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Introducing a hierarchical structure for fabrication of a high performance steel
Materials Chemistry and Physics 129 (2011) 1096–1103
Contents lists available at ScienceDirect
Materials Chemistry and Physics journal homepage: www.elsevier.com/locate/matchemphys
Introducing a hierarchical structure for fabrication of a high performance steel A.Y. Chen a,b , H.H. Ruan b , J.B. Zhang c , X.R. Liu a , J. Lu d,∗ a
School of Material Science & Engineering, Shanghai Institute of Technology, 201400, China Department of Mechanical Engineering, The Hong Kong Polytechnic University, Hong Kong, China c Baosteel Technology Centre, Shanghai 201900, China d Department of Manufacturing Engineering and Engineering Management, City University of Hong Kong, Hong Kong, China b
a r t i c l e
i n f o
Article history: Received 6 January 2011 Received in revised form 22 May 2011 Accepted 25 May 2011 Keywords: Nano materials Multilayered structure Plastic deformation Mechanical properties
a b s t r a c t The multifunctional diversities existing in nature provide clues to speculate the structure–property–function relationships. A hierarchically structured steel is designed by using principles derived from nature and fabricated in situ by a one-step method of surface mechanical attrition treatment (SMAT). The microstructure of the processed steel is characterized by multilayered structure with hard nanocrystalline surface and compliant inner-layer, in particular with a smooth mechanical gradient induced by dual-phase constituents and multiscale grain size distribution. The hierarchically structured steel exhibits simultaneously high stiff, strong and large ductility, which originate from the joint deformation mechanisms of distinct reinforcing layers. The four layers present their own unique deformation mechanisms, including second-phase hardening, transformation induced plasticity and twin strengthening. The unique spatial form of gradation can release stress concentration and improve energy-dissipation leading to exceptional mechanical properties compared with the uniform materials. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Steels are used far more extensively than any other metals, polymers, or ceramics as high-performance structural materials [1]. Now, the most famous iron and steel enterprises are facing the challenges to develop novel steels with exceptional strength and ductility driven by the requests of the automotive and aeronautic industries for safety and lightweight. The strengthening technologies of the conventional steels are involved in addition of alloying elements, fabrication of multiple phase structures [2–5], and refinement of grain size down to nanocrystalline (NC) or ultra-fined grain (UFG) scale [6,7], etc. Unfortunately, the strength and ductility tend to be mutually exclusive, which means higher strength is generally accompanied by a lower plasticity. The deterioration of plasticity is accelerated with the decrease of grain size down to NC/UFG scale [8]. Nature is the origins in design of multifunctional materials, which can provide clues to explore the structure–property–function relationships [9]. Many biological materials, such as bone [10,11], nacre [12], or scale [13], exhibit stiffness as well as high toughness resulting from their lamellar structures to combine the multiscale components with multiple
∗ Corresponding author at: Department of Manufacturing Engineering and Engineering Management, City University of Hong Kong, Hong Kong, China. E-mail address: [email protected] (J. Lu). 0254-0584/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.matchemphys.2011.05.068
functions. For example, the armour of fish scale consists of four different reinforcing layers with decreased gradation of modulus and hardness from the outer to inner layer. The inner layer is soft, while the high strength is achieved by nanoscaled grains in surface layer with a certain thickness [14]. Gradation in microstructure is the most important common feature in biological structures, which exhibits more damage-resistant than the monolithic counterpart [15]. Such layered biological structures have served as a design methodology applied in inorganic reinforced polymer-matrix artificial composites [16]. Although the notion of mimicking natural structures in man-made multilayered composites has generated enormous interest [17], the complex hierarchical designs in biological composites are extremely difficult to replicate synthetically, especially in covering so many dimensions (from nano to micro). Additionally, the synthetic technologies are generally complex, interface flaw-induced, size-dependent, and difficult for engineering applications, which severely hinder the implementation of hierarchical designs. We applied some of the structural concepts found in biological materials to design and fabricate hierarchical steel. A multilayered structure composed of diverse microstructures spanning over various length scales was deliberately chosen to implement design strategies in terms of tailoring strength and toughness. The hierarchically structured (HiSt) steel is assembled from nanometer scale to micrometer scale with functional gradient variation without interface in one-step by using an advanced technology of surface mechanical attrition treatment (SMAT) [18,19].
A.Y. Chen et al. / Materials Chemistry and Physics 129 (2011) 1096–1103
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304 SS sheets (100 mm × 90 mm × 1 mm) were treated by SMAT. The SMAT was performed by the optimized parameters, as given in Table 1. Both sides of AISI 304 SS sheets were treated under the same conditions. Since the SMAT is performed in the vacuum and at room temperature, the medium contaminations by the element diffusion of GCr15 balls are very limited. Thus, the contamination influence on the deformation mechanism and the properties of the HiSt steel is very small. The samples for the tensile tests were cut into dog-bone-shaped specimens with a gauge length of 30 mm and a cross section of 6 mm × 1 mm. Tensile tests were performed at room temperature at a strain rate of 6.7 × 10−4 s−1 using a MTS Alliance RT/50 Materials Testing System. Four specimens were used to obtain consistent stress–strain curves. Vickers microhardness was measured by applying a load of 50 g for 15 s and taking the average of six separate measurements on each depth. A Philips Xpert X-ray diffractometer (XRD) with Cu K␣ radiation was used to determine the phase constituent and estimate the phase content. The XRD tests were carried out along the cross-sectional direction, in these tests, the samples were carefully prepared by mechanical polishing from the surface to reach the targeted depths. The volume fraction of ˛ martensite phases was estimated from the peak integrated intensities Ih,k,l after background subtraction, as introduced in Ref. [20]. The average amount of ˛ martensite phases is calculated from two separate measurements on each depth. Transmission electron microscopy (TEM) observations were made with a JEM 2010 transmission electron microscope with an operating voltage of 200 kV. The plane-view TEM foils were ion-thinned at low temperature. The fracture surfaces of the specimens were investigated using a HITACHI S-4200 field emission scanning electron microscope (SEM).
3. Results and discussion Fig. 1. The schematic principle of SMAT technology.
3.1. Hierarchical microstructure 2. Experimental 2.1. SMAT technology SMAT is a newly developed technique, which can refine the coarse grain (CG) structure progressively down to the nanometer regime by means of repeated multidirectional impacts with high energy onto the sample surface [18]. The schematic principle of SMAT technology is shown in Fig. 1. The requirements of specific structures and properties shall be realized by manipulating the strain and strain rate at different depths from the surface. The resulted strain and strain rate at different depths from the surface are the driven force to refine the grain size into nano and submicrometer scale, simultaneously associated with phase transformation. The processed steel with dimensions of semi-industrial experiment can be fabricated by step moving manner. 2.2. Experimental procedure The chemical composition of AISI 304 austenite stainless steel (SS) is 17.6Cr, 8.3Ni, 0.04C, 0.19Si, 1.36Mn, 0.006S, 0.020P and balance Fe (all in mass%). The
The hierarchical structures originate from three critical aspects: (1) layered structure with hard surface and compliant interior; (2) hierarchical components spanning over various length scales; (3) most importantly, mechanical gradient between each successive layers. Guided by these strategies, the microstructure of HiSt steel is characterized by a symmetrically layered structure, which consists of four different reinforcing layers, as shown in Fig. 2 (from outer to inner layer): NC layer, dual phase layer, twinning layer and CG layer, respectively. The outermost NC layer provides superior stiffness and hardness by using hard ˛ martensite phase. The design origin of dual phase layer is to provide the important mechanical transition, which consists of hard ˛ martensite and soft austenite phase, resembling the biological structure of strong inorganic platelets embedded in a soft, ductile organic matrix. The third twinning layer is composed of multiscaled twins. The underneath softer, more compliant CG layer offers overall ductility. The strength ( y ) of multilayered composites can be estimated from the volume fraction (Vi ) and the tensile strength ( i ) of different layers, in terms of the “rule-of-mixture” [21,22]: y =
m
Vi i
(1)
i=1 m
Vi = 1
(2)
i=1
Fig. 2. Bioinspired design principle of HiSt steel. The outermost surface is composed of nano ˛ martensite phase to form hardest surface layer. The next is a transitional layer, consisted of dual phase structure with ˛ martensite phase and austentite phase in UFG scale. The UFG and ˛ martensite phase can further increase strength, and UFG austentite phase can deform by transformation-induced plasticity. The third layer is a twin layer, strengthening by twin boundaries. The inner layer is a complaint coarse grained (CG) layer to improve plasticity. From outer to inner, the grain size increases from nanometer to micrometer and the volume fraction of ˛ martensite phase decreases to establish a dual graded structure.
Thus, the higher strength of the HiSt steel shall be achieved by reducing one dimension of building blocks to the nanoscale, and increasing the volume fraction of hard layer. However, the ductility of the multilayered composite is contradictory to the strength, i.e. the stronger materials exhibit more brittle. Two critical designs are introduced to improve the ductility by additional deformation mechanisms, i.e. special deformation mechanisms of each layers and the gradation between the layers. Therefore, a dual gradient structure of grain size and phase constituent distribution is designed to induce mechanical gradient, mimicking the gradation in biological materials. Integrated the relationship of strength and ductility, the thickness of the four layers is selected to be (from outer to inner) 5–20 m, 60–200 m, 100–300 m, and 100–200 m, respectively. The above hierarchical structures are
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A.Y. Chen et al. / Materials Chemistry and Physics 129 (2011) 1096–1103
Fig. 3. Representative TEM images and corresponding SAED patterns of HiSt steel at four layers with depth of (a) 5 m, (b) 100 m, (c) 250 m, (d) 400 m, respectively. The lower inset of (a) is the grain size distribution.
difficult to be assembled in nanoscale and without layer interfaces by using bottom-up methods, but a newly developed technique of SMAT can prepare this complex composite in semi-industrial experiment by one-step manner. Fig. 3 presents the representative transmission electron microscopy (TEM) images and corresponding selected area electron diffraction (SAED) patterns of different layers of the AISI 304 SS sheets after treated by SMAT. The outer layer at 5 m depth from surface is composed of ˛ martensite phase in body centered cubic (BCC), where the grain size distribution is in the range of 2–100 nm and the average grain size is about 10 nm, as shown in Fig. 3a. The microstructures observed at 100 m depth from surface, exhibit a dual phase structure of ˛ martensite phase and austenite phase in face centered cubic (FCC) both with grain size from several nanometers to hundreds of nanometers, as shown in Fig. 3b. The corresponding SAED pattern exhibits mixed rings of ˛ and phase, indicating by the grey rings and orange rings in Fig. 3b, respectively. Fig. 3c shows the microstructure of twinning layer at 250 m depth from surface, which consists of twins with spacing of several hundreds of nanometers. The center layer at 400 m depth from surface is characterized by multiscale dislocation cells (DCs), as shown in Fig. 3d. According to the TEM observations, the HiSt steel is composed of micron-sized grain/DC (1–20 m), submicron grain/DC (0.5–1 m) and NC/UFG (0–0.5 m) with volume fraction of 40%, 30% and 30%, respectively, as given in Fig. 4. Furthermore, the
volume fraction of ˛ martensite phase decreases sharply from 95% in the surface of NC layer to about 6% in the center of CG layer. The grain size and phase constituent both show gradient distribution.
Fig. 4. Dual gradient distribution of the grain size and volume fraction of ˛ martensite phase with depth in the HiSt steel.
Treatment temperature (◦ C)
25 40
1099
98 ∼0.5 250 GGr15 50
8
Vacuum degree (Pa) Quantity of balls (pcs) Ball material
Diameter of ball (mm)
Impact velocity (m s−1 )
80 × 80
Fig. 5. Gradient variation of microhardness with depth of the HiSt steel at crosssection at different tensile strain.
Vibrating frequency (Hz)
Table 1 Processing parameters of SMAT.
Impact area (mm2 )
Treatment time (min)
A.Y. Chen et al. / Materials Chemistry and Physics 129 (2011) 1096–1103
3.2. Mechanical properties The dual gradient structure results in a gradient distribution of hardness, and the smaller grain size and larger volume fraction of ˛ martensite phase make higher hardness. The microhardness, H, decreases with distance from outer to inner, presenting a parabolic shape (in Fig. 5). The highest average microhardness (∼500 HV) of the outer layer is attributed to the nanoscaled ˛ martensite phase. The dual phase layer has a larger grain size and smaller amount of ˛ martensite phase, therefore, a steep variation of microhardness is observed. The combination of high tensile strength and high ductility is the most impressive feature of HiSt steel. The engineering stress–strain curve of HiSt steel shows high yield stress ( 0.2 ) of 600 MPa, more than two times that of the original steel (270 MPa), and ultimate tensile stress ( b ) of 875 MPa, as shown in Fig. 6a. The total elongation to failure (εb ) and uniform elongation (εu ) of the HiSt steel also maintain quite high level of 51% and 45%, respectively. Interestingly, the hierarchical steel significantly deviates from the general strength–ductility rule of the conventional steels (interstitial free (IF) steel, mild steel, C–Mn (CMn) steel and high strength low alloy (HSLA) steel) and advanced high strength steels (AHSS), as given in Fig. 6b. The advanced high strength steels, such as dual-phase (DP) steel, transformation-induced plasticity (TRIP) steel, and complexphase (CP) steel are developed far more extensively driving by the requests of safety and weight saving. Compared with the other steels at the same strength level, for example, the TRIP steel with yield strength of 600 MPa (TRIP 600), the elongation to fracture is about 20%, but that of the HiSt steel is as high as 51%, which indicates the design strategies based on bioinspired concepts are fascinating. 3.3. The strength–ductility mechanisms The multilayered structure and gradient within and between layers work collectively to provide enhanced mechanical performance. The present steel consists of four reinforcing layers, each of them deforms in different mechanisms. The NC structure is selected by the super hard ˛ martensite phase particularly with grain size in several nanometers. The NC materials often exhibit superior strength, but brittleness. In order to balance the strength–ductility relationship of NC materials, the thickness of NC layer with grain size less than several tens of nanometers in HiSt steel is about 20 m, similar to that found in fish scale, which shows a highly stable deformation manner by a circumferential cracking [13,15]. The dual phase layer is an important transitional layer, which embod-
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Fig. 6. Comparisons of mechanical properties between the HiSt steel and other steels. (a) Engineering stress–strain curves of the HiSt steel and original AISI 304 SS. The inserts are the dimension of the tensile sample and the picture after tensile, respectively. (b) Strength–ductility relationship of different structured steels.
ies: (1) to delocalize the deformation of brittle NC layer; (2) to supply the required strength by NC/UFG grain size and hard ˛ phase; (3) to generate TRIP mechanism. The hard martensite phase can be used as precipitated phase to further strengthen, and the soft austenite phase can transform to ˛ martensite phase to induce TRIP effect [23–25]. The third layer, as an adapter with close value of microhardness to the fourth layer, presents twin strengthening. Twin boundaries (TBes) can act as barriers to dislocation motion, thus, provide a strengthening mechanism similar to grain boundaries. Moreover, Suresh et al. [26] had constructed a three-dimensional model to point out that the highly coherent twin structure can not only block dislocation action at the TBes, but also avoid the lost of ductility characteristic of nanocrystalline. Thereby, the twinning layer can not only enhance the strength, but also improve the ductility [5,27,28]. The underlying softer, more compliant CG layer dissipates energy via plasticity. The multiscaled dislocation cells in the fourth layer transform to martensite, but the refined dislocation cells can primarily transform to the martensite with thin lath structure. Therefore, each layer exhibits special deformation behaviour to induce excellent combination of strength and ductility, which is not the case for homogenous systems. Besides the important multilayered structure, the junctions between layers are clearly ‘functionally gradient’, which means that they possess gradual changes in properties, as well as no interfaces between multilayers. The gradient structure can promote
an effective transitional region to reduce stress concentration at the intersection between interfaces. If the hardest NC layer transfers load directly through the NC-CG junction, such as the case of NC coating, this may lead to magnify interfacial tensile shear stress, promoting brittle tendency [29]. The junction of NC layer and dual phase layer possesses a steep gradation relative to their neighbouring twinning layer and CG layer. During tensile deformation, compliant CG layer, twinning layer and dual phase layer deform jointly with graded increase of hardness by the introduced strengthening mechanisms, until their magnitudes are close to that of the original NC layer, as shown in Fig. 5. The hardness variations of HiSt steel exhibit obviously gradient tendency, indicating an important smoothing role of gradient in deformation. A sequence of strengthening mechanisms is suggested based on experimental observations, which can be subdivided into three stages, as illustrated in Fig. 7. In the early stage 1, the shear stress induced by the mismatched strengths of the multilayers can trigger the tangential motions of nanocrytalline, such as GB sliding and grain rotation [30–32]. The tangential motions can result in a sheared zone [32]. The well-defined shear bands are formed at surface, characterized by thin and sheet-like regions, where high plastic strain is localized, as indicated by the white arrows in Fig. 7a. The shear localization induced by GB sliding is a dominant deformation mechanism in nanocrystalline materials, especially for the grain size on the order of 20 nm and lower [31]. Coupling of the boundary motions and grain rotations can lead to grain growth, driven by the reduction of the surface free energy [32]. The TEM observations reveal larger grain size and enlarged grains along the shear localization regions at 20% plastic strain, where the average gain size is increased to about 100 nm with a much wider grain size distribution, as presented in Fig. 8a. The growth of grain size directly results in the decrease of surface hardness, which is in agreement with the result of microhardness variation in Fig. 5. These larger grains are thought to provide larger mean free paths for dislocations and sink of these dislocations because the grain boundaries are the primary sources for dislocation interactions (cross slip, reaction, etc.) [31]. The changes of dislocation density estimated from HRTEM images confirm the dislocation annihilation, where dislocation density in the surface and subsurface is high up to 35–40 × 1016 m−2 , but decreases to about 10–15 × 1016 m−2 at 20% strain, as shown in Fig. 8b. The experimental observations, such as shear bands, grain growth, and reduced dislocation density, show the complex deformations of NC by GB sliding and grain rotation. Meantime, the dislocation-mediated plasticity is dominant in twinning and CG layer under applied stress. Compared Fig. 7b with Fig. 3c, the high density of tangled dislocations is accumulated at twin boundaries. The obvious increases of dislocation density of twinning and CG layer are observed by HRTEM images, as given in Fig. 8b. At stage 2, when the plastic deformation increases, the martensite transformation is induced by plastic strain. For the specimen deformed at 30% strain, the ˛ martensite phase with random orientation in UFG layer and lath shape in CG layer are observed at depth of 50 m and 300 m, respectively, shown in Fig. 7c. The volume fraction of ˛ martensite phase in function of plastic strain varies severely, indicating an obvious strain-induced martensite transformation, as shown in Fig. 8c and d. The amount of marteniste transformation agrees with the change tendency of hardness distribution in Fig. 5. The austenite phase of different sizes transforms to martensite phase to hold out the uniform straining. At stage 3, it is interesting to find that the crack initiation presents in the inner layer but not in brittle NC layer. The side-view from the fracture point (in Fig. 7g) presents many cracks, which initiate and propagate at the interlayer. In addition, the tensile fracture surface is embodied with many microcracks in the twinning/CG layers, which are not observed in the NC/UFG layer (in Fig. 7f and h).
A.Y. Chen et al. / Materials Chemistry and Physics 129 (2011) 1096–1103
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Fig. 7. A schematic illustration depicting the deformation mechanism of the HiSt steel. (a) SEM micrograph of NC layer; (b) TEM image of twin layer at 20% strain; (c) and (d) TEM image of UFC and CG layer at 30% strain, respectively; (e–h) fracture morphologies after tensile test; (f) is the general view of the fracture; (e) is the magnification of transition zone of UFC/twin layer; (f) is the magnification of CG layer and (g) is the side view of tensile sample close to the fracture point.
In contrast to larger equiaxed dimples of CG layer, the transitional zone of UFG/twinning layer exhibits much smaller dimples elongated in opposite direction (in Fig. 7e). This distinct transition of grain size dependence on fracture surface suggests that the multiscale structure can effectively improve the plasticity of NC/UFG materials. It is interesting to find that the nano-sized grains in surface present a consecutive distribution during plastic deformation. However, the grain size of the surface nano-sized grains is increased from the original 10 nm to about 100 nm at 20% strain and further increased to 250 nm after tensile test, and the thickness is reduced
from about 20 to 8 m after tensile test. The growth of nano-sized grains and thinner of the nano layer are the results of the grain boundary sliding and grain rotation. Although a novel engineering material with high performance is achieved by implementing bioinspired concepts, there still has large space to be optimized in microstructures, such as the structure of soft layer, the layer thickness and gradient between layers. Further strengthening by grain refinement and the increased proportion of NC/UFG layer may induce the deterioration of ductility. Therefore, the increased proportion of twinning layer inevitably
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Fig. 8. Microstructures of HiSt steel of different strains. (a) TEM image of 5 m depth at 20% strain, the upper inset is the SAED pattern and the lower is grain size distribution; (b) dislocation density; (c) the volume fraction of ˛ martensite phase with strains and (d) XRD patterns after tensile test.
becomes an optimal selection for its comparable level of higher strength, and still providing substantial ductility. The improvement of strength and ductility can be anticipated by reduction of the coarse grained layer proportion and development of high density nanotwinned structure to meet the higher performance requirement. 4. Conclusions Based on the mechanical design strategies to mimic the structure of fish scale and seashell, a complex and multiscale material is designed and fabricated by using an advanced but simple processing of surface mechanical attrition treatment. The hierarchically structured steel is composed of a symmetrically layered structure, e.g. NC layer, dual phase layer, twinning layer and coarse grained layer from outer to inner, respectively. The junctions between layers present a gradient transition induced by a dual gradient structure of grain size and phase constituent distribution. Coupling multiscaled deformation behaviour over the nanoscale to microscale provides additional toughening mechanisms similar to those acting in natural materials. The hierarchical steel significantly deviates from the general strength–ductility rule of the conventional steels and advanced high strength steels. The impressive combination of strength and ductility of the hierarchical steel is obtained by a synergistic effect of multiple deformation mechanisms. These results indicate that microstructural features of hierarchical steel put bioinspired designs into practice, which should guide the synthesis of engineered structural materials.
Acknowledgements This work was supported by the Research Grants Council of the Hong Kong Special Administrative Region of China under the CityU8/CRF/08 project and by the research project (No. ITP/004/08NP) of the Hong Kong Innovation and Technology Commission (ITC). Financial support from the National Natural Science Foundation of China (Grant No. 51074107) is also acknowledged. This work is also supported by Shanghai Leading Academic Discipline (No. J51504). References [1] M. Militzer, Science 298 (2002) 975. [2] P. Podsiadlo, A.K. Kaushik, E.M. Arruda, A.M. Waas, B.S. Shim, J.D. Xu, H. Nandivada, B.G. Pumplin, J. Lahann, A. Ramamoorthy, N.A. Kotov, Science 318 (2007) 80. [3] K.T. Park, S.Y. Han, B.D. Ahn, D.H. Shin, Y.K. Lee, K.K. Um, Scripta Mater. 51 (2004) 909. [4] P.J. Jacques, Q. Furnémont, F. Lani, T. Pardoen, F. Delannay, Acta Mater. 55 (2007) 3681. [5] O. Bouaziz, S. Allain, C. Scott, Scripta Mater. 58 (2008) 484. [6] Y.H. Zhao, X.Z. Liao, S. Cheng, E. Ma, Y.T. Zhu, Adv. Mater. 18 (2006) 2280. [7] X.H. Chen, J. Lu, L. Lu, K. Lu, Scripta Mater. 52 (2005) 1039. [8] N. Tsuji, Y. Ito, Y. Saito, Y. Minamino, Scripta Mater. 47 (2002) 893. [9] C. Sanchez, H. Arribart, M. Madeleine, G. Guille, Nat. Mater. 4 (2005) 277. [10] R.K. Nalla, J.H. Kinney, R.O. Ritchie, Nat. Mater. 2 (2003) 164. [11] K. Tai, M. Dao, S. Suresh, A. Palazoglu, C. Ortiz, Nat. Mater. 6 (2007) 454. [12] S. Kamat, X. Su, R. Ballarini, M. Nassirou, A.H. Heuer, Nature 405 (2000) 1036. [13] B.J.F. Bruet, J. Song, M.C. Boyce, C. Ortiz, Nat. Mater. 7 (2008) 748. [14] K.D. Jandt, Nat. Mater. 7 (2008) 692. [15] K.J. Koester, J.W. Ager, R.O. Ritchie, Nat. Mater. 7 (2008) 672. [16] L.J. Bonderer, A.R. Studart, L.J. Gauckler, Science 319 (2008) 1069.
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Contents lists available at ScienceDirect
Materials Chemistry and Physics journal homepage: www.elsevier.com/locate/matchemphys
Introducing a hierarchical structure for fabrication of a high performance steel A.Y. Chen a,b , H.H. Ruan b , J.B. Zhang c , X.R. Liu a , J. Lu d,∗ a
School of Material Science & Engineering, Shanghai Institute of Technology, 201400, China Department of Mechanical Engineering, The Hong Kong Polytechnic University, Hong Kong, China c Baosteel Technology Centre, Shanghai 201900, China d Department of Manufacturing Engineering and Engineering Management, City University of Hong Kong, Hong Kong, China b
a r t i c l e
i n f o
Article history: Received 6 January 2011 Received in revised form 22 May 2011 Accepted 25 May 2011 Keywords: Nano materials Multilayered structure Plastic deformation Mechanical properties
a b s t r a c t The multifunctional diversities existing in nature provide clues to speculate the structure–property–function relationships. A hierarchically structured steel is designed by using principles derived from nature and fabricated in situ by a one-step method of surface mechanical attrition treatment (SMAT). The microstructure of the processed steel is characterized by multilayered structure with hard nanocrystalline surface and compliant inner-layer, in particular with a smooth mechanical gradient induced by dual-phase constituents and multiscale grain size distribution. The hierarchically structured steel exhibits simultaneously high stiff, strong and large ductility, which originate from the joint deformation mechanisms of distinct reinforcing layers. The four layers present their own unique deformation mechanisms, including second-phase hardening, transformation induced plasticity and twin strengthening. The unique spatial form of gradation can release stress concentration and improve energy-dissipation leading to exceptional mechanical properties compared with the uniform materials. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Steels are used far more extensively than any other metals, polymers, or ceramics as high-performance structural materials [1]. Now, the most famous iron and steel enterprises are facing the challenges to develop novel steels with exceptional strength and ductility driven by the requests of the automotive and aeronautic industries for safety and lightweight. The strengthening technologies of the conventional steels are involved in addition of alloying elements, fabrication of multiple phase structures [2–5], and refinement of grain size down to nanocrystalline (NC) or ultra-fined grain (UFG) scale [6,7], etc. Unfortunately, the strength and ductility tend to be mutually exclusive, which means higher strength is generally accompanied by a lower plasticity. The deterioration of plasticity is accelerated with the decrease of grain size down to NC/UFG scale [8]. Nature is the origins in design of multifunctional materials, which can provide clues to explore the structure–property–function relationships [9]. Many biological materials, such as bone [10,11], nacre [12], or scale [13], exhibit stiffness as well as high toughness resulting from their lamellar structures to combine the multiscale components with multiple
∗ Corresponding author at: Department of Manufacturing Engineering and Engineering Management, City University of Hong Kong, Hong Kong, China. E-mail address: [email protected] (J. Lu). 0254-0584/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.matchemphys.2011.05.068
functions. For example, the armour of fish scale consists of four different reinforcing layers with decreased gradation of modulus and hardness from the outer to inner layer. The inner layer is soft, while the high strength is achieved by nanoscaled grains in surface layer with a certain thickness [14]. Gradation in microstructure is the most important common feature in biological structures, which exhibits more damage-resistant than the monolithic counterpart [15]. Such layered biological structures have served as a design methodology applied in inorganic reinforced polymer-matrix artificial composites [16]. Although the notion of mimicking natural structures in man-made multilayered composites has generated enormous interest [17], the complex hierarchical designs in biological composites are extremely difficult to replicate synthetically, especially in covering so many dimensions (from nano to micro). Additionally, the synthetic technologies are generally complex, interface flaw-induced, size-dependent, and difficult for engineering applications, which severely hinder the implementation of hierarchical designs. We applied some of the structural concepts found in biological materials to design and fabricate hierarchical steel. A multilayered structure composed of diverse microstructures spanning over various length scales was deliberately chosen to implement design strategies in terms of tailoring strength and toughness. The hierarchically structured (HiSt) steel is assembled from nanometer scale to micrometer scale with functional gradient variation without interface in one-step by using an advanced technology of surface mechanical attrition treatment (SMAT) [18,19].
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304 SS sheets (100 mm × 90 mm × 1 mm) were treated by SMAT. The SMAT was performed by the optimized parameters, as given in Table 1. Both sides of AISI 304 SS sheets were treated under the same conditions. Since the SMAT is performed in the vacuum and at room temperature, the medium contaminations by the element diffusion of GCr15 balls are very limited. Thus, the contamination influence on the deformation mechanism and the properties of the HiSt steel is very small. The samples for the tensile tests were cut into dog-bone-shaped specimens with a gauge length of 30 mm and a cross section of 6 mm × 1 mm. Tensile tests were performed at room temperature at a strain rate of 6.7 × 10−4 s−1 using a MTS Alliance RT/50 Materials Testing System. Four specimens were used to obtain consistent stress–strain curves. Vickers microhardness was measured by applying a load of 50 g for 15 s and taking the average of six separate measurements on each depth. A Philips Xpert X-ray diffractometer (XRD) with Cu K␣ radiation was used to determine the phase constituent and estimate the phase content. The XRD tests were carried out along the cross-sectional direction, in these tests, the samples were carefully prepared by mechanical polishing from the surface to reach the targeted depths. The volume fraction of ˛ martensite phases was estimated from the peak integrated intensities Ih,k,l after background subtraction, as introduced in Ref. [20]. The average amount of ˛ martensite phases is calculated from two separate measurements on each depth. Transmission electron microscopy (TEM) observations were made with a JEM 2010 transmission electron microscope with an operating voltage of 200 kV. The plane-view TEM foils were ion-thinned at low temperature. The fracture surfaces of the specimens were investigated using a HITACHI S-4200 field emission scanning electron microscope (SEM).
3. Results and discussion Fig. 1. The schematic principle of SMAT technology.
3.1. Hierarchical microstructure 2. Experimental 2.1. SMAT technology SMAT is a newly developed technique, which can refine the coarse grain (CG) structure progressively down to the nanometer regime by means of repeated multidirectional impacts with high energy onto the sample surface [18]. The schematic principle of SMAT technology is shown in Fig. 1. The requirements of specific structures and properties shall be realized by manipulating the strain and strain rate at different depths from the surface. The resulted strain and strain rate at different depths from the surface are the driven force to refine the grain size into nano and submicrometer scale, simultaneously associated with phase transformation. The processed steel with dimensions of semi-industrial experiment can be fabricated by step moving manner. 2.2. Experimental procedure The chemical composition of AISI 304 austenite stainless steel (SS) is 17.6Cr, 8.3Ni, 0.04C, 0.19Si, 1.36Mn, 0.006S, 0.020P and balance Fe (all in mass%). The
The hierarchical structures originate from three critical aspects: (1) layered structure with hard surface and compliant interior; (2) hierarchical components spanning over various length scales; (3) most importantly, mechanical gradient between each successive layers. Guided by these strategies, the microstructure of HiSt steel is characterized by a symmetrically layered structure, which consists of four different reinforcing layers, as shown in Fig. 2 (from outer to inner layer): NC layer, dual phase layer, twinning layer and CG layer, respectively. The outermost NC layer provides superior stiffness and hardness by using hard ˛ martensite phase. The design origin of dual phase layer is to provide the important mechanical transition, which consists of hard ˛ martensite and soft austenite phase, resembling the biological structure of strong inorganic platelets embedded in a soft, ductile organic matrix. The third twinning layer is composed of multiscaled twins. The underneath softer, more compliant CG layer offers overall ductility. The strength ( y ) of multilayered composites can be estimated from the volume fraction (Vi ) and the tensile strength ( i ) of different layers, in terms of the “rule-of-mixture” [21,22]: y =
m
Vi i
(1)
i=1 m
Vi = 1
(2)
i=1
Fig. 2. Bioinspired design principle of HiSt steel. The outermost surface is composed of nano ˛ martensite phase to form hardest surface layer. The next is a transitional layer, consisted of dual phase structure with ˛ martensite phase and austentite phase in UFG scale. The UFG and ˛ martensite phase can further increase strength, and UFG austentite phase can deform by transformation-induced plasticity. The third layer is a twin layer, strengthening by twin boundaries. The inner layer is a complaint coarse grained (CG) layer to improve plasticity. From outer to inner, the grain size increases from nanometer to micrometer and the volume fraction of ˛ martensite phase decreases to establish a dual graded structure.
Thus, the higher strength of the HiSt steel shall be achieved by reducing one dimension of building blocks to the nanoscale, and increasing the volume fraction of hard layer. However, the ductility of the multilayered composite is contradictory to the strength, i.e. the stronger materials exhibit more brittle. Two critical designs are introduced to improve the ductility by additional deformation mechanisms, i.e. special deformation mechanisms of each layers and the gradation between the layers. Therefore, a dual gradient structure of grain size and phase constituent distribution is designed to induce mechanical gradient, mimicking the gradation in biological materials. Integrated the relationship of strength and ductility, the thickness of the four layers is selected to be (from outer to inner) 5–20 m, 60–200 m, 100–300 m, and 100–200 m, respectively. The above hierarchical structures are
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Fig. 3. Representative TEM images and corresponding SAED patterns of HiSt steel at four layers with depth of (a) 5 m, (b) 100 m, (c) 250 m, (d) 400 m, respectively. The lower inset of (a) is the grain size distribution.
difficult to be assembled in nanoscale and without layer interfaces by using bottom-up methods, but a newly developed technique of SMAT can prepare this complex composite in semi-industrial experiment by one-step manner. Fig. 3 presents the representative transmission electron microscopy (TEM) images and corresponding selected area electron diffraction (SAED) patterns of different layers of the AISI 304 SS sheets after treated by SMAT. The outer layer at 5 m depth from surface is composed of ˛ martensite phase in body centered cubic (BCC), where the grain size distribution is in the range of 2–100 nm and the average grain size is about 10 nm, as shown in Fig. 3a. The microstructures observed at 100 m depth from surface, exhibit a dual phase structure of ˛ martensite phase and austenite phase in face centered cubic (FCC) both with grain size from several nanometers to hundreds of nanometers, as shown in Fig. 3b. The corresponding SAED pattern exhibits mixed rings of ˛ and phase, indicating by the grey rings and orange rings in Fig. 3b, respectively. Fig. 3c shows the microstructure of twinning layer at 250 m depth from surface, which consists of twins with spacing of several hundreds of nanometers. The center layer at 400 m depth from surface is characterized by multiscale dislocation cells (DCs), as shown in Fig. 3d. According to the TEM observations, the HiSt steel is composed of micron-sized grain/DC (1–20 m), submicron grain/DC (0.5–1 m) and NC/UFG (0–0.5 m) with volume fraction of 40%, 30% and 30%, respectively, as given in Fig. 4. Furthermore, the
volume fraction of ˛ martensite phase decreases sharply from 95% in the surface of NC layer to about 6% in the center of CG layer. The grain size and phase constituent both show gradient distribution.
Fig. 4. Dual gradient distribution of the grain size and volume fraction of ˛ martensite phase with depth in the HiSt steel.
Treatment temperature (◦ C)
25 40
1099
98 ∼0.5 250 GGr15 50
8
Vacuum degree (Pa) Quantity of balls (pcs) Ball material
Diameter of ball (mm)
Impact velocity (m s−1 )
80 × 80
Fig. 5. Gradient variation of microhardness with depth of the HiSt steel at crosssection at different tensile strain.
Vibrating frequency (Hz)
Table 1 Processing parameters of SMAT.
Impact area (mm2 )
Treatment time (min)
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3.2. Mechanical properties The dual gradient structure results in a gradient distribution of hardness, and the smaller grain size and larger volume fraction of ˛ martensite phase make higher hardness. The microhardness, H, decreases with distance from outer to inner, presenting a parabolic shape (in Fig. 5). The highest average microhardness (∼500 HV) of the outer layer is attributed to the nanoscaled ˛ martensite phase. The dual phase layer has a larger grain size and smaller amount of ˛ martensite phase, therefore, a steep variation of microhardness is observed. The combination of high tensile strength and high ductility is the most impressive feature of HiSt steel. The engineering stress–strain curve of HiSt steel shows high yield stress ( 0.2 ) of 600 MPa, more than two times that of the original steel (270 MPa), and ultimate tensile stress ( b ) of 875 MPa, as shown in Fig. 6a. The total elongation to failure (εb ) and uniform elongation (εu ) of the HiSt steel also maintain quite high level of 51% and 45%, respectively. Interestingly, the hierarchical steel significantly deviates from the general strength–ductility rule of the conventional steels (interstitial free (IF) steel, mild steel, C–Mn (CMn) steel and high strength low alloy (HSLA) steel) and advanced high strength steels (AHSS), as given in Fig. 6b. The advanced high strength steels, such as dual-phase (DP) steel, transformation-induced plasticity (TRIP) steel, and complexphase (CP) steel are developed far more extensively driving by the requests of safety and weight saving. Compared with the other steels at the same strength level, for example, the TRIP steel with yield strength of 600 MPa (TRIP 600), the elongation to fracture is about 20%, but that of the HiSt steel is as high as 51%, which indicates the design strategies based on bioinspired concepts are fascinating. 3.3. The strength–ductility mechanisms The multilayered structure and gradient within and between layers work collectively to provide enhanced mechanical performance. The present steel consists of four reinforcing layers, each of them deforms in different mechanisms. The NC structure is selected by the super hard ˛ martensite phase particularly with grain size in several nanometers. The NC materials often exhibit superior strength, but brittleness. In order to balance the strength–ductility relationship of NC materials, the thickness of NC layer with grain size less than several tens of nanometers in HiSt steel is about 20 m, similar to that found in fish scale, which shows a highly stable deformation manner by a circumferential cracking [13,15]. The dual phase layer is an important transitional layer, which embod-
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Fig. 6. Comparisons of mechanical properties between the HiSt steel and other steels. (a) Engineering stress–strain curves of the HiSt steel and original AISI 304 SS. The inserts are the dimension of the tensile sample and the picture after tensile, respectively. (b) Strength–ductility relationship of different structured steels.
ies: (1) to delocalize the deformation of brittle NC layer; (2) to supply the required strength by NC/UFG grain size and hard ˛ phase; (3) to generate TRIP mechanism. The hard martensite phase can be used as precipitated phase to further strengthen, and the soft austenite phase can transform to ˛ martensite phase to induce TRIP effect [23–25]. The third layer, as an adapter with close value of microhardness to the fourth layer, presents twin strengthening. Twin boundaries (TBes) can act as barriers to dislocation motion, thus, provide a strengthening mechanism similar to grain boundaries. Moreover, Suresh et al. [26] had constructed a three-dimensional model to point out that the highly coherent twin structure can not only block dislocation action at the TBes, but also avoid the lost of ductility characteristic of nanocrystalline. Thereby, the twinning layer can not only enhance the strength, but also improve the ductility [5,27,28]. The underlying softer, more compliant CG layer dissipates energy via plasticity. The multiscaled dislocation cells in the fourth layer transform to martensite, but the refined dislocation cells can primarily transform to the martensite with thin lath structure. Therefore, each layer exhibits special deformation behaviour to induce excellent combination of strength and ductility, which is not the case for homogenous systems. Besides the important multilayered structure, the junctions between layers are clearly ‘functionally gradient’, which means that they possess gradual changes in properties, as well as no interfaces between multilayers. The gradient structure can promote
an effective transitional region to reduce stress concentration at the intersection between interfaces. If the hardest NC layer transfers load directly through the NC-CG junction, such as the case of NC coating, this may lead to magnify interfacial tensile shear stress, promoting brittle tendency [29]. The junction of NC layer and dual phase layer possesses a steep gradation relative to their neighbouring twinning layer and CG layer. During tensile deformation, compliant CG layer, twinning layer and dual phase layer deform jointly with graded increase of hardness by the introduced strengthening mechanisms, until their magnitudes are close to that of the original NC layer, as shown in Fig. 5. The hardness variations of HiSt steel exhibit obviously gradient tendency, indicating an important smoothing role of gradient in deformation. A sequence of strengthening mechanisms is suggested based on experimental observations, which can be subdivided into three stages, as illustrated in Fig. 7. In the early stage 1, the shear stress induced by the mismatched strengths of the multilayers can trigger the tangential motions of nanocrytalline, such as GB sliding and grain rotation [30–32]. The tangential motions can result in a sheared zone [32]. The well-defined shear bands are formed at surface, characterized by thin and sheet-like regions, where high plastic strain is localized, as indicated by the white arrows in Fig. 7a. The shear localization induced by GB sliding is a dominant deformation mechanism in nanocrystalline materials, especially for the grain size on the order of 20 nm and lower [31]. Coupling of the boundary motions and grain rotations can lead to grain growth, driven by the reduction of the surface free energy [32]. The TEM observations reveal larger grain size and enlarged grains along the shear localization regions at 20% plastic strain, where the average gain size is increased to about 100 nm with a much wider grain size distribution, as presented in Fig. 8a. The growth of grain size directly results in the decrease of surface hardness, which is in agreement with the result of microhardness variation in Fig. 5. These larger grains are thought to provide larger mean free paths for dislocations and sink of these dislocations because the grain boundaries are the primary sources for dislocation interactions (cross slip, reaction, etc.) [31]. The changes of dislocation density estimated from HRTEM images confirm the dislocation annihilation, where dislocation density in the surface and subsurface is high up to 35–40 × 1016 m−2 , but decreases to about 10–15 × 1016 m−2 at 20% strain, as shown in Fig. 8b. The experimental observations, such as shear bands, grain growth, and reduced dislocation density, show the complex deformations of NC by GB sliding and grain rotation. Meantime, the dislocation-mediated plasticity is dominant in twinning and CG layer under applied stress. Compared Fig. 7b with Fig. 3c, the high density of tangled dislocations is accumulated at twin boundaries. The obvious increases of dislocation density of twinning and CG layer are observed by HRTEM images, as given in Fig. 8b. At stage 2, when the plastic deformation increases, the martensite transformation is induced by plastic strain. For the specimen deformed at 30% strain, the ˛ martensite phase with random orientation in UFG layer and lath shape in CG layer are observed at depth of 50 m and 300 m, respectively, shown in Fig. 7c. The volume fraction of ˛ martensite phase in function of plastic strain varies severely, indicating an obvious strain-induced martensite transformation, as shown in Fig. 8c and d. The amount of marteniste transformation agrees with the change tendency of hardness distribution in Fig. 5. The austenite phase of different sizes transforms to martensite phase to hold out the uniform straining. At stage 3, it is interesting to find that the crack initiation presents in the inner layer but not in brittle NC layer. The side-view from the fracture point (in Fig. 7g) presents many cracks, which initiate and propagate at the interlayer. In addition, the tensile fracture surface is embodied with many microcracks in the twinning/CG layers, which are not observed in the NC/UFG layer (in Fig. 7f and h).
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Fig. 7. A schematic illustration depicting the deformation mechanism of the HiSt steel. (a) SEM micrograph of NC layer; (b) TEM image of twin layer at 20% strain; (c) and (d) TEM image of UFC and CG layer at 30% strain, respectively; (e–h) fracture morphologies after tensile test; (f) is the general view of the fracture; (e) is the magnification of transition zone of UFC/twin layer; (f) is the magnification of CG layer and (g) is the side view of tensile sample close to the fracture point.
In contrast to larger equiaxed dimples of CG layer, the transitional zone of UFG/twinning layer exhibits much smaller dimples elongated in opposite direction (in Fig. 7e). This distinct transition of grain size dependence on fracture surface suggests that the multiscale structure can effectively improve the plasticity of NC/UFG materials. It is interesting to find that the nano-sized grains in surface present a consecutive distribution during plastic deformation. However, the grain size of the surface nano-sized grains is increased from the original 10 nm to about 100 nm at 20% strain and further increased to 250 nm after tensile test, and the thickness is reduced
from about 20 to 8 m after tensile test. The growth of nano-sized grains and thinner of the nano layer are the results of the grain boundary sliding and grain rotation. Although a novel engineering material with high performance is achieved by implementing bioinspired concepts, there still has large space to be optimized in microstructures, such as the structure of soft layer, the layer thickness and gradient between layers. Further strengthening by grain refinement and the increased proportion of NC/UFG layer may induce the deterioration of ductility. Therefore, the increased proportion of twinning layer inevitably
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Fig. 8. Microstructures of HiSt steel of different strains. (a) TEM image of 5 m depth at 20% strain, the upper inset is the SAED pattern and the lower is grain size distribution; (b) dislocation density; (c) the volume fraction of ˛ martensite phase with strains and (d) XRD patterns after tensile test.
becomes an optimal selection for its comparable level of higher strength, and still providing substantial ductility. The improvement of strength and ductility can be anticipated by reduction of the coarse grained layer proportion and development of high density nanotwinned structure to meet the higher performance requirement. 4. Conclusions Based on the mechanical design strategies to mimic the structure of fish scale and seashell, a complex and multiscale material is designed and fabricated by using an advanced but simple processing of surface mechanical attrition treatment. The hierarchically structured steel is composed of a symmetrically layered structure, e.g. NC layer, dual phase layer, twinning layer and coarse grained layer from outer to inner, respectively. The junctions between layers present a gradient transition induced by a dual gradient structure of grain size and phase constituent distribution. Coupling multiscaled deformation behaviour over the nanoscale to microscale provides additional toughening mechanisms similar to those acting in natural materials. The hierarchical steel significantly deviates from the general strength–ductility rule of the conventional steels and advanced high strength steels. The impressive combination of strength and ductility of the hierarchical steel is obtained by a synergistic effect of multiple deformation mechanisms. These results indicate that microstructural features of hierarchical steel put bioinspired designs into practice, which should guide the synthesis of engineered structural materials.
Acknowledgements This work was supported by the Research Grants Council of the Hong Kong Special Administrative Region of China under the CityU8/CRF/08 project and by the research project (No. ITP/004/08NP) of the Hong Kong Innovation and Technology Commission (ITC). Financial support from the National Natural Science Foundation of China (Grant No. 51074107) is also acknowledged. This work is also supported by Shanghai Leading Academic Discipline (No. J51504). References [1] M. Militzer, Science 298 (2002) 975. [2] P. Podsiadlo, A.K. Kaushik, E.M. Arruda, A.M. Waas, B.S. Shim, J.D. Xu, H. Nandivada, B.G. Pumplin, J. Lahann, A. Ramamoorthy, N.A. Kotov, Science 318 (2007) 80. [3] K.T. Park, S.Y. Han, B.D. Ahn, D.H. Shin, Y.K. Lee, K.K. Um, Scripta Mater. 51 (2004) 909. [4] P.J. Jacques, Q. Furnémont, F. Lani, T. Pardoen, F. Delannay, Acta Mater. 55 (2007) 3681. [5] O. Bouaziz, S. Allain, C. Scott, Scripta Mater. 58 (2008) 484. [6] Y.H. Zhao, X.Z. Liao, S. Cheng, E. Ma, Y.T. Zhu, Adv. Mater. 18 (2006) 2280. [7] X.H. Chen, J. Lu, L. Lu, K. Lu, Scripta Mater. 52 (2005) 1039. [8] N. Tsuji, Y. Ito, Y. Saito, Y. Minamino, Scripta Mater. 47 (2002) 893. [9] C. Sanchez, H. Arribart, M. Madeleine, G. Guille, Nat. Mater. 4 (2005) 277. [10] R.K. Nalla, J.H. Kinney, R.O. Ritchie, Nat. Mater. 2 (2003) 164. [11] K. Tai, M. Dao, S. Suresh, A. Palazoglu, C. Ortiz, Nat. Mater. 6 (2007) 454. [12] S. Kamat, X. Su, R. Ballarini, M. Nassirou, A.H. Heuer, Nature 405 (2000) 1036. [13] B.J.F. Bruet, J. Song, M.C. Boyce, C. Ortiz, Nat. Mater. 7 (2008) 748. [14] K.D. Jandt, Nat. Mater. 7 (2008) 692. [15] K.J. Koester, J.W. Ager, R.O. Ritchie, Nat. Mater. 7 (2008) 672. [16] L.J. Bonderer, A.R. Studart, L.J. Gauckler, Science 319 (2008) 1069.
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