Solidification behaviour of austenitic stainless steels during welding and directed energy deposition

Abstract

The effect of cooling rate on the solidification behaviour of austenitic stainless steels during high energy density welding and directed energy deposition (DED) has been reviewed. Precedent studies demonstrated the confinement of austenite–ferrite duplex region and the susceptibility of specific alloy compositions on the Schaeffler diagram to alteration of solidification mode at high cooling rates during the high energy density welding. Meanwhile, mitigated cooling conditions have dominated during the DED process. The instances of microstructural fluctuations owing to cooling rate variation have been compiled. The incorporation of DED steels into the implicated Schaeffler diagrams demonstrated reliable predictions at high cooling rates. The printability of austenitic stainless steels during the DED process has been discussed in terms of solidification cracking susceptibility.

Keywords

Welding additive manufacturing directed energy deposition stainless steels solidification Schaeffler diagram microstructure

Introduction

Metal additive manufacturing (MAM) realises intricate microstructure and property engineering [1 –12]. According to ISO/ASTM52900-15 [13], MAM has been classified into two categories: (i) directed energy deposition (DED) and (ii) powder bed fusion (PBF). A laser or electron beam coupled with a wire or powder feeding module generates a melt pool in the DED while fusing a powder bed in the PBF selectively. PBF has been reputed for selective laser melting (SLM) and selective laser sintering (SLS) [14]. In contrast, DED has been introduced as laser engineered net shaping (LENS), laser metal deposition (LMD), direct metal deposition (DMD) and laser powder fusion (LPF) [15,16].

A considerable ratio of MAM concerns stainless steels [17 –23]. Solidification behaviour has a significant influence on the printability and properties of additively manufactured stainless steels and has frequently been interpreted in terms of the primary solidifying phases, consequent evolutions, and the attendant phase compositions. Saboori et al. [24] reviewed the effect of solidification parameters on the microstructure and mechanical properties of type 316L 1 stainless steels produced by the DED process. More recently, Astafurov and Astafurova [25] reviewed solidification-driven variations in the phase composition of austenitic stainless steels produced by various MAM processes. A common feature of these processes is a high cooling rate and its variation during additive manufacturing, which manipulates solidification behaviour and results in the variation of solidification cells and dendrites size, ferrite content and microsegregation of alloying elements generally. Therefore, additively manufactured stainless steels are accosted by intriguing manipulation of solidification behaviour at high cooling rates and inevitable microstructural inhomogeneities owing to process-induced and inherent cooling rate variations.

Welding literature progressively describes the solidification behaviour of stainless steels at high cooling rates [26 –28], the phase composition effects on the properties of weld metals, such as the beneficial effect of delta-ferrite on the hot cracking susceptibility [29 –32], and its deleterious effects on the high-temperature microstructure stability, mechanical properties and corrosion susceptibility [33 –37]. Considering many similarities between welding and MAM processes, welding literature has already devised a charter path for MAM progress [38,39]. Regarding the solidification behaviour, Smith et al. [40] applied welding models for studying the solidification microstructure of 304L and 316L steels produced by the DED process. Hwa et al. [41] reported similar cooling rate variation effects in the DED and autogenous laser welding of 316L steel. Nevertheless, an integrated review of the cooling rate effects on the solidification behaviour of stainless steels during welding and additive manufacturing has not been reported yet.

The cooling rate of stainless steels during high energy density welding and additive manufacturing has been estimated either by computation [42,43] or experimentally 2 by measuring cellular or dendritic arms spacing in the solidified microstructure [44]. In general, the cooling rate of austenitic stainless steels varies between 10⁺² and 10⁺⁵ K s⁻¹ during autogenous, high energy density welding and the DED [44 –46]. Based on the typical range of cooling rates, this paper reviews the cooling rate effects on the solidification behaviour of austenitic stainless steels during the high energy density autogenous welding and the laser-based DED.

Solidification mode of stainless steels

The solidification of stainless steels has been categorised into four modes according to the type of the primary phases solidifying from the melt:

mode A (primary austenite with no ferrite solidified from the melt)

mode AF (primary austenite with ferrite solidified secondarily from the melt)

mode FA (primary ferrite with austenite solidified secondarily from the melt)

mode F (primary ferrite with no austenite solidified from the melt).

Suutala et al. [47] first categorised the solidification mode of stainless steel GTA weld metals in three types: A (austenite and austenite–ferrite), B (ferrite–austenite) and C (ferrite). Then, classification into four modes (A, AF, FA, F) was introduced [48]. Katayama et al. [49] classified it into five modes (A, AE, E, FE, F), E denoting the eutectic ferrite–austenite solidification. Elmer [50] referred to the five modes as A, AF, E, FA and F. However, the E mode has not been considered lately while four modes (A, AF, FA, F) 3 have been referred to more extensively. Schematics illustrating the microstructure evolution in accordance with these modes have been presented frequently [28,47,51 –54]. The solidification process has widely been studied regarding the Fe–Cr–Ni ternary system [55,56], where pseudo-binary vertical sections of the ternary phase diagram at constant Fe predicted solidification mode and subsequent phase transformations reasonably. Figure 1 shows the solidification modes on the vertical section of the Fe–Cr–Ni phase diagram at 70-wt% Fe [57]. Commercial stainless steels often span the peritectic–eutectic triangle on the pseudo-binary diagram, and alloys of low and high Cr-equivalent (Cr_eq)/Ni-equivalent (Ni_eq) ratios on either side of isoplete solidify with the formation of austenite and ferrite as the primary solidifying phase, respectively. The partition ratio between austenite and melt is about 0.75 for Cr and 0.95 for Ni thus, melt enriched by Cr and marginally by Ni, which favours ferrite formation at the later stage of the AF mode. However, the partition ratios are displaced in the FA mode mutually; thus Ni enrichment of melt stimulates austenite formation in the later stage of solidification [52].

Figure 1

The solidification mode of stainless steels on the pseudo-binary vertical section of the Fe–Cr–Ni phase diagram at 70-wt% Fe, calculated using Thermocalc. Adopted from [57] under the terms and conditions of the Creative Commons Attribution (CC BY) license, courtesy of MDPI.

The solidified microstructure of stainless steels encounters subsequent transformation of ferrite into austenite during cooling in the solid state after solidification, depending on the phase composition and cooling condition. Hence, there is no further transformation for fully austenitic microstructure solidified in A mode, but the transformation of ferrite occurs in alloys solidified with AF mode [58 –60]. Alloys solidified in FA and F modes contain higher amounts of ferrite thus their transformations occur diversely, which yield microstructures consisting of austenite and distinct types of ferrite, e.g. globular, vermicular, skeletal, lacy, lathy and acicular [61 –66]. Transformation of ferrite to Widmanstätten austenite (WA) and massive austenite has also been reported in stainless steel weld metals [67 –69].

The solidification mode and microstructure of stainless steel weld metals are conventionally ascertained using the Schaeffler constitution diagram based on Cr_eq and Ni_eq 4 [70]. Figure 2 shows regions of the different solidification modes and the characteristic microstructures of stainless steels weld metals on the Schaeffler diagram [49]. Other constitution diagrams with modified equations have also been reported [71,72]. According to Suutala et al. [47], the solidification mode and microstructure of stainless steel weld metals can also be ascertained from Cr_eq/Ni_eq ratio 5 : (i) A/AF mode at Cr_eq/Ni_eq < 1.48, resulting in austenitic microstructure with ferrite at boundaries, if any, (ii) FA mode at 1.48 < Cr_eq/Ni_eq < 1.95, resulting in austenite with vermicular and lathy ferrite (LF) residue, (iii) F mode at Cr_eq/Ni_eq > 1.95, resulting in austenite with lathy structure. The critical ratio for the transition from AF to FA (1.48) increases slightly by increasing the cooling rate during arc welding [48].

Figure 2

Regions of different solidification modes and characteristic microstructures of stainless steels weld metals are presented on the Schaeffler diagram. Solidification mode abbreviations read type A (= mode A): single-phase austenite, type AE (= mode AF): primary austenite with eutectic ferrite–austenite, type FE (= mode FA): primary ferrite with eutectic ferrite–austenite, type F (= mode F): single-phase ferrite. Characteristic microstructure abbreviations (encircled) read FA (fully austenite), IF (intercellular eutectic ferrite), EF (eutectic ferrite), VF (vermicular (skeletal) ferrite), LF, AF (acicular ferrite), WA, FF (fully ferrite). Adopted from [49] with permission from the Welding Research Institute of Osaka University.

Behaviour of alloy clouds during welding

Katayama and Matsunawa [73] studied the solidification behaviour of more than 30 kinds of Fe–Cr–Ni alloys and commercial stainless steels, e.g. 310S, 304, 316, 347, 321 and 309S, with a variety of equivalencies and the cooling rates during (i) autogenous pulsed laser welding with an energy input of 5–37 J pulse⁻¹ (Nd:YAG: 200 W, 1, 2 and 3.6 ms pulse durations, in air or Ar atmosphere), (ii) autogenous laser welding (continuous wave (CW) CO₂: 500–2000 W, 8.33–250 mm s⁻¹) and (iii) autogenous GTA welding (100 A, 12.5 V, 2.5 mm s⁻¹). Pulsed laser welds of 310S steel (Cr_eq = 25.94, Ni_eq = 22.11, Cr_eq/Ni_eq = 1.17) presented fully austenitic cellular microstructure with PCAS decreasing by increasing cooling rate. Based on the measured PCAS (0.5–2.2 µm), the cooling rate was estimated to be about 5 × 10⁴–5 × 10⁵ K s⁻¹. Pulsed laser welds of 304 steel (Cr_eq = 18.84, Ni_eq = 11.14, Cr_eq/Ni_eq = 1.69) presented fully austenitic cellular microstructure, in contrast to a GTA weld metal which presented 5 vol.-% skeletal ferrite. Ferrite content decreased from 4.0 vol.-% to 0.4 vol.-% by increasing the cooling rate from 3 × 10³ to 5 × 10⁴ K s⁻¹ during CW CO₂ laser welding. Figure 3 demonstrates (a) solidification microstructures based on metallography and X-ray diffraction (XRD) and (b) presumed solidification modes at a cooling rate of 2 × 10⁵ K s⁻¹ for pulsed laser-welded alloys. Alloys showing 8 vol.-% ferrite on the Schaeffler diagram became fully austenitic, while alloys of higher than 20–30 vol.-% ferrite became fully ferritic (see red lines). Hence austenite–ferrite (duplex) region with FA and especially AF modes was narrowed. Alteration of FA mode to A/AF mode by increasing cooling rate from 220 to 1550 K s⁻¹ during GTA was demonstrated for an alloy of Cr_eq = 20.6, Ni_eq = 14.8, Cr_eq/Ni_eq = 1.39 and so for 304 steel [73].

Figure 3

Fe–Ni–Cr alloys and commercial stainless steels studied by Katayama and Matsunawa are presented on the Schaeffler diagram: (a) identified microstructures and (b) presumed solidification mode at a cooling rate of 2 × 10⁵ K s⁻¹. Fine and coarse dashed lines represent the Schaeffler ferrite contours and solidification mode regions for GTA, respectively. Adopted from [73] with permission from the Laser Institute of America.

Lippold [74] and Lienet and Lippold [75] reported the solidification behaviour of 35 kinds of commercial stainless steels in the Cr_eq/Ni_eq range from 1.41 to 2.08, autogenously welded by pulsed laser of 8–17.5 J pulse⁻¹ energy input (Nd:YAG: 80 and 140 W, 8 Hz, 8 ms pulse duration, 1.6 mm s⁻¹, in He gas atmosphere). An average cooling rate of 3 × 10⁵ K s⁻¹ was determined based on the PDAS measurement. Figure 4 shows the solidification mode of the studied steels on the Schaeffler diagram. Steels consisting of less than about 3 vol.-% and more than about 7 vol.-% ferrite on the Schaeffler diagram solidified in A and F modes, respectively. Many of the steels presented mixed modes but in a narrowed region. Based on the Hammar–Svensson (HS) equivalencies 6 , a transition in the solidification mode from AF to FA was found in the Cr_eq/Ni_eq range from 1.59 to 1.69, which indicates a further increase of the critical Cr_eq/Ni_eq for AF to FA transition in pulsed laser welding in comparison to the value of 1.55 determined for GTA welding [48].

Figure 4

Solidification mode of commercial stainless steels studied by Lippold and Lienet is presented on the Schaeffler diagram. Adopted from tabulated data in [74,75] with permission from the American Welding Society.

Berjemo et al. [76] applied autogenous laser welding with energy inputs of 75 and 110 J mm⁻¹ (Ytterbium fiber CW: 750 W–10 mm s⁻¹, 2200 W–20 mm s⁻¹) on 17 kinds of stainless steels in the Cr_eq/Ni_eq range from 1.22 to 1.88 (HS eqs: 1.23–2.04) and Cr_eq + Ni_eq of about 35% and 45%. Table 1 shows cooling rates determined by the SDAS measurement for 304 steel. The solidification mode of the studied weld metals is shown in Figure 5. Like the aforementioned laser welding, single-phase austenite and single-phase ferrite regions developed at the expense of duplex microstructures.

Figure 5

Solidification mode of commercial stainless steels studied by Berjemo et al. is presented on the Schaeffler diagram. Adopted from tabulated data in [76] with permission from Elsevier.

Table 1

Cooling rates estimated for laser-welded 304 steels based on SDAS measurement. Adopted from [76] with permission from Elsevier.

Energy input (J mm⁻¹)	SDAS (µm)	Cooling rate (K s⁻¹)
75	1.76	13056
110	2.43	4125

David et al. [77] studied cooling rate variation effects on the microstructure of 15 kinds of commercial stainless steels during autogenous pulsed laser welding with different energy inputs (Nd:YAG: 110–120 W, 190–200 W, 1 ms pulse duration, 2.1–42.3 mm s⁻¹, in Ar atmosphere). Figure 6 shows the studied steels on the Schaeffler diagram. For each type, two heats of different compositions (A and B) within the standard range were studied. In types 304, 308 and 347 steels, ferrite decreased gradually with increasing cooling rate to approach a fully austenitic microstructure at the highest cooling rate. In the type 309 steels, ferrite initially increased then decreased with no transition to a fully austenitic microstructure. In the type 312 steels, ferrite increased with increasing cooling rate to a fully ferritic microstructure at the highest cooling rate. Different heats of 310 steel solidified in A mode whatever the cooling rate, but significant microstructural variations occurred among 316 welds. Implications for Scheffler diagrams were presented at high cooling rates.

Figure 6

Commercial stainless steels studied by David et al. are presented on the Schaeffler diagram. Adopted from [77] with permission from the American Welding Society.

Behaviour of selected commercial stainless steels

Vitek et al. [78] applied autogenous laser welding (CW CO₂: 9000 W, 13, at 25 and 63 mm s⁻¹, in He gas atmosphere) on three types of commercial stainless steels (308, 310 and 312). Figure 7 shows the approximate compositions of these steels on a vertical section of the Fe–Cr–Ni phase diagram. The 312 (Cr_eq = 30.83, Ni_eq = 9.85, Cr_eq/Ni_eq = 3.13) and 308 (Cr_eq = 21.66, Ni_eq = 13.21, Cr_eq/Ni_eq = 1.64) steels solidify with F and FA modes, respectively, but acquire fully austenitic and duplex ferritic–austenitic microstructures according to phase diagram, while presented duplex ferritic–austenitic microstructures with different ferrite contents after conventional arc welding. The 310 steel (Cr_eq = 27.95, Ni_eq = 25.62, Cr_eq/Ni_eq = 1.09) solidifies in A mode with fully austenitic microstructure at room temperature according to the phase diagram, which retained a fully austenitic microstructure after conventional arc welding. The cooling rate was assumed to vary with scan speed and the intrinsic cooling condition. Higher scan speed provides a higher cooling rate and vice versa. The melt cools at the highest rate in the root and the lowest rate in the crown. The 312 steel showed a fully ferritic microstructure at the highest cooling rate and a duplex ferritic–austenitic microstructure at the lowest cooling rate. It was concluded that the solidification mode is promoted from FA to F at the highest cooling rate, and the transformation of ferrite to austenite is suppressed subsequently. The solidification mode of 308 steel was altered at the highest cooling rate as rationalised by a fully austenitic microstructure at the highest cooling rate and a duplex austenitic–ferritic microstructure at the lowest cooling rate. Partitioning of Cr and Ni decreased remarkably by increasing the cooling rate, which rendered the partitionless crystallisation of austenite at the highest cooling rate. The 310 steel showed a fully austenitic microstructure at all cooling rates.

Figure 7

The approximate composition of the commercial stainless steels studied by Vitek et al. Adopted from [78] with permission from Springer.

Zacharia et al. [79] applied autogenous laser welding (Nd:YAG: 120 and 200 W, 2.1–42.3 mm s⁻¹, in Ar atmosphere) on five types of commercial stainless steels as shown in Table 2. Calculated cooling rates at the solidification temperature varied, with scan speed, in the range of 0.56–3.2 × 10⁵ and 1.0–8.8 × 10⁵ K s⁻¹ for high and low welding powers, respectively. All alloy types represented duplex microstructures after arc welding, while laser welding altered solidification microstructures into single-phase austenite or ferrite.

Table 2

Commercial stainless steels studied by Zacharia et al. Adopted from [79] with permission from Springer.

				Microstructure
Steel	Cr_eq	Ni_eq	Cr_eq/Ni_eq	2.1 mm s⁻¹	42.3 mm s⁻¹
310	26.40	20.53	1.29	Austenite	Austenite
308	20.25	13.33	1.52	Duplex	Austenite
316	22.27	14.74	1.51	Duplex	Austenite
304	19.24	11.07	1.74	Duplex	Austenite
312	29.72	12.96	2.29	Duplex	Ferrite

Molian [80] applied autogenous laser welding (CW CO₂: 1200 W, 13, 21.2 mm s⁻¹, in He–Ar gas atmosphere) on a 304 steel. A fully austenitic structure with high dislocation density and negligible microsegregation was found. Zambon and Bonollo [81] reported laser welding (CW CO₂:7000 W, 38.3 mm s⁻¹) of 304 (Cr_eq = 19.24, Ni_eq = 10.82, Cr_eq/Ni_eq = 1.78) and 316 (Cr_eq = 19.00, Ni_eq = 12.56, Cr_eq/Ni_eq = 1.51) steels. Based on the observation of austenite with 1–2 vol.-% ferrite, solidification mode of 304 and 316 steels was, respectively, assumed FA and AF a priori. Akgun and Inal [82] carried out autogenous laser welding (CW CO₂: 5000 W, 7 mm s⁻¹, in Ar atmosphere) on 304L steel (Cr_eq = 19.48, Ni_eq = 9.51, Cr_eq/Ni_eq = 2.05). A cooling rate in the 10⁴–10⁵ K s⁻¹ range was determined based on the SDAS measurement. XRD identified austenite with ferrite in the range of 4.3–9.9 vol.-%, increasing with increasing the distance from the surface. Cooling rate variation was assumed to alter ferrite content, but the addition of nitrogen to shielding gas was unlikely to do so. Based on microstructure and thermodynamic assessment, the FA mode was rationalised. El-Batahgy [83] conducted autogenous laser welding with energy input in the range of 40–480 J mm⁻¹ (CW CO₂, 2000–5000 W, 8.3-50 mm s⁻¹) on 304L(Cr_eq = 18.93, Ni_eq = 10.1, Cr_eq/Ni_eq = 1.87), 316L (Cr_eq = 19.35, Ni_eq = 12.05, Cr_eq/Ni_eq = 1.61) and 347 (Cr_eq = 18.73, Ni_eq = 11.80, Cr_eq/Ni_eq = 1.59) steels. The microstructure of all laser welds was austenite with a few percent of ferrite, resulting from primary ferrite mode presumably.

Farid and Molian [84] compared high brightness photolytic iodine laser (PIL) (spot size 0.02 mm, 10–20 W) with CW Nd:YAG (0.4 mm spot size, 10–20 W) and CW CO₂ (0.25 mm spot size, 200–600 W) lasers for lap joints of thin sheets (0.1 mm thickness) of 316 steel (Cr_eq/Ni_eq = 1.33). High brightness PIL provided the highest cooling rate. Table 3 shows the microstructural characteristics determined for different weld metals. A higher cooling rate stimulated single-phase austenite with finer cell size in the PIL welds.

Table 3

Characteristics of 316L thin sheets welded by various laser-welding processes. Adopted from [84] with permission from Kluwer Academic Publishers.

Feature	PIL	CW:Nd:YAG	CW-CO2
Cell size (µm)	0.05	4	0.6
Microstructure	fully austenite	austenite + 3–4 vol.-% ferrite	austenite + 3–4 vol.-% ferrite
Mode	A	AF	AF

McPherson et al. [85] reported microstructure of 316LN steel (Cr_eq = 20.99, Ni_eq = 12.52, Cr_eq/Ni_eq =1.68), autogenously laser welded (Nd:YAG: 7.6 mm, 3500 W, 10 mm s⁻¹, in He atmosphere). A mixture of cellular austenite with ferrite and austenite with LF was identified, in addition to the nonequilibrium Cr₂N phase formed owing to a high cooling rate. Celen et al. [86] applied pulsed laser welding (Nd:YAG: 769–200 W, 5.8–7.7 mm s⁻¹) on thin sheets of 304 steel (Cr_eq = 21.13, Ni_eq = 15.40, Cr_eq/Ni_eq = 1.37), and reported fully austenitic microstructure. Kell et al. [87] applied electron backscattering diffraction (EBSD) to study the microstructure of 316L sheets, autogenously laser welded (Coherent CO₂: 125–1000 W, 0.017–0.1 mm s⁻¹, under compressed air flow). AF solidification mode was assumed a priori. EBSD revealed a fully austenitic microstructure and a single crystallographic orientation of a cell colony. Energy-dispersive spectroscopy (EDS) indicated Cr, Mo, Mn, S and P microsegregation at cell boundaries but Ni, O and Si were homogeneously distributed. However, ferrite was identified when a uniformly distributed, square beam was applied. Late studies on dissimilar alloys and high-power lasers applied for welding the thick plates, beyond the scope of this paper, are no longer considered.

Behaviour of Fe–Cr–Ni model alloys

Elmer et al. [88] studied the cooling rate effects on the solidification mode of seven high-purity Fe–Cr–Ni alloys, autogenously welded by electron beam (2000 W, 6.3–5000 mm s⁻¹). Figure 8 shows the approximate compositions of their alloys in a vertical section of the Fe–Cr–Ni phase diagram. A high cooling rate stimulated A mode in the hypereutectic alloys (1–3) with Cr_eq/Ni_eq in the range from 1.15 to 1.51, and F mode in the hypoeutectic alloys (5–7) with Cr_eq/Ni_eq in the range from 1.84 to 2.18. Solidification of the nearly eutectic alloy (4) with a Cr_eq/Ni_eq = 1.60 evolved from FA mode to AF mode by increasing the cooling rate. The development of solidification mode from competitive epitaxial growths of austenite and ferrite at the melt periphery has been explained. The hypereutectic alloy (1) with a Cr_eq/Ni_eq = 1.15 solidified with the epitaxial growth of austenite at both low and high cooling rates, hence realising A mode always. Solidification of the hypoeutectic alloy (7) with Cr_eq/Ni_eq = 2.18 initiated with the epitaxial growth of both ferrite and austenite at the melt periphery, but ferrite dominated immediately. The nearly eutectic alloy (4) provided a duplex (ferrite–austenite) microstructure at the melt periphery, and solidification initiated with the epitaxial growth of both austenite and ferrite. However, the preliminary AF cells evolved into the FA cells within a few tens of micrometers at the lower cooling rate while the AF cells dominated through at a higher cooling rate. Consistent with those transitions, Inoue et al. [89] reported that the preliminary epitaxial growth of austenite is followed by ferrite nucleation owing to Cr enrichment in the melt hereinafter, ferrite continues to grow as a primary solidifying phase. A reverse FA to AF transition occurred by accelerated growth of austenite continuously.

Figure 8

The approximate composition of the alloys studied by Elmer et al. with Cr_eq/Ni_eq ratios of 1.15 (1), 1.39 (2), 1.51 (3), 1.60 (4), 1.84 (5), 2.01 (6) and 2.18 (7). Adopted from [88] with permission from Springer.

Brooks et al. [90] reported cooling rate effects on the solidification mode and microsegregation of alloying elements in the autogenous high energy density (laser and electron beams) weld overlays on GTA weld metal of three Fe–Cr–Ni alloys. Figure 9 represents the solidification modes and microstructures of the laser weld overlays and GTA weld metals. Both laser and electron beams and the different scan speeds produced similar microstructures. All studied alloys solidified in FA mode during GTA welding, consisting of austenite and skeletal ferrite. A high cooling rate altered the solidification mode for an alloy of Cr_eq/Ni_eq 7 = 1.51 from FA mode to A mode in the high energy density weld (HEDW) metals, representing a cellular austenitic structure with strong Cr segregation at cell boundaries. There was no alteration of FA mode for an alloy of Cr_eq/Ni_eq = 1.68, but refined skeletal ferrite was retained at the cell cores with strong Cr and Ni microsegregation. However, solidification mode promoted to F mode in an alloy of Cr_eq/Ni_eq = 1.85, consisting of a relatively homogeneous austenitic microstructure which was assumed to result from the massive transformation of ferrite to austenite in the solid state.

Figure 9

Microstructure and solidification mode of the GTA and HEDW overlays: (a) Fe–21.1Cr–14.2Ni alloy (Cr_eq/Ni_eq = 1.51); (b) Fe–21.9Cr–13.0Ni alloy (Cr_eq/Ni_eq = 1.68); (c) Fe–22.8Cr–11.9Ni alloy (Cr_eq/Ni_eq = 1.85). Si = 0.01; C = 0.006–0.016; N = 0.001–0.004. Reproduced from [90] with permission from Springer.

Selection of solidification mode

Siefert and David [28] categorised three models for modification of stainless steel weld metals solidification at high cooling rates, being quoted by theme: (i) partitionless crystallisation from undercooled melt associated with the kinetic predominance of austenite; (ii) dendrite tip undercooling down to the extended austenite liquidus temperature at rapid solidification and (iii) massive transformation of primary ferrite to austenite in the solid state during cooling after solidification. Vitek et al. [78] first proposed that a high undercooling realises partitionless crystallisation of both ferrite and austenite, based on comparable T₀ temperatures of austenite and ferrite in near-eutectic stainless steels with FA mode. Then, austenite predominates kinetically in the subsequent growth stage. A similar argument was proposed by Akgun and Inal [82], which applies a thermodynamic model to calculate T₀ temperatures and indicates thermodynamic stability of ferrite over austenite in a 304L steel solidifying with FA mode. Consistent with the second model, Brooks et al. [90] attributed austenite predominance to dendrite tip undercooling, comparing ferrite liquidus temperature with extended austenite liquidus temperature for an undercooled near-eutectic alloy solidifying with FA mode. Taking the melt undercooling into account, Figure 10 shows further explanation. A near-eutectic alloy (C_I) solidifies with FA mode under near-equilibrium condition, whereas sufficient undercooling at the dendrite tip realises austenite crystallisation before ferrite, based on the undercooled liquidus (dotted lines). This realises the alteration of solidification mode from FA to AF mode. In contrast, alloy C_II solidifies with ferrite as the primarily solidifying phase at the same condition.

Figure 10

Schematic pseudo-binary vertical section of Fe–Cr–Ni phase diagram at constant Fe with undercooled austenite (γ_L) and ferrite (δ_L) liquidus lines. Adopted from [90,91] with permission from the Institute of Metals, the ASM International, Springer and Materials Research Society.

The competitive formation of austenite and ferrite has experimentally been studied in the levitated and chill-cast Fe–Cr–Ni alloy droplets solidified at high cooling rates. Based on ferrite prevalence in the smaller and more rapidly cooled particles, preferred homogeneous nucleation of ferrite at high undercooling was assumed for alloys that solidify with austenite as the primary phase under near-equilibrium conditions [92,93]. This fact was confirmed by the recalescence phenomenon [52,94,95] and thermodynamic assessments subsequently [96 –99]. Mizukami et al. [100] measured melt undercooling at the surface of 304 steel droplets (Cr_eq/Ni_eq = 1.81) chill-cast at various cooling rates. The maximum initial undercooling increased with increasing the cooling rate. The steel solidified with primary ferrite mode under near-equilibrium conditions. At a high cooling rate, however, solidification initiated with the formation of cellular austenite at the surface and evolved to dendritic ferrite at the interior subsequently. The phase selection was rationalised by computing dendrite tip temperature for each phase, assuming the steady-state dendritic growth into the undercooled melt. With a great number of nucleation sites and a criterion that the one observed has a higher tip temperature, a kinetical predominance of austenite at a high cooling rate was rationalised. Similarly, Lippold [74] justified that a high solidification rate inherent to weld metal decreases dendrite tip radius and tip temperature, thus altering undercooling at the dendrite tip based on the dendrite growth theory. The dendrite tip temperature of ferrite decreases more rapidly than austenite with an increasing solidification rate. Assuming that the preferred primary phase will be that with a higher dendrite tip temperature as a phase selection criterion, predominance of austenite over ferrite above a critical solidification (cooling) rate was interpreted. The value of the critical solidification (cooling) rate for transition to primary austenite increases with increasing Cr_eq/Ni_eq ratio.

The dendrite tip temperature has been computed as an interface response function according to dendrite growth theory [101]. This function has been applied for the construction of solidification microstructure selection maps in laser welds generally [102], and specifically for the prediction of the transition from ferrite to austenite in stainless steels laser welds, as well as for the construction of their solidification microstructure selection maps at various solidification (cooling) rates [103 –105]. An increase in the solidification rate is shown to promote the transition to primary austenite at a higher Cr_eq/Ni_eq ratio.

Solidification mode of stainless steels produced by the DED

Table 4 summarises the solidification mode and resultant microstructure of some austenitic stainless steels produced by the DED process under different conditions [40,106 –122]. In these studies, a precise determination of microstructure constituents, e.g. austenite, ferrite and second-phase precipitates, has been well appreciated. However, solidification mode has been chiefly ascertained based on the alloy composition on the Schaeffler diagrams a priori. In a few studies, the solidification mode has been adequately recognised based on the characteristic concentration profiles of Cr and Ni across cells. For instance, Figure 11 represents the concentration profiles of Cr and Ni in the solidified microstructure of two DED stainless steel blocks: (a) 316L steel solidified in AF mode and (b) 304L steel solidified in FA mode. There has been no ferrite at cell boundaries by random. Therefore, these profiles instead indicate A and F modes. Concentration profiles of Ni and Cr change abruptly when the microanalyzer beam captures a retained ferrite particle at the boundary. Characteristic microsegregation profiles of Cr and Ni for unambiguous recognition of the different solidification modes have already been deserved in welding literature [123]. For the circumstances of no discernible concentration profiles, one may apply a criterion recommended by Iamboliev et al. [124] for the identification of FA mode from AF mode. In pulsed laser spot welds of 304L steel solidified in FA mode, many epitaxial grains formed on a single partially melted grain at the melt periphery with some austenite < 001 > directions varying along cell boundaries, in contrast to 310S steel solidified in AF mode.

Figure 11

Cr and Ni concentration profiles across solidification cells with no ferrite at the cell boundaries located at the Cr maxima: (a) 316L steel consistent to AF mode and (b) 304L steel consistent to FA mode. Reproduced from [40] with permission from Springer.

Table 4

Solidification mode and microstructure of selected DED stainless steels a .

Steel	Cr_eq	Ni_eq	Cr_eq/Ni_eq	Processing condition and build size b	Solidification mode and microstructure	Ref.
316L c	21.67	11.71	1.85	800 W, 15 mm s⁻¹,N₂ atmosphere,substrate preheated at 250 °C, cubes 20 × 20 × 20 mm³	AF mode proposed, cellular growth, average cooling rates of 9 × 10³ and 5.8 × 10³ K s⁻¹ on cold and preheated substrates, respectively. XRD identified ferrite in the deposits of cold substrate	[106]
316L c	20.79	13.28	1.57	900 W, 15 mm s⁻¹, scanned at 67° and 90°, cubes 20 × 20 × 20 mm³	AF mode proposed, cellular growth, cooling rate varied with build thickness and scan regim from 4.85 × 10³–2.05 × 10⁴ K s⁻¹. XRD identified ferrite in the specimen of higher cooling rate	[107]
316L c	19.75	11.67	1.69	Power and scan speed NA, N₂ shielding and chamber filling, blocks 10 × 10 × 20 mm³	AF mode proposed as stimulated by nitrogen pick up, cellular growth, scanning electron microscopy (SEM) revealed ferrite in both shielding and chamber filling conditions but XRD revealed ferrite in the latter only	[108]
316L c	19.01	11.47	1.66	400 W, 17 mm s⁻¹, blocks 13 × 41 × 41 mm³	AF mode consistent to Ni and Cr concentration profiles d , cellular growth, SEM revealed austenite with secondary ferrite (0.2 vol.-%), Cr and Ni microsegregations at cell boundaries	[40]
316L	20.90	13.23	1.58	600 W, 6.35 mm s⁻¹, blocks 30 × 30 × 70 mm³	A/AF mode can be find out, cellular growth, XRD revealed austenite with no ferrite, STEM-EDS indicated identical Ni, Cr and Mo segregation patterns	[109]
316L c	20.55	13.65	1.51	34–45 W, 0.4-0.6 mm s⁻¹, up rods (d1.16–1.36 mm)	AF mode proposed, optical microscopy and XRD identified austenite + ferrite, ferrite increases toward centre	[110]
316L	19.15	11.85	1.62	250–516 W, 16.9 mm s⁻¹, multiple pillars and cubes	AF mode proposed, cellular growth, EBSD revealed austenite + < 2 vol.-%ferrite, cooling rate estimated about 10³ K s⁻¹	[111]
316L c	20.45	13.78	1.48	650 W, 12 mm s⁻¹, blocks 20 × 55 × 122 mm³	A mode can be find out, cellular and dendritic growth, EBSD and XRD revealed austenite with no ferrite, SEM-EDS showed Cr, Mo microsegregations	[112]
316L	19.00	12.62	1.51	600 W, 15 mm s⁻¹, N₂ atmosphere, blocks 3 × 20 × 30 mm³	mainly FA mode presented, nitrogen stimulated AF mode occasionally, electric and magnetic field affected mode	[113]
316L c	19.90	14.90	1.34	2 kW, 8.33 mm s⁻¹, blocks 50 × 55 × 160 mm³, 30 × 70 × 125 mm³	AF mode proposed, dendritic growth, OM, SEM and XRD identified austenite + eutectic ferrite owing to heat retaining powder isolation	[114]
316L	19.13	11.74	1.63	360 W, 16.3 mm s⁻¹, multiple cylinder, disk, hexagon	AF mode proposed, cellular growth, eddy current indicated 0.2–0.5 vol.-% ferrite	[115]
316L c	22.41	12.26	1.83	400 W, 15 mm s⁻¹, cubes 20 × 20 × 20 mm³	AF mode proposed, cellular growth, EBSD and TEM showed austenite + ferrite, TEM-EDS indicated Cr enrichment and Ni depletion in ferrite, XRD identified ferrite	[116]
316L c	21.00	13.50	1.56	180 W, 6–22 mm s⁻¹, air vent wall 50 mm height	A mode proposed, cellular growth, SEM and EBSD indicated austenite with no ferrite, cooling rate estimated about 10³–4.5 × 10³ K s⁻¹	[117]
316 c	19.43	13.94	1.39	200–350 W, 3-8 mm s⁻¹, multiple walls of 3 mm thickness	AF mode proposed, cellular growth, XRD revealed austenite with no ferrite	[118]
316L c	20.8	13.1	1.59	450–850 W, 21 mm s⁻¹, walls 10 × 75 × 150 mm³	Both AF and FA proposed, considerable skeletal and LF presented	[119]
304L c	19.33	10.79	1.79	400 W, 17 mm s⁻¹, blocks 13 × 41 × 41 mm³	FA mode consistent to Ni and Cr microsegregations d , cellular growth, SEM revealed austenite with skeletal ferrite (1.1 vol.-%), Cr microsegregation at cell boundaries and Ni at cell cores	[40]
304 c	20.59	13.97	1.47	power NA, 8.5–34 mm s⁻¹, Blocks (9.5 mm wide, 76 mm long), Shells (0.5 mm thickness)	FA mode consistent with Ni and Cr concentration profiles d , austenite with vermicular and acicular ferrite, cellular growth, the cooling rate of shells measured about 3.5 × 10³ K s⁻¹	[91]
304 c	20.53	10.98	1.87	3 kW, 840 mm s⁻¹, horizontal rods (d10 mm)	AF mode proposed, EBSD demonstrated austenite + 10 vol.-% ferrite, EDS indicated enrichment by Cr, depletion by Ni	[120]
304L	20.10	11.19	1.80	380 W (LP), 12.7 mm s⁻¹, blocks 20 × 20 × 80 mm³ 3800 W (HP), 8.5 mm s⁻¹, blocks 15 × 50 × 75 mm³	FA mode proposed, WDS indicated similar Cr enrichment in ferrite and Ni in austenite, but ferritscope revealed 2 vol.-% ferrite for LP and 1.5 vol.-% ferrite for HP, cooling rate estimated 10⁴ K s⁻¹ for LP and 10³ K s⁻¹ for HP	[121]
304L	19.79	10.91	1.81	2300 W (LP), 8.5 mm s⁻¹, 4000 W (HP), 10.6 mm s⁻¹, walls 11 × 70 × 110 mm³	FA mode proposed, skeletal ferrite (0.9–1.3 vol.-%) identified in austenite for both LP and HP	[122]

Solidification mode has been taken into account in these selected studies.

Stainless steel substrate at ambient temperature and Ar inert gas atmosphere have been used, unless otherwise specified.

Constitution diagrams and equivalencies, e.g. Schaeffler, DeLong, WRC1992, have been applied for predicting solidification mode a priori.

Characteristic concentration profiles have been applied (Figure 11), which enables the unambiguous determination of solidification mode.

Brooks et al. [91] pioneered the first comparison of the solidification behaviour of the DED builds with autogenous laser welds for 304L steels. According to the approximate compositions presented on the pseudo-binary phase diagram (Figure 12), both alloys solidify with ferrite as the primary phase under near-equilibrium conditions. However, alloy #2 (Cr_eq/Ni_eq = 1.53) solidified with austenite as the primary phase without residual ferrite during laser welding. Alloy #3 (Cr_eq/Ni_eq = 1.78) represented refined austenite grains without residual ferrite in the laser welds but likely solidified with F mode and then transformed to massive austenite in the solid state. The solidification microstructures of DED blocks of both alloys were more complicated than laser welding counterparts by varying amounts of residual ferrite. A DED block of alloy #2 consisted of residual ferrite with vermicular and acicular morphologies along with cell cores, resulting from FA mode and subsequent transformation in the solid state. Scanning transmission electron microscopy (STEM)-EDS analysis was consistent with FA mode. In a somewhat mixed mode, grains consisting of vermicular ferrite likely solidified in a peritectic reaction with austenite growing from the melt. In contrast, grains consisting of acicular ferrite likely solidified as ferrite completely. Subsequent diffusional transformation of ferrite in the solid state was completed somewhere. Meanwhile, the microstructure of shell builds was similar to blocks, and scan speed had a negligible effect on the microstructure.

Figure 12

The approximate composition of the 304L DED steels studied by Brooks et al. Reproduced from [91] with permission from the Materials Research Society.

A DED block of alloy #3 solidified in the FA mode but represented residual ferrite inhomogeneously. Based on the microsegregation intensity and phase composition, both diffusional and massive transformation of primary ferrite was assumed in the solid state. The lack of microsegregation was attributed to a massive partitionless transformation, whereas compositional differences between the residual ferrite and austenite indicated long-range diffusional transformation. In conclusion, the DED was not as contentious as laser welding in altering solidification mode but gave more complicated solidification microstructures.

Wang et al. [122] reported the predominance of FA mode based on the skeletal ferrite in the DED walls of a 304L steel with a Cr_eq/Ni_eq ratio of 1.81. These walls were built at melting powers of 2300 and 4000 W. Grain size and grain shape were sensitive to cooling rate in terms of the melting power and the distance from the substrate, with larger and equiaxed grains at the high power and at the top. However, comparable ferrite numbers (0.9) were found for both conditions. Melia et al. [121] also reported the predominance of FA mode for 304L DED builds of a comparable Cr_eq/Ni_eq (1.80), printed at powers of 380 and 3800 W. Residual ferrite increased slightly from 1.5 to 2.0 vol.-% at high power. Wavelength-dispersive spectroscopy (WDS) indicated microsegregation of Cr to ferrite, Ni to austenite and P at the austenite–ferrite interface. A lower cooling rate increased microsegregation at the higher melting power.

Systematic studies on the type 316L DED builds recognised austenite as the primary solidifying phase (A and AF modes) [106 –119,125 –136]. A series of builds represented austenitic microstructures with none or negligible amounts of ferrite [109,112,115,117, 118,130 –133]. However, it was not clear how much ferrite solidified secondarily and subsequently transformed to austenite during cooling. A reference to a similar argument in welding literature aims at better understanding. Inoue et al. [59] distinguished A and AF modes for two stainless steels weld metals solidified with quenching under tin melt at cooling rates greater than 10⁴ K s⁻¹. No alteration of solidification microstructure, except for grain boundary migration, and no enhancement of alloying elements redistribution was found in the A mode. However, spherical or elongated ferrite particles, solidified in succession to the primary austenite cells in the AF mode, only slightly transformed into austenite during very rapid cooling, with a bit change in shape. Therefore, one may propose the partial transformation of secondarily solidifying ferrite during subsequent cooling in the solid state for the DED blocks upon cooling at relatively slower rates. Barika et al. [109] have recently reported an elegant multiscale study on the microstructural features of a 316L (Cr_eq/Ni_eq = 1.58) DED block with no retained ferrite. STEM-EDS maps indicated cell boundary enrichment by Cr, Ni and Mo, consistent to the primary austenite solidification mode. Furthermore, considerable amounts of nanometer-sized inclusions such as Cr–Mo-enriched precipitates and Si oxides were identified in the microstructure.

Another series of 316L DED blocks have shown considerable amounts of residual ferrite in the cellular austenitic microstructure, which indicates the predominance of AF mode more clearly [107,108,110,134 –139]. An earlier study reported ferrite retention at about 3–10 vol.-% owing to the microsegregation of Cr and Mo at cell boundaries [134]. The microstructure was sensitive to the cooling rate, changing with the specimen thickness in the build direction. Yadollahi et al. [135] revealed about 9 vol.-% ferrite by EBSD in the middle of a DED block with no ferrite at the bottom. These results clarify various amounts of ferrite in AF mode depending on the cooling rate. Recent studies also provide insight into the effect of process-driven cooling rate effects on the microstructure [107,108]. For instance, different scanning patterns resulted in different ferrite contents, where ferrite was retained in a sample produced by 90° rotation per layer in contrast to another produced by 60° rotation. The authors assumed that different cooling rates produced different ferrite contents. Manvatkar et al. [128] computed the decrease in the cooling rate with increasing build thickness and demonstrated good agreement between the computed cooling rate and the PCAS variation. Zheng et al. [136] also estimated the cooling rate variation with increasing build thickness by measuring PCAS.

Moreover, manipulated solidification behaviour of stainless steels has also been reported in modulated DED processes such as high-power laser melting with an electromagnetic stirring module, where mixed AF + FA modes have been identified [113,114].

DED steels on the implicated Schaeffler diagrams

David et al. [77] and Nakao et al. [44] have shown that a high cooling rate implicates the conventional Schaeffler diagram, which is nevertheless used to ascertain the solidification mode and microstructures of weld metals and DED builds based on the Cr_eq and Ni_eq. Both papers reported the so-called implicated Schaeffler diagrams constructed for autogenous laser welds of stainless steels, but cooling rates were estimated by computational and experimental PCAS measurement differently. Despite differences in the magnitudes of the proposed cooling rates, the same trend could be found in both studies. Since stainless steels solidify at high cooling rates during the DED process, it should be more reasonable to ascertain the solidification mode and microstructure from those implicated Schaeffler diagrams, which take the cooling rate effects into account. Figure 13 boards some DED stainless steels on the implicated Schaeffler diagrams constructed at three different cooling rates, ca. 1.9 × 10³, 6.0 × 10³ and 3.9 × 10⁴ K s⁻¹ [44]. Hopefully, the implicated Schaeffler diagram provided for a cooling rate of 6 × 10³ K s⁻¹ predicted the solidification behaviour of DED stainless steels more reasonably. However, one may apply each of these implicated diagrams depending on the approximate cooling rate of specific energy input, build thickness, etc.

Figure 13

DED stainless steels incorporated to the implicated Schaeffler diagrams constructed for high cooling rates: (a) 1.9 × 10³ K s⁻¹, (b) 6.0 × 10³ K s⁻¹ with reported ferrite percents given inside parenthesis and (c) 3.9 × 10⁴ K s⁻¹. The implicating Schaeffler diagrams were adopted from [44] with permission from the Japan Welding Society.

Printability of DED stainless steels

Mukherjee et al. [140] introduced the term printability as a measure of the relative ease of metallic alloy additive manufacturing. Printability of 316L steel using different MAM processes has been compared in terms of susceptibilities for residual stress, distortion, compositional change and lack of fusion defects [141]. Comparably, the term weldability interprets welding behaviour and weld metal performance in service. As a measure of weldability, hot cracking of stainless steels has widely been discussed in welding literature [29 –32,49,51,58,142 –150]. Hot cracking has also been reported in the high energy density welds [49,73,74,78] and the DED builds of stainless steels [151 –154].

The American Welding Society (AWS) describes hot cracking as any type of cracking at temperatures near the completion of solidification [155]. One type of hot cracking is solidification cracking which occurs owing to the retention of liquid film at grain boundaries in the later stage of fusion zone solidification. Microsegregation of some alloying and harmful impurity elements such as P and S decreases the melting point of the aftermost melt thus augmenting hot cracking. Stainless steels hot cracking susceptibility is ascertained from weldability diagrams, which demarcate crack and no-crack regions on a plot of total impurities content (P + S) versus Cr_eq/Ni_eq ratio. Conventional alloys with low S + P contents (cf. < 0.02-wt%) resist adequately against hot cracking during arc welding apart from Cr_eq/Ni_eq ratio. However, a higher impurity content promotes cracking susceptibility when Cr_eq/Ni_eq ratio decreases below a threshold. The threshold corresponds to the critical Cr_eq/Ni_eq ratio for the transition of solidification mode from AF to FA. It varies with cooling rate at about 1.49–1.55 in arc welding [48] and 1.59–1.68 in pulsed laser welding [74,75]. Indeed, hot cracking occurs in primary austenite mode but steels solidifying in primary ferrite mode with 5–10 vol.-% residual ferrite are practically passive to hot cracking. Yu et al. [152] adapted contemporary weldability diagrams for predicting hot cracking susceptibility of 316L blocks produced by the DED process. In general, predictions were unsatisfactory. A high-Si alloy suffered hot cracking apart from low P + S and solidification in FA mode. The effect of high nitrogen was timely interpreted to promote hot cracking in the DED steels. The nitrogen atmosphere, stimulating primary austenite solidification, promotes hot cracking susceptibility of the DED stainless steels as weld metals [154,156,157].

Nonetheless, there are instances of extraordinary resistance against hot cracking during high energy density welding [80,84,86] as many DED stainless steels solidify with primary austenite mode safely. Instances of negligible microsegregations have also been reported [78,80,124]. Therefore, it is thought that a high cooling rate decreases microsegregations at solidifying grain boundaries. According to [158,159], increasing of solidification rate stimulates solute trapping and renders partitionless solidification at most.

In addition to hot cracking, the formation of the deleterious sigma phase during the aging of ferrite in duplex stainless steels and martensite cracking in low alloy steels promote failure susceptibility. Suutala et al. [63] delimited these structure-sensitive properties on the Schaeffler diagram by representing the middle part of the Bystram diagram [160], modified by reference to the cooling rate stimulated, hot cracking susceptibility. Figure 14 sets the DED stainless steels to the modified Bystram diagram. The formation of the FeCr (sigma) phase in a DED steel is anticipated. Those steels solidifying with A and AF modes show hot cracking at high P + S (cf. > 0.04 wt-%) and high N content, while those solidifying with FA mode have been more resistant. There has not been any report of martensite cracking in the DED stainless steels yet. Many of the accredited DED stainless steels are inside the triangular SAFE ZONE, wherein alloy compositions and process conditions have practically been justified for the safe microstructure evolution.

Figure 14

DED stainless steels [from the same references cited in Figure 13(a)] are presented on the modified Bystram diagram. The triangular safe zone was adopted from [63] with permission from Springer.

Conclusions

Stainless steels with Cr_eq/Ni_eq ratios in the range of 1.5–1.7 which solidify with the primary ferrite mode under near-equilibrium conditions are much vulnerable to intriguing alteration of solidification mode to primary austenite at high cooling rates. Consequently, acute microstructure heterogeneities are realised. However, those steels that solidify with the primary austenite mode in the near-equilibrium conditions retain the same mode at high cooling rates. However, the amount of residual ferrite varies by cooling rate variation considerably. Mitigated conditions result in arms spacing, phase composition and solute redistribution variations.

Process-induced and inherent cooling rate variations during high energy density welding and the DED process manipulate the microstructure of stainless steels intriguingly. Although cooling rate effects in the DED are less stringent than welding, substantial cooling rate variation effects on the arms spacing and ferrite content along the build direction have been reported.

An implicated Schaeffler diagram, constructed for stainless steel laser welds at a cooling rate of 6.0 × 10³ K s⁻¹, predicted the solidification microstructure of the DED stainless steels much better than the original constitution diagram which is, nevertheless, used conventionally.

Available structure-sensitive properties of the DED stainless steels, such as solidification cracking susceptibility, enunciated the optimistic printability on the modified Bystram diagram.

Footnotes

Disclosure statement

No potential conflict of interest was reported by the authors.

1

All stainless steel compositions comply with the American Iron and Steel Institute (AISI) designation system throughout this paper.

2

Empirical relations such as d = 80(Ṫ)^−0.33 have been obtained in stainless steels gas tungsten arc (GTA) and tungsten inert gas (TIG) weld metals between measured dendritic arms spacing (d) and experimental cooling rates (Ṫ) from temperature measurement by thermocouple. Owing to difficulties with temperature measurement in tiny melt pools of high energy density welding and MAM, the empirical relation is often extrapolated at high cooling rates. Both primary dendritic arms spacing (PDAS) and secondary dendritic arms spacing (SDAS) have been applied for dendritic microstructures with different constants, and primary cellular arms spacing (PCAS) has been implemented for cellular microstructures.

3

These modes are applied in this paper, but the term primary austenite is applied to emphasize the solidification mode with austenite as the primarily solidifying phase (the leading one) irrespective of any secondarily solidifying ferrite whenever A and AF modes are not to be discerned essentially. Similarly, the term primary ferrite will be sued with no discernment between FA and F modes to emphasize primarily solidifying ferrite irrespective of any secondarily solidifying austenite. Noteworthy, those terms have occasionally been used for A and F modes in the literature. For a schematic corresponding to the present terminology see [51].

4

Schaeffler equivalencies: Cr_eq = Cr+Mo+1.5Si+0.5Nb and Ni_eq = Ni+0.5Mn+30C, in wt-%. Unless otherwise specified, the Schaeffler equivalencies apply hereafter.

5

Cr_eq = Cr+Mo+1.5Si+0.5Nb+2Ti and Ni_eq = Ni+0.5Mn+30C+30(N-0.06), in wt-%.

6

Cr_eq = Cr+1.37Mo+1.5Si+2Nb+3Ti and Ni_eq = Ni+0.31Mn+22C+14.2N+Cu, in wt-%.

7

Cr_eq = Cr+Mo+1.5Si+0.5Nb; Ni_eq = Ni+0.5Mn+30C+30N, in wt-% (DeLong equivalencies)

References

Dutta-Majumdar

, Manna

Laser material processing. Int Mater Rev. 2011;56:341–388.

Murr

, Gaytan

, Ramirez

, . Metal fabrication by additive manufacturing using laser and electron beam melting technologies. J Mater Sci Technol. 2012;28:1–14.

, Meiners

, Wissenbach

, . Laser additive manufacturing of metallic components: materials, processes and mechanisms. Int Mater Rev. 2012;57:133–164.

Frazier

Metal additive manufacturing: a review. J Mater Eng Perform. 2014;23:1917–1928.

Körner

Additive manufacturing of metallic components by selective electron beam melting-a review. Int Mater Rev. 2016;61:361–377.

Herzog

, Seyda

, Wycisk

, . Additive manufacturing of metals. Acta Mater. 2016;117:371–392.

Sames

, List

, Pannala

, . The metallurgy and processing science of metal additive manufacturing. Int Mater Rev. 2016;61:315–360.

Zhang

, Wu

, Guo

, . Additive manufacturing of metallic materials: a review. J Mater Eng Perform. 2018;27:1–13.

Harun

WSW

, Kamariah

MSIN

, Muhamad

, . A review of powder additive manufacturing processes for metallic biomaterials. Powder Tech. 2018;327:128–151.

10.

Fayazfar

, Salarian

, Rogalsky

, . A critical review of powder-based additive manufacturing of ferrous alloys: process parameters, microstructure and mechanical properties. Mater Des. 2018;144:98–128.

11.

DebRoy

, Wei

, Zuback

, . Additive manufacturing of metallic components—process, structure and properties. Prog Mater Sci. 2018;92:112–224.

12.

Poorganji

, Ott

, Kelkar

, . Review: materials ecosystem for additive manufacturing powder bed fusion processes. JOM. 2020;72:561–576.

13.

ISO/ASTM52900-15. Standard terminology for additive manufacturing—general principles—terminology. West Conshohocken (PA): ASTM International; 2015; www.astm.org.

14.

Gong

, Ye

, Chi

, . Research status of laser additive manufacturing for metal: a review. J Mater Res Tech. 2021;15:855–884.

15.

Thompson

, Bian

, Shamsaei

, . An overview of direct laser deposition for additive manufacturing; part I: transport phenomena, modeling and diagnostics. Addit Manuf. 2015;8:36–62.

16.

Shamsaei

, Yadollahi

, Bian

, . An overview of direct laser deposition for additive manufacturing; part II: mechanical behavior, process parameter optimization and control. Addit Manuf. 2015;8:12–35.

17.

Adeyemi

, Akinlabi

, Mahamood

RM.

Powder bed based laser additive manufacturing process of stainless steel: a review. Mater Today: Proceed. 2018;5:18510–18517.

18.

Krakhmalev

, Fredriksson

, Svensson

, . Microstructure, solidification texture, and thermal stability of 316 L stainless steel manufactured by laser powder bed fusion. Metals (Basel). 2018;8:643.

19.

Bajaj

, Hariharan

, Kini

, . Steels in additive manufacturing: a review of their microstructure and properties. Mater Sci Eng A. 2020;772:138633.

20.

Jin

, Zhang

, Jin

, . Wire arc additive manufacturing of stainless steels: a review. Appl Sci. 2020;10:1563.

21.

Zai

, Zhang

, Wang

, . Laser powder bed fusion of precipitation-hardened martensitic stainless steels: a review. Metals (Basel). 2020;10:255.

22.

Haghdadi

, Laleh

, Moyle

, . Additive manufacturing of steels: a review of achievements and challenges. J Mater Sci. 2021;56:64–107.

23.

Michla

JRJ

, Nagarajan

, Krishnasamy

, . Conventional and additively manufactured stainless steels: a review. Trans Indian Inst Met. 2021;74:1261–1278.

24.

Saboori

, Aversa

, Marchese

, . Microstructure and mechanical properties of AISI L produced by directed energy deposition-based additive manufacturing: a review. Appl Sci. 2020;10:3310.

25.

Astafurov

, Astafurova

Phase composition of austenitic stainless steels in additive manufacturing: a review. Metals (Basel). 2021;11:1052.

26.

David

, Vitek

JM.

Correlation between solidification parameters and weld microstructures. Int Mater Rev. 1989;34:213–245.

27.

David

, Babu

, Vitek

JM.

Welding: solidification and microstructure. JOM. 2003;55:14–20.

28.

Siefert

, David

SA.

Weldability and weld performance of candidate austenitic alloys for advanced ultrasupercritical fossil power plants. Sci Tech Weld Join. 2014;19:271–294.

29.

Arata

, Matsuda

, Katayama

Solidification crack susceptibility in weld metals of fully austenitic stainless steels (report I): fundamental investigation on solidification behavior of fully austenitic and duplex microstructures and effect of ferrite on microsegregation. Trans JWRI. 1976;5:135–151.

30.

Lippold

, Savage

WF.

Solidification of austenitic stainless steel weldments: part III—the effect of solidification behavior on hot cracking susceptibility. Weld J. 1982;61:388s–396s.

31.

Brooks

, Thompson

AW.

Microstructural development and solidification cracking susceptibility of austenitic stainless steel welds. Int Mater Rev. 1991;36:16–44.

32.

Shankar

, Gill

TPS

, Mannan

, . Solidification cracking in austenitic stainless steel welds. Sadhan (A). 2003;28:359–382.

33.

Nakao

, Nishimoto

, Ishizaki

Effects of delta-ferrite on sensitisation of austenitic stainless steel weld metal. Weld Int. 1992;6:523–530.

34.

David

, Vitek

, Alexander

DJ.

Embrittlement of austenitic stainless steel welds. J Nondest Eval. 1996;15:129–136.

35.

Bilmes

, Gonzalez

, Llorete

, . Effect of δ ferrite solidification morphology of austenitic stainless steel weld metal on properties of welded joints. Weld Int. 1996;10:797–808.

36.

Mills

WJ.

Fracture toughness of type 304 and 316 stainless steels and their welds. Int Mater Rev. 1997;42:45–82.

37.

Luppo

, Hazarabedian

, Ovejero-Garcia

Effects of delta ferrite on hydrogen embrittlement of austenitic stainless steel welds. Corr Sci. 1999;41:87–103.

38.

Oliveira

, Santos

, Miranda

RM.

Revisiting fundamental welding concepts to improve additive manufacturing: from theory to practice. Prog Mater Sci. 2020;107:100590.

39.

Wei

, Bhadeshia

HKDH

, David

, . Harnessing the scientific synergy of welding and additive manufacturing. Sci Tech Weld Join. 2019;24:361–366.

40.

Smith

, Sugar

, Marchi

, . Microstructural development in DED stainless steels: applying welding models to elucidate the impact of processing and alloy composition. J Mater Sci. 2021;56:762–780.

41.

Hwa

, Kumai

, Devine

, . Microstructural banding of directed energy deposition-additively manufactured 316L stainless steel. J Mater Sci Tech. 2021;69:96–105.

42.

Manvatkar

, De

, DebRoy

Heat transfer and material flow during laser assisted multi-layer additive manufacturing. J Appl Phys. 2014;116:124905.

43.

Yan

, Liu

, Tang

, . Review on thermal analysis in laser-based additive manufacturing. Opt Laser Tech. 2018;106:427–441.

44.

Nakao

, Nishimoto

, Zhang

WP.

Effects of rapid solidification by laser surface melting on solidification modes and microstructures of stainless steel. Trans Japan Weld Soc. 1988;19:20–26.

45.

Mukherjee

, DebRoy

A digital twin for rapid qualification of 3D printed metallic components. Appl Mater Today. 2019;14:59–65.

46.

Berjemo

MAV

, Karlsson

, DebRoy

NOLAMP14; 2013.p. 3.

47.

Suutala

, Takalo

, Moisio

The relationship between solidification and microstructure in austenitic and austenitic-ferritic stainless steel welds. Metall Trans A. 1979;10A:512–514.

48.

Suutala

Effect of solidification conditions on the solidification mode in austenitic stainless steels. Metall Trans A. 1983;14A:191–197.

49.

Katayama

, Fujimoto

, Matsunawa

Correlation among solidification process, microstructure, microsegregation and solidification cracking susceptibility in stainless steel weld metals (materials, metallurgy & weldability). Trans JWRI. 1985;14:123–138.

50.

Elmer

JW.

The influence of cooling rate on the microstructure of stainless steel alloys [dissertation]. Massachusetts (MA): Massachusetts Institute of Technology; 1988.

51.

Brooks

, Thompson

, Williams

JC.

A fundamental study of the beneficial effects of delta ferrite in reducing weld cracking. Weld J. 1984;63:71s–83s.

52.

Koseki

, Flemings

MC.

Solidification of undercooled Fe–Cr–Ni alloys: part II. Microstructural evolution. Metall Mater Trans A. 1996;27A:3226–3240.

53.

Lippold

, Kotecki

DJ.

Welding metallurgy and weldability of stainless steels. Hoboken (NJ): John Wiley & Sons; 2005.

54.

Welding metallurgy. Hoboken (NJ): John Wiley & Sons; 2005.

55.

Rivlin

, Raynorv

GV.

Phase equilibria in iron ternary alloys: critical evaluation of constitution of chromium–iron–nickel system. Int Metals Rev. 1980;1:21–38.

56.

Kundrat

, Elliott

JE.

Phase relationships in the Fe–Cr–Ni system at solidification temperatures. Metall Trans A. 1988;19A:899–908.

57.

Großwendt

, Becker

, Röttger

, . Impact of the allowed compositional range of additively manufactured 316L stainless steel on processability and material properties. Materials (Basel). 2021;14:4074.

58.

Takalo

, Suutala

, Moisio

Austenitic solidification mode in austenitic stainless steel welds. Metall Trans A. 1979;10A:1173–1181.

59.

Inoue

, Koseki

, Okita

, . Solidification and transformation behaviour of austenitic stainless steel weld metals solidified as primary austenite: study of solidification and subsequent transformation of Cr–Ni stainless steel weld metals (1st report). Weld Int. 1997;11:876–887.

60.

Suutala

Effect of manganese and nitrogen on the solidification mode in austenitic stainless steel welds. Metall Trans A. 1982;13A:2121–2130.

61.

Takalo

, Suutala

, Moisio

Influence of ferrite content on its morphology in some austenitic weld metals. Metall Trans A. 1976;7A:1591–1592.

62.

Lippold

, Savage

WF.

Solidification of austenitic stainless steel weldments: part 2—the effect of alloy composition on ferrite morphology. Weld J. 1980;59:48s–58s.

63.

Suutala

, Takalo

, Moisio

Ferritic–austenitic solidification mode in austenitic stainless steel welds. Metall Trans A. 1980;11A:717–725.

64.

Suutala

, Takalo

, Moisio

Single-phase ferritic solidification mode in austenitic-ferritic stainless steel welds. Metall Trans A. 1979;10A:1183–1190.

65.

Brooks

, Williams

, Thompson

AW.

Microstructural origin of the skeletal ferrite morphology of austenitic stainless steel welds. Metall Trans A. 1983;14A:1271–1281.

66.

Inoue

, Koseki

, Okita

, . Solidification and transformation behaviour of austenitic stainless steel weld metals solidified as primary ferrite: study of solidification and subsequent transformation of Cr–Ni stainless steel weld metals (2nd report). Weld Int. 1997;11:937–949.

67.

Leone

, Kerr

HW.

The ferrite to austenite transformation in stainless steels. Weld J. 1982;61:13s–21s.

68.

Inoue

, Koseki

, Okita

, . Solidification and transformation behaviour of Cr–Ni stainless steel weld metals with ferritic single-phase solidification mode: study of solidification and transformation of Cr–Ni stainless steel weld metals (4th report). Weld Int. 1998;12:282–296.

69.

Perricone

, DuPont

, Anderson

, . An investigation of the massive transformation from ferrite to austenite in laser-welded Mo-bearing stainless steels. Metall Mater Trans A. 2011;42A:700–716.

70.

Kotecki

, Siewert

TA.

WRC-1992 constitution diagram for stainless steel weld metals: a modification of the WRC-1988 diagram. Weld J. 1992;71:171s–178s.

71.

Olson

DL.

Prediction of austenitic weld metal microstructure and properties. Weld J. 1985;64:281s–295s.

72.

Berjemo

MAV.

Predictive and measurement methods for delta ferrite determination in stainless steels. Weld J. 2012;91:113s–121s.

73.

Katayama

, Matsunawa

Solidification microstructure of laser welded stainless steels. ICALEO, vol. 44; 1984. p. 60–67.

74.

Lippold

JC.

Solidification behavior and cracking susceptibility of pulsed-laser welds in austenitic stainless steels. Weld J. 1994;73:129s–139s.

75.

Lienert

, Lippold

JC.

Improved weldability diagram for pulsed laser welded austenitic stainless steels. Sci Tech Weld Join. 2003;8:1–9.

76.

Berjemo

MAV

, DebRoy

, Hurtig

, . Towards a map of solidification cracking risk in laser welding of austenitic stainless steels. Physica D. 2015;78:230–239.

77.

David

, Vitek

, Hebble

TL.

Effect of rapid solidification on stainless steel weld metal microstructures and its implications on the Schaeffler diagram. Weld J. 1987;66:289s–300s.

78.

Vitek

, Dasgupta

, David

SA.

Microstructural modification of austenitic stainless steels by rapid solidification. Metall Trans A. 1983;14A:1833–1841.

79.

Zacharia

, David

, Vitek

, . Heat transfer during Nd-Yag pulsed laser welding and its effect on solidification structure of austenitic stainless steels. Metall Trans A. 1989;20A:957–967.

80.

Molian

PA.

Solidification behaviour of laser welded stainless steel. J Mater Sci Let. 1985;4:281–283.

81.

Zambon

, Bonollo

Rapid solidification in laser welding of stainless steels. Mater Sci Eng A. 1994;178:203–207.

82.

Akgun

, Inal

OT.

Laser surface melting and alloying of type 304 L stainless steel. J Mater Sci. 1995;30:6097–6104.

83.

El-Batahgy

AM.

Effect of laser welding parameters on fusion zone shape and solidification structure of austenitic stainless steels. Mater Lett. 1997;32:155–163.

84.

Farid

, Molian

PA.

High-brightness laser welding of thin-sheet 316 stainless steel. J Mater Sci. 2000;35:3817–3826.

85.

McPherson

, Samson

, Baker

, . Steel microstructures in autogenous laser welds. J Laser Appl. 2003;15:200–210.

86.

Celen

, Karadeniz

, Ozden

Effect of laser welding parameters on fusion zone morphological, mechanical and microstructural characteristics of AISI 304 stainless steel. Mat-Wiss u Werkstofftech. 2008;39:845–850.

87.

Kell

, Tyrer JR , Higginson

, . Microstructural characterization of autogenous laser welds on 316L stainless steel using EBSD and EDS. J Microscopy. 2005;217:167–173.

88.

Elmer

, Allen

, Eagar

TW.

Microstructural development during solidification of stainless steel alloys. Metall Trans A. 1989;20A:2117–2131.

89.

Inoue

, Koseki

, Okita

, . Epitaxial growth and phase formation of austenitic stainless steel weld metals near fusion boundaries: study of solidification and transformation of Cr–Ni stainless steel weld metals (3rd report). Weld Int. 1998;12:195–206.

90.

Brooks

, Baskes

, Greulich

FA.

Solidification modeling and solid-state transformations in high-energy density stainless steel welds. Metall Trans A. 1991;22A:915–926.

91.

Brooks

, Headley

, Robino

CV.

Microstructures of laser deposited 304L austenitic stainless steel. Mater Res Soc Symp Proc. 2000;625:21–30.

92.

Kelly

, Cohen

, Vander Sande

JB.

Rapid solidification of a droplet-processed stainless steel. Metall Mater Trans A. 1984;15A:819–833.

93.

Wright

, Bae

, Kelly

, . The microstructure and phase relationships in rapidly solidified type 304 stainless steel powders. Metall Mater Trans A. 1988;19A:2399–2405.

94.

Koseki

, Flemings

MC.

Solidification of undercooled Fe–Cr–Ni alloys: part I. Thermal behavior. Mater Trans A. 1995;26A:2991–2999.

95.

Koseki

, Flemings

MC.

Solidification of undercooled Fe–Cr–Ni alloys: part III. Phase selection in chill casting. Metall Mater Trans A. 1997;28A:2385–2395.

96.

Löser

, Herlach

DM.

Theoretical treatment of the solidification of undercooled Fe–Cr–Ni melts. Metall Mater Trans A. 1992;23A:1585–1591.

97.

Löser

, Volkman

, Herlach

DM.

Nucleation and metastable phase formation in undercooled Fe–Cr–Ni melts

Mater Sci Eng A. 1994;178:163–166.

98.

Volkman

, Löser

, Herlach

DM.

Nucleation and phase selection in undercooled Fe–Cr–Ni melts: part I. Theoretical analysis of nucleation behavior. Metall Mater Trans A. 1997;28A:453–460.

99.

Volkman

, Löser

, Herlach

DM.

Nucleation and phase selection in undercooled Fe–Cr–Ni melts: part II. Containerless solidification experiments. Metall Mater Trans A. 1997;28A:461–469.

100.

Mizukami

, Suzuki

, Umeda

Initial stage of rapid solidification of 18-8 stainless steel. Mater Sci Eng A. 1993;173:363–364.

101.

Trivedi

, Kurz

Dendritic growth. Int Mat Rev. 1994;39:49–74.

102.

Kurz

, Trivedi

Microstructure and phase selection in laser treatment of materials. Trans ASME. 1992;114:450–458.

103.

Fukumoto

, Kurz

The δ to γ transition in Fe–Cr–Ni alloys during laser treatment. ISIJ Int. 1997;37:677–684.

104.

Fukumoto

, Kurz

Prediction of the δ to γ transition in austenitic stainless steels during laser treatment. ISIJ Int. 1998;38:71–77.

105.

Fukumoto

, Kurz

Solidification phase and microstructure selection maps for Fe–Cr–Ni alloys. ISIJ Int. 1999;39:1270–1279.

106.

Moheimani

, Iuliano

, Saboori

The role of substrate preheating on the microstructure, roughness, and mechanical performance of AISI 316L produced by directed energy deposition additive manufacturing. Int J Add Manuf Tech. 2022;119:7159–7174.

107.

Saboori

, Piscopo

, Lai

, . An investigation on the effect of deposition pattern on the microstructure, mechanical properties and residual stress of 316L produced by directed energy deposition. Mater Sci Eng A. 2020;780:139179.

108.

Aversa

, Saboori

, Librer

, . The role of directed energy deposition atmosphere mode on the microstructure and mechanical properties of 316L samples. Addit Manuf. 2020;34:101274.

109.

Barkia

, Aubry

, Ashtiani

, . On the origin of the high tensile strength and ductility of additively manufactured 316L stainless steel: multiscale investigation. J Mater Sci Tech. 2020;41:209–218.

110.

Weng

, Gao

, Jiang

, . A novel strategy to fabricate thin 316L stainless steel rods by continuous directed energy deposition in Z direction. Addit Manuf. 2019;27:474–481.

111.

Zheng

, Haley

, Yang

, . On the evolution of microstructure and defect control in 316L SS components fabricated via directed energy deposition. Mater Sci Eng A. 2019;764:138243.

112.

Mukherjee

Effect of build geometry and orientation on microstructure and properties of additively manufactured 316L stainless steel by laser metal deposition. Materialia. 2019;7:100359.

113.

, Sun

, Wang

, . Effects of electromagnetic field on the laser direct metal deposition of austenitic stainless steel. Opt Laser Tech. 2019;119:105586.

114.

Guo

, Zou

, Huang

, . Study on microstructure, mechanical properties and machinability of efficiently additive manufactured AISI 316L stainless steel by high-power direct laser deposition. J Mater Process Tech. 2017;240:12–22.

115.

Yang

, Yee

, Zheng

, . Process–structure–property relationships for 316L stainless steel fabricated by additive manufacturing and its implication for component engineering. J Therm Spray Tech. 2017;26:610–626.

116.

Ziętala

, Durejko

, Polański

, . The microstructure, mechanical properties and corrosion resistance of 316L stainless steel fabricated using laser engineered net shaping. Mater Sci Eng A. 2016;677:1–10.

117.

Marya

, Singh

, Marya

, . Microstructural development and technical challenges in laser additive manufacturing: case study with a 316L industrial part. Metall Mater Trans B. 2015;46B:1654–1665.

118.

Lima

MSF

, Sankaré

Microstructure and mechanical behavior of laser additive manufactured AISI 316 stainless steel stringers. Mater Des. 2014;55:526–532.

119.

Koduri

, Henderson

, Costello

, . Microstructural observations of 316L stainless steel laser powder depositions. In: Marquis

FDS

, editor. Powder materials: current research and industrial pratices III. Chicago, Il;The Minerals, Metals and Materials Society; 2003. p. 229–238.

120.

Huang

, Zhang

, Dai

, . Mechanical properties of 304 austenite stainless steel manufactured by laser metal deposition. Mater Sci Eng A. 2019;758:60–70.

121.

Melia

, Nguyen

HDA

, Rodelas

, . Corrosion properties of 304L stainless steel made by directed energy deposition additive manufacturing. Corr Sci. 2019;152:20–30.

122.

Wang

, Ta

, Beese

AM.

Effect of processing parameters on microstructure and tensile properties of austenitic stainless steel 304L made by directed energy deposition additive manufacturing. Acta Mater. 2016;110:226–235.

123.

Brooks

, Williams

, Thompson

AW.

STEM analysis of primary austenite solidified stainless steel welds. Metall Trans A. 1983;14A:23–31.

124.

Iamboliev

, Katayama

, Matsunawa

Interpretation of phase formation in austenitic stainless steel welds. Weld J. 2003;82:337s–347s.

125.

Dutta-Majumdar

, Pinkerton

, Liu

, . Microstructure characterisation and process optimization of laser assisted rapid fabrication of 316L stainless steel. Appl Surf Sci. 2005;247:320–327.

126.

Amine

, Newkirk

, Liou

An investigation of the effect of laser deposition parameters on characteristics of multilayered 316 L deposits. Int J Adv Manuf Technol. 2014;73:1739–1749.

127.

Lei

, Xie

, Zhou

, . Comparative study on microstructure and corrosion performance of 316 stainless steel prepared by laser melting deposition with ring-shaped beam and Gaussian beam. Opt Laser Technol. 2019;111:271–283.

128.

Manvatkar

, Gokhale

, Jagan Reddy

, . Estimation of melt pool dimensions, thermal cycle, and hardness distribution in the laser-engineered net shaping process of austenitic stainless steel. Metall Mater Trans A. 2011;42A:4080–4087.

129.

Sklyar

, Turichin

, Klimova

, . Microstructure of 316L stainless steel components produced by direct laser deposition. Steel in Trans. 2016;46:883–887.

130.

Zhang

, Wang

, Liu

, . Characterization of stainless steel parts by laser metal deposition shaping. Mater Des. 2014;55:104–119.

131.

Feenstra DR , Cruz

, Gao

, . Effect of build height on the properties of large format stainless steel 316L fabricated via directed energy deposition. Addit Manuf. 2020;34:101205.

132.

Chen

, Liu

, Song

A multiscale investigation of deformation heterogeneity in additively manufactured 316L stainless steel. Mater Sci Eng A. 2021;820:141493.

133.

Yan

, Zou

, Cheng

, . Revealing relationships between heterogeneous microstructure and strengthening mechanism of austenitic stainless steels fabricated by directed energy deposition (DED). J Mater Res Tech. 2021;15:582–594.

134.

Hofmeister

, Griffith

, Ensz

, . Solidification in direct metal deposition by LENS processing. JOM. 2001;53:30–34.

135.

Yadollahi

, Shamsaei

, Thompson

, . Effects of process time interval and heat treatment on the mechanical and microstructural properties of direct laser deposited 316L stainless steel. Mater Sci Eng A. 2015;644:171–183.

136.

Zheng

, Zhou

, Smugeresky

, . Thermal behavior and microstructural evolution during laser deposition with laser-engineered net shaping: part I. Numerical calculations. Metall Mater Trans A. 2008;39A:2228–2236.

137.

Chen

, Chen

, Cheng

, . Microstructure and properties of hybrid additive manufacturing 316L component by directed energy deposition and laser remelting. J Iron Steel Res Int. 2020;27:842–848.

138.

Gray

III , Livescu

, Rigg

, . Structure/property (constitutive and spallation response) of additively manufactured 316L stainless steel. Acta Mater. 2017;138:140–149.

139.

Brown

, Adams

, Balogh

, . In situ neutron diffraction study of the influence of microstructure on the mechanical response of additively manufactured 304L stainless steel. Metall Mater Trans A. 2017;48A:6055–6069.

140.

Mukherjee

, Zuback

, De

, . Printability of alloys for additive manufacturing. Sci Rep. 2016;6:19717.

141.

Mukherjee

, DebRoy

Printability of 316 stainless steel. Sci Tech Weld Join. 2019;24:412–419.

142.

Brooks

JA.

Weldability of high N, high Mn austenitic stainless steel. Weld Res Int. 1975;4:189s–195s.

143.

Brooks

JA.

The effects of phosphorus, sulfur and ferrite content on weld cracking of type 309 stainless steel. Weld Res Int. 1978;7:139s–143s.

144.

Kujanpää

, Suutala

, Takalo

, . Correleation between solidification cracking and microstructure in austenitic and austenitic-ferritic stainless steel welds. Weld Res Int. 1979;9:55–76.

145.

Kujanpää

, Suutala

, Takalo

, . Solidification cracking-estimation of the susceptibility of austenitic-ferritic stainless steel welds. Metal Const. 1980;6:282–285.

146.

, Thompson

, McCarthy

, . Microstructure evolution and solidification cracking in austenitic stainless steel welds. Weld J. 2018;97:301s–314s.

147.

Katayama

Solidification phenomena of weld metals: solidification cracking mechanism and cracking susceptibility (3rd report). Weld Int. 2001;15:627–636.

148.

Korinko

, Malene

SH.

Considerations for the weldability of types 304L and 316L stainless steel. Prac Fail Analys. 2001;1:61–68.

149.

Brooks

JA.

Weld solidification and cracking behavior of free-machining stainless steel. Weld J. 2003;82:51s–64s.

150.

Kou

Solidification and liquation cracking issues in welding. JOM. 2003;55:37–42.

151.

, Cui

, Xiao

Crack detection during laser metal deposition by infrared monochrome pyrometer. Materials (Basel). 2020;13:5643.

152.

, Rombouts

, Maes

Cracking behavior and mechanical properties of austenitic stainless steel parts produced by laser metal deposition. Mater Des. 2013;45:228–235.

153.

Chen

, Lin

, Wang

, . The hot cracking mechanism of 316L stainless steel cladding in rapid laser forming process. Rare Met Mater Eng. 2003;32:183–186.

154.

Song

, Li

, Deng

, . Cracking mechanism of laser cladding rapid manufacturing 316L stainless steel. Key Eng Mater. 2010;419:413–416.

155.

Lippold

JC.

Welding metallurgy and weldability. Hoboken (NJ): John Wiley & Sons; 2015.

156.

Hafez

, Ghanem

, Morsy

MA.

The influence of shielding gases on solidification structures and grain size of AISI 304 stainless steel fiber laser welds. Laser Manuf Mater Process. 2019;6:345–355.

157.

Cieslak

, Ritter

, Savage

WF.

Solidification cracking and analytical electron microscopy of austenitic stainless steel weld metals. Weld J. 1982;61:1s–8s.

158.

Boettinger

, Coriell

, Sekerka

RF.

Mechanisms of microsegregation-free solidification. Mater Sci Eng. 1984;65:27–36.

159.

Aziz

Non-equilibrium interface kinetics during rapid solidification: theory and experiment. Mater Sci Eng. 1988;98:369–372.

160.

Bystram

MC.

Some aspects of stainless alloy metallurgy and their application to welding problems. Br Weld J. 1956;3:41–46.