WO2020222695A1

WO2020222695A1 - Process for producing a steel workpiece by additive powder bed fusion manufacturing, and steel workpiece obtained therefrom

Info

Publication number: WO2020222695A1
Application number: PCT/SE2020/050438
Authority: WO
Inventors: Zhijian James SHEN; Yuan Zhong
Original assignee: Shen Zhijian James; Yuan Zhong
Priority date: 2019-04-30
Filing date: 2020-04-30
Publication date: 2020-11-05
Also published as: SE1950517A1

Abstract

The present disclosure relates to a process for producing a steel workpiece. The process comprises the steps: adding one or more active metals or active metal precursors as secondary phases to a steel powder to produce a steel powder comprising active metal; and subjecting the steel powder to an additive powder bed fusion manufacturing process, thereby producing the steel workpiece; wherein by active metal it is meant any metal where the sacrificial reaction of oxygen with the active metal is favoured over the reaction of oxygen with the iron in the steel powder at the concentrations and conditions prevailing during the additive manufacturing process, and wherein the active metals or active metal precursors are added to the steel powder in a quantity sufficient to react substantially completely with total oxygen present during the additive powder bed fusion manufacturing process. The disclosure further relates to use of a steel powder comprising active metal in such a process, as well as a steel workpiece obtained by such a process.

Description

Process for producing a steel workpiece by additive powder bed fusion manufacturing, and steel workpiece obtained therefrom

TECHNICAL FIELD

The present disclosure relates to a process for producing a steel workpiece by additive powder bed fusion manufacturing. The disclosure further relates to use of a steel powder comprising active metals that oxidize in such a process, as well as a steel workpiece obtained by such a process.

BACKGROUND ART

Metal additive manufacturing, colloquially termed 3D printing, is capable of being used to fabricate a variety of components having structures of such complexity that they are unviable or impossible to produce by traditional casting or subtractive manufacturing (machining) methods. Metal additive manufacturing therefore has the potential to find utility in a large range of fields including medical implants, engines and automobile components.

Selective laser melting (SLM) is an additive manufacturing process utilizing a focused laser beam to melt and fuse metal powder together in a localised melt pool, thus forming solid metal components in one step. To date, the metals that have been demonstrated in SLM processes include steel, cobalt-chrome, inconel, aluminium and titanium powders.

A known drawback of SLM processes is however that components manufactured by the technique tend to have inconsistent and unreliable mechanical behaviour. This limitation hinders the more widespread implementation of SLM as a manufacturing technique.

There remains a need for a means for metal additive manufacturing that allows for the production of components with reliably excellent mechanical properties. SUMMARY OF THE INVENTION

The inventors of the present invention have identified shortcomings with prior art means of metal additive manufacturing. The inventors have identified that the inconsistent and unreliable mechanical behaviour of components manufactured by SLM is due at least partly to the presence of residual oxygen in the additive manufacturing apparatus during the additive manufacturing process. Although metal additive manufacturing is typically carried out under inert conditions or under vacuum (depending on the additive manufacturing method utilized), it is all but impossible to avoid the presence of at least some residual oxygen both from the processing chamber and the precursor powder. This residual oxygen may manifest as an oxide layer at the grain boundaries of the produced component, leading to weakened grain boundaries and degraded mechanical properties.

It is an object of the present invention to achieve a means of overcoming, or at least alleviating, the above-mentioned shortcomings. In particular, it is desired to provide a means of additive manufacturing that avoids or reduces the formation of oxides at the grain boundaries of a manufactured component.

These objects are achieved by a process for producing a steel workpiece according to the appended independent claim.

The steel workpiece produced by the process may have an elongation at fracture of greater than 80% as determined by the method of ASTM Test E8-16a, i.e. it has exceptional ductility.

The process comprises the steps of adding one or more active metals and/or active metal precursors as secondary phases to a steel powder to produce a steel powder comprising active metal; and subjecting the steel powder to an additive powder bed fusion manufacturing process, thereby producing the steel workpiece. The active metals and/or active metal precursors are added to the steel powder in a quantity sufficient to react substantially completely with total oxygen present during the additive powder bed fusion manufacturing process.

By steel workpiece it is meant any finished or unfinished component producible by metal additive manufacturing of a steel powder. For example, in order to produce a finished component from the workpiece, one or more further steps such as machining or coating may be required.

The inventors have discovered that the provision of active metal(s) in the steel powder consistently leads to the produced workpieces having exceptional ductile properties as defined above. Moreover, the produced workpieces also have consistently excellent strength, both yield strength and tensile strength.

Without wishing to be bound by theory, it is thought that the presence of one or more active metals in the steel powder allows the active metals to be sacrificially oxidized by oxygen present during the additive manufacturing process. The resulting active metal oxide particles are dispersed as nanoinclusions throughout the metal, accompanied with the consumption of oxygen contamination inside the grain boundaries and the elimination of large oxide precipitation that otherwise may form at the grain boundaries. This means that the grain boundaries are substantially free of oxide and are strengthened in comparison to oxide-coated boundaries. Instead of weakening the produced component, it is thought that the dispersed active metal oxides strengthen the metal component, and the resulting workpiece can thus be considered to be formed of an oxide-dispersion strengthened (ODS) alloy. By adding the active metals and/or active metal precursors to the steel powder it can be ensured that sufficient amount of active metals is present to react substantially completely with the oxygen present during the process, and that the desired chemical composition and amount of nanoinclusions can be controlled essentially independently of the composition of the steel powder and alloyed metals therein.

By an active metal it is meant any metal where the sacrificial reaction of oxygen with the active metal is favoured over the reaction of oxygen with the iron in the steel powder at the concentrations and conditions prevailing during the additive manufacturing process.

According to a further aspect, the objects of the invention are achieved by the use of a steel power as defined in the appended independent claim. The steel powder comprises one or more active metals and/or active metal precursors added as secondary phases, and is used in an additive powder bed fusion manufacturing process. Such use may be under an oxygen concentration sufficient to ensure substantially complete oxidation of active metals or active metal precursors. Such use may produce a steel workpiece having an elongation at fracture of greater than 80%, as determined by the method of ASTM Test E8-16a.

According to yet a further aspect, the objects of the invention are achieved by a steel workpiece as defined in the appended independent claim. The steel workpiece is produced by the process as defined herein.

These further aspects have the same advantages as described in relation to the inventive process as described herein, i.e. improved reliability, exceptional ductility and excellent strength of the produced steel workpieces.

The following features are applicable to the process, the use and the steel workpiece as defined in the appended independent claims, unless otherwise stated.

The steel workpiece may have an elongation at fracture of greater than 80%, as determined by the method of ASTM Test E8-16a, or may have an elongation of at least 85%, or in some cases an elongation of at least 90%.

The additive powder bed fusion manufacturing process may be selective laser melting, electron beam melting, selective laser sintering, selective heat sintering, or direct metal laser sintering. Processes whereby the metal powder is locally melted during processing, such as selective laser melting or electron beam melting, are preferred. Selective laser melting is especially preferred.

The steel powder may be any appropriate steel powder, such as an austenitic steel powder. Austenitic steel powders of the 300 series, such as 316L steel powders, are preferred.

The active metal may be the same as a metal alloyed in the steel powder. Alternatively, it may be one or more metals that are different to those metals alloyed in the steel powder. The active metals may be selected from Si, Cr, Mo, Ti, Ta and Nb and are added as elemental metals or active metal precursors. Silicon is preferred. The active metal precursors may be hydrides, carbides, nitrides, silicides or intermetallic compounds of the active metals.

The active metals and/or active metal precursors are added in a quantity sufficient to react substantially completely with oxygen present during the additive powder bed fusion manufacturing process. This assists in ensuring grain boundaries that are substantially free of large oxide particles in the workpiece and thus assists in improving the mechanical properties of the workpiece.

The steel powder may further comprise a nanoparticle metal oxide. The nanoparticle metal oxide may preferably be yttrium oxide nanoparticles. By adding a metal oxide nanoparticle secondary phase to the steel powder, a stronger oxide-dispersion strengthened (ODS) alloy may be obtained in the resulting workpiece, and the mechanical properties of the workpiece may be enhanced.

The steel workpiece may have a yield strength of at least 500 MPa, such as at least 550 MPa, as determined by the method of ASTM Test E8-16a. Such yield strengths are in excess of those typically obtained in steel workpieces from additive manufacturing. It is well established that there is typically a trade-off between strength and ductility, and steel workpieces combining both exceptional ductility and excellent strength are highly unusual.

The steel workpiece may have a tensile strength of at least 580 MPa, such as at least 590 MPa or at least 600 MPa, as determined by the method of ASTM Test E8-16a.

The additive powder bed fusion manufacturing process may comprise a step of rotating a scanning direction between adjacent layers. For example, the scanning direction of each layer may be offset from the immediately preceding layer by from about 45° to about 90°, such as about 67°. The microstructure and mechanical properties of the workpiece may be customized by offsetting the scanning direction in each layer in such a manner.

The oxygen present in the additive manufacturing process may be controlled to provide an oxygen concentration in inert gas from about 0.01 % to about 2.0 % by volume, preferably from about 0.05 % to about 1.0 %, even more preferably from about 0.1 % to about 0.8 %. By controlling the oxygen concentration, it can be ensured that the active metal oxide nanoinclusions are formed in the desired quantities, whilst avoiding the formation of large oxide particles and oxygen contamination at the grain boundaries.

In some embodiments, the active metals and/or active metal precursors are added to the steel powder in a quantity sufficient to react substantially completely with total oxygen present during the additive powder bed fusion manufacturing process.

Alternatively, or in addition, the oxygen present in the additive manufacturing process may be controlled to provide a total oxygen concentration sufficient to ensure substantially complete oxidation of the active metals and/or active metal precursors during the additive powder bed fusion manufacturing process.

That is to say that, in the workpiece, there may be a slight stoichiometric excess of active metals and/or active metal precursors, alternatively a slight stoichiometric excess of oxygen as oxygen contamination, or alternatively the active metals and/or active metal precursors may be stoichiometrically balanced with the total oxygen present.

Further objects, advantages and novel features of the present invention will become apparent to one skilled in the art from the following detailed description.

BRIEF DESCRIPTION OF THE DRAWINGS

For a fuller understanding of the present invention and further objects and advantages of it, the detailed description set out below should be read together with the accompanying drawings, in which the same reference notations denote similar items in the various diagrams, and in which:

Figure la schematically illustrates the layout of the specimens showing the dimension and laser scanning strategy.

Figure lb is a photograph illustrating the as-prepared tensile test specimens.

Figure lc is a photograph illustrating the machined specimens with gouge length of 10 mm and gauge cross-section of lxl mm for tensile tests.

Figure Id is a photograph illustrating the tensile test facility for small sized specimens. Figure 2a is a graph illustrating the cell spacing, density and Vickers hardness of the various small-scale specimens.

Figure 2b is SEM images of the etched surfaces showing the cell spacing decreases as scanning speed increases. Figure 3a shows a SEM image of PH45 at a first magnification.

Figure 3b shows a SEM image of PH90 at a first magnification.

Figure 3c shows a SEM image of PH45 at a second magnification.

Figure 3d shows a SEM image of PH90 at a second magnification.

Figure 3e shows a figurative illustration of the cell growth in the various layers of PH45. Figure 3f shows a figurative illustration of the cell growth in the various layers of PH90.

Figure 4a shows an EBSD orientation map of PH45 in a first direction.

Figure 4b shows an EBSD orientation map of PH45 in a second direction.

Figure 4c shows the inverse pole figures of the top surface of PH45.

Figure 4d shows an EBSD orientation map of PH90 in a first direction. Figure 4e shows an EBSD orientation map of PH90 in a second direction.

Figure 4f shows the inverse pole figures of the top surface of PH90.

Figure 5 shows the tensile engineering stress-strain curve of SLM specimens PH90 (blue), PH45 (red) and PV45 (green).

Figure 6a figuratively illustrates the fracture tip of PH45, with an OM image inserted in the figure.

Figure 6b is a SEM image illustrating the fracture tip of PH45 in a side view of fracture tip showing cell direction and melt pool boundaries.

Figure 6c is a SEM image illustrating the fracture tip of PH45 in a top view showing different fracture mode at rupture. Figure 6d is a SEM image illustrating the fracture tip of PH45 in a magnified images showing details.

Figure 6e figuratively illustrates the fracture tip of PH90, with an OM image inserted in the figure.

Figure 6f is a SEM image illustrating the fracture tip of PH90 in a side view of fracture tip showing cell direction and melt pool boundaries.

Figure 6g is a SEM image illustrating the fracture tip of PH90 in a top view showing different fracture mode at rupture.

Figure 6h is a SEM image illustrating the fracture tip of PH90 in a magnified images showing details.

Figure 7 shows the derived true stress-strain curves and the strain hardening curves of different specimens (PH90, PH45 and PV45).

Figure 8a is an EBSD image revealing the grain boundaries and the sub-grain boundaries.

Figure 8b is a SEM image of the etched surface at the same site showing the cellular structure and the melt pools.

Figure 8c is a magnified SEM image of site c as marked in Figure 8a.

Figure 8d is a magnified SEM image of site d as marked in Figure 8a.

Figure 8e is a magnified SEM image of site e as marked in Figure 8a.

Figure 8f is a schematic drawing of the hierarchical structures in AM materials.

DETAILED DESCRIPTION

Additive powder bed fusion manufacturing processes involve the layer-by-layer build-up of a workpiece by depositing a layer of metal powder, such as steel, and then precision melting selected areas of the powder layer using a heat source. For example, in selective laser melting (SLM) processes the heat source is a laser, typically focused by a lens and precision-directed using mirrors. A further layer of powder is then deposited and the melting step repeated in order to provide the next layer of metal fused upon the original layer. The deposition and melting steps are then alternatingly repeated until the desired metal workpiece has been completed. The workpiece is removed from the surrounding bed of non-fused powder and may then be subjected to further treatment in order to provide a finished component.

The processing parameters such as laser power, scanning speed, hatch spacing and layer thickness are known to affect the mechanical properties of the produced workpiece, and these parameters should be tuned for each powder in order to provide the desired optimized mechanical properties. For example, the laser power may be from about 150 W to about 300 W, preferably about 200W. The scanning speed may be from about 500 mm/s to about 7000 mm/s, preferably from about 700 to about 900 mm/s. The hatch spacing (line separation distance) may be from about 0.01 mm to about 0.15, preferably from about 0.06 mm to about 0.15 mm. The layer thickness may be from about 0.01 mm to about 0.03 mm, preferably about 0.02 mm.

The microstructure and mechanical properties of the workpiece may also be customized by changing the scanning direction in each layer. For example, the scanning direction of each layer may be offset from the immediately preceding layer by from about 45° to about 90°, such as preferably about 67°.

The present invention is based upon the discovery by the inventors that providing one or more active metals in the steel powder allows workpieces to be produced that have exceptional ductility and excellent strength. By an active metal it is meant any metal where a sacrificial reaction of oxygen with the active metal is favoured over the reaction of oxygen with the iron in the steel powder at the concentrations and conditions prevailing during the additive manufacturing process. Suitable active metals include, but are not limited to alloying metals such as Si, Cr, Mo, Ti, Ta and Nb. One or more of these metals may be present as an alloy metal in the steel powder.

The active metals and/or precursors of the active metals are added as secondary phases to the steel powder. By precursor of the active metal, it is meant a metal compound that undergoes thermal decomposition and sacrificial oxidation under the conditions prevailing during the additive manufacturing process to provide an oxide of the active metal. Advantages of using active metal precursors include, but are not limited to, easy to achieve homogeneous and fine distribution of the active metals in metal powder, reduction of oxidation risk during powder mixing process, and in-situ forming of more reactive active metals during the additive manufacturing process. Such precursors may include, but are not limited to hydrides, carbides, nitrides, silicides or intermetallic compounds of the active metals.

Although metal additive manufacturing is typically performed under inert conditions (inert atmosphere or vacuum), there is typically some residual oxygen present during the additive manufacturing process, for example in the inert atmosphere (which may e.g. have up to 10000 ppm oxygen) or associated with the steel powder. It is thought that this residual oxygen reacts with the steel powder during the additive manufacturing process, forming an oxide layer at the metal grain boundaries and thus weakening the produced workpiece. Without wishing to be bound by theory, it is thought that the presence of one or more active metals in the steel powder allows the active metals to be sacrificially oxidized by oxygen present during the additive manufacturing process. The resulting active metal oxides are dispersed as

nanoparticles throughout the metal and are not localised at the grain boundaries as large oxide particles or oxygen contamination. This means that the grain boundaries are

substantially free of oxide when a steel powder comprising one or more than one active metals is used, and the grain boundaries are therefore strengthened in comparison to oxide- coated boundaries. Instead of weakening the produced component, it is thought that the dispersed active metal oxide somewhat strengthens the metal component, and the resulting workpiece can thus be considered to be formed of an oxide-dispersion strengthened (ODS) alloy.

The active metals or active metal precursors are added as a secondary phase to the steel powder. Some quantity of active metals may also be alloyed in the steel powder. The active metals and/or precursors thereof is added or alloyed in quantities sufficient to react with substantially total oxygen present in the additive manufacturing process. This may be determined by the skilled person by using steel powders having a variety of concentrations of active metals (and/or precursors) and determining which powder provides workpieces having optimal mechanical properties. Alternatively, a theoretical quantity of active metal required may be determined from the stoichiometry of the reaction of the active metals with oxygen, provided that the amount of oxygen present is known and under control. The oxygen present in the additive manufacturing process may be residual oxygen only, or oxygen may be controllably added to the process to provide a desired oxygen concentration higher than the residual concentration.

The steel powder may comprise from about 0.1% to about 1.5% by weight of one or more active metals in total, such as from about 0.1% to about 0.9%. The total percentage by weight active metals is the sum of all active metals added as secondary phase and/or alloyed in the steel powder, relative to the total weight of the steel powder including all added secondary phases.

The invention will now be described in more detail with reference to certain exemplifying embodiments and the drawings. However, the invention is not limited to the exemplifying embodiments discussed herein and/or shown in the drawings, but may be varied within the scope of the appended claims. Furthermore, the drawings shall not be considered drawn to scale as some features may be exaggerated in order to more clearly illustrate certain features.

Examples

Example 1: Small-scale specimens (40 mm x 4 mm)

All specimens were prepared using a gas atomized spherical SS316L powder with particle size ranging from 10 to 45pm (Carpenter powder products AB, Torshalla, Sweden) as the precursor powder. The powder alloy comprises approximately 1% by weight silicon as the active metal.

The specimens were prepared by a commercial SLM system EOS M270 (EOS GmbH, Krailling, Germany). The experimental details of the various specimens are shown in Figs la-d with the processing parameters listed in Table 1.

Figure la schematically illustrates the layout of the specimens showing the dimension and laser scanning strategy, where BD represents building direction and TD represents transverse direction. Figure lb is a photograph illustrating the as-prepared tensile test specimens. Figure lc is a photograph illustrating the machined specimens with gouge length of 10 mm and gauge cross-section of lxl mm for tensile tests. Figure Id is a photograph illustrating the tensile test facility for small sized specimens. Table 1. Processing parameters of the test specimens and the calculated relative density (theoretical density of SS316L is 8000 kg/m³)

Specimen Laser Scanning Hatch Layer Energy Rotation Relative

No. power speed (v) spacing thickness density angle (°) density

(W) (mm/s) (d)(mm) (h)(mm) (J/mm³) (%)

Cl 195 7000 0.01 0.02 139 45 95.3

C2 195 4250 0.02 0.02 114.5 45 98.6

C3 195 1700 0.05 0.02 114.5 45 99

C4 195 850 0.1 0.02 114.5 45 99.2

C5 195 566 0.15 0.02 114.5 45 99.5

PH45 195 850 0.1 0.02 114.5 45 99.2

PH90 195 850 0.1 0.02 114.5 90 99.3

PV45 195 850 0.1 0.02 114.5 45 99

The laser melt traces were characterized on the etched surface by Light Optical Microscopy (LOM). Etched surfaces of both as-prepared specimens and the specimens after tensile tests together with the fracture surface were observed by JEOL JSM-7000F field emission scanning electron microscopy (SEM) (JEOL, Tokyo, Japan). Prior to microstructure observation, the specimens were mounted, polished and etched in 50 ml HF solution (HF:HN03:H20 = 2:6:42) for 5 mins. EBSD was performed on a HKL Nordlys orientation imaging microscope system (Oxford Instruments, Oxford, UK) equipped on a TESCAN MIRA 3LMH with a step size less than 1.5 pm.

The average cell spacing as illustrated in Fig. 2 was calculated by randomly choosing 10 sites with more than 1000 cells counted in each etched specimen. The densities of the specimens were checked by Archimedes method. The as-prepared specimens were machined to dog- bone shaped tensile test specimens as seen in Fig. lc. The gouge length was 10 mm and the cross-section size of tested part was around lxl mm. An extensometer was used to measure the elongation during tests as seen in Fig. Id. The tensile test directions with reference to as- prepared specimens were also indicated in Fig. la, where BD represents building direction and TD represents transverse direction. The reported values in this study for tensile properties were average value of 3 tests. Vickers Hardness tests were carried out at RT using a

Zwick/Roell ZHV indenter (Zwick/Roenhjll, Ulm, Germany) with a dual time of 10s. Five tests were performed for each specimen. A high load of 10 kgf was used to test the overall hardness and a low load of 100 gf was used to determine the local hardness.

According to solidification theory, AG/R (where AG is the temperature gradient and R is the solidification rate) determines the solidification mode (planar, dendritic or cellular) and as well as the dendritic arm spacing or the cell spacing. Both AG and R are related to the cooling rate, which can be tuned by the scanning speed (v) in SLM process. In this work, we firstly tuned the laser scanning speed and line spacing (d) to control the cell spacing of the sub-grain cellular structure. The energy density (w=P/(vdh), P is the laser power and h is the layer thickness) was kept at a roughly same level for all the samples.

Figures 2a and 2b illustrate the effect of the laser scanning speed on the cell spacing and consequently on the mechanical properties. Figure 2a is a graph illustrating the cell spacing, density and Vickers hardness of the various small-scale specimens. Figure 2b is SEM images of the etched surfaces showing the cell spacing decreases as scanning speed increases.

It has already been proved that the dislocation network appears together with the element segregation at the cell boundaries in our recent study. Therefore, the dislocation network characterization which needs great amount of TEM work can be replaced by observing element segregated cell boundaries on the polished & etched surface in SEM. Figure 2a shows the variation of cell spacing and the bulk density. The cell spacing decreased significantly as the increase of the laser scanning speed. Meanwhile, the bulk density of the specimens was also influenced by the different combination of v and d. Almost full density (with a relative density of 99.2%) without obvious macro defects was achieved in specimen C4.

Hardness test results were carried out with a relatively high load of lOkgf in order to avoid the influence of the local structure. The hardness increased with increasing density until it reaches 232(HV10) for sample C4 and then it dropped in C5 with even higher density (99.5%). The density was still the most critical issue of concern for many applications. But when the density variation becomes a negligible factor, for example in C4 and C5, the balance between a higher density and a smaller cell spacing should be considered. In other words, the scanning speed should be increased when defects are already controlled at a low level, which leads to a reduced cell spacing, higher hardness and probably higher strength.

The laser scanning strategy affects the arrangement of the cellular structure. One of the important parameters is the rotation of laser scanning direction between successive layers. In order to investigate this property, we produced two batches of specimens with different rotation angles of 45° (PH45) and 90° (PH90).

Figures 3a-3f illustrate the effect of the laser scanning strategy on the morphology and the arrangement of the cells. Figure 3a shows a SEM image of PH45 at a first magnification. Figure 3b shows a SEM image of PH90 at a first magnification. Figure 3c shows a SEM image of PH45 at a second magnification. Figure 3d shows a SEM image of PH90 at a second magnification. Figure 3e shows a figurative illustration of the cell growth in the various layers of PH45. Figure 3f shows a figurative illustration of the cell growth in the various layers of PH90.

No obvious difference on density or macro defect was observed in both specimens. However the length and the arrangement of the cells differed significantly. The cells were noticed frequently crossing the melt pool boundaries in PH45 (Fig. 3a, c), which resulted in

'continuous' longer cells. In contrast, most cells stopped at the melting pool boundaries and the trend of forming longer cells was hindered in PH90 (Figure 3b, d). The growth direction of the cellular structure is known to incline to the local AG and may also be influenced by the Marangoni convection in the melt pool [22, 23] and recent modeling work proved heat flow direction determines the solidification texture of sub-grain structure [18]. Following the grain growth, the cells have the chance to grow epitaxially at the melt pool boundaries and then change the mode to competitive growth away from the boundaries. The cells were more likely to form longer columnar if epitaxial growth is triggered, as illustrated in Fig. 3e. When the laser rotated dramatically (90°), the driving force for cells to incline to the scanning direction overcomes the tendency for epitaxial growth, which results in the non-continuous short cells in Fig. 3f. This difference of cell arrangement could presumably cause significant difference in mechanical properties.

Grain structure is also an important factor of consideration. EBSD analysis (Fig. 4a-f) was done on the two specimens. Figure 4a shows an EBSD orientation map of PH45 in a first direction. Figure 4b shows an EBSD orientation map of PH45 in a second direction. Figure 4c shows the inverse pole figures of the top surface of PH45. Figure 4d shows an EBSD orientation map of PH90 in a first direction. Figure 4e shows an EBSD orientation map of PH90 in a second direction. Figure 4f shows the inverse pole figures of the top surface of PH90.

Both specimens consisted of the columnar grains growing in the building direction. PH90 showed an ordered 'mosaic-like' pattern trapped in adjacent melt tracks (Fig. 4b) while the melt tracks were difficult to be distinguished in PH45 (Fig. 4a). While both specimens showed texture with preferred crystallographic orientation (101) along the building direction, PH45 showed slightly stronger texture as indicated by the IPF mapping in Figure 4c, f.

PH45 had slightly higher strength but a relative lower ductility compared with PH90. The average yield strength of both specimens was high due to the hierarchical structure. The tensile strength and yield strength for PH45 were 612 MPa and 474 MPa while those of PH90 were 555 MPa and 434 MPa, respectively. The average elongations at rupture (fracture) for PH45 and PH90 were 30% and 34%, respectively, which were lower than that of the counterpart fabricated by traditional methods. PH45 showed better tensile strength than PH90, which attributed to several microstructure features including grain morphology, texture and the dislocation network. The former two factors are often discussed in metals while the last one, dislocation network, does not exist in metals from most of the manufacture processes, therefore rarely discussed.

The difference of the cell arrangement resulted in different fracture modes under tensile.

Figures 6a-h illustrate the performance of the PH45 and PH9 samples during tensile testing. Figure 6a figuratively illustrates the fracture tip of PH45, with an OM image inserted in the figure. Figure 6b is a SEM image illustrating the fracture tip of PH45 in a side view of fracture tip showing cell direction and melt pool boundaries. Figure 6c is a SEM image illustrating the fracture tip of PH45 in a top view showing different fracture mode at rupture. Figure 6d is a SEM image illustrating the fracture tip of PH45 in a magnified images showing details. Figure 6e figuratively illustrates the fracture tip of PH90, with an OM image inserted in the figure. Figure 6f is a SEM image illustrating the fracture tip of PH90 in a side view of fracture tip showing cell direction and melt pool boundaries. Figure 6g is a SEM image illustrating the fracture tip of PH90 in a top view showing different fracture mode at rupture. Figure 6h is a SEM image illustrating the fracture tip of PH90 in a magnified images showing details. In the figures, the red lines refer to the melt pool boundaries and the yellow arrows indicate the growth direction of the cells.

A uniform deformation of cells across the melt pools were identified in PH45 (Fig. 6b) while distinct shorter cells were found in PH90 (Fig. 6f). The uniform deformation featured a delamination fracture (arrow in Fig. 6c) and the fracture of long cells (Fig. 6d) was directly observed. One should notice the frequently formed large crater on the fracture surface was absent indicating PH45 had high defect tolerance. By comparison, the delamination feature disappeared in PH90 and many large craters presented on the fracture surface (Fig. 6g). Larger stress was needed to tear apart cells as the dislocations were well pinned by the cell boundaries. Dislocation glided easier in PH90 when relative fewer dislocation walls were blocking their movements. Therefore, a most significant strengthening effect and a higher strength were expected in PH45. Similar phenomenon at the scale of grain level has long been known: more dislocation walls are encountered and a better strength is obtained when columnar grains are tensile tested perpendicular to the long columnar axis. [24, 25]

Specimen PV45 was prepared by the same process parameters and strategies as PH45 but with length of the specimen standing in building direction. The yield strength is 442 MPa and the tensile strength is 547 MPa, both values were lower than those of specimen PH45.

However, the average elongation at rupture (fracture) of PV45 reached 87% (with one specimen reaching more than 100%), which was much longer than that of PH45. This combination of strength and elongation of SLM SS316L was better than most of 316L fabricated by various methods. [26, 27]

The obtained strength was attributed to the presence of dislocation network structure. In-situ TEM observation of the compression tests proved the cell boundaries delayed the dislocation movement during deformation. [28] Removing cell structure by annealing leaded to a dramatic decrease in yield strength. [25, 29-31] The yield strength is related to cell spacing instead of the grain size following a Hall-Petch like relation. [12] Therefore, the tremendous dislocation cell boundaries played a more important role than the much fewer grain boundaries in determination of the strength. Similar relation as Hall-petch law between the local hardness and the cell spacing was found.

Many factors influence the ductility including grains, texture and also cellular structure. The grain morphology and size affects the free-path length of the dislocations during deformation. The grains and their texture also influence the deformation twinning generation. The tremendous pre-existing dislocations at the cell boundaries limit the capacity of strain hardening but meanwhile enable stable plastic flow. It is difficult to identify which factor takes the leading role in determining the ductility of SLM SS316L from the present experiments. However, we can get a clue from the dramatic ductility difference between PH45 and PV45.

Figure 7 shows the derived true stress-strain curves and the strain hardening curves of different specimens (PH90, PH45 and PV45). The strain hardening rate curves and the true stress-strain curves were derived to help understanding the ductility variation. PV45 has an obvious strain hardening rate recovery region. Twinning occur primary in grains close to <111> parallel to the tensile axis according to Schmid's law. But no preferred orientation <111> //tensile axis was observed in both samples according to the IPF mapping results (Fig. 4). Relative more grain boundaries along tensile axis in PH45 than PV45 increased the twinning stress and partially suppressed the twinning formation. [32, 33] It concluded that the grain size influenced the twinning generation and further the strain hardening capacity. In addition, the cell boundaries coupled with bundles of dislocations served as the nucleation sites for twinning. The cell boundaries also stabilized the dislocation network until large strain but couldn't fully block the dislocation motion at high stress levels. This resulted in a stable plastic flow and delayed the onset of necking. However, it was difficult to quantify the anisotropy level of cellular structure in different tensile directions. On the other hand, the grain boundary was known to be able to fully block the dislocation motion and deteriorate the ductility. [34] The columnar shaped grains generated much more high-angle boundaries perpendicular to building direction (PH45) than along building direction (PV45). The strain hardening capacity was thus lowered and ductility dropped dramatically in PH45. The present result proved that the grain size and the dislocation network both influenced the ductility, although further experiments are needed to clarify the leading factor. Figures 8a-f illustrate relations between the cellular structure, the grains and the melt pools. Figure 8a is an EBSD image revealing the grain boundaries and the sub-grain boundaries.

Figure 8b is a SEM image of the etched surface at the same site showing the cellular structure and the melt pools. Figure 8c is a magnified SEM image of site c as marked in Figure 8a. Figure 8d is a magnified SEM image of site d as marked in Figure 8a. Figure 8e is a magnified SEM image of site e as marked in Figure 8a. Figure 8f is a schematic drawing of the hierarchical structures in AM materials: Melt pool boundaries (blue dashed line), High angle grain boundaries (red line), low angle grain boundary (green line) and cell boundaries (black line).

The relation between hierarchical structures has not been investigated before. The cells were formed due to cellular grain growth under high AG/R value combined with element segregation at the solidification front. Therefore, the grain growth and the convection in the melt pools both influenced the arrangement of the cells. Here we introduced a new method by comparing the grain boundaries (shown by EBSD mapping) and the cell boundaries together with the melt pool boundaries (revealed by SEM on etched surface) at the same site (Fig. 8a, 8b). Different cases were marked in Figs. 8a-f and also summarized in Table 2. In general, cells were always similar in the same sub-grain (with low-angle grain boundaries) without crossing melt pool boundary but might be different in other cases. The cells were similar in adjacent sub-grains (case 3) due to the sub-grains texture caused by large AG at local site, in different melt pools (case 4) because of the epitaxial cell growth at melt pool boundaries. On the other hand, the cells more likely changed their arrangement if non- epitaxial grain growth occurred at the melt pool boundaries (case 6), which corresponds to the situation in PH90. In a word, any process parameter that changed the AG or the melt pool features influenced the cell formation and further the mechanical behavior.

Melt pool boundaries formed by layer by layer process are more like to accumulate defects.

[35] Scanning strategy should enable just enough overlapping of melt pools to minimize the amount of melt pool boundaries. Grain boundary is believed to have great influence on strength and ductility in traditional fabricated materials. The influence of grain boundary on ductility is still active but the tremendous cell boundaries vague its impact on strength. The cell boundaries benefit both strength and ductility at certain tensile axis. Careful control of the cellular structure makes it possible to fabricate customized materials by SLM. Table 2. The cell continuity regarding to sub-grains and melt pools (cases marked in Figs. 8a-f)

Same sub-grain Different sub-grains

Same melt pool Similar cells (case 1) Similar cells (case 2) or different cells (case 3)

Different melt Similar cells (case 4) or different Different cells (case 6)

pools cells (case 5)

In summary, the examples above assist in in understanding the sub-grain cellular dislocation network and its influence on the mechanical properties, and demonstrate how these properties can be manipulated using this understanding. The features of cellular structure under different processing conditions were predicted and proved by comprehensive microstructure and mechanical characterizations. At high density levels, minimizing the cell spacing increased the average hardness. The arrangement of dislocation network determined the fracture mode and impacted the tensile properties. A careful arrangement of continuous longer cells was demonstrated to generate a combination of superior strength and good ductility along the building direction.

Example 2: Larger-scale samples

For the as-built PV/PH specimens in Example 1, the produced specimen cross section is only l*4mm and the tensile test sample thickness is less than 1mm. In such small specimens the defects will play an unduly large role during the tensile tests and will result in unduly low tensile properties. In order to ameliorate the role of specimen size in the measured mechanical properties, a new series of specimens were prepared (S-series). The as-built S- series specimens were 8mm in diameter and the diameter of tensile test specimens (gauge length part) is 3mm (polished). S-series specimens are therefore larger and defects will not be so critical to the measured mechanical properties. In essence, S-series specimens provide a more realistic indication of the true mechanical properties obtainable using our additive manufactured 316L steel. All S-series specimens were prepared using the same steel powder as in Example 1. The laser power used was 195 W, a layer thickness of 0.02 mm was used, and a rotation angle offset of 45° was used between adjacent layers. Scanning speed and line spacing were varied as shown in Table 3 below. Table 3. Scanning parameters and obtained mechanical properties of S-series specimens

Sample Mechanical properties Scanning parameters

Yield strength Tensile Elongation Scanning speed / Line spacing /Mpa strength / Mpa /% mm/s / mm

51 491.4 594.7 60.9 1700 0,05

52 521,1 599.6 67.9 850 0,1

53 521,0 614,3 82,6 566 0,15

54 536.5 601,5 79,3 566 0,08

55 535,0 589,9 83.5 700 0,08

57 516,4 583.8 68,8 1000 0,08

58 528,9 596.2 80,7 566 0,1

59 521,1 603.3 74,1 566 0,12

510 529.6 634,1 82.6 800 0,15

511 543,3 601.7 85,5 700 0,1

It can be seen that the greater specimen sizes used in the S-series provides more consistent mechanical properties, and that the scanning parameters can be tuned to optimize the mechanical properties obtained. Elongation at fracture in excess of 70%, and in many cases in excess of 80% were obtainable by appropriate choice of scanning parameters.

Example 3: Fresh powder

The specimens of Examples 1 and 2 were produced using SS 316L powder recycled from previous additive manufacturing studies, and therefore potentially comprising a substantial degree of oxidation on the powder surface. In order to explore the properties obtainable under optimal conditions, a series of specimens were prepared using fresh SS 316L powder. The powder was otherwise the same as that used in Examples 1 and 2 and comprises silicon as the active metal. Specimens were prepared using 316L powder only (Normal), or with varying quantities of a nanoparticle (approx diameter 800 nm) Y2O3 secondary phase (ODS-1 and ODS-2). The laser power used was 195 W, a layer thickness of 0.02 mm was used, and a rotation angle offset of 67° was used between adjacent layers. Scanning speed and line spacing were varied as shown in Table 4 below.

Table 4. Scanning parameters and obtained mechanical properties of "fresh" -series specimens

Sample Scanning parameters Secondary Mechanical properties

phase %

(w/w)

Scanning Line Yield Tensile Elongation speed / spacing / strength / strength / /% mm/s mm MPa MPa

Normal 900 0.15 0 552 661 83.2

ODS-1 800 0.08 1 574 627 90.5

ODS-2 700 0.06 2 553 597 95.7

It can be seen that under optimized conditions a steel workpiece having an elongation in excess of 80%, a yield strength in excess of 550 MPa and a tensile strength in excess of 590 MPa may be obtained. It is though that the superior mechanical properties obtained are due to a powder having less initial oxidation, thus providing more evenly dispersed and smaller nanoinclusions, as well as fewer defects. Addition of a nanoparticle oxide secondary phase to the steel powder may improve the mechanical properties of the obtained workpiece, especially the elongation-to-failure. Elongation-to-failure as high as 95.7% was obtained using SS 316L powder in combination with 2 % w/w nanoparticle Y2O3 powder.

Example 4: Active metal precursor added powder

The powder mixtures were prepared by adding an active metal precursor, silicon nitride nanopowder, in weight percentage of 1.0% and 1.5% and mixed with the fresh SS 316L powder. Weight percentage is relative to the total weight of the steel powder including all added secondary phases. Specimens were prepared using 316L powder only (Normal), or the powder mixtures with the addition of silicon nitride of 1.0% ( ODS-3) and 1.5% (ODS-4). The laser power used was 195 W, a layer thickness of 0.02 mm was used, and a rotation angle offset of 67° was used between adjacent layers. Scanning speed and line spacing were varied as shown in Table 5 below. The oxygen concentration was regulated to 0.07% for reference 316L sample (Normal) and 0.1% for two active metal precursor added samples (ODS-3 and ODS-4).

Table 5. Scanning parameters and obtained mechanical properties of active metal precursor added specimens

Sample Scanning parameters Secondary Mechanical properties

phase %

(w/w)

Normal 900 0.15 0 552 661 83.2

ODS-3 800 0.08 1.0 592 655 82.5

ODS-4 700 0.06 1.5 565 623 80.2

It can be seen that under optimized conditions a steel workpiece having an elongation in excess of 82%, a yield strength in excess of 590 MPa and a tensile strength in excess of 650 MPa may be obtained. It is though that the superior mechanical properties obtained are due to the formation of homogeneously distributed oxide nanoinclusions with increased amount.

Claims

1. A process for producing a steel workpiece, the process comprising the steps:

- adding one or more active metals and/or active metal precursors as secondary phases to a steel powder to produce a steel powder comprising active metal; and

- subjecting the steel powder to an additive powder bed fusion manufacturing process, thereby producing the steel workpiece; wherein by active metal it is meant any metal where the sacrificial reaction of oxygen with the active metal is favoured over the reaction of oxygen with the iron in the steel powder at the concentrations and conditions prevailing during the additive manufacturing process, and wherein the active metals and/or active metal precursors are added to the steel powder in a quantity sufficient to react substantially completely with total oxygen present during the additive powder bed fusion manufacturing process.

2. The process according to claim 1, wherein the additive powder bed fusion manufacturing process is selective laser melting, electron beam melting, selective laser sintering, selective heat sintering, or direct metal laser sintering.

3. The process according to any one of claims 1-2, wherein the steel powder is 316L steel powder.

4. The process according to any one of the preceding claims, wherein the active metals are selected from Si, Cr, Mo, Ti, Ta and Nb.

5. The process according to any one of the preceding claims, wherein the active metal precursors are hydrides, carbides, nitrides, silicides or intermetallic compounds of the active metals.

6. The process according to any one of the preceding claims, wherein one or more active metals are added in a quantity sufficient to react substantially completely with total oxygen present during the additive powder bed fusion manufacturing process.

7. The process according to any one of claims 1-5, wherein one or more active metal precursors are added in a quantity sufficient to react substantially completely with total oxygen present during the additive powder bed fusion manufacturing process.

8. The process according to any one of the preceding claims, wherein the steel powder further comprises a nanoparticle metal oxide, and wherein the nanoparticle metal oxide is preferably yttrium oxide nanoparticles.

9. The process according to any one of the preceding claims, wherein the steel workpiece has an elongation at fracture of greater than 80%, and/or the steel workpiece has a yield strength of at least 550 MPa, and/or the steel workpiece has a tensile strength of at least 600 MPa, as determined by the method of ASTM Test E8-16a.

10. The process according to any one of the preceding claims, wherein the additive powder bed fusion manufacturing process comprises a step of rotating a scanning direction between adjacent layers.

11. The process according to any one of the preceding claims, wherein the oxygen present in the additive manufacturing process is controlled to provide an oxygen concentration in inert gas from about 0.01% to about 2 %, preferably from about 0.05 % to about 1.0 %, even more preferably from about 0.1 % to about 0.8 %.

12. Use of a steel powder comprising one or more active metals and/or active metal precursors added as secondary phases in an additive powder bed fusion manufacturing process for producing a steel workpiece.

13. The use according to claim 12, wherein the steel workpiece has an elongation at fracture of greater than 80%, as determined by the method of ASTM Test E8-16a.

14. A steel workpiece produced by a process according to any one of claims 1-11.