Effect of Volumetric Energy Density and Part Height on the Material Properties of Low‐Alloyed Steels Manufactured by Laser‐Based Powder Bed Fusion of Metals

The layer‐by‐layer manufacturing approach in laser powder bed fusion of metals (PBF‐LB/M) leads to heat accumulation in the workpiece with increasing part heights. The effect of this heat accumulation on the resulting material properties has, however, only barely been studied for low‐alloyed steels. The goal of this work is to analyze the influence of different PBF‐LB/M‐specific boundary conditions like varying part heights and volumetric energy densities (VED) on the resulting material properties. It isfound that lower part regions possess similar hardness (380–410 HV1) and retained austenite values (7%–8%), independent of the applied VED. Higher energy inputs lead to higher retained austenite contents of up to 20% due to an incomplete transformation upon cooling. This rise in retained austenite content is also linked to a decreased material hardness down to 320 HV1. In higher part regions, this effect is reversed as the retained austenite content starts to decrease for the highest investigated VED. This is caused by the in situ preheating temperatures caused by heat accumulation, which favor a bainitic transformation. The part height‐specific properties indicate that the microstructure formation forms through a continuous transformation in lower part regions and through an isothermal transformation in higher regions.


Introduction
Low-alloyed steels have recently been the subject of investigations in laser-based powder bed fusion of metals (PBF-LB/M).Due to their good weldability, case-hardening steels are suited very well for the PBF-LB/M process.In combination with the extraordinary design freedom, highly sophisticated products can be generated by additive manufacturing.3] In the first work, Kamps [4] investigated the processing of 16MnCr5 by means of PBF-LB/M.The microstructure was described as fine-grained and resulted in an average hardness of around 330 HV1.More thorough investigations by Schmitt et al. [5] found a martensitic-bainitic microstructure with shares of ferrite.This conclusion on the underlying microstructure was performed based on the high process-specific cooling rates, which can be in the order of 10 3 to 10 5 K s À1 [6,7] in PBF-LB/M.Aumayr et al. [1] investigated the material properties of the low-alloyed steel Böhler E185 AMPO when processed by means of PBF-LB/M.This material could be processed successfully without larger defects.The resulting microstructure was described as bainitic-martensitic.It was also possible to process the tempering steel 30CrNiMo8 by means of PBF-LB/ M. [3] Due to the higher carbon concentration (%0.3 wt%) of this material, the substrate was preheated to 300 °C.The as-built tensile strength was almost as high as the one of the quenched and tempered reference specimens.In their work, Schmitt et al. [8] investigated the influence of different platform preheating temperatures on the material properties of 16MnCr5.For the lowest temperatures, a mostly ferritic structure was observed.

Increasing the platform temperatures leads to a transition toward
The layer-by-layer manufacturing approach in laser powder bed fusion of metals (PBF-LB/M) leads to heat accumulation in the workpiece with increasing part heights.The effect of this heat accumulation on the resulting material properties has, however, only barely been studied for low-alloyed steels.The goal of this work is to analyze the influence of different PBF-LB/M-specific boundary conditions like varying part heights and volumetric energy densities (VED) on the resulting material properties.It isfound that lower part regions possess similar hardness (380-410 HV1) and retained austenite values (7%-8%), independent of the applied VED.Higher energy inputs lead to higher retained austenite contents of up to 20% due to an incomplete transformation upon cooling.This rise in retained austenite content is also linked to a decreased material hardness down to 320 HV1.In higher part regions, this effect is reversed as the retained austenite content starts to decrease for the highest investigated VED.This is caused by the in situ preheating temperatures caused by heat accumulation, which favor a bainitic transformation.The part height-specific properties indicate that the microstructure formation forms through a continuous transformation in lower part regions and through an isothermal transformation in higher regions.a bainitic-ferritic and even pearlitic microstructure of the PBF-LB/M specimens.Wang et al. [9] processed the low-alloyed steel 24CrNiMo by means of PBF-LB/M.The resulting microstructure was found to be bainitic and the material possesses a high microhardness in the as-built state.Wei et al. [10] found that 24CrNiMo steels transform into bainite for high energy inputs due to heat accumulation during PBF-LB/M.The work of Bartels et al. [11] indicates that a bainite-like microstructure is formed when processing Bainidur AM by means of PBF-LB/M.In contrast to the hardened specimens, the additively manufactured material is characterized by a good tempering stability.This thermal stability indicated that the underlying microstructure is more thermally stable than martensite, which is typically associated with a hardness drop-off on tempering.Similar results were observed when processing Bainidur AM by means of DED-LB/M. [12]he microstructure formation of steels upon cooling from the austenitization temperature can be distinguished into continuous cooling and isothermal transformation. [13,14]A continuous cooling is typically defined by specific cooling rates. [15]artensite (fast), bainite (intermediate), and pearlite/ferrite (slow cooling) forms depend on the underlying cooling rates. [16]he isothermal transformation is characterized by a cooling to a defined temperature followed by an isothermal holding at this temperature for a defined time. [17]Thereby, the austenite is transformed in different microstructural constituents depending on the holding time at these temperatures.Martensite is typically formed through continuous cooling while bainite is formed through an isothermal transformation. [18]Transferring the fundamentals regarding the microstructure formation to PBF-LB/M, the ambient temperature of the workpiece gains importance.This temperature can be affected through either a (high-temperature) preheating [19,20] or heat accumulation [21,22] caused by the continuous energy input during build-up.
In their work, Mertens et al. [23] studied different preheating temperatures from 100 to 400 °C for the PBF-LB/M process.When preheating the substrate to a temperature of 400 °C, the martensite transformation is suppressed.This results in a homogeneously distributed bainitic microstructure.Saewe et al. [24] also applied a high-temperature preheating to improve the processability of the high-speed steel M50.The increased substrate temperatures affected the microstructure formation and resulted in a transition from an epitaxial (200 °C) toward a columnar dendritic microstructure (500 °C).The findings by Paravicini Bagliani [25] et al. also indicate that the degree of the retained austenite of low-alloyed steels is strongly affected by the temperature to which is the material is quenched.Already minor differences in temperature can significantly affect the retained austenite content of the final workpiece.Kunar et al. [26] further found that the partitioning time plays a major role regarding the formation of retained austenite.Regarding heat accumulation, Mohr et al. [27] found that the part height is highly influential on the phase formation.High volume energy densities (VEDs) and short interlayer times (ILT) favor an overheating of the to-be-produced parts with increased part heights.The high energy inputs further result in deeper melt pool depths, thus favoring keyhole porosity.This effect could be at least partially countered by reducing the energy input either by adjusting the VED or by increasing the interlayer times.An approach to counter this overheating by modulating the laser power was performed by Ettaieb et al. [28] for the titanium alloy Ti-6Al-4V.Consequently, the fluctuations of the melt pool size due to self-reinforcing overheating could be countered when using adjusted laser powers.Furthermore, a significant difference in subgrain size up to a factor of six depending on the applied parameter combinations was obtained within the investigations of Mohr et al. [27] .The corresponding effect on temperature development was proven in an additional work by Mohr et al. using thermography. [29]It is shown that the temperature within the workpiece can vary between approximately 150 and 600 °C depending on the parameter combination, chosen interlayer times, and part height.Williams et al. [30] developed an in situ thermography approach for determining the temperature development during the PBF-LB/M process.They found that the interlayer time significantly affects the surface temperature of the workpiece.Low ILTs can result in a drastic temperature increase for the same applied process parameters.From these studies, it can be assumed that the microstructure formation of low-alloyed steels could be altered extremely during the PBF-LB/M process as the different transformation intervals of the steels will be passed.Schmitt et al. [31] studied the influence of different part heights of up to 15 mm on the hardness of 16MnCr5 in their respective work.They found that the hardness decreased with increasing part heights and attributed this to grain coarsening effects.Furthermore, they stated that the microstructure transitioned from a martensitic toward a ferritic one.
The goal of this work is to study the influence of varying part heights and VEDs on the microstructure formation and the corresponding material properties of the low-alloyed steel Bainidur AM in PBF-LB/M.It is assumed that the transformation behavior is significantly altered by the energy input and transitions from a continuous cooling for small parts and low VEDs toward an isothermal transformation for larger parts and high VEDs.The investigations will be performed using the low-alloyed case-hardening steel Bainidur AM (1.7980; 18MnCrMoV4-8-7).By analyzing the geometry-specific as-built properties, we aim at expanding the application field of additively manufactured case-hardening steels to the use for construction parts due to their good combination of ductility and strength.

Experimental Section
All experiments were performed on an AconityMINI machine (Aconity 3D, Germany).Bainidur AM (Deutsche Edelstahlwerke Speciality Steel GmbH, Germany) was used as powder material with a nominal particle size from 15 to 45 μm.The particle size distribution (d 10% = 18.28 μm, d 50% = 32.72 μm, d 90% = 46.58μm) was determined using a CamsizerX2 (Microtrac Retsch GmbH, Germany).Table 1 lists the chemical composition of the Bainidur AM powder material according to the supplier's certificate.The moderate carbon content of 0.22 wt% indicates a good weldability of the material.Apart from that, the powder is characterized by minor concentrations of Si, Mn, Cr, and Mo, all in the range of around 1 wt%.Figure 1 presents the morphology of the Bainidur AM powder determined by means of scanning electron microscopy.The powder is mainly spherical with shares of potato-shaped particles.
The carbon content of the powder material was also validated using an elemental CS-Analyzer (ELTRA GmbH, Germany).A carbon content of 0.219 wt% was determined, meaning that no statistically relevant fluctuations of the carbon content could be observed.Furthermore, the powder was dried in a vacuum furnace at 110 °C for 12 h prior to the experiments to get rid of undesirable powder moisture, which might hinder the flowability and thus the coating operation.

Process Parameters
The parameter window was deduced using cubic specimens with a dimension of 10 Â 10 Â 10 mm 3 in previous studies. [11]From these results, three different volumetric energy densities (VEDs) were selected for the investigations on the influence of the part height on the corresponding material properties.The studied parameter combinations are listed in Table 2.
Laser spot size was maintained constant at around 105 μm.A constant layer thickness of 60 μm was used.The scanning direction was rotated by 67 °after every layer.All experiments were performed under argon gas atmosphere.

Sample Preparation and Analysis
The additively manufactured specimens were cut in half, embedded in an epoxy resin, grinded toward the center, and polished for analyzing the relative part density.Relative part density was determined using the same approach as presented in the study of Bartels et al. [11] Etching was performed with a nital solution (<5%) for the preliminary investigations.A Qness indentation tester Q10 Aþ (ATM Qness GmbH, Germany) was used for determining the Vickers Hardness (HV1).The microstructure was analyzed using optical light microscopy (Leica DM4 M) and scanning electron microscopy (Tescan Vega and Mira3 SEM).Retained austenite (RA) content was determined by means of X-Ray diffraction (XRD) using a D8 discover system (Bruker Corporation, Billerica, USA).
Tensile testing was performed using specimens according to DIN 50125:2009-07 using shape C with an inner diameter of 4 mm.Specimens with and edge length of 10 mm in xand y-direction and a total height of 60 mm were manufactured additively.Afterward, the parts were machined to the final geometry of the tensile sample.The experiments were performed on a QUASAR 100 (SCHÜTZ þ LICHT Prüftechnik GmbH, Germany).Elongation at break was determined using an extensometer.

Part-Height-Specific Microstructure
Specimens with different part heights were manufactured to assess the influence of the number of layers on the microstructure formation.Five different heights (10, 15, 25, 35, and 60 mm) were studied.The edge length in xand y-direction was maintained constant at 10 mm each.No additional minimum layer times were defined.Figure 2 displays the exemplary build jobs fabricated on the AconityMINI machine with different VEDs and part heights.
For these experiments, the hardness was measured along build direction.The distance between two measurement points in z-direction was set to 1 mm.At least three hardness measurements were performed at each z-position.Microstructural characteristics were obtained through optical light microscopy and SEM.Klemm-I etching was applied to reveal the microstructure.Additional XRD analyses were performed in five regions to obtain the corresponding retained austenite content, which is then correlated with the material hardness.To assess the tempering stability of the additively manufactured specimens, samples were heat-treated in a furnace of type N 30/85SHA (Nabertherm GmbH, Germany).The heat treatment parameters were chosen  based on a previous study. [11]Tempering was performed at 600 °C since the as-built material was known to withstand this temperature without significant hardness losses.Oxidation of the samples was avoided by preventing the experiments in a nitrogen gas atmosphere.

Results and Discussion
The applied process parameters were extracted from a previous study on the processability and the microstructure formation on Bainidur AM.Based on these findings, three different parameter combinations resulting a low, medium, and high VED were selected.In the first step (Section 3.1), the influence of different part heights on the microstructure formation will be analyzed.These investigations are followed by a study on the role of the volume energy density (Section 3.2) regarding the microstructure formation for specimens with a part height of 60 mm.

Influence of Different Part Heights on Material Properties
The part-height-specific microstructural properties are assessed for specimens with five different part heights up to 60 mm for a medium VED.All specimens were manufactured on the same build plate within one build job.The samples could be manufactured without larger noticeable defects, as indicated in Figure 3.
It is noticeable that the intensity of the etching increases along the build direction.This correlates with the increased weld depth which was observed for the 60 mm specimen.The corresponding energy input is higher than the energy output from the specimens due to heat flux or surface cooling.Furthermore, this overheating will also act as an in situ preheating during the fabrication process, as shown by Mohr et al. in their respective work. [29]As known from isothermal transformation diagrams for steels, elevated temperatures and holding times affect the resulting microstructure significantly.The average penetration depth of the single layers into the lower layer increases, as can be seen in Figure 3 (green, orange, and red boxes).This is the consequence of an increased preheating temperature of the previously manufactured layers, which results in a higher melt pool temperature and thus melt pool size during PBF-LB/M.It is therefore assumed, that the material properties, in this case the hardness, will also be affected by the altered transformation behavior.Figure 4 shows the results on the hardness gradient and the corresponding retained austenite content for the different part heights.
The material hardness decreased continuously for a higher number of manufactured layers.All specimens possess the highest hardness in the bottom region of the specimen (bottom 10 mm), typically in the range between 390 and 410 HV1.At a part height of around 30-35 mm, a hardness of around 360-370 HV1 was determined both for the samples with a part height of 35 mm and 60 mm.As the trend is similar for all specimens and independent of the part height, a thermallyinduced effect is assumed.Scanning electron microscopy (SEM) was used in the next step to analyze the microstructure within the bottom and top regions of the Bainidur AM samples.Figure 5 shows the respective microstructure in the different regions of the specimens.
The bottom area of the specimen is characterized by two distinct regions, the fusion zone and the heat-affected zone.Whereas the fusion zone is characterized by finely dispersed carbides within the laths similar to lower bainite or (tempered) martensite, the heat-affected zone is more similar to a degenerated upper bainitic-like structure due to the cementite between the ferritic structures.A more detailed characterization of the bainitic phase can be found in the study of Bartels et al. [11] However, the top regions are no longer characterized by this distinction between the fusion and heat-affected zone.Larger shares of austenitic isles can be identified (Figure 5a) that appear like granular bainite without precipitated carbides within the grains.This could be explained by the slower cooling rates in the top regions and the higher expected temperatures, which would favor the transformation into a globular bainitic or at least bainite-like structure. [32]Further regions can also be identified that seem like a degenerated upper bainitic structure (Figure 5b).However, this structure possesses larger areas of austenitic isles compared to the bottom region of the specimen (Figure 5c).The top region is further characterized by a most likely complete absence of carbides.However, a validation of the absence of these (if present) nanoscaled carbides in the different regions of the specimens (top and bottom) would require the use of a transmission electron microscope.
Due to the low amount of carbon and other austenite-stabilizing alloying elements, a ferritic transformation can be expected.The microstructure in low-alloyed steels was already proven to be mostly bainitic-martensitic for small specimens. [11]orrespondingly, the isothermal time-temperature-transformation (TTT) diagram was calculated for the low-alloyed steel Bainidur AM using JMatPro.The diagram is shown in Figure 6.A quenching temperature of 920 °C was chosen.Since PBF-LB/M is known to produce a very fine microstructure, an average grain size of 9 μm was chosen for the calculation.
The TTT diagram reveals that an isothermal transformation during PBF-LB/M requires a global temperature of at least 350 °C.Furthermore, the beginning of the bainitic transformation can start for extremely low holding times at around 4 s.At a temperature of around 450 °C, a complete isothermal transformation can take place within approximately 100 s.Both lower and higher temperatures will again require longer holding times to support a complete transformation of the austenite into bainite during cooling.A literature review shows that the part height significantly affects the surface temperature of the specimen.The findings of Mohr et al. [29] indicated that surface temperatures between 300 and 400 °C can be obtained during the PBF-LB/M process when applying a comparable medium VED (see Figure A1).Correlating this information with the obtained RA content, the resulting material hardness, and the microstructural characteristics obtained through SEM, an incomplete bainitic transformation is concluded.The incomplete transformation correlates well with the obtained microstructural characteristics (see Figure 5) even though  shares of (tempered) martensite and other phases cannot be concluded completely at this state.This assumption is based on two main points: on the one hand, the preheating temperature is too high to support the formation of martensite or bainite through continuous cooling.On the other hand, the holding times are too short to form 100% bainite by an isothermal transformation.
To validate this assumption, additional tempering investigations were performed.The hardness was measured on one-half of the specimen in the as-built state, while the other half was  exposed to a tempering step at 600 °C for 1 h before determining the hardness.Furthermore, the retained austenite content was determined both in the as-built state and in the tempered state.Figure 7 presents the propagation of the obtained hardness and the corresponding retained austenite content.The hardness in the lower half of the specimens is barely affected by the tempering heat treatment and falls between 375 HV1 and 400 HV1 until a part height of around 20 mm is reached.XRD measurements reveal an RA content of around 7% in the as-built state, which falls to below 2% after tempering.Here, a transformation of the retained austenite into bainitic structures is most likely taking place.After passing a part height of 25 mm, a difference between the material hardness in the as-built and tempered state can be identified.While the hardness of the as-built specimens drops to around 360 HV1, the hardness of the tempered samples remains constant around 375 HV1.This effect is also reflected in the retained austenite content, which increases to around 11% in the as-built state.Tempering led to a transformation of the austenite, resulting in an increased hardness.Moving to the upper half of the specimens, an even more significant difference could be identified.While the hardness decreases continuously in the as-built state and reaches a minimum of 320 HV1, a complementary increase of the retained austenite (up to 15% and 20%) content was determined.After tempering, the retained austenite content dropped to around 7% in both cases.Furthermore, the material hardness exceeded a minimum hardness of 350 HV1 in the tempered state.Through-hardening of the additively manufactured specimens leads to an average hardness of around 470 HV1. [11] can further be seen that the bottom region possesses a finegrained structure that appears like lower bainite.After tempering, this microstructure is coarsened and slightly resolved due to the high temperatures.These results support the assumption that the microstructure during PBF-LB/M is initially formed based on a continuous transformation at the absence of a platform preheating.This is supported by the fact that the material hardness in the lower regions is only barely affected by the cyclic energy input during the manufacturing process (see Figure 4).When building larger parts, an incomplete transformation of the austenite during cooling is present in the upper regions of the parts (see Figure 7).One reason for this is that the increased preheating temperatures during PBF-LB/M affect the transformation behavior, which is highly depending on the holding temperatures and times.Here, it is most likely that an isothermal transformation of the material takes place within the specimen.The increased retained austenite content in the top regions can then be explained by either too short holding times or too low holding temperatures during PBF-LB/M to assure a complete transformation into e.g., bainite.This is assisted by the finding on the tempering behavior, as the hardness increases by around 30 HV1 in the top regions when tempering the material at 600 °C for 1 h.This is contrary to the hardness in the bottom region, which was not affected by the tempering process.Correspondingly, the bottom region can either be tempered martensitic structure, a bainite-like structure, or a combination of both.The elevated temperatures of the specimen in the higher part regions might then result in locally altered chemical compositions, which would then affect the stability of the retained austenite by element saturation and diffusion.A longer holding time might help to further reduce the RA content, as shown in the study of Saha Podder and Bhadeshia. [33]

Influence of Different VEDs on Material Properties
Based on the previous findings, two different approaches will be followed.On the one hand, the energy input will be decreased by applying a lower VED.The goal of this parameter set is to investigate whether the continuous transformation and the associated increase of the retained austenite content can be prolonged with lower energy inputs.On the other hand, the energy input will be increased.The aim is to find out whether the higher VED results in an earlier overheating of the structure.Correspondingly, this overheating should result in higher holding temperatures and longer holding times during build-up, thus resulting in a more complete transformation of the retained austenite.All specimens were built with a part height of 60 mm. Figure 8 presents the cross-sections of the additively manufactured specimens along build direction for a low, medium, and high VED.
A high VED results in a reduced relative part density along build direction even though the bottom segments were characterized by a high relative part density.The specimens manufactured using the medium VED possess a mostly homogeneous density along build direction.A slight decrease could be identified along build direction.In contrast, the lowest VED is characterized by a homogenous relative part density for all regions within the specimens.The decrease in relative part density for increasing VEDs can therefore be attributed to a continuous overheating which must take place during build-up.This overheating can favor keyhole formation during PBF-LB/M, thus resulting in keyhole porosity, which is known for its spherical pores. [34]dditionally, the specimens manufactured with a low VED and high VED were etched for analyzing the microstructure.The obtained results are presented in Figure 8. Again, three different sections were studied per specimen.A clear difference in the etching behavior can be observed for the specimens manufactured with the two different VEDs.Whereas the weld track boundaries can be easily identified for the parameter set with a low VED in all regions of the sample, the geometry of the weld tracks cannot be determined any longer for elevated part heights when applying a high VED.It can be seen for the low VED that the average weld penetration depth increases in higher regions within the part.This is attributed to the elevated preheating temperatures of the underlying structure, which results in the larger weld track depths since the overall thermal energy is higher.
The different regions of the specimens are further characterized by a different etching behavior.For the lowest VED, a brownish etching can be identified within all regions of the sample.The darker regions at the edges of the weld tracks most likely resemble a martensitic or bainitic-ferritic microstructure. [35]This correlates well with the findings presented in a previous work for this class of low-alloyed steels. [11]Throughout all three images in Figure 8a, only minor whitish areas, which would resemble retained austenite, [36] can be identified within the cross-sections.A slight transition toward a bluish color is however visible in the top regions.The change in etching color can be explained by two factors.First, the lower regions are continuously tempered throughout the manufacturing process. [37]Second, the heat agglomeration during build-up affects the initial microstructure formation in the top regions. [29]This effect can be especially identified when considering the development of the etching colors for parameter VED High in Figure 8b.Here, a promoted bluish etching can already be observed within the center region of the specimen even though the bottom regions are similar independent of the applied VEDs.Furthermore, the amount of the brighter white regions presenting retained austenite tends to increase. [36]Moving further toward the top, the single weld tracks that form the final geometry can no longer be distinguished from one another.Here, a switch toward a blue-grayish microstructure is evident.This supports the assumption, that an excessive in situ preheating temperature is present during the build-up of larger components in PBF-LB/M.To better evaluate the effect of this in situ heat agglomeration on the material properties, the hardness was determined along build direction for the different VEDs. Figure 10 presents the findings on the material hardness as well as the corresponding retained austenite content.The material hardness in the bottom region is the highest for the lowest VED and decreases with increasing VED.Possible reasons for this a fine-grain-hardening effects.Lower VEDs are typically characterized by a faster scanning or lower laser power, both favoring higher cooling rates.The faster cooling again results in a finer grain, an effect commonly observed in PBF-LB/M.As for the medium VED, the lowest VED results in a continuous decrease in material hardness.The lowest VED is however characterized by a very slow decrease in material hardness.Even after manufacturing a part with a height of 60 mm, a hardness of around 380 HV1 was obtained in the top region.This correlates well with the measured RA contents of these specimens, which continuously increased from 7% in the bottom region to around 10% in the top region.The slower increase backs the assumption that the initial transformation of the austenite is performed through continuous cooling.The highest VED results in the fastest decrease of the material hardness.A minimal hardness of 320 HV1 was already reached after a part height of approximately 35 to 40 mm in the center region of the sample.However, the top region of the specimen is characterized by a higher hardness of up to 400 HV1, which is similar to the original hardness in the bottom region.
The evaluation of the material hardness also correlates with the retained austenite content.XRD measurements reveal the highest RA contents of around 17% in the center region of the specimen.The top region, however, is characterized by a lower RA content of 6%, similar to the bottom region, which also correlates well with the high material hardness.This supports the assumption that the higher holding temperatures for the highest VED positively affect the transformation of the austenite to a bainitic or bainite-like microstructure.Overall, the material hardness develops differently for the three investigated VEDs for an exemplary part height of 60 mm.This further manifests the theory that the phase formation needs to be distinguished into a continuous and an isothermal transformation.The retained austenite content is sensitive to the underlying quenching temperature. [25]Increasing temperatures due to heat accumulation could then affect the transformation of the retained austenite upon cooling.Considering the propagation of the retained austenite content along with increasing part height, a differentiation into the two mechanisms isothermal and continuous transformation is possible.A potential differentiation into these two mechanisms is shown in Figure 11.The different regions were defined based on the corresponding hardness and retained austenite contents (see Figure 10) as well as the underlying microstructure (see Figure 9).Since the AconityMINI machine does not provide a port for an infrared camera, the temperature values were approximated based on the work by Mohr et al. [29] In lower regions of the specimens, a continuous transformation of the austenite takes place.With increased part heights, the process intrinsic preheating temperatures of the specimens rise.This temperature rise is independent of the applied VED.These increased temperatures result in an isothermal holding of the workpiece.Since the transformation from austenite to, e.g., bainite is a time and temperature-dependent process, both the holding time as well as the holding temperature are decisive.Four different transformation mechanisms could be observed within the specimens: 1) Complete continuous transformation, 2) incomplete continuous transformation, 3) incomplete isothermal transformation, and 4) complete isothermal transformation.At this point, it needs to be mentioned that a clear differentiation between the mechanism (2) and ( 3), which describe an incomplete continuous and an incomplete isothermal transformation, is hardly possible.Applying the highest VED results in the fastest decrease of the material hardness, as previously shown in Figure 10.Based on the isothermal transformation diagram of Bainidur AM (see Figure 6) and the work of Mohr et al., [29] the underlying temperatures can be assigned to the different transformation regions.The highest hardness in the early stages (1) can be associated with an almost complete transformation of the austenite during cooling.This is mirrored in the microstructure (see Figure 11e), which appears lower-bainitic or temperedmartensitic. Results on the bainitic-martensitic microstructure in this region can be found in previous work. [11]After that, the second stage (2) is characterized by an increasingly incomplete transformation.The findings from the literature indicate that temperatures below 350 °C should be obtained in these regions.This temperature is below the bainite finish temperature (%360 °C) and above the martensite finish temperature of the material (%240 °C).Correspondingly, larger shares of retained austenite are present.The third region (3) is characterized by elevated temperatures that reach the ones required for the formation of lower bainite.However, the short holding times and the late onset of the transformation of the austenite into bainite (see Figure 2) might not be sufficient to completely transform the austenite into bainite during holding. [38,39]Thereby, the heat flux toward the substrate, especially between consecutive layers, will hinder the austenite from an almost complete bainitic transformation.The resulting microstructure (see Figure 11f ) is correspondingly characterized by high ratios of retained austenite, which are responsible for the reduced material hardness in this region.With increasing part height, a close to complete isothermal transformation is achieved (4).This is due to the expectable elevated temperatures that exceed 420 °C.According to the isothermal TTT of Bainidur AM, the required holding time falls in the range of approximately 100 s to undergo a complete transformation of the austenite.This helps in explaining the reduced austenite content and increased material hardness for higher part heights of the specimens manufactured with VED high.SEM images of this microstructure reveal a pronounced globular structure (see Figure 11g).This structure is characterized by a reduced share of the blocky austenite compared to the previous region.The globular structure might therefore be granular bainite, which forms at these elevated temperatures for comparable materials. [40]However, due to the high process-specific cooling rates, other constituents like a softer martensite that is formed upon the final cooling of the entire part cannot be ruled out.Reducing the VED shifts the underlying transformation mechanisms toward higher part heights and delays the onset of the isothermal transformation.For a medium VED (see Figure 11b), the fourth region, which is characterized by a close-to-complete isothermal transformation, is no longer evident.It can, however, be expected that this mechanism would occur when further increasing the part height.This is underlined by the fact that only regions (1) and ( 2) are present for the lowest VED.Here, the in situ preheating temperature is too low to even transition toward an isothermal holding.This is manifested by the comparably low retained austenite contents and the associated slow hardness drop-off.The findings by Mohr et al. further back this assumption since different maximum surface temperatures were obtained for comparable VEDs. [29]Processing the material with VED High leads to maximum surface temperatures of almost 500 °C after 1000 layers (= 60 mm).Applying a parameter comparable to VED medium lowers the maximum surface temperature to around 400 °C while VED Low would result in a temperature in the range of 300 °C.These temperature correlate with the expectable intervals for the austenite-to-bainite transformation within this work.An exemplary scheme of the propagation of the expectable surface temperatures is presented in Appendix A (see Figure A1).These findings support the assumption that the microstructure is initially formed through continuous transformation processes since only a minor heat accumulation is observed.With an increasing energy input and corresponding heat accumulation during build-up, the main transformation mechanism progresses toward isothermal holding.However, the transformation cannot be completed fully in lower part regions since the underlying holding temperature is not sufficient.This is associated with the lowest hardness values and the highest retained austenite contents.Moving toward larger parts, an almost complete isothermal transformation is observed.To manifest these findings, future research should focus on the application of a high-temperature preheating system to already start the build job within a temperature range at which an isothermal transformation takes place.

Tensile Properties of Additively Manufactured Bainidur AM
Furthermore, the tensile properties of the additively manufactured samples were analyzed.Vertically fabricated cubes were machined into tensile specimens.Five parts were manufactured within one build job for each parameter set.The obtained properties are listed in Table 3.
All specimens possess a similar ultimate tensile strength of 1200 AE 22 MPa (Low VED), 1222 AE 14 MPa (Medium VED),  and c) medium VED.d) Shows the arising shift of the transformation toward lower part heights when applying a higher VED.e,f,g) show the microstructure that is formed in the regions 1 (e), 2 þ 3 (f ), and 4 (f ).The underlying temperatures were estimated based on the finding by Mohr et al. [29] and 1203 AE 21 MPa (High VED), respectively.However, the yield strengths and elongation at breaks varied significantly throughout the different parameter combinations.All samples broke in the top third of the specimen, which could be expected due to the reduced material hardness in this region.The specimens manufactured with the lowest VED possess an excellent ductility of around 15.2 AE 1.7% and a yield strength of 997 AE 62 MPa.A medium VED resulted in an average ductility of 13.7 AE 0.6% and a yield strength of 900 AE 55 MPa.The highest VED also resulted the highest yield strength of 1097 AE 61.6 MPa.However, no precise information about the ductility can be provided at this point as the specimens tore outside of the region of the extensometer.The obtained ductility falls in the range of 5.4 AE 2.4%, which falls below the ductility of hardened steels.One potential reason for the reduced ductility with increasing VEDs can be found in the gradual overheating of the structures and the corresponding pore formation in higher regions of the part (see Figure 8).This explains the ductile breaking of the samples for the low and medium VED which developed toward a brittle fracture for the high VED due to the increased porosity in the top regions.Furthermore, the different tensile behaviors could be explained by mechanisms like work hardening. [41]This effect could be promoted differently depending on the underlying microstructure.During load, the retained austenite might transform into other microstructural constituents that possess a higher strength.This work hardening could explain the difference between yield strength and tensile strength for the different VEDs.Materials with a higher retained austenite content might favor the deformation-induced transformation of the retained austenite into fresh martensite, which is characterized by a higher strength. [42]Furthermore, the fine distribution of the retained austenite in different shares might have beneficiary effects on the material properties like ductility.Higher austenite contents could help to explain the differences in ductility, [26] as observed in this work.However, additional work on this is required for better understanding the tensile properties.Future work will address important issues like dimple formation, which can tremendously affect the fracture behavior. [43]These investigations will focus on the most promising process parameters to analyze the influence of parameters like varying grain size and austenite contents, which can be achieved by a tailored heat treatment.By comparing quenched and tempered specimens with bainitized and as-built ones, the goal is to identify the influence of the microstructure on the underlying fracture mechanisms.Overall, the tensile properties of Bainidur AM exceed the ones of additively manufactured 16MnCr5 [5] and are in the range of the quench and tempering steel 30CrNiMo8. [3]Even though all VEDs result in an inferior ultimate tensile strength compared to the results for 42CrMo4 shown by Damon et al. [44] and Shi et al., [45] a significantly higher elongation at break can be observed when using either a low or medium VED.

Conclusion
This work presents investigations on the influence of the part height and the energy input on the material properties of additively manufactured Bainidur AM.Increasing the part height results in a continuous overheating of the specimen.This overheating again affects the transformation behavior of low-alloyed steels during PBF-LB/M.The main findings of our work are: Increased part heights negatively affect porosity during buildup due to overheating, which is indicated by the increased weld penetration depth in higher part regions.
The process-intrinsic preheating results in a hardness decrease due to an incomplete transformation of the austenite during cooling for all VEDs.Higher retained austenite contents result in hardness drop-off by up to 25%.
Low VEDs result in a slow rise of the preheating temperatures of the additively manufactured specimens during build-up, leading to a mostly consistent hardness along the build direction with only a minor decrease in the average hardness.
High VEDs lead to a fast rise of the intrinsic preheating temperatures and thereby result in an altered microstructure associated with an increased retained austenite content and hardness drop-off.
Ultimate tensile strength is similar for all parameter combination.The main differences can be identified in the respective yield strength and ductility, which are affected by the applied VED and most likely the underlying fracture mechanisms.
Therefore, it is concluded that large specimens form their respective microstructure initially through a continuous transformation in the lower regions close to the substrate.In higher regions, the heat accumulation results in an elevated preheating.Consequently, the transformation behavior is affected, indicating that an isothermal transformation occurs in the upper regions of the specimens.Controlling the energy input during PBF-LB/M allows to manipulate this transformation behavior to generate parts with almost homogeneous material properties in the asbuilt state.

Appendix A
Figure A1 presents the expected surface temperatures.The VEDs used within this work were corrected since Mohr et al. [29] applied a layer thickness of 50 μm within their experiments.After correction, similar VEDs were obtained for the different parameters.Since a different material was used (316L), minor changes in the peak surface temperatures should be expected (e.g., due to altered heat conduction and energy absorption).
Figure A1.Estimated surface temperatures based on the findings by Mohr et al. [29] The VEDs applied within this work were corrected to a lower layer thickness (from 60 μm down to 50 μm) to match the formula.

Figure 1 .
Figure 1.Morphology of the powder material Bainidur AM with a nominal size from 15 to 45 μm for a) 100Â magnification and b) 500Â magnification.

Figure 2 .
Figure 2. Exemplary build job with cuboid specimens for analyzing the build-height-specific microstructure for a) solid cuboids processed with the three different VEDs and b) cubes with different part heights from 10 to 60 mm.Tensile specimens were manufactured from the same build job as in (b) with all specimens possessing a part height of 60 mm.

Figure 3 .
Figure 3. Cross-sections for specimens with different part heights manufactured using a medium VED.

Figure 4 .
Figure 4. Development of the material hardness depending on the respective part height for VED Medium.

Figure 5 .
Figure 5. Microstructure in the (a,b) top region of the specimen showing a) a globular structure with austenite isles and b) a degenerated upper bainitic-like structure as well as the region in (c,d) the bottom of the specimen characterized by c) a degenerated upper bainitic-like structure in the HAZ and d) a fine lower bainitic-or martensite-like structure in the fusion zone for specimens with a part height of 60 mm and a medium VED.

Figure 6 .
Figure 6.Isothermal time-temperature-transformation diagram for the low-alloyed steel Bainidur AM calculated using JMatPro.An average grain size of 9 μm and a quenching temperature of 920 °C were selected for calculation.

Figure 7 .
Figure 7. Experimentally determined a) material hardness and corresponding retained austenite content along build direction after tempering at 600 °C.b,c) Shows the microstructure in the bottom region in the as-built and tempered state.The specimens were manufactured on the AconityMINI machine using the medium (54.1 J mm À3 ) VED.

Figure 9 .
Figure 9. Cross-sections of different regions within the part manufactured with a) low VED and b) high VED.The specimens were etched using a Klemm-I reagent.

Figure 10 .
Figure 10.Experimentally determined values of material hardness (top) and retained austenite (bottom) content along build direction for different VEDs.The specimens were manufactured on the AconityMINI machine.The standard deviation is not shown for better perceptibility.

Figure 11 .
Figure 11.Differentiation into the four different transformation mechanisms during PBF-LB/M which are affected by the part height and the applied energy input, which can be distinguished for a a) high VED, b) low VED,and c) medium VED.d) Shows the arising shift of the transformation toward lower part heights when applying a higher VED.e,f,g) show the microstructure that is formed in the regions 1 (e), 2 þ 3 (f ), and 4 (f ).The underlying temperatures were estimated based on the finding by Mohr et al.[29]

Table 1 .
Chemical composition of the Bainidur AM powder batch according to the supplier's certificate.

Table 2 .
Investigated process parameter range for the fabrication of the specimens.

Table 3 .
Tensile properties of the additively manufactured specimens for three different VEDs (low, medium, and high).*Ductility of the specimens processed with the highest VED could not be obtained from the tensile test as the breaking occurred outside of the region observable with the extensometer.