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Strengthening and hardening mechanisms of additively manufactured stainless steels: The role of cell sizes
Scripta Materialia 177 (2020) 17–21
Contents lists available at ScienceDirect
Scripta Materialia journal homepage: www.elsevier.com/locate/scriptamat
Strengthening and hardening mechanisms of additively manufactured stainless steels: The role of cell sizes Zan Li a, Bo He b,∗, Qiang Guo a,∗ a b
State Key Laboratory of Metal Matrix Composites, Shanghai Jiao Tong University, Shanghai, 200240, China Department of Materials Science & Engineering, Shanghai University of Engineering Science, Shanghai, 201620, China
a r t i c l e
i n f o
Article history: Received 30 July 2019 Revised 1 October 2019 Accepted 5 October 2019
Keywords: Additive manufacturing Stainless steel Deformation mechanism Cell size Micro-pillar compression
a b s t r a c t We designed micromechanical experiments to examine the deformation behaviors of austenitic stainless steels fabricated by additive manufacturing. Micro-pillars containing different dislocation cell sizes were produced from the bulk specimen, and were mechanically tested for direct assessment of cell size reliance of yield and plastic deformation processes. The results highlight the cruciality of dislocation density in determining the yield strength of additively manufactured stainless steels, challenging the previous viewpoint that the strength scales with the cell size. However, dislocation nucleation and hardening of these micro-pillars were found to be cell size-correlated, with easier dislocation nucleation for micro-pillars containing smaller cells. © 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
After more than thirty years of development, additive manufacturing (AM) has become an important manufacturing process owing to its unique design freedom and rapid build capability [1,2]. For metals, the fast and integrative processing feature of AM technology as compared to conventional metal processing routes is particularly attractive and thus has garnered a lot of research attention [3–5]. The additively manufactured metals, moreover, typically exhibit microstructures that are distinct from those of metals and alloys produced by traditional fabrication approaches such as casting and forging. For instance, in the case of the most widely studied and practically applied AM metal, 316L stainless steels (SS), the highly non-equilibrium process and significant solute partition during the AM procedure favor the formation of an ultrafine cellular dislocation structure in the austenite matrix [6–8], which render AM 316L SS an exceptional combination of strength and ductility. A fundamental understanding of the structure-property relationship in AM 316L SS is vital to tailor their mechanical performances and guide their further property improvement. One critical factor in determining the strength of AM 316L SS is the nature of their intrinsic cellular dislocation structure, which is usually decorated with elemental segregations and precipitates. Our previous analysis suggested the strengthening effect of these segregations/precipitates may be small [8]. Yet a Hall-Petch typed strengthening behavior, where the yield strength scales with the average cell size, is usually believed to account for the high yield ∗
Corresponding authors. E-mail addresses: [email protected] (B. He), [email protected] (Q. Guo).
https://doi.org/10.1016/j.scriptamat.2019.10.005 1359-6462/© 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
strength of AM 316L SS [8,9], and in situ transmission electron microscopy (TEM) experiments unraveled that the pre-existing dislocation cells have an ability to impede dislocation motion, resulting in high yield strength [10]. However, considering the fact that the dislocation cells in AM 316L SS are not traditionally deformation-induced dislocation walls, as they possess a lower energy configuration [8,11], the simple analogy between cellular dislocation walls in AM metals and conventional grain boundaries remains to be verified. Indeed, it was recently reported that, the enhanced strength of as-built AM 316L SS is mainly associated with the high dislocation densities of these dislocation cells rather than the cell size [12]. Such a controversy in strengthening mechanisms of AM 316L SS may stem from the heterogeneity in their inherent microstructures caused by different thermomechanical history during fabrication. Therefore, the role of particular cellular structures/configurations on the mechanical behavior should be pinpointed, in order to reveal the underlying deformation mechanisms of AM 316L SS. Here, we present a strategy that allows a site-specific study of the mechanical properties of AM 316L SS fabricated by laser powder-bed-fusion (L-PBF) technology. By virtue of an in situ micro-mechanical experiment, we directly assessed the detailed deformation processes as well as the mechanical behaviors of single crystalline AM 316L SS micro-pillars with different representative cell sizes. Our results challenge the conventional viewpoint that the yield strengths of AM 316L SS are primarily dominated by the cell size, and unravel the determining role of initial dislocation density on the high strength of AM 316L SS. Additionally, with
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Z. Li, B. He and Q. Guo / Scripta Materialia 177 (2020) 17–21 Table 1 Cell size, Schmid factor, and shear strength improvement summary of AM 316L SS micro-pillars at various selected regions. The increment in resolved shear flow stress was statistically calculated by considering the discrete strain events at strains higher than 0.03, where a relatively constant increment for each burst event was observed (Fig. 4b). The error bars correspond to the standard deviations (± SD) of the measurements.
Fig. 1. (a) Cross sectional SEM image of etched AM 316L SS. Fusion boundaries were indicated by white arrows. Regions (R-1 to R-6) containing various cell sizes were used for micro-pillar preparations, and were encircled by yellow dashed lines. (b)(g) are SEM images and statistically calculated cell size distributions (insets) for regions R1 to R6, respectively. Note that the cell size of R6 was estimated from the spacing of the elongated cellular structure, which is normally believed to be the characteristic length scale associated with the cell boundary-dislocation interactions. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
a closer inspection of the discrete strain events displayed in the stress-strain curves, we found that the dislocation nucleation and the strain hardening processes of AM 316L SS micro-pillars are cell size-correlated, with easier dislocation nucleation for micro-pillars containing smaller cells. The present study will further the understanding of the underlying mechanisms that govern the excellent mechanical responses of AM 316L SS. The 316L SS rectangular plate was fabricated in an argon environment by a commercial BLT S210 PBF machine equipped with a 400 W fiber laser. The initial plasma-atomized 316L SS powders had a particle size ranging from 20 to 80 μm, with a mean value of ∼30 μm. We performed a series of control experiments to optimize the processing parameters for 316L SS, where the details had been reported in our previous work [13]. A cubic sample was cut from the as-printed plate by electro-discharge machining (EDM). Its surface was polished with 360–40 0 0 grit metallographic silicon carbide papers, followed by polishing with 3 μm and 1 μm diamond suspensions. A modified Carpenter’s reagent was used for etching to enhance the image contrast for scanning electron microscopic (SEM, FEI Scios) observations [13]. Fig. 1a displays the representative SEM image of the etched surface, where fusion boundaries and the ultrafine cellular structure can be clearly seen. We chose six representative regions (R1 – R6) with cell size ranging from 330 nm to 590 nm, as shown in Fig. 1b-1g, for the subsequent micro-pillar preparation. In particular, two micro-pillars with di-
Pillars
Cell size (nm)
R-1 R-2 R-3 R-4 R-5 R-6
435 438 336 451 502 590
± ± ± ± ± ±
63 70 46 76 77 118
Schmid factor
0.492 0.472 0.496 0.495 0.494 0.491
± ± ± ±
0.006 0.002 0.005 0.001
± 0.001
Increment in resolved shear flow stress (MPa) 11.5 ± 2.2 10.8 ± 1.0 9.6 ± 2.5 11.6 ± 0.7 11.9 ± 1.7 12.7 ± 2.4
ameter and height of ∼4.5 μm and ∼9 μm respectively, were fabricated from each region by a focused ion beam (FIB, FEI Scios) system (Fig. 2a). The taper angle of these micro-pillars was controlled within 2° in order to reduce the inhomogeneous deformation associated with the variations in gauge cross-section areas [14]. The actual cell size in pillars fabricated from regions 1 to 6 was measured to be 445 ± 36 nm, 430 ± 21 nm, 345 ± 36 nm, 463 ± 46 nm, 492 ± 42 nm and 595 ± 75 nm, respectively (for details please refer to Fig. S1). We employed electron backscatter diffraction (EBSD, TEAMTM EBSD Analysis System, EDAX) to examine the grain orientation of each micro-pillar and found that the two micro-pillars that contained identical cell sizes had very similar grain orientations (Fig. 2b). The Schmid factor for each micro-pillar was estimated by averaging the Schmid factors of five randomly selected positions that were ∼ 1 μm distance away from the milling trench, as indicated in Fig. 2b. Note that there are local misorientations inside one individual grain, which may be induced by the strong thermal gradient during the rapid solidification and extensive thermomechanical cycles during the layer-by-layer building processes [8]. The average cell size and Schmid factor of each region are documented in Table 1, from which a small Schmid factor deviation can be observed, in spite of the existence of local misorientation gradients. In situ compression tests of the micro-pillars were conducted in SEM using a cradle-based nanoindentation system (NanoFlip, Nanomechanics Inc.), with an 8 μm-diameter conductive diamond flat punch tip. The compression tests were carried out at room temperature under the displacement controlled mode to a maximum strain of ∼ 0.1, at a nominal strain rate of 5 × 10−4 s−1 . Fig. 3a shows the representative compressive engineering stress vs. strain responses of micro-pillars fabricated from different regions. Because of the intermittent nature of the compressive flow and the difficulty in accurately defining the onset of plastic deformation, the yield strengths of the micro-pillars were identified by the stress at which the first strain jump (burst) occurred, using a strain jump of ε = 0.002 as the threshold [14]. Surprisingly, as revealed in Fig. 3a, the yield strengths of micro-pillars possessing various cell sizes were found to be fairly independent of the cell size and no salient trend was found. This can be better appreciated by plotting the yield strength versus cell size figure (Fig. 3b), where the Hall-Petch typed dependence of yield strength on the inverse square root of cell size (d−1/2 ) was not observed [8,9]. We further considered the Schmid factor for the operating slip system in each micro-pillar, where the calculated resolved shear yield strength (τ y ) versus d−1/2 (Fig. 3c) also displayed a weak (if any) linear correlation (dashed line; coefficient of determination for the best fit, R = 0.32). The fitted slope of the τ y vs. d−1/2 data (k) in metals is generally regarded as a measure of dislocation pinning capacity at boundaries [15], and the fitted k ≈ 1365 MPa nm1/2 from Fig. 3c was substantially lower than that of strengthening
Z. Li, B. He and Q. Guo / Scripta Materialia 177 (2020) 17–21
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Fig. 2. (a) SEM image of the as-milled micro-pillars with the selected regions (R1-R6) superimposed. The inset shows the enlarged view of a couple of representative micropillars for compression test. A rectangular trench, with dimension of 30 μm × 10 μm × 5 μm (length × width × depth) was milled next to the pillars to ensure a clear view of the pillars during mechanical testing. The tilt anlge was 15° during the observation. (b) The corrsponding EBSD inverse-pole figure map superimposing high-angle (grain misorientation > 10°, black line) and low-angle (grain misorientation 2°−10°, white line) grain boundaries. The row of 5 crosses illustrates the locations for Schmid factor measurement of a micro-pillar in R-1.
Fig. 3. (a) Representative stress-strain curves of micro-pillars containing various cell sizes. (b) Yield strength versus cell size of micro-pillars, showing their weak (if any) correlation. (c) Resolved shear yield strength versus cell size and a best linear fit of the data (dashed line). For comparison, the Hall-Petch fitting plot of bulk austenitic 316L SS is also shown [8]. (d) A plot of resolved shear yield strength as a function of the square root of dislocation cell volume fraction (Vf 1/2 ). The dashed line is the best linear fit of the data.
mechanism by grain boundaries measured from bulk austenitic 316L SS fabricated by conventional methods (k ≈ 2674 MPa nm1/2 [8]). This indicated that the dislocation cells in AM 316L SS were relatively weak dislocation barriers and cannot be treated equivalent to typical grain boundaries in terms of strengthening. On the contrary, by calculating the volume fraction (Vf ) of dislocation cells
in each region (Fig. 1b–g) (see Supplementary text. S1, Fig. S2 and S3 for details of the measurement), we found that τ y exhibits a well-fitted linear correlation with Vf 1/2 (R = 0.91) (Fig. 3d), thus obeying the well-known Taylor-strengthening power law for forest dislocations [16]. That being said, in contradiction to previous hypothesis of a strong dependence of yield strengths on cell sizes, we
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Z. Li, B. He and Q. Guo / Scripta Materialia 177 (2020) 17–21
Fig. 4. (a) Statistical distribution of displacement (normalized by the Burgers vector, b) of a discrete strain event versus the resolved shear flow stress at which it occurred, for micro-pillars containing various cell sizes. The general ‘higher flow stress-larger displacement’ is observed, in spite of some relatively large scattered data points (e.g., a displacement of 1150 b at the resolved shear stress of 297 MPa is included for pillars in R2). (b) A plot of shear flow strength improvement of a burst versus the strain value at which it occurs for micro-pillars with various cell sizes. The general trend is depicted by an arrow.
demonstrated that it is likely to be the dislocation density, rather than the cell size, that dominated the yield strength of AM 316L SS. This proposal is in line with a recent study of heat-treated AM 316L SS, where the hardness of the sample was significantly reduced after dislocation recovery, although the dislocation cell size stayed unchanged [17]. It is worth mentioning here that, the yield strengths of micropillars (between 550 and 620 MPa) were found to be comparable with those of the corresponding bulk samples (552–635 MPa [13]). As a matter of fact, this lack of specimen size effect was expected, because the sizes of typical dislocation obstruction features (i.e., dislocation cells) were approximately one order of magnitude smaller than the pillar diameter. This much smaller internal microstructural length scale as compared with the external one (i.e., pillar diameter) indicates that the deformation of these micropillars was governed by dislocation-internal boundary (cell boundary) interactions instead of free surface effects [18,19]. Specifically, the numerous dislocation cells in the pillars would intrinsically act as dislocation trapping and/or nucleation sites [10], making it less likely for dislocation source-truncation or exhaustion hardening [14,20]. However, unlike macroscopic samples characterized by smooth plastic flow, the AM 316L SS micro-pillars showed lots of strain jumps during the hardening processes, which were likely to associate with internal dislocation avalanches. It had been well demonstrated that dislocation activity during the avalanche is usually dominated by a single-slip system [14,21], and indeed, as revealed in Fig. S4, the post-deformation micro-pillars were shown to contain numerous single slip bands. These observations suggest that, although dislocation cells may trap dislocations [10], such trapping activity may not be strong enough to stimulate a massive mobile dislocation cross-slip, thus leading to a discrete plastic flow in AM 316L SS micro-pillars. This agrees well with our results of reduced k obtained in Fig. 3c and the previous experimental observations [10], where dislocation cell walls were found to have the ability to hinder but not entirely block the dislocation motion. To further understand the dislocation activities in AM 316L SS after yielding (i.e., the hardening process) so that the role of cell size on plastic deformation can be more precisely elucidated, we made an in-depth study of the plastic flow intermittency. In particular, it is known that the statistical attributes of strain bursts are believed to closely correlate with the intrinsic microstructural features, e.g., the density and morphology of pre-existing dislocations
[21,22]. As shown in Fig. 4a, a plot of the displacement magnitude for all strain bursts of each micro-pillar (normalized by Burgers vector of austenitic 316L SS, b = 0.257 nm) versus the corresponding resolved shear flow stress revealed a general trend that the discrete displacement increases with increasing resolved shear flow stress. Furthermore, the distribution of these burst events didn’t show a cell size dependence, and we highlighted this in Fig. 4a by comparing the data distributions of the two extremes, i.e., micropillars containing the smallest (red triangle symbol) and largest (blue diamond symbol) cell sizes. This observation again suggested that the cell boundaries in this study were not strong dislocation obstacles as low- and high-angle grain boundaries, which would usually cause a distinct correlation between strain burst size and characteristic dislocation obstruction length scales (grain sizes) [23–25]. In addition, we calculated the increments in resolved shear flow stress after each strain burst event as a function of the engineering strain at which the burst occurred (Fig. 4b). Two important findings can be immediately drawn. First, the strength improvements for early burst events were larger and they gradually attenuated with increasing strain for almost all tested samples, eventually reaching a nearly constant value after an engineering strain of ∼ 0.03 was achieved. We propose that this observation stemmed from the dislocation cells-mobile dislocation interactions during plastic deformation. Specifically, as evidenced by the lack of local misorientations across the cell boundaries and their unique characteristic of thermal tolerance, dislocation cells in the as-printed 316L SS are energetically stable [8,11]. Upon straining, mobile dislocations were trapped and piled up at these cell boundaries, which would cause stress concentrations and favor dislocation nucleation/emission from these boundaries [10]. The strong trapping activities of dislocation cells were also demonstrated in our previous post-mortem TEM analysis [8,13]. Such dislocation trapping induced stress concentrations would subsequently give rise to a reduced stress increment that was needed to compensate for the dislocation avalanches and sustain further plastic deformation, leading to the evolution of shear flow stress increment as shown in Fig. 4b. Here we consider that the stress increment arose at each strain burst primarily from the resistance to dislocation generation at the cell boundaries, rather than the resistance to dislocation nucleation inside the cells (e.g., from the Frank-Read sources [26]). This is because, the stress
Z. Li, B. He and Q. Guo / Scripta Materialia 177 (2020) 17–21
increment required for the operation of normal Frank–Read sources comes from the back stress rendered by pile-up dislocations [27], which can be expressed as τ = GNb/[2π (1−ν )L]. By assuming that the separation L between a source and dislocation pile-up to be about half the cell size and taking shear modulus G = 75 GPa, and a Poisson’s ratio of ν ∼ 0.3, the calculated shear stress from just two piled dislocation (N = 2) would range from 29 to 58 MPa for samples with a cell size of 30 0 nm–60 0 nm studied here. This value is substantially higher than the experimental measured ones (Fig. 4b), suggesting that the dislocation multiplication in the cell interiors was unlikely the operating mechanism that accounted for the stress increments on the stress vs. strain response. Second, as revealed by the statistically calculated shear stress improvement (Table 1), micro-pillars with smaller cell size generally exhibited lower stress increments after the discrete strain bursts. We believe that this cell size correlated dislocation nucleation/hardening responses arose from the differences in the density of piled-up dislocations at the cell boundaries for a specific shear strain, γ . It is indicated in Orowan equation that (γ = ρ × b × l, where ρ is mobile dislocation density and l is dislocation mean free path) [16], to reach the same γ , higher density of dislocations have to be multiplied for smaller dislocation trapping distances; that is, stress concentrations along cell boundaries would be more intensive for smaller dislocation cells, thereby leading to reduced stress improvement that required for dislocation multiplication. However, a quantitative analysis of the mechanisms deserves further investigations and is out of the scope of the present study. In summary, our inspection of the burst evolution showed that, although the yielding behavior of AM 316L SS was mainly dominated by dislocation densities, the subsequent dislocation nucleation and hardening processes of the micro-pillars were cell size-correlated. Note that the influence of twinning effect on hardening behavior was not considered here, because we didn’t observe extensive twinning activities in the deformed pillars (Fig. S5). We employed micro-pillar compression experiments to probe the cell size reliance of the yield and plastic deformation processes of AM 316L SS. The greatest advantage of the micromechanical approach as compared to traditional bulk mechanical characterizations is that it enables a site-specific investigation, which is particularly significant for AM metals that intrinsically contain heterogeneous microstructures. Our findings provide insight into the mechanisms that govern the superior mechanical properties of AM 316L SS. The experimental strategy adopted here also provides a universal pathway to advance the understanding of the microstructureproperty relationship for other rapidly developing AM metallic materials that have highly heterogeneous microstructures. The authors would like to acknowledge the financial support from the National Natural Science Foundation of China (Grant Nos. 51801120, 51771111), the Ministry of Science and Technology
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of China (Nos. 2016YFE0130200, 2017YFB0703100), Shanghai Science and Technology Committee Innovation Grant (Grant Nos. 17JC140 060 0, 17JC140 0603, 1752071240 0), and Distinguished Professor Program of Shanghai University of Engineering Science. Declaration of Competing Interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Supplementary materials Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.scriptamat.2019.10. 005. References [1] D.W.Rosen Gibson, B. Stucker, Additive Manufacturing Technologies, Springer, New York, 2014. [2] S.H. Huang, P. Liu, A. Mokasdar, L. Hou, Int. J. Adv. Manuf. Tech. 67 (2013) 1191. [3] J.J. Lewandowski, M. Seifi, Annu. Rev. Mater. Res. 46 (2016) 151. [4] D. Herzog, V. Seyda, E. Wycisk, C. Emmelmann, Acta Mater. 117 (2016) 371. [5] W.E. Frazier, J. Mater. Eng. Perform. 23 (2014) 1917. [6] Z. Wang, A.M. Beese, Acta Mater. 131 (2017) 410. [7] Z. Sun, X. Tan, S.B. Tor, C.K. Chua, NPG Asia Mater. 10 (2018) 127. [8] Y.M. Wang, T. Voisin, J.T. McKeown, J. Ye, N.P. Calta, Z. Li, Z. Zeng, Y. Zhang, W. Chen, T.T. Roehling, R.T. Ott, M.K. Santala, P.J. Depond, M.J. Mattews, A.V. Hamza, T. Zhu, Nat. Mater. 17 (2018) 63. [9] Y. Zhong, L.F. Liu, S. Wikman, D.Q. Cui, Z.J. Shen, J. Nucl. Mater. 470 (2016) 170. [10] L. Liu, Q. Ding, Y. Zhong, J. Zou, J. Wu, Y.L. Chiu, J. Li, Z. Zhang, Q. Yu, Z. Shen, Mater. Today. 21 (2018) 354. [11] M. Shamsujjoha, S.R. Agnew, J.M. Fitz-Gerald, W.R. Moore, T.A. Newman, Metall. Mater. Trans. A. 49 (2018) 3011. [12] T.R. Smith, J.D. Sugar, C.S. Marchi, J.M. Schoenung, Acta Mater. 164 (2019) 728. [13] Z. Li, T. Voisin, J.T. McKeown, J. Ye, T. Braun, C. Kamath, W.E. King, Y.M. Wang, Int. J. Plasticity. 120 (2019) 395. [14] M.D. Uchic, P.A. Shade, D.M. Dimiduk, Annu. Rev. Mater. Res. 39 (2009) 361. [15] Z. Li, L. Zhao, Q. Guo, Z. Li, G. Fan, C. Guo, D. Zhang, Scripta Mater. 131 (2017) 67. [16] M.A. Meyers, Mechanical Metallurgy Principles and Applications, Prentice-Hall, New Jersey, 1984. [17] P. Krakhmalev, I. Yadroitsava, G. Fredriksson, I. Yadroitsev, S. Afr. J. Ind. Eng. 28 (2017) 12. [18] Y. Hu, Q. Guo, L. Zhao, Z. Li, G. Fan, Z. Li, D.B. Xiong, Y. Su, D. Zhang, Scr. Mater. 146 (2018) 236. [19] D.C. Jang, J.R. Greer, Scr. Mater. 64 (2011) 77. [20] C. Chisholm, H. Bei, M.B. Lowry, J. Oh, S.A. Syed Asif, O.L. Warren, Z.W. Shan, E.P. George, A.M. Minor, Acta Mater. 60 (2012) 2258. [21] F.F. Csikor, C. Motz, D. Weygand, M. Zaiser, S. Zapperi, Science 318 (2007) 251. [22] K.Y. Xie, S. Shrestha, Y. Cao, P.J. Felfer, Y. Wang, X. Liao, J.M. Cairney, S.P. Ringer, Acta Mater. 61 (2013) 439. [23] T. Hu, L. Jiang, H. Yang, K. Ma, T.D. Topping, J. Yee, M. Li, A.K. Mukherjee, J.M. Schoenung, E.J. Lavernia, Acta Mater. 94 (2015) 46. [24] J.Y. Zhang, G. Liu, J. Sun, Int. J. Plasticity. 50 (2013) 1. [25] P.J. Imrich, C. Kirchlechner, C. Motz, G. Dehm, Acta Mater. 73 (2014) 240. [26] J.J. Gilman, J. Appl. Phys. 30 (1959) 1584. [27] K.S. Ng, A.H.W. Ngan, Acta Mater. 56 (2008) 1712.
Contents lists available at ScienceDirect
Scripta Materialia journal homepage: www.elsevier.com/locate/scriptamat
Strengthening and hardening mechanisms of additively manufactured stainless steels: The role of cell sizes Zan Li a, Bo He b,∗, Qiang Guo a,∗ a b
State Key Laboratory of Metal Matrix Composites, Shanghai Jiao Tong University, Shanghai, 200240, China Department of Materials Science & Engineering, Shanghai University of Engineering Science, Shanghai, 201620, China
a r t i c l e
i n f o
Article history: Received 30 July 2019 Revised 1 October 2019 Accepted 5 October 2019
Keywords: Additive manufacturing Stainless steel Deformation mechanism Cell size Micro-pillar compression
a b s t r a c t We designed micromechanical experiments to examine the deformation behaviors of austenitic stainless steels fabricated by additive manufacturing. Micro-pillars containing different dislocation cell sizes were produced from the bulk specimen, and were mechanically tested for direct assessment of cell size reliance of yield and plastic deformation processes. The results highlight the cruciality of dislocation density in determining the yield strength of additively manufactured stainless steels, challenging the previous viewpoint that the strength scales with the cell size. However, dislocation nucleation and hardening of these micro-pillars were found to be cell size-correlated, with easier dislocation nucleation for micro-pillars containing smaller cells. © 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
After more than thirty years of development, additive manufacturing (AM) has become an important manufacturing process owing to its unique design freedom and rapid build capability [1,2]. For metals, the fast and integrative processing feature of AM technology as compared to conventional metal processing routes is particularly attractive and thus has garnered a lot of research attention [3–5]. The additively manufactured metals, moreover, typically exhibit microstructures that are distinct from those of metals and alloys produced by traditional fabrication approaches such as casting and forging. For instance, in the case of the most widely studied and practically applied AM metal, 316L stainless steels (SS), the highly non-equilibrium process and significant solute partition during the AM procedure favor the formation of an ultrafine cellular dislocation structure in the austenite matrix [6–8], which render AM 316L SS an exceptional combination of strength and ductility. A fundamental understanding of the structure-property relationship in AM 316L SS is vital to tailor their mechanical performances and guide their further property improvement. One critical factor in determining the strength of AM 316L SS is the nature of their intrinsic cellular dislocation structure, which is usually decorated with elemental segregations and precipitates. Our previous analysis suggested the strengthening effect of these segregations/precipitates may be small [8]. Yet a Hall-Petch typed strengthening behavior, where the yield strength scales with the average cell size, is usually believed to account for the high yield ∗
Corresponding authors. E-mail addresses: [email protected] (B. He), [email protected] (Q. Guo).
https://doi.org/10.1016/j.scriptamat.2019.10.005 1359-6462/© 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
strength of AM 316L SS [8,9], and in situ transmission electron microscopy (TEM) experiments unraveled that the pre-existing dislocation cells have an ability to impede dislocation motion, resulting in high yield strength [10]. However, considering the fact that the dislocation cells in AM 316L SS are not traditionally deformation-induced dislocation walls, as they possess a lower energy configuration [8,11], the simple analogy between cellular dislocation walls in AM metals and conventional grain boundaries remains to be verified. Indeed, it was recently reported that, the enhanced strength of as-built AM 316L SS is mainly associated with the high dislocation densities of these dislocation cells rather than the cell size [12]. Such a controversy in strengthening mechanisms of AM 316L SS may stem from the heterogeneity in their inherent microstructures caused by different thermomechanical history during fabrication. Therefore, the role of particular cellular structures/configurations on the mechanical behavior should be pinpointed, in order to reveal the underlying deformation mechanisms of AM 316L SS. Here, we present a strategy that allows a site-specific study of the mechanical properties of AM 316L SS fabricated by laser powder-bed-fusion (L-PBF) technology. By virtue of an in situ micro-mechanical experiment, we directly assessed the detailed deformation processes as well as the mechanical behaviors of single crystalline AM 316L SS micro-pillars with different representative cell sizes. Our results challenge the conventional viewpoint that the yield strengths of AM 316L SS are primarily dominated by the cell size, and unravel the determining role of initial dislocation density on the high strength of AM 316L SS. Additionally, with
18
Z. Li, B. He and Q. Guo / Scripta Materialia 177 (2020) 17–21 Table 1 Cell size, Schmid factor, and shear strength improvement summary of AM 316L SS micro-pillars at various selected regions. The increment in resolved shear flow stress was statistically calculated by considering the discrete strain events at strains higher than 0.03, where a relatively constant increment for each burst event was observed (Fig. 4b). The error bars correspond to the standard deviations (± SD) of the measurements.
Fig. 1. (a) Cross sectional SEM image of etched AM 316L SS. Fusion boundaries were indicated by white arrows. Regions (R-1 to R-6) containing various cell sizes were used for micro-pillar preparations, and were encircled by yellow dashed lines. (b)(g) are SEM images and statistically calculated cell size distributions (insets) for regions R1 to R6, respectively. Note that the cell size of R6 was estimated from the spacing of the elongated cellular structure, which is normally believed to be the characteristic length scale associated with the cell boundary-dislocation interactions. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
a closer inspection of the discrete strain events displayed in the stress-strain curves, we found that the dislocation nucleation and the strain hardening processes of AM 316L SS micro-pillars are cell size-correlated, with easier dislocation nucleation for micro-pillars containing smaller cells. The present study will further the understanding of the underlying mechanisms that govern the excellent mechanical responses of AM 316L SS. The 316L SS rectangular plate was fabricated in an argon environment by a commercial BLT S210 PBF machine equipped with a 400 W fiber laser. The initial plasma-atomized 316L SS powders had a particle size ranging from 20 to 80 μm, with a mean value of ∼30 μm. We performed a series of control experiments to optimize the processing parameters for 316L SS, where the details had been reported in our previous work [13]. A cubic sample was cut from the as-printed plate by electro-discharge machining (EDM). Its surface was polished with 360–40 0 0 grit metallographic silicon carbide papers, followed by polishing with 3 μm and 1 μm diamond suspensions. A modified Carpenter’s reagent was used for etching to enhance the image contrast for scanning electron microscopic (SEM, FEI Scios) observations [13]. Fig. 1a displays the representative SEM image of the etched surface, where fusion boundaries and the ultrafine cellular structure can be clearly seen. We chose six representative regions (R1 – R6) with cell size ranging from 330 nm to 590 nm, as shown in Fig. 1b-1g, for the subsequent micro-pillar preparation. In particular, two micro-pillars with di-
Pillars
Cell size (nm)
R-1 R-2 R-3 R-4 R-5 R-6
435 438 336 451 502 590
± ± ± ± ± ±
63 70 46 76 77 118
Schmid factor
0.492 0.472 0.496 0.495 0.494 0.491
± ± ± ±
0.006 0.002 0.005 0.001
± 0.001
Increment in resolved shear flow stress (MPa) 11.5 ± 2.2 10.8 ± 1.0 9.6 ± 2.5 11.6 ± 0.7 11.9 ± 1.7 12.7 ± 2.4
ameter and height of ∼4.5 μm and ∼9 μm respectively, were fabricated from each region by a focused ion beam (FIB, FEI Scios) system (Fig. 2a). The taper angle of these micro-pillars was controlled within 2° in order to reduce the inhomogeneous deformation associated with the variations in gauge cross-section areas [14]. The actual cell size in pillars fabricated from regions 1 to 6 was measured to be 445 ± 36 nm, 430 ± 21 nm, 345 ± 36 nm, 463 ± 46 nm, 492 ± 42 nm and 595 ± 75 nm, respectively (for details please refer to Fig. S1). We employed electron backscatter diffraction (EBSD, TEAMTM EBSD Analysis System, EDAX) to examine the grain orientation of each micro-pillar and found that the two micro-pillars that contained identical cell sizes had very similar grain orientations (Fig. 2b). The Schmid factor for each micro-pillar was estimated by averaging the Schmid factors of five randomly selected positions that were ∼ 1 μm distance away from the milling trench, as indicated in Fig. 2b. Note that there are local misorientations inside one individual grain, which may be induced by the strong thermal gradient during the rapid solidification and extensive thermomechanical cycles during the layer-by-layer building processes [8]. The average cell size and Schmid factor of each region are documented in Table 1, from which a small Schmid factor deviation can be observed, in spite of the existence of local misorientation gradients. In situ compression tests of the micro-pillars were conducted in SEM using a cradle-based nanoindentation system (NanoFlip, Nanomechanics Inc.), with an 8 μm-diameter conductive diamond flat punch tip. The compression tests were carried out at room temperature under the displacement controlled mode to a maximum strain of ∼ 0.1, at a nominal strain rate of 5 × 10−4 s−1 . Fig. 3a shows the representative compressive engineering stress vs. strain responses of micro-pillars fabricated from different regions. Because of the intermittent nature of the compressive flow and the difficulty in accurately defining the onset of plastic deformation, the yield strengths of the micro-pillars were identified by the stress at which the first strain jump (burst) occurred, using a strain jump of ε = 0.002 as the threshold [14]. Surprisingly, as revealed in Fig. 3a, the yield strengths of micro-pillars possessing various cell sizes were found to be fairly independent of the cell size and no salient trend was found. This can be better appreciated by plotting the yield strength versus cell size figure (Fig. 3b), where the Hall-Petch typed dependence of yield strength on the inverse square root of cell size (d−1/2 ) was not observed [8,9]. We further considered the Schmid factor for the operating slip system in each micro-pillar, where the calculated resolved shear yield strength (τ y ) versus d−1/2 (Fig. 3c) also displayed a weak (if any) linear correlation (dashed line; coefficient of determination for the best fit, R = 0.32). The fitted slope of the τ y vs. d−1/2 data (k) in metals is generally regarded as a measure of dislocation pinning capacity at boundaries [15], and the fitted k ≈ 1365 MPa nm1/2 from Fig. 3c was substantially lower than that of strengthening
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Fig. 2. (a) SEM image of the as-milled micro-pillars with the selected regions (R1-R6) superimposed. The inset shows the enlarged view of a couple of representative micropillars for compression test. A rectangular trench, with dimension of 30 μm × 10 μm × 5 μm (length × width × depth) was milled next to the pillars to ensure a clear view of the pillars during mechanical testing. The tilt anlge was 15° during the observation. (b) The corrsponding EBSD inverse-pole figure map superimposing high-angle (grain misorientation > 10°, black line) and low-angle (grain misorientation 2°−10°, white line) grain boundaries. The row of 5 crosses illustrates the locations for Schmid factor measurement of a micro-pillar in R-1.
Fig. 3. (a) Representative stress-strain curves of micro-pillars containing various cell sizes. (b) Yield strength versus cell size of micro-pillars, showing their weak (if any) correlation. (c) Resolved shear yield strength versus cell size and a best linear fit of the data (dashed line). For comparison, the Hall-Petch fitting plot of bulk austenitic 316L SS is also shown [8]. (d) A plot of resolved shear yield strength as a function of the square root of dislocation cell volume fraction (Vf 1/2 ). The dashed line is the best linear fit of the data.
mechanism by grain boundaries measured from bulk austenitic 316L SS fabricated by conventional methods (k ≈ 2674 MPa nm1/2 [8]). This indicated that the dislocation cells in AM 316L SS were relatively weak dislocation barriers and cannot be treated equivalent to typical grain boundaries in terms of strengthening. On the contrary, by calculating the volume fraction (Vf ) of dislocation cells
in each region (Fig. 1b–g) (see Supplementary text. S1, Fig. S2 and S3 for details of the measurement), we found that τ y exhibits a well-fitted linear correlation with Vf 1/2 (R = 0.91) (Fig. 3d), thus obeying the well-known Taylor-strengthening power law for forest dislocations [16]. That being said, in contradiction to previous hypothesis of a strong dependence of yield strengths on cell sizes, we
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Fig. 4. (a) Statistical distribution of displacement (normalized by the Burgers vector, b) of a discrete strain event versus the resolved shear flow stress at which it occurred, for micro-pillars containing various cell sizes. The general ‘higher flow stress-larger displacement’ is observed, in spite of some relatively large scattered data points (e.g., a displacement of 1150 b at the resolved shear stress of 297 MPa is included for pillars in R2). (b) A plot of shear flow strength improvement of a burst versus the strain value at which it occurs for micro-pillars with various cell sizes. The general trend is depicted by an arrow.
demonstrated that it is likely to be the dislocation density, rather than the cell size, that dominated the yield strength of AM 316L SS. This proposal is in line with a recent study of heat-treated AM 316L SS, where the hardness of the sample was significantly reduced after dislocation recovery, although the dislocation cell size stayed unchanged [17]. It is worth mentioning here that, the yield strengths of micropillars (between 550 and 620 MPa) were found to be comparable with those of the corresponding bulk samples (552–635 MPa [13]). As a matter of fact, this lack of specimen size effect was expected, because the sizes of typical dislocation obstruction features (i.e., dislocation cells) were approximately one order of magnitude smaller than the pillar diameter. This much smaller internal microstructural length scale as compared with the external one (i.e., pillar diameter) indicates that the deformation of these micropillars was governed by dislocation-internal boundary (cell boundary) interactions instead of free surface effects [18,19]. Specifically, the numerous dislocation cells in the pillars would intrinsically act as dislocation trapping and/or nucleation sites [10], making it less likely for dislocation source-truncation or exhaustion hardening [14,20]. However, unlike macroscopic samples characterized by smooth plastic flow, the AM 316L SS micro-pillars showed lots of strain jumps during the hardening processes, which were likely to associate with internal dislocation avalanches. It had been well demonstrated that dislocation activity during the avalanche is usually dominated by a single-slip system [14,21], and indeed, as revealed in Fig. S4, the post-deformation micro-pillars were shown to contain numerous single slip bands. These observations suggest that, although dislocation cells may trap dislocations [10], such trapping activity may not be strong enough to stimulate a massive mobile dislocation cross-slip, thus leading to a discrete plastic flow in AM 316L SS micro-pillars. This agrees well with our results of reduced k obtained in Fig. 3c and the previous experimental observations [10], where dislocation cell walls were found to have the ability to hinder but not entirely block the dislocation motion. To further understand the dislocation activities in AM 316L SS after yielding (i.e., the hardening process) so that the role of cell size on plastic deformation can be more precisely elucidated, we made an in-depth study of the plastic flow intermittency. In particular, it is known that the statistical attributes of strain bursts are believed to closely correlate with the intrinsic microstructural features, e.g., the density and morphology of pre-existing dislocations
[21,22]. As shown in Fig. 4a, a plot of the displacement magnitude for all strain bursts of each micro-pillar (normalized by Burgers vector of austenitic 316L SS, b = 0.257 nm) versus the corresponding resolved shear flow stress revealed a general trend that the discrete displacement increases with increasing resolved shear flow stress. Furthermore, the distribution of these burst events didn’t show a cell size dependence, and we highlighted this in Fig. 4a by comparing the data distributions of the two extremes, i.e., micropillars containing the smallest (red triangle symbol) and largest (blue diamond symbol) cell sizes. This observation again suggested that the cell boundaries in this study were not strong dislocation obstacles as low- and high-angle grain boundaries, which would usually cause a distinct correlation between strain burst size and characteristic dislocation obstruction length scales (grain sizes) [23–25]. In addition, we calculated the increments in resolved shear flow stress after each strain burst event as a function of the engineering strain at which the burst occurred (Fig. 4b). Two important findings can be immediately drawn. First, the strength improvements for early burst events were larger and they gradually attenuated with increasing strain for almost all tested samples, eventually reaching a nearly constant value after an engineering strain of ∼ 0.03 was achieved. We propose that this observation stemmed from the dislocation cells-mobile dislocation interactions during plastic deformation. Specifically, as evidenced by the lack of local misorientations across the cell boundaries and their unique characteristic of thermal tolerance, dislocation cells in the as-printed 316L SS are energetically stable [8,11]. Upon straining, mobile dislocations were trapped and piled up at these cell boundaries, which would cause stress concentrations and favor dislocation nucleation/emission from these boundaries [10]. The strong trapping activities of dislocation cells were also demonstrated in our previous post-mortem TEM analysis [8,13]. Such dislocation trapping induced stress concentrations would subsequently give rise to a reduced stress increment that was needed to compensate for the dislocation avalanches and sustain further plastic deformation, leading to the evolution of shear flow stress increment as shown in Fig. 4b. Here we consider that the stress increment arose at each strain burst primarily from the resistance to dislocation generation at the cell boundaries, rather than the resistance to dislocation nucleation inside the cells (e.g., from the Frank-Read sources [26]). This is because, the stress
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increment required for the operation of normal Frank–Read sources comes from the back stress rendered by pile-up dislocations [27], which can be expressed as τ = GNb/[2π (1−ν )L]. By assuming that the separation L between a source and dislocation pile-up to be about half the cell size and taking shear modulus G = 75 GPa, and a Poisson’s ratio of ν ∼ 0.3, the calculated shear stress from just two piled dislocation (N = 2) would range from 29 to 58 MPa for samples with a cell size of 30 0 nm–60 0 nm studied here. This value is substantially higher than the experimental measured ones (Fig. 4b), suggesting that the dislocation multiplication in the cell interiors was unlikely the operating mechanism that accounted for the stress increments on the stress vs. strain response. Second, as revealed by the statistically calculated shear stress improvement (Table 1), micro-pillars with smaller cell size generally exhibited lower stress increments after the discrete strain bursts. We believe that this cell size correlated dislocation nucleation/hardening responses arose from the differences in the density of piled-up dislocations at the cell boundaries for a specific shear strain, γ . It is indicated in Orowan equation that (γ = ρ × b × l, where ρ is mobile dislocation density and l is dislocation mean free path) [16], to reach the same γ , higher density of dislocations have to be multiplied for smaller dislocation trapping distances; that is, stress concentrations along cell boundaries would be more intensive for smaller dislocation cells, thereby leading to reduced stress improvement that required for dislocation multiplication. However, a quantitative analysis of the mechanisms deserves further investigations and is out of the scope of the present study. In summary, our inspection of the burst evolution showed that, although the yielding behavior of AM 316L SS was mainly dominated by dislocation densities, the subsequent dislocation nucleation and hardening processes of the micro-pillars were cell size-correlated. Note that the influence of twinning effect on hardening behavior was not considered here, because we didn’t observe extensive twinning activities in the deformed pillars (Fig. S5). We employed micro-pillar compression experiments to probe the cell size reliance of the yield and plastic deformation processes of AM 316L SS. The greatest advantage of the micromechanical approach as compared to traditional bulk mechanical characterizations is that it enables a site-specific investigation, which is particularly significant for AM metals that intrinsically contain heterogeneous microstructures. Our findings provide insight into the mechanisms that govern the superior mechanical properties of AM 316L SS. The experimental strategy adopted here also provides a universal pathway to advance the understanding of the microstructureproperty relationship for other rapidly developing AM metallic materials that have highly heterogeneous microstructures. The authors would like to acknowledge the financial support from the National Natural Science Foundation of China (Grant Nos. 51801120, 51771111), the Ministry of Science and Technology
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of China (Nos. 2016YFE0130200, 2017YFB0703100), Shanghai Science and Technology Committee Innovation Grant (Grant Nos. 17JC140 060 0, 17JC140 0603, 1752071240 0), and Distinguished Professor Program of Shanghai University of Engineering Science. Declaration of Competing Interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Supplementary materials Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.scriptamat.2019.10. 005. References [1] D.W.Rosen Gibson, B. Stucker, Additive Manufacturing Technologies, Springer, New York, 2014. [2] S.H. Huang, P. Liu, A. Mokasdar, L. Hou, Int. J. Adv. Manuf. Tech. 67 (2013) 1191. [3] J.J. Lewandowski, M. Seifi, Annu. Rev. Mater. Res. 46 (2016) 151. [4] D. Herzog, V. Seyda, E. Wycisk, C. Emmelmann, Acta Mater. 117 (2016) 371. [5] W.E. Frazier, J. Mater. Eng. Perform. 23 (2014) 1917. [6] Z. Wang, A.M. Beese, Acta Mater. 131 (2017) 410. [7] Z. Sun, X. Tan, S.B. Tor, C.K. Chua, NPG Asia Mater. 10 (2018) 127. [8] Y.M. Wang, T. Voisin, J.T. McKeown, J. Ye, N.P. Calta, Z. Li, Z. Zeng, Y. Zhang, W. Chen, T.T. Roehling, R.T. Ott, M.K. Santala, P.J. Depond, M.J. Mattews, A.V. Hamza, T. Zhu, Nat. Mater. 17 (2018) 63. [9] Y. Zhong, L.F. Liu, S. Wikman, D.Q. Cui, Z.J. Shen, J. Nucl. Mater. 470 (2016) 170. [10] L. Liu, Q. Ding, Y. Zhong, J. Zou, J. Wu, Y.L. Chiu, J. Li, Z. Zhang, Q. Yu, Z. Shen, Mater. Today. 21 (2018) 354. [11] M. Shamsujjoha, S.R. Agnew, J.M. Fitz-Gerald, W.R. Moore, T.A. Newman, Metall. Mater. Trans. A. 49 (2018) 3011. [12] T.R. Smith, J.D. Sugar, C.S. Marchi, J.M. Schoenung, Acta Mater. 164 (2019) 728. [13] Z. Li, T. Voisin, J.T. McKeown, J. Ye, T. Braun, C. Kamath, W.E. King, Y.M. Wang, Int. J. Plasticity. 120 (2019) 395. [14] M.D. Uchic, P.A. Shade, D.M. Dimiduk, Annu. Rev. Mater. Res. 39 (2009) 361. [15] Z. Li, L. Zhao, Q. Guo, Z. Li, G. Fan, C. Guo, D. Zhang, Scripta Mater. 131 (2017) 67. [16] M.A. Meyers, Mechanical Metallurgy Principles and Applications, Prentice-Hall, New Jersey, 1984. [17] P. Krakhmalev, I. Yadroitsava, G. Fredriksson, I. Yadroitsev, S. Afr. J. Ind. Eng. 28 (2017) 12. [18] Y. Hu, Q. Guo, L. Zhao, Z. Li, G. Fan, Z. Li, D.B. Xiong, Y. Su, D. Zhang, Scr. Mater. 146 (2018) 236. [19] D.C. Jang, J.R. Greer, Scr. Mater. 64 (2011) 77. [20] C. Chisholm, H. Bei, M.B. Lowry, J. Oh, S.A. Syed Asif, O.L. Warren, Z.W. Shan, E.P. George, A.M. Minor, Acta Mater. 60 (2012) 2258. [21] F.F. Csikor, C. Motz, D. Weygand, M. Zaiser, S. Zapperi, Science 318 (2007) 251. [22] K.Y. Xie, S. Shrestha, Y. Cao, P.J. Felfer, Y. Wang, X. Liao, J.M. Cairney, S.P. Ringer, Acta Mater. 61 (2013) 439. [23] T. Hu, L. Jiang, H. Yang, K. Ma, T.D. Topping, J. Yee, M. Li, A.K. Mukherjee, J.M. Schoenung, E.J. Lavernia, Acta Mater. 94 (2015) 46. [24] J.Y. Zhang, G. Liu, J. Sun, Int. J. Plasticity. 50 (2013) 1. [25] P.J. Imrich, C. Kirchlechner, C. Motz, G. Dehm, Acta Mater. 73 (2014) 240. [26] J.J. Gilman, J. Appl. Phys. 30 (1959) 1584. [27] K.S. Ng, A.H.W. Ngan, Acta Mater. 56 (2008) 1712.