A review on the science of plastic deformation in laser-based additively manufactured steel

Abstract

Additive manufacturing (AM) is believed to be a breakthrough technology that may potentially revolutionize various industries, including aerospace, biomedical, and automotive industries. The metallic materials processed by AM, especially by laser-based AM, generally demonstrate different mechanical behaviors as compared to their wrought/cast counterparts. Among all the metallic materials, steel has been the most widely used one with an enormous amount of production. Steel can be processed to have varied phases (martensite, ferrite, bainite, austenite, etc.) with different mechanical behaviors. Depending on the stacking fault energy (SFE), the plasticity of steel can be achieved through the slip of dislocations, the formation of martensite, or the generation of deformation twins. In this review, we focus on the plasticity of steel made by laser-based AM and pay less attention to the effect of processing parameters on other properties (e.g., density, surface roughness and solidification cracks). This review helps to deepen the understanding of the deformation behaviors of laser-based AM fabricated steel, which can be potentially applied in different industrial fields.

Introduction

Additive manufacturing (AM) is a novel processing technology, which has been proposed around 40 years ago. It has undergone rapid development over the last decade. The AM technique enables the flexible design of sophisticated functional components or customized products directly from digital files and building 3D parts in an incrementally layer-wise manner under the guidance of digital models, enlarging the manufacturing freedom of some complex components such as the turbine parts [1]. Moreover, as compared to conventional manufacturing (CM) techniques such as subtractive manufacturing, the AM technique requires no expensive machining or forming processes (i.e., punches, dies, or molds) [2].

Laser is usually applied as a heating source for metal AM. In general, there are two representative laser-based AM strategies including laser-directed energy deposition (DED, Fig. 1a) and laser powder-bed fusion (L-PBF, Fig. 1b) [3]. DED is an AM technique that uses a focused laser beam to fuse powder or wire feedstocks delivered to the substrate, generating a tiny melt pool and depositing material layer by layer [4, 5]. DED has several unique advantages including alloy design, high fabrication flexibility, and site-specific deposition; thus, its applicability extends from AM to remanufacturing [6], repairing [7], and surface coating [8]. However, owing to the low precision of DED products, further machining is required to achieve the desired structural and accuracy requirements [3]. L-PBF is another AM process technology that has shown promising application in printing alloys [9, 10] and ceramics [11]. The build process includes a cycle of: (i) depositing powder layers onto the substrates or on previously deposited layer, (ii) melting the powder by the focused laser beam selectively according to each layer profile, (iii) lowering one-layer thick of the platform and (iv) recoating a new powder layer [12]. Near-net-shaping fabrication directly from powders enables L-PBF to be employed in end applications or at least reduces the required post-processing to some extent [12, 13]. L-PBF is typically used on small- and medium-sized products, but thanks to the multi-laser scanning technique and high-speed scanner, the build volume of the largest L-PBF system reached the maximum side length of 800 mm [14], allowing for the printing of components longer than 1 m.

Figure 1

Schematic illustration of a the laser-directed energy deposition (DED) and b the laser powder-bed fusion (L-PBF) [3]

The deposited layer thickness of L-PBF is much smaller than that for DED, e.g., the layer thickness of Ti64 deposited by DED (Optomec 750) is about 200–1000 $\mu\mu$m [15], while it is about 20–100 $\mu\mu$m for L-PBF (Realizer SLM50) [16]. A successful L-PBF requires a high flowability of the powder particles and thus the uniform laying of powders [17]. Compared with DED, the laser power used in L-PBF is lower but with a considerably finer laser beam and a high quality, which can enable both sufficient heat transfer to the melt and a higher printing accuracy [3].

To date, there is a large body of reviews on the topics of laser-based AM metallic materials, covering diverse topics including materials [18,19,20,21,22], powders [23], residual stresses [24], defects [25], processing parameters [26], metallurgy [27], microstructure and associated mechanical behaviors [28, 29], and all encompass reviews [9, 30]. However, to the best of our knowledge, there are few comprehensive and in-depth reviews on the deformation mechanisms of laser-based AM metallic materials. The understanding of the science of deformation of laser-based AM metallic materials serves as the basis for the development of high-performance structural components for engineering applications. Steel is the most widely used structural material in various industries all over the world (e.g., about 2.87 billion tons of crude steel were produced in 2021 [31]), which demonstrates the diverse microstructure resulting from the complex phase transformation, leading to intriguing deformation behavior that incorporates almost all deformation mechanisms. Therefore, the present review focuses on the science of deformation of laser-based AM steel.

Depending on the alloying and thermal–mechanical processing, steel generally consists of multiple phases including ferrite, martensite, bainite, austenite, etc., allowing the operation of multiple deformation mechanisms. The dislocation activity such as the dislocation–dislocation interaction and dislocation–boundaries interaction is the most common deformation mechanism in steel, irrespective of the constituting phases. The presence of the austenite phase in steel introduces other strengthening mechanisms, including twinning-induced plasticity (TWIP, contributed by deformation twins) and transformation-induced plasticity (TRIP, resulted from martensitic transformation) effects. The competing role of the dislocation plasticity, deformation twins, and martensitic transformation depends on the stacking fault energy (SFE) [32]. The TRIP effect is preferred when SFE is below 20 mJ/m² [33], the TWIP effect is dominant when SFE is in the range of 20 ~ 40 mJ/m² [34] and when SFE is beyond 40 mJ/m², the dislocation slip is predominant [35]. The SFE, which depends on the chemical compositions and temperature [36], determines the separation of two partial dislocations in the crystalline material with a close-packed structure [37]. Generally, the higher SFE results in the smaller separation distance between two partial dislocations and thus thinner stacking faults (SFs), and vice versa [37]. In addition to the SFE, the operation of different deformation mechanisms is also affected by microstructural characteristics such as crystallographic orientation [38], grain size [39] and external loading conditions, including stress [40], strain [41], strain rate [42], and temperature [43].

Here we are aimed to reveal the plastic deformation mechanisms in the laser-based AM steel. The typical microstructural features of the laser-based AM steel will be presented by comparing to the steel processed by the conventional processing route. The deformation mechanisms including the dislocation plasticity, deformation twins, and martensitic transformation will be discussed based on the microstructure in different steel grades processed by the laser-based AM technique. Particularly, the deformation reviewed in this work is confined in the time-independent class at room temperature; for the time-dependent deformation (i.e., fatigue and fracture) in AM metals, we refer the reader to other reviews [29, 44,45,46,47]. Additionally, the metal matrix composites (MMCs) are not included in this review. Although there are several steel grades reinforced with hard second-phase particles manufactured especially for DED, covering all of them in detail would take up much too much space in this review. We recommend the reader to another review for the MMCs [48]. We expect that this review will deepen the understanding of the microstructure–mechanical properties relation of laser-based AM steel and provide insights on the microstructure-guided design of high-performance laser-based AM steel for a broad industrial application.

General microstructure in laser-based AM steels

Like the conventionally processed steel, the mechanical properties of laser-based AM steel are also governed by the microstructure, which is, however, completely different from that of the forged/casted counterpart, as shown in Fig. 2. Figure 2a-c presents the typical microstructure of the L-PBF 316L [49]. Slightly elongated grains along the build direction (BD) are observed (Fig. 2a): <110> direction tends to align with the BD, and <001> follows the laser scanning path along x- and y-directions. In the fine-scale STEM/HAADF maps (Fig. 2b and c), the brighter wall contrast origins from the local chemistry heterogeneity (i.e., Cr and Mo segregation) [50, 51]. The direction of elongation of cells is found to be along <001> direction (Fig. 2c), which is consistent with a cell-like dendritic growth structures, i.e., solidification cell structures. Figure 2d-f presents the typical microstructure of the DED 316L [52]. Figure 2(d) presents the track interfaces and solidification boundaries in DED 316L. A lot of equiaxed grains lie along the track center with the formation of columnar grains in the center between two tracks (Fig. 2e). The maximum thermal gradient direction shifts in the process of re-solidification generated by track overlapping and explains the morphological complexity of the grains, and the equiaxed grains would be columnar grains that develop almost vertically beneath the track center [53]. Additionally, columnar grains between two tracks grow toward the track centers rather than aligning with the BD, and when depositing a subsequent track, partial remelting changes the original growth direction [52]. Figure 2(f) depicts the sub-grain microstructures. The transverse cross-sections normal to columnar microstructures should be viewed as thin equiaxed microstructures [54], and so these microstructures are referred to as cells indiscriminately [55]. The columnar microstructures are more likely to form at track edges where remelting occurs [52]. Different from the sophisticated microstructure of the laser-based AM 316L, the wrought 316L exhibits a simple equiaxed microstructure with an average grain size of ~ 29 $\mu\mu$m, as shown in Fig. 2g and h [56].

Figure 2

The typical microstructure of 316L produced by L-PBF (a–c [49]), DED (d–f [52]), and CM (g, h [56]). a Electron back-scattered diffraction (EBSD) inverse pole figure (IPF) for typical grain structure of L-PBF 316L. The right-hand pole figures are for the top face. b The sub-grain structure and the end-on cells at higher magnification showing in the inset. c The elongated sub-grain structures. d The approximate center- and edge-lines (dashed lines) of a DED track; e the grain morphology in the track vicinity; f the colony boundaries (dashed lines) and sub-grain microstructures. g EBSD images and h optical microscopy of the wrought 316L

The unique microstructure of laser-based AM steel is a result of the complex thermal history of the laser-based AM process, involving highly localized melting, rapid cooling, and re-heating of the material repeatedly during deposition of adjacent layers and tracks, a thermal condition termed as intrinsic heat treatment (IHT) [57,58,59]. Consequently, the steel processed by laser-based AM experiences cyclic liquid-to-solid and/or solid-state phase transformation. Usually, a lower heat input possible to obtain fully dense parts is used to minimize the magnitudes of residual stress [24]. Additionally, the smaller size of the melt pool during laser-based AM led to smaller dendrite spacing and grain size, and fewer elemental segregation [26]. As a result, fine grains, solidification textures, and extremely non-equilibrium phases/composition structures can be found in laser-based AM steel [9].

The rapid solidification associated with laser-based AM will result in fine solidification cells [60, 61]. These cells exhibit similar crystallographic orientations. A bundle of such cells together form a grain, up to hundreds of microns in size, i.e., a volume of material bound by high-angle grain boundaries (HAGBs) [18]. The dislocation cell structure is another fingerprint of laser-based AM steel that is characterized by high-density dislocations at the cell boundaries (CBs) accompany by segregating elements or precipitates [51]. The cell structure is formed due to the rapid cooling during the solidification with a relatively high thermal gradient (G) and the crystal growth rate (R) ratio [62]. A larger G and R produce a smaller cell size [63, 64]. The solidification CBs are decorated by high-density dislocations that formed due to the tension–compression deformation cycles caused by the localized heating/cooling heterogeneity [65,66,67]. Lots of dislocation cells inside the grains make a strain gradient, which is transformed to the low-angle grain boundaries (LAGBs) [68]. Additionally, the steel powders used for laser-based AM are usually atomized by liquid or gas, generating a thin oxide layer surrounding the powder surface [69, 70]. The decomposition of the oxide layer during the melting process provides the extra oxygen into the melt pool during deposition, generating nano-sized oxide inclusions [50, 71]. Additionally, the gas impurities in the printing atmosphere may also include the oxygen for the formation of the oxide inclusions. Consequently, it is very difficult to completely get rid of the oxide inclusions in the laser-based AM components [62]. The aforementioned solidification columnar grains, dislocation cells, and oxide inclusions together form the hierarchically heterogeneous microstructure [50, 72,73,74,75], which can span nearly six orders of magnitude (10^–3 ~ 10^–9 m) [50].

The laser-based AM steel has various phases, including ferrite, austenite, martensite, bainite, etc. The solid-state phase transformation in steels is governed by temperature and chemical compositions and both are affected by the laser-based AM process. For example, the 316L made by L-PBF typically exhibits a complete austenitic microstructure without solid-state phase transformation [76]. However, for the lower cooling rate of the DED process, some $\delta$-ferrite (~ 10.9%) is observed at the grain boundaries (GBs) [77]. Additionally, the transformation of martensite into bainite is observed in the bottom and middle regime of laser-based AM H13 [78] and 24CrNiMoY [79] samples due to the thermal history caused by the deposition of the subsequent layers. A certain amount of retained austenite (RA) phase has been observed in some laser-based AM steel grades (e.g., precipitation hardening (PH) steel [80,81,82], maraging steel [83,84,85,86], tool steel [87, 88], and martensitic stainless steel [60, 89, 90]), which typically have a fully martensitic structure in conventional processing. The RA is supposed to be generated mainly under the rapid solidification state, which leads to element segregation and the formation of thermal residual stress (high-density dislocations), thereby chemically [91] and mechanically [92] stabilizing the austenitic phase (cf. Section 3.3 on 316L). Additionally, it is claimed that the governing length scale for thermally induced martensitic transformation is the solidification cells, with the dense dislocation walls being sufficient in resisting the displacive transformation [61]. Besides, reversion and growth of austenite are observed in laser-based AM martensitic 420 because of the in situ partitioning of the austenitic stabilizer element (i.e., C atoms) triggered by the IHT [89]. In laser-based AM steel grade, the aforementioned metastable RA and austenite with low SFE (i.e., ASS and HMnS) will lead to deformation twinning and/or martensitic transformation during deformation [80].

Additionally, the powder blends adhesion to laser-based AM provides a promising approach for adjustment of the composition and chemistry of the alloy flexibly, e.g., by deliberately adding oxide particles to fabricate oxide dispersion strengthened (ODS) steel [75, 93, 94], by elaborately adjusting the Al- or C-content in high-manganese austenitic steel to tune the SFE [91, 95, 96], and by tuning the Ti content in Maraging steel to modify the fraction of RA [97], and thus, enables tailored mechanical properties.

Plastic deformation mechanisms in laser-based AM steels

Generally, the microstructure of the laser-based AM steel has many unique features, such as the hierarchical structures, the solidification cells, the various boundaries/sub-boundaries, the metastable structures, and the nanoscale precipitates and inclusions, as shown in Fig. 3. During deformation, the hierarchical heterogeneous structures can interact with dislocation and provide back stress hardening, the dense dislocations, CBs, GBs, phase boundary (PB), inclusions and precipitations, which are all effective barriers for dislocation motion and thus strengthen the AM steel. Especially, the austenite phase specifically RA from the reverse transformation during the processing of martensitic steel grades and the preexisted austenite with low SFE from ASS and HMnS can provide TRIP and TWIP effects. Therefore, in this review, we attempt to discuss the relevant deformation mechanisms of laser-based AM steel from the aspect of its unique microstructural features.

Figure 3

Review over typical microstructures produced by CM, L-PBF, and DED, as well as the associated deformation mechanisms in different laser-based AM steels in this review

Dislocation activity in laser-based AM steel

Dislocations are line defects in crystalline materials, which are the main carrier of plastic deformation. The dislocation activities including dislocation storage, nucleation, motion, multiplication, and pinning will occur at the submicron to nanoscale during the plastic deformation [98]. Numerous strengthening strategies have been developed to enhance the strength by inhibiting the dislocation activity, including the presence of numerous GBs to impede the motion of dislocations [99], the introduction of solid solution or second hard phases to block the slip of dislocations [100], and also the generation of forest dislocations by plastic straining to prevent further dislocation activities [101].

Based on the dislocation theory, the yield strength ${\sigma }_{y}$ of steels can be described considering the contributions of the individual strengthening mechanisms: $\sigma_{y} = \sigma_{f} + \sigma_{GB} + \sigma_{CB} + \sigma_{SS} + \sigma_{BS} + \sigma_{\rho } + \sigma_{P1} + \sigma_{P2}\sigma_{y} = \sigma_{f} + \sigma_{GB} + \sigma_{CB} + \sigma_{SS} + \sigma_{BS} + \sigma_{\rho } + \sigma_{P1} + \sigma_{P2}$ [102], which are summarized in Table 1. As mentioned previously in Sect. 2, the dislocation cell structure, nano-oxide inclusions, and hierarchical heterogeneous microstructure are unique structures that exist in almost all laser-based AM steel grades; additionally, the laser-based AM have some unique advantages for the design and processing of precipitation strengthening materials, and thus, we review the dislocation activity in laser-based AM steel from the following three aspects.

Table 1 The dislocation slip-related strengthening mechanisms in laser-based AM metals

G is shear modulus, M= 0.4 for BCC and 0.2 for FCC, b is Burgers vector,$\rho\rho$ is dislocation density, ${K}_{i}$ is the strengthening coefficient, ${c}_{i}$ is the volume fraction for element i, k is Hall–Petch factor, d is grain size, f is the volume fraction of precipitation, r is the radius of precipitate, L is the spacing of precipitates on slip plane. Reedited from [102].

Dislocation–dislocation cell interaction

The ${\sigma }_{f}$, ${\sigma }_{GB}$, and ${\sigma }_{SS}$ are all grain size-dependent parameters [104, 105, 109] and are similar to traditionally processed steels, and thus, we do not discuss them here. The strengthening effect of the solidification cell structure is widely discussed [50, 64, 109,110,111,112]. Some of them assume the strengthening effect of CB can be related to the Hall–Petch relation like the GB does [50, 110]. However, others suggest the CB alone cannot be simply resembled as GB and is not suitable using the Hall–Petch relation [64, 109], due to low misorientation angles between adjacent cells [109], and the CB is a weaker barrier for dislocation motion than GB [64, 113, 114]. The observed cell spacing is previously thought to scale with the strength of L-PBF 316L [115]. Others, however, have noted that the size of cells within the same sample is non-uniform, with an average spacing varying from 0.2 to 1.0 $\mu$m, making it difficult to construct a single-size scaling law [116]. The non-uniform distribution of the cell structures is observed with the volume fraction varying from sample to sample, which could be connected with the processing parameters and sample geometry [117]. Importantly, the EBSD IPF orientation mapping in SEM revealed no misorientations across the majority of CBs (Fig. 4e), implying that the CB should be treated differently from a typical interface like HAGB and dislocation wall [50]. Additionally, the dislocations that can transmit through thin CBs are blockage by thick CBs, which has recently been discovered in AiSi10Mg alloys [118]. These findings suggest that the strengthening ability of the cell structure is not well understood and remains to be studied.

Figure 4

Deformation structures with different tensile strain levels of L-PBF 316L. a The cell structures at ~ 3% strain, with little change of the cell size and shape. The dislocations trapped by the cell walls are observed (inset). b The intersections of deformation twins with cells after ~ 12% strain and these cells are slightly elongated. c The cell structures and deformation twins after ~ 36% strain. d IPF map (top) of a cell region at ~ 3% strain, with associated IQ (middle) and KAM (bottom) maps. e The misorientation angle variation obtained from the point-to-origin (top) and point-to-point (bottom), across the cells marked by the solid line at ~ 3% strain [50]

Plenty of residual stresses is generated and stored during the laser-based AM process due to the generation of thermal strain coupled with the reduction in yield strength during deposition [65,66,67]. Generally, depending on the processing parameters, the density of these stored dislocations is in the range of 10^14~16 m⁻² [65, 72, 80, 119, 120], which is the main contributor to the high-yield strength of laser-based AM steels according to the classic Taylor hardening law [120]. The stored high-density dislocations are mainly concentrated at the solidification CBs (Fig. 2b insert) accompanied with segregation of alloying element, thus forming the low-energy configuration, namely dislocation cell structure [51]. This kind of dislocation cell structure is widely observed in laser-based AM steels [51, 121].

In conventional manufactured low-alloyed steel, the size of dislocation cells in deformed part is closely related to the flow stress [122]. Similarly, the presence of dislocation cell structure in laser-based AM steels is responsible for the improved tensile strength [51]. The strengthening behavior of the dislocation cell structure [51], and the dislocation size-related strengthening mechanism [123] are confirmed by the micropillar test. The dislocation nucleation is easier for smaller cell sizes, which can be ascribed to the different densities of dislocations accumulated at the CB under a particular shear strain, ${\tau }_{s}$ [123]. According to the Orowan equation (${\tau }_{s}$= $\rho \bullet b\bullet l$, where l is the mean free path [124]), the higher dislocation density has to be achieved to reach the same ${\tau }_{s}$, for smaller dislocation slip distance. Considering the sensitivity of the cell size to the processing parameters, the mechanical properties of laser-based AM steel can be optimized by controlling microstructure through tuning the printing strategies (i.e., tuning the cooling rate by adjusting the scanning speed, cell sizes in the range of 0.2 ~ 1 $\mu$m are processed) [51]. The defects inherent to the misorientation of solidification cells can act as dislocation sources [51]. The pinning effect of particles and the low misorientation between solidification cells effectively stabilize the dislocation cell structure to maintain the cell configuration during the entire plastic deformation (Figs. 4a-c) and enable the steady increase in work hardening rate for plastic flow, delaying the onset of plastic instability according to the Considere criterion ($\sigma =d\sigma /d\varepsilon$, where $\sigma$ and $\varepsilon$ are true stress and true strain, respectively) [51, 125]. Additionally, in laser-based AM austenitic steel with low SFE, the dislocation cell structure cannot fully impede the dislocation motion [51]. This is because the perfect dislocation can widely dissociate into two partial dislocations and move forward with increased applied stress when they are trapped by dislocation cells. Consequently, the presence of dislocation cells enhances the strength without sacrificing ductility, which makes the dislocation cell an idea “modulator” to allow the development of an excellent combination of strength and ductility [51].

Precipitation hardening

Precipitation hardening is realized by introducing finely dispersed precipitates to impede the motion of dislocations and has been widely used to enhance the strength of many metallic systems [100]. The strengthening mechanism, hardened by either precipitates or dispersoids, is derived from the interaction of dislocations with the second phase [126]. Generally, the interaction depends on the amount, strength, dimensions, spacing, and the detailed behavior of the precipitate, all of which differs among different alloy systems [108]. Generally, the precipitation hardening could be described using either shearing or looping mechanisms, depending on the nature of the precipitate and the crystallographic relationship with the matrix [127]. Considering the laser-based AM steel, the precipitates in maraging steel grades [81, 84, 85, 121, 128,129,130,131,132] and oxide inclusions in ODS steel grades [75, 93, 133,134,135,136] provide a significant strengthening effect. Additionally, the formation of tiny secondary phase in the laser-based AM steel, for example, the formation of bainite in DED H13 [78], L-PBF 24CrNiMoY [79], and DED 24CrNiMo [137] may also provide a strengthening effect depending on the interaction between the matrix and second phase and will be discussed in the next section.

Maraging steel demonstrates ultra-high-strength and excellent toughness, originating from a soft ultralow carbon martensitic matrix together with different types of nanoscale intermetallic precipitates [127]. It is reported that the IHT of the DED process is sufficient to initiate the early stage of precipitation of grade 300 maraging steel [138]. However, owing to some relatively sluggish precipitation such as Mo precipitates, the precipitation triggered by primary IHT is limited [83], making an extra heat treatment or special design of the laser-based AM process necessary [139]. Interestingly, in situ precipitation of Ni–Al in Fe–Ni-Al maraging steel is triggered by IHT during DED without any further post-heat treatment [131]. The low lattice mismatch between the martensite and NiAl phases enables a high nucleation rate, delivering an extremely high number density of Ni–Al precipitates triggered by IHT [131]. Moreover, nanostructured maraging steel with excellent mechanical properties (1300 MPa and 10% elongation) is fabricated by in situ optimization of martensite and Ni–Ti nanoprecipitation through tuning thermal history of the DED process, as shown in Fig. 5 [139]. Since these precipitation phases triggered by IHT have a low lattice mismatch with the martensite matrix and are about several nanometer sizes, which can be regarded as shearable precipitates by dislocation, i.e., these precipitates harden the material by shearing mechanism [126, 140]. These excellent works indicate the possibility of producing components that are precipitation strengthened by exploiting the IHT of the DED process without any post-heat treatment.

Nanoscale inclusions in laser-based AM ODS steels can be obtained through in situ oxidation in the melting phase under an oxygen-containing atmosphere or during the consolidation of mechanically alloyed powders [93, 134, 141, 142]. Various nanoscale oxides are usually formed during the laser-based AM due to a high possibility of powder contamination by oxygen and a low oxygen solubility of steels, which can contribute to the strength through Orowan bypass strengthening, but also provide high resistance to grain coarsening at elevated temperatures through Zener pinning effect [93]. The unexpected formation of ultra-fine oxides paves a new pathway for the design of ODS steel specifically for laser-based AM, which demonstrates excellent mechanical properties at high temperatures [135]. The oxide inclusions usually lack a well-defined crystallographic structure [143], suggesting that the interface between the steel matrix and oxide inclusions is incoherent [144]. Moreover, the diameter of the oxide inclusions (~ 100 nm) in laser-based AM ODS steels is usually larger than a shearable precipitate [108, 140, 145]. Thus, dislocations are unable to slip across such dispersoids in ODS materials by a shearing mechanism and have to bow around the obstacles for deformation to proceed.

Heterogeneous deformation-induced hardening

The hierarchically heterogeneous microstructure enables the laser-based AM steels with an excellent combination of high strength and good ductility [50, 146,147,148,149]. The microstructural spatial length-scale characteristics influence the evolution of geometrically necessary dislocation (GND)-type hardening and the plastic localization and thus influence the overall mechanical behavior [112]. The GNDs, accommodating local deformation gradients [150,151,152] generated by thermally induced volume changes during the laser-based AM process [109], are sessile dislocations. The deformation gradients generated due to inhomogeneous deformation of constituting phases during the plastic deformation of laser-based AM steel are accommodated by GNDs accumulation and pile-up against the domain boundaries, allowing compatible deformation of the heterogeneous structures [150]. The presence of GND generates the long-range back stress in the soft domain to compensate for the applied shear stress and make the soft domain stronger [153]. Such back stresses inhibit the forward dislocation glide and facilitate the reverse motion, resulting in low yield stress at the reversed loading. This is the so-called Bauschinger effect, which gives rise to kinematic hardening, elevating both strength and tensile ductility of heterogeneous materials [154,155,156,157]. Particularly, the high-density GBs at the molt pool boundary and the thick CBs inside each grain of the laser-based AM 316L belong to the hard domains [113]. However, the effectiveness of back stress hardening associated with the CB is lower than that of GB due to the low storage ability of GNDs [64, 113, 114], as illustrated in Fig. 6.

Figure 5

APT analysis of precipitation in laser-based AM maraging steel. a The as-printed microstructure, the Ti atom maps of the austenite (left) and martensite (right) of the b soft region, and c hard region in a. The precipitates only form upon IHT in the martensite of the hard region [139]

The asymmetric tension–compression response is observed in laser-based AM steel with heterogeneous structures [158]. In laser-based AM 316L, the cell structure of heterogeneous dislocations leads to the development of long-range internal stress with forward stress in the hard cell wall and back stress in the soft cell interior, which can resist the forward slip of dislocations and contribute to the reverse slip within the cell interiors, resulting in lower yield stress under reverse loading [159]. Considering the thermal sensitivity of microscale (i.e., intra-, and inter-granular) residual stress, the microscale residual stress can be relieved by heat treatment, which can reduce the tension–compression asymmetries and alter the work hardening behavior.

Deformation twinning in laser-based AM austenitic steels

In addition to the complex dislocation activity, the generation for deformation twinning in laser-based AM austenitic steels (i.e. high-manganese steel, HMnS [91, 95, 160,161,162,163,164] and austenitic stainless steel (ASS) [50, 62, 119, 120, 146, 165,166,167,168]) becomes highly possible with the decrease in SFE. The underlying mechanism for the improved work hardening behavior with the generation of deformation twins is frequently ascribed to the dynamic Hall–Petch effect [160]. Specifically, the twin boundaries (TBs) sub-divides the original coarse grains into twinned and untwined areas and decreases the dislocation mean free path, improving the work hardening behavior and delaying the onset of plastic instability (i.e., necking) [119]. Unlike GBs and phase interfaces, TBs have lower energy and a stronger capacity for dislocation storage [169]. The dislocations can glide along or pass through with the TBs through dissociation into partial dislocations, which can improve both the strength and ductility simultaneously [169].

The strong twinning activity of the laser-based AM 316L is confirmed by TEM analysis as shown in Fig. 7 (a-d). A lot of SFs are observed in the as-built 316L, which indicates its low SFE [120, 146]. It is well known that mechanical twinning is facilitated at the expense of dislocation glide due to the low SFE value [170], which induces large deformation plasticity [171]. Thus, the formation of SFs in the as-built 316L favors mechanical twinning, which improves the tensile ductility. The high initial dislocation density together with the mechanical twins contributes to the high strength of the laser-based AM 316L since TBs can effectively impede the motion of dislocation by reducing the mean free path of dislocations [120, 146, 149, 172].

Figure 6

a The shear strain mismatch compensated by GNDs at different boundaries, b the different hetero-deformation-induced (HDI) stress, i.e., back stress, induced by GB and CB with elevated strain [113]

The SFE of 316L is estimated to be about 40–60 mJ/m² [173], which is a range unfavorable to the formation of deformation twins [174]. However, the N atoms can be easily introduced during the laser-based AM process, especially when using N atomized powder and/or printing in an atmosphere with N, which can lower the SFE, and therefore facilitate the dissociation of dislocations and the formation of deformation twinning [165]. N is widely believed to reduce the SFE, promoting the twinning activity during deformation in high N-containing steels [175, 176]. The N influences structural development by decreasing the SFE and increasing the internal friction: with the increased N content, the structures become finer, the planarity of glide becomes sharper, and therefore, the onset of deformation twinning shifts to higher stresses and lower strains, i.e., the more significant contribution of deformation twinning to the total strain [176].

The propensity of deformation twinning in a specific grain is dependent on its crystallographic orientation, which is well reported in conventionally manufactured TWIP steel [177, 178]. The grain orientation-dependent TWIP behavior in L-PBF 316L is verified by TEM analysis, as shown in Fig. 8 [166]. No twins are formed within the < 001 > -oriented grain (Fig. 8a). However, separated single-twin lamellae are commonly observed in the non- <001> -oriented grains with sub-microns in width (Fig. 8a and b). Therefore, the deformation twinning is unfavorable for grains with orientations close to <001> alignment along the strain direction. By comparing the Schmid factor between the leading and trailing partial dislocations, the relationship between texture and loading direction can be elucidated [179]. Extensive SFs arise when the Schmid factor of the leading partial dislocation is higher than the Schmid factor of the trailing partial dislocation [179,180,181], permitting the production of $\varepsilon$-martensite and twins [180].

Figure 7

The SFs formed in the laser-based AM austenitic steels, a 316L processed by DED [146], b 316L processed by L-PBF[120], c and d are SAD patterns of regions ‘1’ and ‘2’ in b, e the bright-field and f HRTEM image of SFs and twins in L-PBF 304L [147]

The distinct crystallographic-orientation-dependent strain hardening behaviors of laser-based AM 316L are shown in Fig. 9. The different mechanical behavior can be ascribed to the orientation-dependent TWIP effect [167]. The grain with a higher Taylor factor along the loading direction tends to develop deformation twins (i.e., the grain lies in <110> and <111> orientation) [182]. Note that the Taylor factor depicts the ratio of flow stress to the critical resolved shear stress of polycrystalline material [183]. The grains with higher Taylor factors require higher applied stress to initiate the dislocations during plastic deformation, leading to greater dislocation proliferation and pile-ups, which is conducive to the deformation twinning [166]. Based on the Taylor factor theory, in-process engineering of microstructure into a <110> texture in laser-based AM 316L is investigated by controlling the melt pool shape, e.g., by modifying the processing parameters to form a deeper and narrower melt pool, the columnar grains grow perpendicular to the curved melt pool border toward the top-center of the melt pool during solidification and tend to tilt by $45^\circ$ from the building direction following the heat transfer path and thus produces a <011> crystallographic texture in the direction of construction [166], or by varying the workpiece geometry (e.g., strut diameter) [168]. The excellent ductility is thus achieved with the assistance of the TWIP effect originating from the <110> texture formed during laser-based AM. These works lay a foundation for the excellent strength and ductility of laser-based AM parts through the TWIP effect associated with texture engineering.

Figure 8

The grain orientation-dependent deformation twinning in L-PBF 316L during the tensile test. a TEM image showing the deformed sample with the adjacent < 001 > and < 321 > grains. The GB and single-twin lamellae are indicated. b The SFs and nano-twins along the < 011 > zone axis in the TEM image [166]

The deformation twinning in laser-based AM 316L is usually accompanied by other deformation mechanisms, including dislocation slips, cell wall evolution, and cell–twin, twin–twin interaction, as shown in Fig. 10 [50, 62, 120, 166]. Plenty of dislocations is stored at the TBs after large tensile strain (Fig. 10a), indicating that the TBs are effective barriers for impeding the motion of dislocation and thus substantially strengthen the L-PBF 316L. Additionally, owing to the low excess energy and high symmetry of twins, a large number of mobile dislocations can be generated and stored at TBs [120]. The dislocation/twin interactions in the L-PBF 316L are revealed in Fig. 10b-d [166]. Plenty of multiple deformation twin lamellae and dislocations are discovered within the <011> -textured austenite grain (Fig. 10b). The deformation twins and dislocations likely interact with each other during plastic deformation. The multiple nano-twin lamellae in Fig. 10b and the atomic scale high-resolution TEM morphology of the nano-twin lamellae and SFs are displayed in Fig. 10c and d, respectively. The formation of SFs composed of partial dislocations is due to the dislocation–twin interactions [184], which indicates that deformation twinning is the dominant mechanism in the <011> -textured sample, and the simultaneously enhanced strength and ductility could be ascribed to the formation of nano-twins [166]. A progressive deformation twinning mechanism is revealed in Fig. 10e-h, including twins nucleate at ~ 3% strain from HAGBs and penetrate the cell walls and LAGBs to subdivide these structures to promote cell–twin interaction (Fig. 10e, f), then with the formation of different twinning systems in the individual austenite grain at strain of ~ 12% (Fig. 10g), the twin–twin interactions proceed, thus forming a massive 3D network obstacle to dislocation motion with the increased strain (Fig. 10h), offering a progressive work hardening [50].

Figure 9

a Engineering stress–strain curves with different textures of laser-based AM 316L, b, c, and d represent the true stress–strain hardening rate over the true strain with <100> , <110> , and <111> textures, respectively [167]

Deformation-induced martensitic transformation

Generally, in austenitic or austenite-containing steel grade, the metastable austenite will transform to $\varepsilon$- and ${\alpha }^{^{\prime}}$-martensite progressively during the plastic deformation, providing the TRIP effect to improve the work hardening and thus simultaneous enhancement of strength and uniform elongation [185]. In laser-based AM steel grade, the metastable austenite comes from two aspects: the preexisted austenite with low SFE in austenitic steel grades (i.e., ASS and HMnS), and the RA in martensitic steel grades formed from reverse transformation due to the thermal history of AM process or post-heat treatment. The effectiveness of the TRIP effect in enhancing the work hardening behavior can be explained from two mechanisms, including the generation of new dislocations resulting from the accommodation of the transformation shape strain of fresh martensite [186], and the dynamic strain partitioning between the soft matrix phase and the hard fresh martensite [187]. The occurrence of martensitic transformation depends on the SFE of austenite [188]. The ${\alpha }^{^{\prime}}$-martensite can be generated in individual $\varepsilon$-martensite for the austenite with a very low SFE (e.g. blew 12.4 mJ/m² in Fe–Cr-Ni alloy) [189]. The $\varepsilon$ phase can act as a transient phase of ${\alpha }^{^{\prime}}$-martensite transformation [190], but it is not a prerequisite condition, because the ${\alpha }^{^{\prime}}$-martensite can also nucleate at the shear band intercepts in the form of twins or SFs [191].

The strain-induced martensite transformation in conventional 316L during plastic deformation is observed in the study [192]. However, the TRIP effect in laser-based AM 316L is substantially suppressed [146, 193], which indicates the mechanical stability of austenite in laser-based AM 316L is higher than that of conventional 316L. The stability of austenite is affected by the grain size [194], chemical composition (especially C and Mn) [195], and initial dislocations [196]. The effect of grain size on mechanical stability is studied by Naghizadeh et al. [194], and they reported that when grain size decreased below ~ 50$\mu m\mu m$ m, the susceptibility to martensitic transformation reduced. It has been reported that the SF probability and shear band formation decreased with the decrease in grain size [197], which reduce deformation-induced martensite nucleation [198, 199]. In addition, the austenite grains in laser-based AM steel have a high density of dislocations, LAGBs, and fine cell structures, which reduce the dislocation slip and deformation twinning, leading to a reduction in ${\alpha }^{^{\prime}}$-martensite nucleation sites and an improvement in austenitic stability [193]. In the as-built laser-based AM steel, chemical heterogeneity most likely occurred as a result of solute micro-segregation caused by dendritic and/or cellular solidification, where solute atoms accumulated in dendrite interfaces or cell walls [95]; thus, the solute-poor dendrite/cell interior would have higher SFE due to the absence of SFE-lowering elements (i.e., Mn, Si, and Cr), which would have decreased the nucleation of deformation-induced martensite [200, 201].

According to the calculation of SFE based on the chemical compositions (SFE =  − 53 + 6.2 (wt% Ni) + 0.7 (wt% Cr) + 3.2 (wt% Mn) + 9.3 (wt% Mo) [200], the SFE of L-PBF 304L is estimated to be around 19.2 mJ/m² [147]. According to the general correlation between the SFE and the operation of plastic modes, it is expected that the TRIP effect shall dominate during the plastic deformation of L-PBF 304L. The L-PBF 304L [147] has an improved yield strength while maintaining the uniform elongation compared with the wrought [202] and annealed [203] counterpart, as shown in Fig. 11a and b. Specifically, the TRIP effect plays an important role at the deformation stages II and III (Fig. 11b). The progressive deformation at the deformation stage II is induced by the transformation of ${\alpha }^{^{\prime}}$-martensite, which impedes the dislocation slip, thus increasing the work hardening rate. The enhanced work hardening behavior in the deformation stage III is mainly contributed by the TRIP effect, which consumes other microstructures such as deformation twins, leaving the ${\alpha }^{^{\prime}}$-martensite as the main constituting phase after deformation (Fig. 11c), though some $\varepsilon$-martensite and austenite phases can be occasionally detected (Fig. 11d and e). Besides, the newly formed $\alpha \mathrm{^{\prime}}$-martensite contains high-density dislocations (Fig. 11c), which is ascribed to the displacive shear transformation of $\gamma \to \alpha \mathrm{^{\prime}}$[204]. Considering the almost completely martensitic transformation, the TRIP effect makes a significant contribution in accommodating the plastic deformation at high strain levels [147].

Figure 10

The exampled deformation twinning with other deformation mechanisms. a High-density dislocations stored at TBs [120], b-d The dislocation–twin interactions in a <011> -textured austenite grain [166], b plenty of multiple deformation twin lamellae (indicated by white arrows) and dislocations in the <011> -textured austenite grain, c the multiple nano-twin lamellae (indicated by white arrows) in b, d the atomic scale high-resolution TEM morphology of the nano-twin lamellae and SFs (indicated by black arrows). (e–h) a progressive deformation twinning mechanism [50], e twins nucleate from HAGBs and penetrate the cell walls and LAGBs at ~ 3% tensile strain, f an enlarged view of e, the LAGBs are indicated by black arrows, g the intersections of deformation twin with the dislocation cells and other twins at ~ 12% tensile strain, the HAGB is indicated by the white arrow, and the different twin sets are labeled as Set 1 and Set 2 by white dotted lines, h the intersections between cell walls and high-density deformation twins after ~ 36% tensile strain. The TB and cell orientations are marked by dark and white dashed lines, respectively

Since the SFE increases with temperature, the deformation mechanism transits from TRIP to dislocation slip when the deformation temperature exceeds the critical temperature, ${M}_{d30}$. Note that ${M}_{d30}$ is defined as the temperature of 50% austenite-to-martensite transformation occurring at a true strain of 30% [205]. It can be estimated by the empirical equation based on the alloy composition (in wt%) [147, 203, 206]:

$$M_{d30} = 413 - 462\left( {C + N} \right) - 9.2{\text{Si}} - 8.1{\text{Mn}} - 13.7{\text{Cr}} - 9.5{\text{Ni}} - 18.5{\text{Mo}}$$

(1)

The calculated ${M}_{d30}$ can be largely varied in laser-based AM 304L due to the different N content [203, 206]. Specifically, the TRIP effect can be effectively inhibited for the DED-processed 304L containing 0.09wt% N, with the ${M}_{d30}$ calculated to be around − 3.6 $^\circ{\rm C}$ [203], according to Eq. (1), but is substantially activated for the L-PBF fabricated 304L containing only 0.016wt% N, with ${M}_{d30}$ of 52.3 $^\circ{\rm C}$ [147].

Texture in austenitic steels may have a significant impact on the TRIP effect because the mechanical driving force for martensitic transformation is determined by the combined effects of stress state and crystallographic texture [207,208,209], affecting the rate of deformation-induced martensitic transformation concerning plastic strain. For example, the polycrystalline steels with a predominant brass texture (i.e., {110} <112>) or copper texture (i.e., {112} <111>) had lower driving forces than that with cube texture (i.e., {001} < 00>), calculated using a criterion derived from Schmid's law, and thus, required higher applied stresses for martensitic transformation [208].

The deformation-induced martensitic transformation in L-PBF 304L is studied under uniaxial tensile loading on the aspects of initial crystallographic texture and the evolution of deformation texture [179]. The extent of deformation-induced martensite depends on the crystallographic texture [180, 181, 191, 210, 211]. The texture and phase evolution during deformation of the as-built 304L is illustrated in Fig. 12 [179]. The crystallographic texture is almost random in an as-built state, but a strong <111> -texture evolves upon deformation with an increasing strain (Fig. 12a and b). This can be ascribed to a single slip of the activation of the {111} <110> slip system in FCC during tensile deformation [212]. Under the <111> -texture evolution, most grains are favorably oriented for partial dislocation splitting (Fig. 12c), with the leading partial dislocation having a greater Schmid factor than the trailing partial dislocation, allowing the extended SFs and $\varepsilon\varepsilon$-martensite formation, which facilitate the $\alpha \prime\alpha \prime$-martensite formation [179].

Figure 11

a The engineering stress–strain curves of the L-PBF [147], wrought [202] and annealed [203] 304L. b The true stress–work hardening rate–true strain curves of the L-PBF 304L [147]. (c-e) Deformed structures of the laser-based AM 304L adjacent to the fracture region [147]. c The high-density dislocations in the α^'-martensite are shown in the BF image. d The co-existence of austenite and α^'-martensite phases is shown in BF TEM and the inset SAED images. e The existence of ε-martensite is shown in the BF TEM image

The transformation also influences the austenite's deformation texture. Generally, in the absence of transformation, tensile and compressive deformation will produce separate textures in FCC metals, i.e., <111> -texture for tensile plastic deformation due to the straining aligning with the slip direction, while <220> -texture for compressive deformation because of the straining direction normal with the slip plane [212]. However, in both tension and compression of the L-PBF 17-4PH, a moderate <111> -texture is observed in the austenite aligning with the straining direction, which indicated that the texture evolution is influenced by preferential martensitic transformation, i.e., grains with (200) planes perpendicular to the straining direction transform first, but with <111> oriented grains showing higher resistance to martensitic transformation and remaining austenite to a large extent till the end of the straining, in the absence of the loading condition [211].

The ultra-high strength of ~ 1800 MPa combined with an excellent total elongation of ~ 25% is obtained in the tempered laser-based AM martensitic 420 [60]. The TRIP effect is found to be responsible for the excellent ductility, which is closely related to the volume fraction of the RA. The heat treatment increases the volume fraction of austenite from 3.5% up to 36%, which in turn provides the enhanced TRIP effect during deformation and results in an ultra-high strength combined with excellent ductility compared with the as-built and CM 420 (Table 2). The post-heat treatment can be omitted by designing the thermal history, i.e., the IHT associated with the laser-based AM process. The in situ diffusional processes and C partitioning initiated by IHT can trigger the subsequent austenite reversion or growth of RA, which leads to a high fraction of 57 ± 8 vol% RA in the DED part [89].

Table 2 Mechanical properties of the martensitic 420 in varied processing conditions

The RA in the range of 0 ~ 97% [18, 80, 92, 211, 214,215,216,217,218,219,220,221,222,223,224,225,226] is measured in laser-based AM 17–4 PH, which is a typical martensitic grade if treated by a conventional thermal–mechanical processing route. The different fraction of RA can be ascribed partly to the variation in the alloying element compositions in the feedstock powders, since the standard alloy specification [227] allows rather wide compositional ranges for Cr, Ni, and Cu, and do not prescribe the compositions of the O and N interstitial alloying elements. The differences in N compositions present in atomized powders lead to a wide range level of RA in the as-built 17–4 PH and thus significantly affect the mechanical properties [214, 222, 226].

High-density SF in nanoscale RA is observed in the microstructure of the L-PBF 17–4 PH [80]. Owing to the low SFE of the RA phase, the direct transformation of $\gamma \to {\alpha }^{^{\prime}}$ is favored, which leads to a continuous TRIP effect in the whole plastic regime, and thus excellent work hardening and mechanical properties (flow stress of 1300 MPa and total elongation of 28%). The microstructure of L-PBF 17–4 PH fabricated with N-atomized powder displays a high amount of RA, leading to discontinuous yielding driven by the deformation-induced martensitic transformation during uniaxial tensile deformation, and displays a 300 MPa higher strength than that fabricated with Ar-atomized powders [226]. Additionally, a relationship that [011] γ//[001]α between the RA and transformed martensite is observed which is following the Bain orientation relationship [228].

TRIP and TWIP effects in laser-based AM HMnS

The HMnS is typically alloyed with 0 ~ 1% C, 15 ~ 30% Mn, 0 ~ 3% Si, and 0 ~ 3% Al to achieve the SFE values in the range of 15 ~ 45 mJ/m² at room temperature [229]. Secondary alloying elements including Ti, Nb, N, Cu, Cr, and/or V are frequently adopted to enhance the strength of HMnS. The SFE in this range is accountable for activation of the TRIP and TWIP effects and thus HMnS show very high work-hardening capability that can be tailored within a wide range by adjusting the SFE and microstructure [230,231,232,233]. Additionally, the high work-hardening capability of HMnS due to activation of both TWIP and TRIP effects is desirable for applications in energy absorption and filigree lattice structures [95, 161]. Compared to the ingot- and strip-cast process, the rapid cooling associated with the laser-based AM process reduces elemental segregation during fabrication of HMnS, and thus, energy-intensive post-heat treatment can be circumvented, which makes HMnS a promising candidate for laser-based AM [91, 95, 96, 109, 160,161,162,163,164, 234].

The microstructure of HMnS fabricated by L-PBF is featured with a fine cellular dendritic structure, indicating a short segregation length and the absence of macrosegregation [161]. Such segregation behavior can be utilized to tune local SFE variation of HMnS. Additionally, the formation of $\alpha \prime\alpha \prime$-martensite in $\varepsilon\varepsilon$-martensite laths is also observed. The $\gamma \to \varepsilon\gamma \to \varepsilon$ transformation is induced due to the rapid cooling of L-PBF, and the $\alpha \prime\alpha \prime$-martensite is restricted in the $\varepsilon$-martensite laths following the $\gamma \to \varepsilon \to \alpha \prime\gamma \to \varepsilon \to \alpha \prime$ transformation reaction [235]. The HMnS fabricated by L-PBF shows a higher strength and lower elongation than the conventionally produced HMnS. Nevertheless, the HMnS produced by the above different strategies exhibit the same deformation characteristics, i.e., high work-hardening capacity [161]. The high-density dislocations and the presence of $\alpha \prime\alpha \prime$ and $\varepsilon\varepsilon$ martensite contribute to the higher strength, while the initial defects such as porosity and impurities impair the ductility.

The L-PBF-processed HMnS exhibits mechanical anisotropy [161]. Since the elongated grains formed along the build direction (BD) due to the rapid solidification during L-PBF, the effective grain size normal to BD ($0^\circ$ sample) is smaller than that along BD ($90^\circ$ sample), and thus, according to the Hall–Petch effect higher yield strength was obtained with $0^\circ$ sample than $90^\circ$ sample. Additionally, the varying contribution of deformation twinning to the working hardening can be estimated based on the equivalent critical uniaxial stress for initiation of deformation twinning. Such critical uniaxial stress can be calculated by using the Byun model ${\sigma }_{twin}=6.14\frac{{\gamma }_{sf}}{{b}_{SP}}$ [236], where ${\gamma }_{sf}$=10 mJ/m² is the SFE and ${b}_{SP}=14.5$ nm [237] is the Burgers vector of Shockley partial dislocation. The calculated ${\sigma }_{twin}$ of the $0^\circ$ sample is only slightly higher than the yield strength, which indicates the initiation of deformation twinning at lower plastic strain and thus a higher the accommodation of plastic strain by deformation twinning, compared with the $90^\circ$ sample.

In the DED-processed X30Mn23 microstructure, notable Mn/C enrichment at the interdendritic regions suggests the increased SFE in these regions and accordingly decreased SFE in the intradendritic regions [238]. The depleted Mn/C content in intradendritic regions results in the thermally induced $\varepsilon$-martensitic transformation during cooling (Fig. 13a and b) [239]. The TRIP and TWIP effects can be simultaneously activated during the deformation process depending on the region with varying SFE, i.e., numerous mechanical twins formed in the interdendritic regions and deformation-induced ε-martensite observed in the intradendritic regions after a fracture (Fig. 13c and d). Additionally, the Mn/C segregation is fully suppressed by alloying sufficient Al content. The alloy X30MnAl23-1 with an increased Al content demonstrates a low SFE in the pure TWIP range, exhibiting the most predictable deformation behavior and highest specific energy absorption upon compression. That is due to the contribution of the TWIP effect to accommodate plastic deformation and the suppressed formation of brittle α´- and ε-martensite (or TRIP effect).

Figure 12

The texture and phase evolution with increasing strain. The evolved texture at a 0.26 and b 0.42 true strain, the formation of SFs and ε-martensite at c 0.26 and d 0.42 true strain, e the details of the favorably oriented grain at the true strain of 0.26 for the ε-martensite formation [179].

The mechanical properties of laser-based AM HMnS can be tailored by controlling the deformation mechanisms through tuning the SFE via change of chemical compositions (Fig. 14a and b), e.g., by blending HMnS powder with Al [91, 95], or C [96]. The HMnS within the X30MnAl23-# (# indicates the fraction of Al, with # ≤ 2 wt%) system specific for AM applications is successfully designed using a methodology combining experimental and theoretical alloy screening strategy [95]. The alloy design by considering the SFE and DED is carried out to develop promising TWIP/TRIP HMnS in the X30MnAl23-# system. The initiation of the TWIP/TRIP effect in the laser-based AM fabricated HMnS can be successfully predicted by the SFE maps. (Fig. 14a)

Figure 13

a EBSD-IQ shows the thermally induced ε-martensite (yellow) formed besides austenite in DED X30Mn23, the white dotted line indicates the prior melt pool boundary. b The EDX mapping illustrates the Mn element enriched in the interdendritic regions. c EBSD-IQ and the overlaid binary EDX-Mn mappings of DED X30Mn23 at fracture. In the dark-shaded areas, Mn segregation is prevalent and thus characterized by higher SFE. Σ3 and ε-martensite are shown in blue and yellow, respectively [95]

Figure 14

a SFE-based deformation map [95], b SFE map associated with simultaneous change of C, Mn, and Al contents [238], c the X30MnAl21-x equilibrium phase diagram, and d Scheil–Gulliver simulation of the Al-0 to Al-5 solidification sequence and phase-dependent C, Al, and Mn distribution in Al-5 solidification [91]

The addition of the strongly bcc-stabilizing element such as Al to X30Mn21 steel is found to significantly affect the texture and microstructure evolution during the laser-based AM process and thus the work hardening behavior [91]. By increasing the Al content to 4 wt%, the solidification mode transits from a fully fcc mode to a bcc–fcc mode, resulting in substantial texture randomization and grain refinement [91]. During primary $\delta$-bcc solidification, the C and Mn elements in the remaining melt between the $\delta$-bcc cells/dendrites are enriched continuously due to elemental segregation (Fig. 14d), which promotes the stabilization of the fcc phase and causes secondary fcc solidification. The bcc–fcc solidification mechanism combined with the bcc-to-fcc solid-state transformation facilitates nucleation processes and suppresses epitaxial growth, resulting in a randomly distributed texture. The TRIP-dominated alloy (Al-0) demonstrates a high work-hardening, whereas fully austenitic alloys (Al-1, Al-2, Al-3) show large plasticity dominated by the TWIP effect. Further, Al addition (Al-4, Al-5) promotes a duplex bcc–fcc microstructure and a strongly increased yield strength and reduced work hardening owing to the refinement of grain and suppression of TWIP effect. Moreover, the suppressed TRIP effect and refined grain in the high Al HMnS (Al-4, Al-5) all contribute to the increased ductility and energy absorption capacity. Therefore, the mechanical properties of laser-based AM components can be tailored through the tuning of Al content as it governs SFE, the solidification kinetics, as well as the constituting phases.

The HMnS within the X(30 + #)Mn22 (# indicates the blended C fraction) system specific for AM applications is investigated by powder blends in [96]. The C powder is used to precisely adjust the C fraction in solid solution. The high C-content can fully dissolve into the austenite matrix without the formation of carbide, enabling flexible adjustment of microstructure and mechanical properties. The C-content variation enables a tunable SFE, and the controlled solid solution strengthening and thus tailored yield strength and deformation mechanisms, which result in a controllable work-hardening behavior. The SFE increases with an increase of C-content, resulting in the shift from TRIP to TWIP effect, leading to a decreased work-hardening rate. The optimum work hardening behavior appears at X90, which can be ascribed to the fully suppressed TRIP effect and the pronounced operation of the TWIP behavior. A further increase in C-content is detrimental owing to the reduction of the TWIP effect. Moreover, the addition of C can also promote increased yield strength and pronounced dynamic strain aging, which is typical for Al-free HMnS [240]. Therefore, the optimization of the different deformation mechanisms can be achieved by careful tuning chemical compositions and printing strategies, developing high-performance laser-based AM steels for broad industrial applications.

Conclusions

Steels are mankind’s most widely used structural material in various industries, which demonstrates the diverse microstructure due to the complex phase transformation, leading to intriguing deformation behavior that incorporates almost all deformation mechanisms. The understanding of the science of deformation of laser-based AM metallic materials serves as the basis for the development of high-performance structural components for engineering applications. The present work thoroughly reviews the deformation mechanisms (i.e. dislocation slip, deformation twinning, and deformation-induced martensitic transformation) with various microstructures of steels prepared by laser-based AM, and the following conclusions are derived:

• 1. The laser-based AM steels with inherent hierarchical structures in different length scales including the HAGB, LAGB, dislocation cell structure, nanoscale precipitation, and inclusion demonstrate dominated dislocation slip plasticity, providing back stress hardening, forest dislocation hardening, and precipitation hardening to increase strength without sacrificing ductility. These microstructure characteristics can be successfully engineered through optimum laser-based AM process, i.e., the dislocation cell structure is very sensitive to the thermal history, the oxide inclusion is very sensitive to the processing atmosphere and also can be fabricated by blending powders, the precipitate has a low mismatch with the matrix (such as NiAl-martensite) can be triggered directed by the IHT during deposition.

• 2. The TWIP behavior of laser-based AM austenitic (or austenite-containing) steel should be influenced by the SFE and the crystallographic texture significantly, and therefore, the TWIP effect can be tailored by elaborately engineering the texture and SFE. The texture-related SF/twin formation can be rationalized by Schmid factor associated leading and trailing partial dislocations, and the orientation-dependent TWIP effect can be explained by the Taylor factor. Owing to the directed solidification behavior associated with laser-based AM, the texture can be successfully engineered in laser-based AM steel.

• 3. The TRIP behavior of laser-based AM steel is closely related to the fraction of RA, the SFE value, and the microstructure characteristic. The fraction of the RA can be influenced by the N content since the N can influence the stability of austenite. Additionally, the texture can influence the martensitic transformation, which can be explained by the Schmid factor theory. Consequently, the TRIP-related deformation mechanism in laser-based AM steel can be engineered by elaborately tuning the SFE value and the microstructure feature, the texture specifically.

• 4. The HMnS is the promising candidate for laser-based AM constructional components for which can tailor the dislocation slip-, TRIP- and TWIP-related deformation mechanisms, thus enabling a large regime of strain hardening, through tuning the SFE and microstructure. The SFE in HMnS can successfully be tailored by powder blends, i.e., by adding C/Al into the HMnS powder, and thus ideally tailoring the deformation mechanisms and mechanical response. Moreover, the solidification model can also be tuned by modification of the SFE, i.e., by adding sufficient Al, the solidification model transited from fcc model to fcc–bcc mode, which affects both the microstructure and texture evolution.

• 5. The powder blends adhesion to AM provides a promising approach for adjustment of the composition and chemistry of the alloy flexibly, for example, by deliberately adding oxide particles to fabricate oxide dispersion strengthened (ODS) steel and by deliberately adjusting the Al- or C-content in high-manganese austenitic steel to tune the SFE and thus enables excellent mechanical properties.