic temperature. In summary, the gradient structure created by SMAT in 304L steel superimposes two gradients: grain size gradient and phase distribution gradient.⁴³⁾

4. Effects of SMAT on Mechanical Properties of Steels, Cu and Cu Alloys

4.1 Superior strength and good ductility

NS materials are known for having high strength, but the ductility is typically poor due to the absence of work hardening in the deformation process.⁶¹⁾ In recent years, GS materials have gained widespread attention for their superior combination of strength and ductility. In general, the improvement of ductility in GS materials is mainly related to the grain growth, back stress hardening, and volume fraction of GS layer.

For the gradient nanograined Cu prepared by surface mechanical grinding treatment (SMGT), the high ductility of the sample was attributed to grain growth caused by the low structural stability of the nanograined Cu.²¹⁾ The investigation on SMAT-ed IF steel showed that the high back stress in the GS layer not only contributes to the synergetic strengthening, while the work hardening caused by high back stress is beneficial for enhancement of ductility.⁶²⁾ Our previous work on SMAT-ed pure copper indicated that there exists an optimum range of volume fractions of GS layer which can produce the highest extra work hardening and extra strengthening.⁶³⁾ In addition, the simulation results (Fig. 9) of Ref. 64) show that uniform elongation increases and yield/ultimate strength decrease slightly as the volume fraction of GS layer decreases. This will help to optimize the strength and ductility of GS materials by controlling the volume fraction of GS layer. On the other hand, the relation between the volume fraction of GS layer and the grain size of the topmost surface has been established, which also can provide a guide for the manufacture of GS materials with an excellent balance of strength and ductility.⁶⁵⁾

Fig. 9

The variation of uniform elongation (a) and yield/ultimate strength (b) with volume fraction of GS layer.⁶⁴⁾

The following sections will mainly focus on the mechanisms of strengthening and improving ductility in the GS materials prepared by SMAT.

4.2 The SMAT treatment at liquid nitrogen temperature

Generally, the SMAT process is carried out at room temperature. However, the strong impact on the sample surface during SMAT can lead to an increase in the temperature of the sample (especially for a long time at room temperature). When the temperature is high enough to cause dynamic recovery during SMAT, the balance between the dislocations generation and the dislocations annihilation can be achieved.

K.A. Darling et al.⁶⁶⁾ first performed SMAT treatment on copper at cryogenic temperature (here SMAT refers to cryogenic milling in the SPEX 8000M mill) and demonstrated its advantage in microstructure refinement compared with SMAT process at room temperature. Another research showed that the SMAT process at liquid nitrogen temperature can effectively suppress the occurring of dynamic recovery and significantly increase the strength of pure Cu without seriously sacrificing its ductility.⁴⁹⁾ This is because the SMAT at cryogenic temperature could increase twin and dislocation density. The twin can provide strengthening and enough work hardening capacity by emitting and blocking the dislocations.⁶⁷⁾ In addition, a thicker GS layer with finer grains is obtained under this process compared with the process at room temperature.⁴⁹⁾ It is noted that although SMAT at cryogenic temperature usually leads to surface grain refinement, e.g. pure Cu,⁴⁹⁾ iron,⁴⁵⁾ 304L,^43,44) 310S,⁴⁴⁾ the thickness of GS layer does not always increase. For example, Ref. 44) shows that the GS layer thickness in the 304L steel treated by SMAT at cryogenic temperature increases compared with that at room temperature, while the trend of 310S steel is just the opposite (Fig. 10).

Fig. 10

Scanning electron microscope (SEM) images of the 304L (a, b) and the 310S (c, d) steels treated at room (a, c) and cryogenic (b, d) temperatures.⁴⁴⁾

It is reported that the yield strength of 310S steel increased from 380 MPa at room temperature to 700 MPa at 77 K. In contrast, for 304L steel, it varies from 220 MPa at room temperature to 380 MPa at 77 K.⁶⁸⁾ Therefore, the high strength and hardness of 310S steel at cryogenic temperature hinder plastic deformation during the cryogenic SMAT process, which results in the formation of a thin GS layer.⁴⁴⁾

The influence of strain rate $\dot{\varepsilon }$ and deformation temperature T on mechanical properties can be expressed by a parameter Z:⁶⁹⁾   

\begin{equation} \mathrm{Z} = \dot{\varepsilon}\exp\left(\frac{Q}{RT}\right) \end{equation} (1)

where Q is the activation energy for diffusion. The Z value (Zener–Holloman parameter) will increase by increasing $\dot{\varepsilon }$ and/or decreasing T. Similar to reducing SFE, this will limit dislocation movement and increase the critical shear stress of the slip. As a result, the dislocation slip is suppressed and the twinning becomes the main deformation mechanism.⁶⁹⁾ Hence, the strength and ductility of the GS materials are expected to improve by increasing the strain rate. A previous study on Cu/Ni multilayered composites prepared by electrodeposition indicates that the strength and ductility of the sample increased simultaneously with the increase of strain rate.⁷⁰⁾ This could be further explored in other GS materials.

4.3 The extra work hardening in the GS materials

Wu et al. reported a superior combination of considerable uniform elongation and the higher yield strength in SMAT-ed IF steel.¹⁸⁾ The investigation reveals a unique extra work hardening in the GS materials, which is critical for improving the ductility of the materials.^18,71) The tensile deformation process of the treated IF steel involves the following three stages.^18,72)

In stage I (Fig. 11(a)), the GS sample only undergoes elastic deformation due to the very small applied strain.

Fig. 11

Schematic of the stress state change during uniaxial tension in the GS sample. (a) Stage I; (b) Stage II; (c) Stage III. The red and yellow arrows represent lateral tensile and compressive stress, respectively.⁷²⁾

In stage II, as the applied strain increases to a certain extent, the coarse-grained center layer begins to plastically deform (Fig. 11(b)), at which time the entire sample is at the elastoplastic deformation stage. The plastically deformed central layer will shrink more in the lateral direction than the elastically deformed GS layers due to different Poisson’s ratios. This indicates that the GS materials have an inherent mechanical incompatibility (unlike the homogeneous materials) during deformation. This causes the uniaxial stress state to transform into the biaxial stress state. In stage III (Fig. 11(c)), the whole sample is in the plastic deformation state. The GS layer of the sample first shrinks rapidly in the lateral direction due to the earlier occurrence of necking instability. However, the mechanical incompatibility occurs again because the surface necking is constrained by the central layer. The largest strain gradient occurs near the interface between them, and this maximum strain gradient will promote the accumulation of GNDs and the development of back stress to produce extra work hardening. In addition, the ductility of GS sample is improved due to the sharp increase in the work hardening rate θ. The phenomenon is caused by the wake high density of dislocations left by the migration process of two interfaces from the surface to the central layer. The investigation also suggests that there may be an optimal GS layer thickness which produces the most significant θ upturn and the most extra work hardening.¹⁸⁾

4.4 The influence of the volume fraction of GS layer and degree of gradient on the mechanical properties of metallic materials

A previous study compared the strength and ductility of pure copper processed by a variety of processes and SMAT (See Fig. 12). The results show that an excellent combination of strength-ductility occurs on pure copper treated by SMAT at cryogenic temperature when the volume fraction of GS layer is between 0.08–0.1.⁶³⁾

Fig. 12

Superior mechanical property. Strength and ductility in the SMAT-ed pure copper samples compared with the pure copper samples processed by a variety of processes.⁶³⁾

Interestingly, 0.1 is also the critical value of θ upturn, which suggests this unique behavior may relate to the volume fraction of GS layer. In addition, it was reported that the GS Cu–30 wt.% Zn processed by multiple-pass friction stir processing (FSP) and rotationally accelerated shot peening (RASP) has a higher volume fraction of GS layer than the one processed by RASP alone. The former exhibits higher synergistic strengthening because the thicker GS layer provides sufficient space for the accumulation of GNDs. This results in the higher tensile strength increment relative to the coarse-grained samples.^31,73)

On the other hand, the degree of gradient has also received attention recently. It is worth noting that the degree of gradient and the volume fraction of GS layer are two different concepts. The former is discussed only in the GS layer range, while the latter refers to the ratio of GS layer volume to total GS sample volume. Cheng et al.⁷⁴⁾ achieved simultaneous improvement of strength and work hardening by increasing the structural gradient parameter s (i.e. the increment in hardness per unit thickness along the gradient direction) of the gradient nanotwinned pure Cu prepared by direct-current electro-deposition method. This can be attributed to the unique morphology of GNDs, which are uniformly distributed inside the grains in the form of concentrated dislocation bundles. The latest research shows that we can achieve optimal mechanical properties by rationally controlling the degree of gradient. The issue of better control of the gradient degree remains to be investigated in future. It should be noted that in this study, all discussions of the influence of gradient degree on mechanical properties are based on constant sample thickness and volume fraction of GS layer.

4.5 The mechanical properties of the Cu alloys with different SFEs treated by SMAT

As mentioned above, the liquid nitrogen temperature not only inhibits dynamic recovery, but also contributes to the generation of twins. On the other hand, the full dislocations are difficult to cross-slip and climb in the metals with low SFEs, which leads to an increase in work hardening.⁷⁵⁾ Therefore, both liquid nitrogen temperature and low SFE may improve mechanical properties. The following sections mainly focus on the Cu alloys with different SFEs treated by SMAT at liquid nitrogen temperature.

A previous investigation showed that the yield strength of SMAT-ed Cu–Zn alloys is improved by decreasing SFE. However, optimal ductility occurs in the alloy with the optimal SFE.⁷⁶⁾ A similar trend was found in another study on SMAT-ed Cu–Al alloys. The SMAT-ed 5 min Cu–4.5 wt.%Al sample with medium SFE shows an optimized combination of strength and ductility.⁷⁷⁾ Taking the Cu–Al alloys for instance, the possible reasons behind this phenomenon will be discussed as follows. Figure 13 shows the reduced tendency of θ of SMAT-ed Cu–2.2 wt.%Al sample is significantly faster than that of the Cu–4.5 wt.%Al sample with the increase of true strain, so the ductility of the Cu–4.5 wt.%Al sample is higher.

Fig. 13

The work hardening rate (θ = dσ/dε) vs. true strain (ε_t) curves (the insert shows corresponding true stress-strain curves).⁷⁷⁾

When the true strain exceeds 22%, the θ of SMAT-ed 5 min Cu–4.5 wt.%Al sample is higher than that of the Cu–6.9 wt.%Al sample. The high work hardening ability maintained at the relatively large plastic strain may be responsible for the high ductility of the Cu–4.5 wt.%Al samples.⁷⁷⁾ Both the high ρ_m/ρ_m0 measured by the stress relaxation experiment and the small K2 value obtained by the Kocks-Mecking model in SMAT-ed 5 min Cu–4.5 wt.%Al sample show that the ductility of the Cu–4.5 wt.%Al sample with optimal SFE is superior.

In addition, the same phenomenon was observed in UFG Cu–Zn alloys.⁷⁸⁾ The lower ductility in the samples with low SFE is caused by extremely small grains and saturated stacking faults, which are difficult to accumulate dislocations and form twins during the deformation process. The underlying mechanism corresponding to the special phenomenon in the SMAT-ed Cu–Al alloys requires further exploration.

According to the discussion mentioned above, the change in yield strength of Cu–Al alloys with different SFEs may be related to the GS layer thickness and back stress. As the SFE decreases, the hardness of the Cu–Al alloys increases. Therefore, the degree of deformation of the low-SFE alloys is less than that of the high-SFE alloys at the same kinetic energy. As a result, the thickness of the GS layer decreases with the SFE decreases. SFE further affects the grain refinement in the GS layer induced by SMAT, i.e. the higher SFE, the smaller grain size in the GS layer is formed due to the presence of more low-angle grain boundaries used for subdividing grains.⁷⁹⁾ The thicker GS layer, the smaller grain size in the GS layer, and a higher tensile strength increment of the SMAT-ed Cu–Al alloy are reached due to fine-grain strengthening.

Finally, the Cu–2.2 wt.%Al sample exhibited the highest tensile strength increment, but the ductility was significantly degraded. While the increase in the strength of the Cu–6.9 wt.%Al sample after SMAT is particularly limited due to the thinner GS layer. It is noted that the SMAT-ed 5 min Cu–6.9 wt.%Al sample has the highest tensile strength based on the high strength of its corresponding annealed sample, although the strength increment is limited compared with the annealed sample.

It is reported that the fine-grain strengthening is not sufficient to explain the increase in tensile strength of GS Cu–Al alloy, and the back stress is also considered to play an important role in increasing the tensile strength of the material.^18,73) The back stress is the long-range internal stress caused by the accumulation of GNDs.^80,81) The GNDs represent extra storage of dislocations required to maintain lattice continuity during the non-uniform plastic deformation.⁸²⁾

Recent research investigated the relationship between GNDs and back stress strengthening in SMAT-ed Cu–Al alloys with different SFEs.⁷⁹⁾ As shown in Fig. 14, the GNDs density increases with the decreasing SFE in annealed Cu–Al alloy samples. The average grain size of the annealed Cu–6.9 wt.% Al alloy (low SFE) is smaller, which makes the dislocations more easily absorbed by the grain boundaries, but the further accumulation is difficult for them. As a result, the GNDs density increment of the SMAT-ed 5 min Cu–Al alloy sample decreases with the decrease of SFE. Therefore, this indicates that the back stress strengthening (also for tensile strength increment) of the Cu–Al alloy decreases with the decreasing SFE.

Fig. 14

Cross-sectional EBSD maps of the GNDs density near the surfaces of the SMAT-5 min and annealed ((a) and (b)) Cu–2.2 wt.% Al, ((c) and (d)) Cu–4.5 wt.% Al, and ((e) and (f)) Cu–6.9 wt.% Al alloys.⁷⁹⁾

SMAT time is also an important parameter affecting the mechanical properties of materials. Within a certain time, as the SMAT time increases, the thickness of the GS layer increases to some extent. The GNDs are mainly distributed in the GS layer, and the density of GNDs increases with the SMAT time. This results in a significant increase in back stress, which can improve the tensile strength.⁸³⁾

On the other hand, the back stress strengthening ability gradually decreases with the increasing SMAT time. This may be due to the high dislocation density in the GS samples under long processing time. The high dislocation density makes it difficult to accumulate more GNDs during the next deformation.⁸³⁾ This is similar to the reason for the reduced GNDs density increment in SMAT-ed Cu–Al alloys with the low SFE.

4.6 Effect of residual stress on mechanical properties of GS materials

The effects of most surface treatment processes on mechanical properties are related to the microstructures. However, the effect of residual stress on mechanical properties is rarely reported. Recently, Ji Hyun Moon et al.⁸⁴⁾ studied the effects of the residual stress and the microstructure on mechanical properties of GS materials processed by UNSM (ultrasonic-nanocrystalline surface modification), respectively. The results in Fig. 15 show that the surface grain refinement and residual stress of the UNSM-ed sample contribute to 80% (38 MPa) and 20% (9 MPa) increase in yield strength, respectively, compared with the initial sample.

Fig. 15

Tensile engineering stress-strain curves of the initial, UNSM, and UNSM + HT (stress released) specimens.⁸⁴⁾

Similarly, the contribution of residual stress to mechanical properties is not negligible for other surface treatment processes (e.g. SMAT) used to fabricate GS materials.⁸⁴⁾ Therefore, residual stress also provides a direction for improving the mechanical properties of GS materials.

5. Summary

This overview mainly focuses on the microstructure in the gradient structured (GS) layer generated by surface mechanical attrition treatment (SMAT) and the related studies on strength and ductility in different GS materials. We can draw the following conclusions from the above statements:

1. (1)    The SMAT is an effective method for manufacturing the gradient nanostructure. In the GS materials fabricated by SMAT, the structures change gradually and continuously from the top surface to interior and do not form significant interface.
2. (2)    With the SMAT time and the number of balls increase, the proportion of α′ in SMAT-ed AISI 304L steel increases. The SMAT at cryogenic temperature could increase twin and dislocation density, which helps to increase the strength and maintain good ductility. The gradient structure in the 304L steel treated by SMAT at cryogenic temperature is a superposition of grain size gradient and phase distribution gradient.
3. (3)    The GS materials fabricated by SMAT not only have high strength, but also effectively improve ductility. The fine-grain strengthening and the back stress strengthening both contribute to the increment of tensile strength of SMAT-ed materials. Moreover, the extra work hardening caused by back stress in the gradient structure is critical for improving the ductility of materials.
4. (4)    A superior combination of strength and ductility can be obtained by reasonably controlling the degree of gradient and the volume fraction of the GS layer.
5. (5)    The investigation on SMAT-ed Cu–Zn alloys showed that the yield strength increases with the SFEs decrease. On the other hand, superior ductility occurs in the alloy with optimal SFE. A similar phenomenon was also observed in the SMAT-ed Cu–Al alloys and the UFG Cu–Zn alloys.
6. (6)    The thickness of GS layer and the density of geometrically necessary dislocations (GNDs) increase with the SMAT time within a certain time range. Thick GS layer provides more effective fine grain strengthening while high-density GNDs results in a significant increase in back stress, both of which help to increase the tensile strength of the sample. On the other hand, the short SMAT time can increase the strength of the material with less sacrifice of ductility. This can be related to a unique bimodal structure consisting of the “hard phase” and the “soft matrix” or “100% peening time”.
7. (7)    The influence of strain rate, volume fraction of GS layer, degree of gradient, SFEs and residual stress on mechanical properties of the GS materials require further investigation.

Acknowledgments

The authors would like to acknowledge financial support by the National Natural Science Foundation of China (NSFC) under Grants No. 51664033, No. 51561015 and No. 51501078.

REFERENCES

• 1)   H.  Gleiter: Acta Mater. 48 (2000) 1–29.
• 2)   C.  Suryanarayana: Bull. Mater. Sci. 17 (1994) 307–346.
• 3)   Y.T.  Zhu and  X.  Liao: Nat. Mater. 3 (2004) 351.
• 4)   M.A.  Meyers,  A.  Mishra and  D.J.  Benson: Prog. Mater. Sci. 51 (2006) 427–556.
• 5)   K.  Lu: Science 328 (2010) 319–320.
• 6)   R.Z.  Valiev,  Y.  Estrin,  Z.  Horita,  T.G.  Langdon,  M.J.  Zechetbauer and  Y.T.  Zhu: JOM 58 (2006) 33–39.
• 7)   Y.T.  Zhu,  T.C.  Lowe and  T.G.  Langdon: Scr. Mater. 51 (2004) 825–830.
• 8)   G.W.  Nieman,  J.R.  Weertman and  R.W.  Siegel: Scr. Metall. 23 (1989) 2013–2018.
• 9)   U.  Erb,  A.M.  El-Sherik,  G.  Palumbo and  K.T.  Aust: Nanostruct. Mater. 2 (1993) 383–390.
• 10)   C.C.  Koch and  Y.S.  Cho: Nanostruct. Mater. 1 (1992) 207–212.
• 11)   S.V.  Petegem and  F.D.  Torre: Scr. Mater. 48 (2003) 17–22.
• 12)   P.G.  Sanders,  J.A.  Eastman and  J.R.  Weertman: Acta Mater. 45 (1997) 4019–4025.
• 13)   R.Z.  Valiev and  T.G.  Langdon: Prog. Mater. Sci. 51 (2006) 881–981.
• 14)   A.P.  Zhilyaev and  T.G.  Langdon: Prog. Mater. Sci. 53 (2008) 893–979.
• 15)   K.  Lu and  J.  Lu: J. Mater. Sci. Technol. 15 (1999) 193–197.
• 16)   E.  Ma and  T.  Zhu: Mater. Today 20 (2017) 323–331.
• 17)   X.  Wu and  Y.  Zhu: Mater. Res. Lett. 5 (2017) 527–532.
• 18)   X.  Wu,  P.  Jiang,  L.  Chen,  F.  Yuan and  Y.T.  Zhu: Proc. Natl. Acad. Sci. USA 111 (2014) 7197.
• 19)   S.  Suresh: Science 292 (2001) 2447–2451.
• 20)   X.H.  Chen,  J.  Lu,  L.  Lu and  K.  Lu: Scr. Mater. 52 (2005) 1039–1044.
• 21)   T.H.  Fang,  W.L.  Li,  N.R.  Tao and  K.  Lu: Science 331 (2011) 1587–1590.
• 22)   Z.B.  Wang,  K.  Lu,  G.  Wilde and  S.V.  Divinski: Acta Mater. 58 (2010) 2376–2386.
• 23)   Z.B.  Wang,  N.R.  Tao,  W.P.  Tong,  J.  Lu and  K.  Lu: Acta Mater. 51 (2003) 4319–4329.
• 24)   Y.  Lin,  J.  Lu,  L.  Wang,  T.  Xu and  Q.  Xue: Acta Mater. 54 (2006) 5599–5605.
• 25)   T.  Balusamy,  S.  Kumar and  T.S.N.S.  Narayanan: Corros. Sci. 52 (2010) 3826–3834.
• 26)   Z.B.  Wang,  N.R.  Tao,  S.  Li,  W.  Wang,  G.  Liu,  J.  Lu and  K.  Lu: Mater. Sci. Eng. A 352 (2003) 144–149.
• 27)   Y.S.  Zhang,  Z.  Han,  K.  Wang and  K.  Lu: Wear 260 (2006) 942–948.
• 28)   D.  Li,  H.N.  Chen and  H.  Xu: Appl. Surf. Sci. 255 (2009) 3811–3816.
• 29)   K.  Lu: Acta Metall. Sinica 51 (2015) 1–10.
• 30)   W.L.  Li,  N.R.  Tao and  K.  Lu: Scr. Mater. 59 (2008) 546–549.
• 31)   Y.F.  Wang,  C.X.  Huang,  M.S.  Wang,  Y.S.  Li and  Y.T.  Zhu: Scr. Mater. 150 (2018) 22–25.
• 32)   D.A.  Hughes and  N.  Hansen: Phys. Rev. Lett. 87 (2001) 135503.
• 33)   C.  Suryanarayana: Prog. Mater. Sci. 46 (2001) 1–184.
• 34)   J.L.  Liu,  M.  Umemoto,  Y.  Todaka and  K.  Tsuchiya: J. Mater. Sci. 42 (2007) 7716–7720.
• 35)   R.  Thevamaran,  O.  Lawal,  S.  Yazdi,  S.J.  Jeon,  J.H.  Lee and  E.L.  Thomas: Science 354 (2016) 312–316.
• 36)   C.  Cui,  X.  Cui,  X.  Li,  K.  Luo,  J.  Lu,  X.  Ren,  J.  Zhou,  C.  Fang,  R.  Farkouh and  Y.  Lu: Int. J. Plastic. 102 (2018) 83–100.
• 37)   G.  Liu,  J.  Lu and  K.  Lu: Mater. Sci. Eng. A 286 (2000) 91–95.
• 38)   N.R.  Tao,  M.L.  Sui,  J.  Lu and  K.  Lua: Nanostruct. Mater. 11 (1999) 433–440.
• 39)   H.Q.  Sun,  Y.N.  Shi,  M.X.  Zhang and  K.  Lu: Surf. Coat. Technol. 202 (2008) 3947–3953.
• 40)   N.R.  Tao,  Z.B.  Wang,  W.P.  Tong,  M.L.  Sui,  J.  Lu and  K.  Lu: Acta Mater. 50 (2002) 4603–4616.
• 41)   X.  Wu,  N.  Tao,  Y.  Hong,  B.  Xu,  J.  Lu and  K.  Lu: Acta Mater. 50 (2002) 2075–2084.
• 42)   K.Y.  Zhu,  A.  Vassel,  F.  Brisset,  K.  Lu and  J.  Lu: Acta Mater. 52 (2004) 4101–4110.
• 43)   M.  Novelli,  P.  Bocher and  T.  Grosdidier: Mater. Charact. 139 (2018) 197–207.
• 44)   M.  Novelli,  J.J.  Fundenberger,  P.  Bocher and  T.  Grosdidier: Appl. Surf. Sci. 389 (2016) 1169–1174.
• 45)   H.A.  Murdoch,  K.A.  Darling,  A.J.  Roberts and  L.  Kecskes: Metals 5 (2015) 976–985.
• 46)   R.  Kalsar and  S.  Suwas: Scr. Mater. 154 (2018) 207–211.
• 47)   W.P.  Tong,  N.R.  Tao,  Z.B.  Wang,  J.  Lu and  K.  Lu: Science 299 (2003) 686–688.
• 48)   K.  Lu and  J.  Lu: Mater. Sci. Eng. A 375–377 (2004) 38–45.
• 49)   Y.  Shen,  C.  Wen,  X.  Yang,  Y.  Pang,  L.  Sun,  J.  Tao,  Y.  Gong and  X.  Zhu: J. Mater. Eng. Perform. 24 (2015) 5058–5064.
• 50)   H.L.  Chan,  H.H.  Ruan,  A.Y.  Chen and  J.  Lu: Acta Mater. 58 (2010) 5086–5096.
• 51)   X.  Liu,  K.  Wu,  G.  Wu,  Y.  Gao,  L.  Zhu,  Y.  Lu and  J.  Lu: Scr. Mater. 124 (2016) 103–107.
• 52)   A.M.  Gatey,  S.S.  Hosmani and  R.P.  Singh: Surf. Eng. 32 (2016) 69–78.
• 53)   W.B.  Liu,  X.  Jin,  B.  Zhang,  D.  Yun and  P.H.  Chen: Materials 12 (2019).
• 54)   Y.  Samih,  B.  Beausir,  B.  Bolle and  T.  Grosdidier: Mater. Charact. 83 (2013) 129–138.
• 55)   G.  Proust,  D.  Retraint,  M.  Chemkhi,  A.  Roos and  C.  Demangel: Microsc. Microanal. 21 (2015) 919–926.
• 56)   B.  Bay,  N.  Hansen,  D.A.  Hughes and  D.  Kuhlmann-Wilsdorf: Acta Mater. 40 (1992) 205.
• 57)   J.  Azadmanjiri,  C.C.  Berndt,  A.  Kapoor and  C.  Wen: CRC Crit. Rev. Solid State Sci. 40 (2015) 164–181.
• 58)   H.W.  Zhang,  Z.K.  Hei,  G.  Liu,  J.  Lu and  K.  Lu: Acta Mater. 51 (2003) 1871–1881.
• 59)   Y.F.  Shen,  X.X.  Li,  X.  Sun,  Y.D.  Wang and  L.  Zuo: Mater. Sci. Eng. A 552 (2012) 514–522.
• 60)   H.  Fujita and  T.  Katayama: Mater. Trans. JIM 33 (1992) 243–252.
• 61)   Y.M.  Wang and  E.  Ma: Acta Mater. 52 (2004) 1699–1709.
• 62)   M.  Yang,  Y.  Pan,  F.  Yuan,  Y.  Zhu and  X.  Wu: Mater. Res. Lett. 4 (2016) 145–151.
• 63)   X.  Yang,  X.  Ma,  J.  Moering,  Z.  Hao,  W.  Wei,  Y.  Gong,  J.  Tao,  Y.  Zhu and  X.  Zhu: Mater. Sci. Eng. A 645 (2015) 280–285.
• 64)   J.  Li and  A.K.  Soh: Int. J. Plast. 39 (2012) 88.
• 65)   J.  Li and  A.K.  Soh: Model. Simul. Mater. Sci. Eng. 20 (2012) 085002.
• 66)   K.A.  Darling,  M.A.  Tschopp,  A.J.  Roberts,  J.P.  Ligda and  L.J.  Kecskes: Scr. Mater. 69 (2013) 461–464.
• 67)   C.  Ye,  A.  Telang,  A.S.  Gill,  S.  Suslov,  Y.  Idell,  Z.  Kai,  J.M.K.  Wiezorek,  Z.  Zhou,  D.  Qian and  S.R.  Mannava: Mater. Sci. Eng. A 613 (2014) 274–288.
• 68)  T. Horiuchi, R. Ogawa, M. Shimada, S. Tone, M. Yamaga and Y. Kasamatsu: Adv. Cryog. Eng. Mater., vol. 28 (Plenum Press, New York, 1982) pp. 93–103.
• 69)   S.  Qu,  X.H.  An,  H.J.  Yang,  C.X.  Huang,  G.  Yang,  Q.S.  Zang,  Z.G.  Wang,  S.D.  Wu and  Z.F.  Zhang: Acta Mater. 57 (2009) 1586–1601.
• 70)   Z.  Fu,  Z.  Zhang,  L.  Meng,  B.  Shu,  Y.  Zhu and  X.  Zhu: Mater. Sci. Eng. A 726 (2018) 154–159.
• 71)   Z.  Zhi,  X.  Li,  D.  Xu,  L.  Lei,  H.  Gao and  T.  Zhu: Extreme Mech. Lett. 8 (2016) 213–219.
• 72)   I.A.  Ovid’Ko,  R.Z.  Valiev and  Y.T.  Zhu: Prog. Mater. Sci. 94 (2018) 462–540.
• 73)   X.L.  Wu,  P.  Jiang,  L.  Chen,  J.F.  Zhang,  F.P.  Yuan and  Y.T.  Zhu: Mater. Res. Lett. 2 (2014) 185–191.
• 74)   Z.  Cheng,  H.  Zhou,  Q.  Lu,  H.  Gao and  L.  Lu: Science 362 (2018) 559.
• 75)   J.  Gubicza,  N.Q.  Chinh,  J.L.  Lábár,  Z.  Hegedűs and  T.G.  Langdon: Mater. Sci. Eng. A 527 (2010) 752–760.
• 76)   B.  Cai,  X.  Ma,  J.  Moering,  H.  Zhou,  X.  Yang and  X.  Zhu: Mater. Sci. Eng. A 626 (2015) 144–149.
• 77)   Z.  Yin,  L.  Sun,  J.  Yang,  Y.  Gong and  X.  Zhu: J. Alloys Compd. 687 (2016) 152–160.
• 78)   Y.H.  Zhao,  X.Z.  Liao,  Z.  Horita,  T.G.  Langdon and  Y.T.  Zhu: Mater. Sci. Eng. A 493 (2008) 123–129.
• 79)   Y.  Zhang,  C.  Yang,  D.  Zhou,  Y.  Zhe,  L.  Meng,  X.  Zhu and  D.  Zhang: Mater. Sci. Eng. A 740–741 (2019) 235–242.
• 80)   D.  Kuhlmann-Wilsdorf and  C.  Laird: Mater. Sci. Eng. 37 (1979) 111–120.
• 81)   C.J.  Bayley,  W.A.M.  Brekelmans and  M.G.D.  Geers: Int. J. Solids Struct. 43 (2006) 7268–7286.
• 82)   H.  Gao and  Y.  Huang: Scr. Mater. 48 (2003) 113–118.
• 83)   X.  Hu,  S.  Jin,  Z.  Hao,  Y.  Zhe,  Y.  Jian,  Y.  Gong,  Y.  Zhu,  S.  Gang and  X.  Zhu: Metall. Mater. Trans. A 48 (2017) 1–8.
• 84)  J.H. Moon, S.M. Baek, S.G. Lee, Y. Seong and A. Amanov: Mater. Res. Lett. 7 (2019) 97–102.