Improving strength-ductility synergy in 301 stainless steel by combining gradient structure and TRIP effect

Journal Pre-proof Improving strength-ductility synergy in 301 stainless steel by combining gradient structure and TRIP effect Q. He, Y.F. Wang, M.S. Wang, F.J. Guo, Y. Wen, C.X. Huang PII:

S0921-5093(20)30233-1

DOI:

https://doi.org/10.1016/j.msea.2020.139146

Reference:

MSA 139146

To appear in:

Materials Science & Engineering A

Received Date: 24 January 2020 Revised Date:

20 February 2020

Accepted Date: 22 February 2020

Please cite this article as: Q. He, Y.F. Wang, M.S. Wang, F.J. Guo, Y. Wen, C.X. Huang, Improving strength-ductility synergy in 301 stainless steel by combining gradient structure and TRIP effect, Materials Science & Engineering A (2020), doi: https://doi.org/10.1016/j.msea.2020.139146. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2020 Published by Elsevier B.V.

Credit Author Statement Q. He: Conceptualization, Methodology, Data curation, Writing-Original draft preparation. Y.F. Wang: Conceptualization, Writing-Review & Editing. M.S. Wang: Data curation, Supervision. F.J. Guo: Data curation, Validation. Y. Wen: Validation. C.X. Huang: Conceptualization, Methodology, Writing-Review & Editing.

Improving strength-ductility synergy in 301 stainless steel by combining gradient structure and TRIP effect Q. He1, Y.F. Wang1,*, M.S. Wang1, F.J. Guo1, Y. Wen1, C.X. Huang1,* 1

School of Aeronautics and Astronautics, Sichuan University, Chengdu 610065, China * Corresponding author: [email protected] (C.X. Huang) [email protected] (Y.F. Wang) Abstract: Here we investigate the microstructure and mechanical properties of gradient-structured 301 stainless steel synthesized by surface mechanical attrition treatment. Microstructure gradient in both martensite content and grain size, from nanostructured (NS) surface with high volume fraction of martensite to coarse-grained (CG) austenite interior, is identified in the mechanical gradient surface layer. Gradient samples

exhibit

excellent

strength-ductility synergy due to

the structure

gradient-induced synergistic hardening and phase transformation induced plasticity (TRIP), i.e., combination of gradient and TRIP effects. Comparative analysis reveals that mechanical gradient can promote the phase transformation process, i.e. encouraging the TRIP effect. Keywords: Stainless steel; Gradient structure; TRIP effect; Synergistic deformation; Strength-ductility synergy.

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1. Introduction Metals with homogeneous microstructure may be high in strength or good in ductility but rarely occupy both simultaneously [1]. For example, the recrystallized coarse-grains (CG) material can display graceful ductility but low yield strength. Although the strength of ultrafine-grained (UFG) or nanostructured (NS) bulks processed by severe plastic deformation can be significantly improved, they usually exhibit negligible uniform elongation under tension [1]. Such paradox demands us to design new structures with optimal strength-ductility synergy. It is well recognized that ductility is primarily controlled by strain hardening efficiency [2]. Therefore, high strain hardening capability is a prerequisite for avoiding the strength-ductility paradox. Heterostructure design by combing multiple domains with significant mechanical incompatibility has attracted interests for its potential in optimizing strength-ductility synergy [3]. Typically, gradient-structured materials composed of NS surface layers and CG interior were extensively reported to possess improved strength and reasonable ductility [4-7]. During straining, the differences in strength and strain hardening capability between neighboring layers induces strain inhomogeneity between them, which is capable of triggering synergistic constraint between layers and introducing gradient strain distribution across layer boundary [5,8]. Accordingly, accumulation of geometrically necessary dislocations (GNDs) in strain gradient zone leads to the development of long-range internal stress, which can promote strengthening and strain hardening simultaneously, i.e., synergistic strengthening and strain hardening [9-11]. Transformation-induced plasticity (TRIP) is one of the most effective methods for improving strain hardening [12-14]. The TRIP process not only suppresses the development of plastic damages by delaying the formation of strain localization, but also introduces strain gradient at phase boundaries in order to maintain strain continuity [15]. GNDs accumulate around phase boundary to accommodate plastic strain gradient, which is expected to further improving the strength and strain hardening [16,17]. The mechanical performance of austenitic stainless steel is susceptible to martensite transformation under plastic deformation [18-21]. Deformation-induced γ→α’-martensite transformation can be significantly affected by the stress/strain status [22-24]. In this paper, the structure gradient-induced synergistic effect and TRIP effect are

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combined in 301 stainless steel (301ss) to optimize the strength-ductility synergy. Since the synergistic interaction between neighboring layers can reshapes the local stress state [17,19], and the martensite transformation can change the mechanical inhomogeneity between layers, it is expected that the gradient effect and TRIP effect in gradient 301ss can affect each other. Such mutual affection is interesting, but, regretfully, is seldom investigated in previous works. 2. Experimental methods The chemical composition of the as-received commercial 301ss plates (in thickness of 1 mm) is 0.098 C, 16.82 Cr, 7.41 Ni, 1.43 Mn, 0.73 Si, 0.017 P, 0.007 S, and the balance of Fe (all in mass %). The as-received sample is characterized by homogenous coarse austenite grains (Fig. 1). Gradient-structured 301ss plates were produced by surface mechanical attrition treatment (SMAT). During the SMAT process, steel shots in diameter of 3 mm were accelerated to a high speed using high-power ultra-sound to impact the annealed plates. Three types of gradient plates were symmetrically treated on both sides for 1 min, 5 min and 10 min, which are labeled as S1min, S5min and S10min samples, respectively.

Fig. 1 Optical image for the as-received 301 stainless steel. Dog bone-shaped tensile specimens with a gauge dimension of 12×2×1 mm3 were machined from as-processed gradient plates. Tensile specimens with only the topmost 0.2-mm-thick layer of S5min sample or the 0.4-mm-thick central layer of S1min sample were prepared by polishing away the other layers of integrated gradient tensile specimens. Uniaxial tensile experiments were carried out at room temperature at a

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strain rate of 5×10-4 s-1. Vickers microhardness measurements were conducted on the cross-section of gradient samples before and after tension, using a load of 50 g for 15 s. The measurements on each type of cross-section were performed at four independent positions. The gradient microstructures were characterized using transmission electron microscopy (TEM). TEM foils in the layers at different depths were obtained by carefully polishing the surface layer and the inner matrix to a thickness of ~50 µm, and then further thinning using ion beam milling. X-ray diffraction (XRD) was utilized to examine the distribution and transition of phases during tension. XRD measurements were performed using a CuKα radiation (lambda = 0.15418 nm), in a scanning range of 40–90° at 0.02° s-1. The volume fraction of retained martensite, Vα’, was calculated from the integrated intensities of diffraction peak (Ihkl) after background subtraction using the following equation [25,26], V =I

I +1.4Iγ

(1)

where Iγ is the average integrated intensity obtained at the (111)γ, (200)γ and (220)γ peaks, and Iα’ is that obtained at the (110)α’, (200) α’ and (211) α’ peaks. 3. Results 3.1 Gradient materials Fig. 2(a) shows the cross-sectional hardness distribution of S1min, S5min and S10min gradient samples. There is large hardness gradient in all of gradient samples, which indicates significant mechanical incompatibility between surface layers and center matrix. Clearly, the mechanical gradient is more pronounced as the processing time increasing. The S10min sample has the largest hardness difference between the top surface layer (as high as 560±20 Hv) and central matrix (~380±11 Hv). Fig. 2(b) is an optical image obtained from the center of S5min sample, showing that the central layer remains the as-annealed equiaxed coarse grains with an average grain size of ~21 µm. For the S1min sample, the parent austenite grains in the surface layer became a laminar microstructure with clear and sharp lamella boundaries and high dislocation density by SMAT treatment (Fig. 2(c)). The selected area electron diffraction (SAED) pattern (Fig. 2(d)) suggests a composited nano-lamellae containing γ matrix, twin and ε-martensite [22,25,27]. The ε-martensite is the intermediate phase during phase transformation from γ matrix to α’-martensite [28]. In addition, some granular

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martensite in size of several tens of nanometer is also detected, as indicated by reds arrow heads in Fig. 2(c). Fig. 2(e) shows the representative microstructure of the surface layer of S5min sample. The grain size statistically measured from tremendous enlarged TEM images is in a range of 30–100 nm. It is obvious that some irregular submicron and nanometer-sized α’-martensite are produced and embedded in the nano-lamellar matrix, as indicated by red arrow heads [29-31]. An extra diffraction of bcc α’-martensite can be identified in the SAED pattern of this microstructure (Fig. 2(f)), as compared with the SAED pattern in Fig. 2(d). Interestingly, the microstructure in the surface layer of S10min sample is obviously different. As shown in Fig. 2(g), there is a large amount of granular α’-martensite that replaces the lamellar structure. According to the SAED pattern (Fig. 2(h)) one can identify that this microstructure is also composed of γ matrix, twin, ε-martensite and α’-martensite. The diffraction spots of the α’ phase are significantly enhanced, indicating extremely large volume fraction of α’-martensite [32]. Fig. 3 presents the results of XRD measurements on as-SMATed gradient samples at different depth, which provides statistical analysis for the phase content. Austenite phase has the diffraction peaks of {111} γ, {200} γ, and {220} γ, and martensite phase shows peaks of {110} α’, {200} α’ and {211} α’ [25,33]. As shown, the intensity of martensite peaks is decreased while the austenite peaks are strengthened as the depth increases (Figs. 3(a-c)), indicating a gradual decrease of martensite volume fraction along depth (Fig. 3(d)). It is obvious that the sample experienced longer attrition time exhibits larger gradient in martensite volume fraction along depth. The volume fraction of martensite in the surface layer of S10min sample is as high as ~75%, implying that most of the γ-austenite has been transformed into α’-martensite by SMAT treatment. Different martensite content (Fig. 3(d)) is one of the main reasons for the difference in hardness gradient among three types of gradient samples (Fig. 2(a)). Importantly, the above microstructure observations clarify that the microstructure gradient of samples exists in both grain size and phase content. This can be attributed to the gradient decrease of accumulative plastic strain and corresponding strain rate from surface to interior during SMAT treatment [34]. Note that the central layer of three types of gradient samples exhibit similar phase content (Fig. 3(d)) but different hardness (Fig. 2(a)), indicating that the center of plate experienced small plastic strain

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but no phase transformation during SMAT treatment [35,36].

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Fig. 2 Three types of as-SMATed gradient samples. (a) Cross-sectional hardness distribution. (b) CG matrix in the central layer of S5min sample. (c) Typical TEM micrograph taken from the surface layer of S1min sample. (d) SAED pattern (with zone _

_

_

axes [110] γ//[110] twin//[1120]ε) corresponding to the lamellar microstructure in (c), showing the co-existence of f.c.c. γ matrix phase, f.c.c. twin and h.c.p. ε-martensite phase. (e) and (f) are the micrograph and associated SAED pattern taken from the surface layer of S5min sample, respectively. (g) and (h) are TEM results of the surface _

layer of S10min sample. (f) and (h) are the SAED patterns with the zone axes [110] γ//[1 _

_

_

10] twin//[1120]ε//[111] α’. All of the above TEM micrographs were taken at the depth of ~50 µm from the top surface layer of gradient samples.

Fig. 3 (a)-(c) XRD spectra measured at varying depth of as-processed S1min, S5min and S10min samples, respectively. (d) Variation of α’-martensite volume fraction along depth. 3.2 Improved strength-ductility synergy

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Fig. 4(a) shows the engineering tensile response of gradient samples. The freestanding 0.2-mm-thick NS surface layer displays extremely high strength (~1760 MPa) but disporting uniform elongation due to the high-volume fraction of martensite (Fig. 3(d)) and severely deformed microstructure (Fig. 2(e)). Table 1 summarizes the tensile properties of gradient samples. Surprisingly, gradient samples exhibit excellent strength-ductility synergy, although there is slight reduction in uniform elongation as increasing SMAT duration. For example, compared with the performance of homogeneous CG matrix (the purple curve), the yield strength of S5min sample is nearly doubled (from 359 MPa to 710 MPa), while the uniform elongation dropped only 28%. For the S1min sample, the yield strength is enhanced by 44% with even a slight increase in uniform elongation simultaneously. Since the improved uniform elongation of the NS surface layer achieved in gradient structure is due to the synergistic constraint from CG interior [5,8,37], the excellent strength-ductility synergy of gradient samples may suggest that the favorable properties of component domains, the high strength of NS surface layer and the graceful ductility of inner CG core, are invoked simultaneously during tension [10]. These indicate that a large mechanical inhomogeneity in gradient sample can introduce strong synergistic constraint between them during tension [17,38]. Fig. 4(b) presents the strain hardening behavior of gradient samples. The evolution of strain hardening rate (Θ) of 301ss can be divided into three stages [39]. The rapid decrease of Θ at the initial stage I is mainly due to the planar slip of partial and dissociated dislocations and the formation of ε-martensite plates [40]. In the second stage, the prevalent of γ→α’-martensite transformation leads to regaining strain hardening, i.e., TRIP effect, resulting in an Θ up-turn [41]. Interestingly, the phenomenon of Θ up-turn was also reported in gradient-structured IF steel and pure Cu samples which have single phase [5,42]. This suggests that the present Θ up-turn in gradient 301ss may be contributed from structure gradient and TRIP effects simultaneously. The final drop of strain hardening, i.e., stage III, is caused by dynamic recovery [36].

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Fig. 4 Tensile response of gradient samples: (a) Engineering stress-strain curves; (b) Θ versus true strain. Dotted lines illustrate the participation of the three Θ stages in homogeneous CG sample. The inset in (b) shows the start of Θ up-turning. Table 1 Mechanical properties of gradient samples, homogeneous CG matrix and freestanding 0.2-mm-thick nanostructured gradient surface layer. Yield strength

Ultimate strength

Uniform elongation

(MPa)

(MPa)

(%)

homogeneous CG

359±11

1174±15

61.9±2.1

S1min

518±12

1202±18

63.4±1.2

S5min

710±13

1236±16

43.5±1.7

S10min

995±17

1305±14

33.2±2.5

1640±16

1760±15

3.01±0.01

Sample

0.2-mm-thick surface layer

The γ→α’-martensitic transformation is generally accompanied by continuous stress and Θ serrations during straining [16,32,43]. It can be seen (both stress-strain and Θ curves) that the magnitude of stress serration in gradient sample seems more serious than that in CG sample. Moreover, the starting strain of stage II in gradient samples, i.e., the begin of Θ up-turning, is earlier than that of homogeneous CG matrix, as compared in the insert (Fig. 4(b)). These suggest that mechanical gradient may play a role in affecting the phase transition process. 3.3 Phase transformation during tension The S5min sample is taken as an example to investigate the microstructure evolution of gradient material during tension. Fig. 5 shows the bright- and corresponding dark-filed TEM micrographs at the same depth of ~50 µm but varying tensile strains. Fig. 5(a) shows the typical microstructure at a tensile strain of ~0.047. The austenite matrix is subdivided by banded and granular (marked by red arrow heads) martensite, forming submicron-sized parallelepiped domains (marked by

9

yellow arrows) within austenite interior. Such parallelepiped domains are the basic microstructural unit for accommodating plastic deformation, which is usually observed in deformed austenite materials [32]. When the applied tensile strain is increased to ~0.15 (Fig. 5(b)), the volume fraction of α’-martensite grains is increased. The parallelepiped domain interior is still austenite mostly but accumulated with extremely high density of dislocations. At the tensile strain of ~0.23 (Fig. 5(c)), the frame of the parallelepiped domain becomes weaker, which is basically converted into α’-martensite. Fig. 5(d) shows the microstructure of the sample after fracture. All domain interiors were transformed into α’-martensite. The SAED pattern shows a discontinuous ring of the bcc structure.

Fig. 5 The evolution of microstructure in the surface layer of S5min sample during tensile test. TEM images and SAED patterns at the depth of ~50 µm at the tensile strains of (a) ~0.047, (b) ~0.15, (c) ~0.23, (d) fracture. The insets in (a), (b) and (c) _

_

_

_

are the SAED patterns with identified zone axis [110] γ//[110] twin//[1120]ε//[111] α’, _

[011] γ//[111] α’ and [111] α’, respectively, while the SAED pattern in the inset of (d) _

is close to zone axis [111] α’.

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Fig. 6 presents the phase transformation characterized using XRD, at the depth of ~50 µm. With the increase of strain, the diffraction intensity of martensite phase is obviously increased. As summarized in Fig. 6(d), for all of the three types of gradient samples, the martensite volume fraction is gradually increased with applied strain and finally reaches almost saturation (~98% α’-martensite volume fraction). This result is in accordance with above microstructure observations (Fig. 5).

Fig. 6 α’-martensite transformation during tensile testing. (a)-(c) XRD spectra measured in gradient samples with varying tensile strains, at the depth of 50 µm. (d) Volume fraction of α’-martensite as a function of tensile strains.

3.4 Evolution of hardness gradient during tension Gradient samples at varying tensile strains, i.e., the lowest (P1), middle (P2) and highest (P3) points at the second strain hardening stage (seeing Fig. 4(b)), were selected for hardness test. As the cross-sectional hardness profiles shown in Figs. 7(a1, b1, c1), in addition to the CG matrix, the surface layers in all of gradient materials exhibit obvious hardness increment (∆H) with straining. This is different from the

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gradient materials with single phase that the NS surface layer experiences almost no hardness increase during tension due to the initial extremely high dislocation density accumulated during SMAT, thus a low strain hardening ability [5]. For the S1min sample (Fig. 7(a2)), the surface layer exhibits a slightly larger ∆H than the central CG matrix. In contrast, ∆H in the NS surface layers of S5min and S10min samples (Figs. 7(b2) and (c2)) are obviously lower than that in the central CG matrix.

Fig. 7 Hardness evolution of gradient samples during tension. (a1, a2) S1min sample. (b1, b2) S5min sample. (c1, c2) S10min sample. The left column shows the plots of the cross-sectional hardness distribution at varying tensile strains. The right column which is fitted by a Gaussian distribution function shows the hardness increment (∆H)

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relative to the as-SMATed state. Fig. 8 presents the evolution of hardness difference between the surface layer and inner core of gradient samples, with increasing applied tension strain. The hardness difference between surface layer and inner core can approximately represents the magnitude of mechanical gradient of gradient sample [9]. Therefore, it can be summarized from the data in Fig. 7 that during tension (i) the mechanical gradient of gradient materials is gradually decreased, (ii) within the strain regime from ~5%-15% mechanical gradient remains almost stable, and (iii) the sample with larger initial mechanical gradient exhibits larger decrease rate in gradient during subsequent tension. In addition, the mechanical gradient of S1min at 6% tensile strain increase in comparison with the starting state, possibly due to the martensite transformation occurring in surface layer but not in inner core.

Fig. 8 The variation of hardness gradient with increasing tension strain. Each datum represents the hardness difference between topmost surface layer and central layer at certain applied strain. 4. Discussion 4.1 Combining gradient and TRIP effects The microstructure observation (Fig. 5) and phase content examination (Fig. 6) confirm a continuous γ→α’-martensitic transformation in the surface layers of three types of gradient samples during tension. Since the density of crystalline defects in surface layer has been accumulated to near saturation during SMAT treatment (S5min and S10min in particular), the large ∆H produced during tension in this layer is largely 13

caused by the increase of α’-martensite content, rather than dislocation hardening [32]. The difference of ∆H in the surface layer among three types of gradient sample (Figs. 7(a2-c2)) can be attributed to the different initial martensite content and transformation rate during strain (Fig. 6(d)). Specifically, higher ∆H can be achieved if the as-processed surface layer has lower initial martensite volume fraction (Fig. 3). In this regard, the change of mechanical gradient (Fig. 8) during tension should also be determined by the difference of phase transformation rate between surface layer and inner CG core. The initial CG interior has lower volume fraction of martensite phase, in which a relatively higher transformation rate is expected. This should be one of the reasons for the fast decrease of hardness gradient in samples (Fig. 8). Importantly, ∆H is an indicator for the retained strain hardening [5]. The large ∆H in surface layer (Fig. 7) indicates that pronounced strain hardening was produced by phase transformation in this layer. At the same time, it has also been widely demonstrated in single phase gradient material that the hetero-deformation of layers with mechanical incompatibility can produce extra strain hardening due to the strain gradient-dependent GNDs piling-up across layer boundary and development of long-range internal stresses [5,8,9,11,44-47], i.e., Hetero-deformation induced (HDI) strain hardening [48,49]. Therefore, the excellent strain hardening (accompanied by significant Θ up-turning) displayed by gradient 301ss samples (Fig. 4(b)) should be a synergistic response for phase transformation-induced hardening and HDI hardening. Furthermore, the inhomogeneous in-layer distribution of new formed martensite and remained austenite phases is effective in partitioning applied strain and thus impeding the propagation of strain localizations, such as shear bands and cracks [1]. Strain partitioning between austenite and martensite phases during straining introduces GNDs accumulation around boundaries in order to remain strain continuity, which is expected to further improving the HDI hardening and strengthening [11,27]. These phase transformation-related effects help with overcoming the quick development of plastic damages in component layers [32]. This point can be indirectly supported by the experimental observations that catastrophic instability (failure) occurred in gradient samples (Fig. 4(a)) soon after the saturation of phase transformation in surface layer (Fig. 6(d)). The above mechanisms suggest that the stable plasticity and excellent strain hardening performances of gradient 301ss samples can be partially attributed to the combination of structure gradient and TRIP effects.

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4.2 Mutual influence of gradient and TRIP effects The steep

hardness

gradient along depth

indicates large mechanical

incompatibilities in yield strength and strain hardening capability between neighboring layers [48]. In the yielding process, the inner softer layer starts plastic deformation firstly, while the harder surface layer still remains elastic. The incompatibility in lateral shrinking rate will lead to mutual constraint between them, with the surface suffers a lateral compressive stress from inner layer, i.e., elastic/plastic interaction [46]. Soon after yielding, local fast shrinking is prone to develop in the hard surface layer due to the relatively low strain hardening efficiency. However, such unstable strain localization can be suppressed by a lateral tensile stress from inner stable matrix [5], introducing the plastically stable/unstable mutual constraint between them. These mutual constraints change the applied uniaxial stress state to a multi-axial state, which helps to improve the level of internal stress [50]. Such hetero-deformation induced high internal stress is effective in encouraging phase transformation [28,51], thereby promoting strain hardening. This is the reason for the earlier start of Θ up-turning and the relatively large flow stress and Θ serrations in tension test (Fig. 4). Moreover, as indicated by the tensile response (Fig. 4(a)), the stable elongation of a surface layer can be exhibited only when it is supported by the CG inner matrix, i.e., in an integrated gradient sample. Therefore, it is the martensite-austenite (NS-CG) microstructure gradient-induced synergistic constraint which maintains the stable plasticity of surface layer and provides the primary opportunity for the continuous occurrence of martensite transformation (Figs. 5 and 6(d)). On the other hand, the continuous martensite transformation in surface layer helps to remain a reasonable mechanical gradient from surface to interior (Figs. 7 and 8), although the fast increase of martensite content and dislocation accumulation in inner CG core plays an unwanted role in reducing the mechanical gradient. This suggests that TRIP and gradient effects affect each other during the whole straining process. A comparison of tensile properties with previous studies on 301ss [52-55] is shown in Fig. 9. This is particularly true that, in the high strength regime, the gradient 301ss usually has higher ductility than those of homogeneous samples. For example, at the yield stress level of ~1000 MPa, the uniform elongation of current gradient 301ss material is higher than 33%, while it is smaller than 10% for most of the

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homogeneous UFG and NS 301ss synthesized by general severe plastic deformation routes [53,54]. This suggests that the present gradient-structured materials have an optimal strength-ductility synergy, which is difficultly achieved in conventional materials with homogeneous microstructure. In particularly, the combination of gradient structure and TRIP effect is effective in producing high strength and high ductility materials.

Fig. 9 Representative tensile properties of 301ss. Hollow symbols are reported data for homogeneous 301ss with coarse, ultrafine and nanoscale grain sizes [52-55], and solid symbols represent other gradient 301ss [55]. Worth to note is that dense twins (Figs. 2(c) and (e)) with a gradient reduction in density along depth can also be introduced in the surface layers during SMAT processing, due to the low stacking fault energy of austenite stainless steel [27,29,36]. This may contribute to graceful strength-ductility synergy in two ways. Firstly, it has been widely confirmed that dense twin boundaries are effective in hindering dislocation slip and promoting dislocation accumulation, which are benefit for strengthening and strain hardening [27,56]. Secondly, the density gradient of twins plays a role in enlarging the mechanical incompatibility along depth, such as the strength and hardness gradients (Fig. 2(a)), which optimize the mechanical performances by promoting the synergistic interaction between the hard surface and soft core, i.e., improving the gradient effect as discussed before. In other words, the

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as-discussed mechanical gradient and the corresponding effects on tension response are caused by several microstructural inhomogeneities including grain size, twin density and phase content gradients [21,27,28,36]. 5. Conclusion The mechanical property and microstructure evolution of gradient-structured 301ss are systematically investigated in this work. The main conclusions are summarized as follows: (i) The 301ss materials synthesized by surface mechanical attrition treatment exhibit large mechanical gradient, and microstructure gradient in both martensite content and grain size from surface to interior. (ii) Due to the combination of structure gradient and TRIP effects, the gradient 301ss materials display excellent strength-ductility synergy. Specifically, structure gradient causes hetero-deformation induced synergistic hardening. Phase transformation continuously occurred until the martensite content reaches saturation, leading to martensite strengthening and strain hardening. (iii) The synergistic constrain from CG interior helps to maintain the plastic stability of hard surface layer and thus provides opportunity for the continuous occurrence of martensite transformation in this layer. The high internal stress caused by synergistic constraint between mechanical incompatible layers plays a role in promoting phase transformation. Acknowledgements This work was supported by the National Natural Science Foundation of China (Nos.11672195 and 51931003) and Sichuan Youth Science and Technology Foundation (2016JQ0047). Reference [1] T.G. Langdon, Twenty-five years of ultrafine-grained materials: Achieving exceptional properties through grain refinement, Acta Mater. 61 (2013) 7035-7059. [2] Y.T. Zhu, X.L. Wu, Ductility and plasticity of nanostructured metals: differences and issues, Mater. Today Nano 2 (2018) 15-20. [3] X.L. Wu, Y.T. Zhu, Heterogeneous materials: a new class of materials with unprecedented mechanical properties, Mater. Res. Lett. 5 (2017) 527-532. [4] T.H. Fang, W.L. Li, N.R. Tao, K. Lu, Revealing extraordinary intrinsic tensile plasticity in

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Declaration of interests ☒ 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. ☒The authors declare the following financial interests/personal relationships which may be considered as potential competing interests:

There are no financial interests/personal relationships which may be considered as potential competing interests.