High-strength Damascus steel by additive manufacturing


Laser additive manufacturing is attractive for the production of complex, three-dimensional parts from metallic powder using a computer-aided design model1,2,3. The approach enables the digital control of the processing parameters and thus the resulting alloy’s microstructure, for example, by using high cooling rates and cyclic re-heating4,5,6,7,8,9,10. We recently showed that this cyclic re-heating, the so-called intrinsic heat treatment, can trigger nickel-aluminium precipitation in an iron–nickel–aluminium alloy in situ during laser additive manufacturing9. Here we report a Fe19Ni5Ti (weight per cent) steel tailor-designed for laser additive manufacturing. This steel is hardened in situ by nickel-titanium nanoprecipitation, and martensite is also formed in situ, starting at a readily accessible temperature of 200 degrees Celsius. Local control of both the nanoprecipitation and the martensitic transformation during the fabrication leads to complex microstructure hierarchies across multiple length scales, from approximately 100-micrometre-thick layers down to nanoscale precipitates. Inspired by ancient Damascus steels11,12,13,14—which have hard and soft layers, originally introduced via the folding and forging techniques of skilled blacksmiths—we produced a material consisting of alternating soft and hard layers. Our material has a tensile strength of 1,300 megapascals and 10 per cent elongation, showing superior mechanical properties to those of ancient Damascus steel12. The principles of in situ precipitation strengthening and local microstructure control used here can be applied to a wide range of precipitation-hardened alloys and different additive manufacturing processes.

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Fig. 1: DED-produced Fe19Ni5Ti (wt%) sample.
Fig. 2: Microstructure characterization at different length scales.
Fig. 3: APT analysis of martensite and austenite in the soft region and hard region.
Fig. 4: The effect of the thermal history.
Fig. 5: Tensile tests of two Fe19Ni5Ti (wt%) steel samples.

Data availability

The authors declare that the data supporting the findings of this study are available within the paper and its supplementary information and extended data files. Raw data are available from the corresponding author upon reasonable request.


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We are grateful to U. Tezins and A. Sturm for their support to the FIB and APT facilities at MPIE, to H. Faul and A. Jansen for their help with tensile tests, and to M. Adamek for his help with dilatometer experiments. A. Kwiatkowski da Silva and P. Bajaj are acknowledged for their input and discussions regarding thermodynamics and additive manufacturing respectively. We thank C. Brunner-Schwer for his support in conducting the DED experiments.

Author information




P.K. performed the microstructure analysis and corresponding data analysis including EDS, EBSD, FIB and APT and the analysis of dilatometer experiments and tensile tests. M.B.W. produced all samples used in this study by DED and acquired the experimental thermal profiles as well as the optical micrographs. E.A.J., A.W., B.G. and D.R. designed the study and acquired funding. P.K. wrote the initial draft. All authors contributed to reviewing and editing the manuscript and discussing and interpreting all the results.

Corresponding author

Correspondence to Philipp Kürnsteiner.

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The authors declare no competing interests.

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Peer review information Nature thanks Claire Davis and the other, anonymous, reviewer(s) for their contribution to the peer review of this work.

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Extended data figures and tables

Extended Data Fig. 1 Electron micrographs of the soft and hard regions.

In the SE micrographs, the hard region appears bright due to the rougher surface emitting more SEs, while the soft region appears darker due to the smooth surface (that is, opposite of how these two regions appear in optical micrographs). a, The interface between the soft and hard regions as well as a zoom to both regions. It is apparent that in the hard region there are two distinct phases: one with a rough surface, which is the martensite with (Ni,Fe)3Ti precipitates, and one with a smooth surface, which is austenite (and does not contain any precipitates). In the soft region, both austenite and martensite have a smooth surface as the martensite does not contain any precipitates in the soft regions. b, Further examples of the rough surfaces in hard regions at higher magnifications. The sample was slightly etched with 5 vol% nital for 10 s at room temperature.

Extended Data Fig. 2 Microstructure characterization.

High-resolution EBSD together with correlative EDS elemental mapping of a location within a hard region, showing that the austenite is enriched in Ti and Ni. a, SE micrograph (left) and EBSD mapping (right) of the same location. b, Overlapping the SE micrograph and the EBSD phase map from a show that austenite has a smooth surface and appears darker than martensite with a rough surface. c, The EDS element mapping shows that the smooth regions (that is, austenite) are enriched in Ti and Ni. These regions represent the interdendritic regions. d, EDS and EBSD are brought together: the area shown in d is marked by light blue dashed boxes in ac. It becomes apparent that austenite has a smooth surface and is enriched in Ti and Ni, whereas martensite has a rough surface and is depleted in Ti and Ni.

Extended Data Fig. 3 Thermodynamic calculation of the driving force for martensite formation.

The Gibbs free energies of single-phase bcc and fcc Fe–Ni–Ti at room temperature for a variable Ti content and a fixed Ni content of 19 wt%. It is apparent that there is a higher energy difference between the fcc and bcc structures (that is, a higher driving force for martensite formation) at lower Ti contents than at higher Ti contents. The two Ti compositions highlighted in the graph are 2.3 wt% Ti and 8.6 wt% Ti, which are typical compositions for the martensite/dendritic region and the austenite/interdendritic region, respectively, and for which the driving force for martensite formation is −2,100 J mol−1 and −1,780 J mol−1, respectively.

Extended Data Fig. 4 Serial sections through an APT reconstruction.

a, The (Ni,Fe)3Ti precipitates are marked by a 10 at% Ti isocomposition surface in the APT reconstruction. b, Consecutive slices through the same dataset from one side of the tip to the other. Each slice is 10 nm thick and all Ti atoms within this slice are shown. This sequence of images illustrates the complex shape and morphology and three-dimensional arrangement of the network of η-phase (Ni,Fe)3Ti precipitates created by IHT during the DED process. cf, Reconstruction of an APT volume that contains small spherical precipitates a few nanometres in diameter in addition to the plate-shaped interconnected network of precipitates. Both precipitate types are η-phase (Ni,Fe)3Ti and are marked by means of 10 at% Ti isocomposition surfaces. The plate-shaped network is depicted in dark green and the small spherical precipitates are depicted in light blue. In c, the whole dataset is shown, and in df, only a thin slice of 5 nm thickness is shown. In addition to the isocomposition surfaces, Ti atoms are shown in df. Panels e and f are enlarged sections of the image shown in c.

Extended Data Fig. 5 Composition of the η-phase precipitates.

a, A Ti atom map of a 5-nm-thick slice through the reconstructed volume from the martensitic phase in the hard region. The top part shows the atom map only; in the bottom part, precipitates are also highlighted by a set of isocomposition surfaces encompassing regions containing more than 10 at% Ti (dark green). b, A proximity histogram, that is, composition profile as a function of the distance to this isocomposition surface30, calculated for all imaged precipitates in the dataset. The average Ti content fits the expected 25 at% almost perfectly. Fe replaces some of the Ni from the prototype Ni3Ti phase, rendering it a (Ni,Fe)3Ti phase with approximately 6–7 at% Fe and 66–67 at% Ni.

Extended Data Fig. 6 Pseudobinary phase diagram for the Fe19Ni-xTi (at%) alloy.

The phase diagram was calculated using the Thermo-Calc software in conjunction with the TCFE7 database. The dashed line at 5 at% Ti highlights the phases that can be expected for the Fe19Ni5Ti (at%) steel used in this study: liquid, fcc A1 austenite, η-phase Ni3Ti, Laves phase and bcc A2 ferrite/martensite.

Extended Data Fig. 7 Determination of Ms.

A dilatometer curve acquired on a DED-produced Fe19Ni5Ti (wt%) sample (with no pause) is depicted. The double tangent method was used to determine Ms as 195 °C. Further dilatometer specimen from the same DED sample as well as DED samples with a 90-s pause time were measured. On all measurements, the martensite start temperatures are within 10 °C.

Extended Data Fig. 8 Experimental time–temperature profiles.

a, Experimental time–temperature profiles acquired with a pyrometer on the surface of the sample during the DED build at different pause times after each fourth layer. It becomes apparent that without pauses, the temperature increases during the entire fabrication and only drops notably when a pause time is introduced. The dashed orange line corresponds to Ms. b, Optical micrographs of the samples built with the corresponding pause times.

Extended Data Fig. 9 Tensile curves.

a, The testing direction is parallel to the laser scan direction. Tensile tests show a substantial improvement in strength as well as ductility when a pause is introduced in the manufacturing process. The pause allows the material to partially transform to martensite and then allows the IHT to trigger (Fe,Ni)3Ti precipitates in the martensite. The results show a few tensile specimens that fracture prematurely at low strains, which is due to additive-manufacturing-process-related defects31,32. These outliers rather represent the additive manufacturing process and show that there is potential for future process optimization. The specimens containing fewer defects and therefore higher strength and ductility show the actual potential of the newly designed maraging steel. In Fig. 5, we show two representative curves for each condition. For the condition ‘90-s pause each layer’, we omitted the one sample fracturing prematurely at 1.7% strain as well as the samples with the highest strength and lowest strength and show the two curves in between. For the condition ‘no pauses’, we omitted the two samples fracturing at the lowest strains of 4.5% and 6.7% as well as the samples with the lowest strength and highest strength. b, The testing direction is parallel to the build direction. Owing to limitations in the size of the DED-produced samples, we used small tensile specimens with a gauge length of 4 mm, a width of 2 mm and a thickness of 0.35 mm to test the tensile properties along the build direction (that is, perpendicular to the layered structure). The tensile specimens were machined with the gauge width parallel to the laser scan direction. There is a notable increase in strength and ductility due to the layered, Damascus-type structure. However, due to the smaller size, compared with the tensile specimen machined along the laser scan direction, a direct comparison between the two is difficult. Both, the Damascus-type layered steel as well as the one that was produced without pauses in between layers show higher strengths along the build direction than in the laser scan direction. While this could be due to the anisotropy of the material, the smaller tensile specimen geometry might also have a role.

Extended Data Fig. 10 Impact toughness.

a, The absorbed energies of subsized V-notch Charpy specimens at three different temperatures of −180 °C, 22 °C and 200 °C. The inset shows the geometry of the used subsized specimens. Charpy specimens were machined along the laser scan direction of the DED sample with the B direction normal to the layers and the b direction parallel to the layers. The values shown in the graph are an average of three specimens at 22 °C and two specimens at −180 °C and 200 °C. b, The values of the absorbed energy in joules in the Charpy V-notch impact testing carried out on subsized specimens shown in the inset in a. Two different normalizing factors are used to convert the results of the subsized specimen to standard specimen (55 × 10 × 8 mm3): the fracture area B × b and the fracture volume B × b2 (see, for example, refs. 33,34,35). It is noted that such normalizing factors are material dependent and there is no literature available on the selection of normalizing factors for additively manufactured maraging steels. The converted values presented in this table should therefore only be regarded as a rough estimate of the impact toughness on standard samples. Nevertheless, the Fe19Ni5Ti (wt%) samples investigated in this study show a high impact toughness compared with 4.9 J (standard V-notch samples) of laser-powder-bed-fusion-produced 18Ni-300 maraging steel in the aged condition (5 h at 480 °C)36.

Supplementary information

Supplementary Video 1

Atom probe tomography reconstruction. Supplementary Video 1 shows an atom probe tomography (APT) reconstruction spinning around the long tip axis (z-axis) showing the 3D view of the dataset containing η-phase precipitates shown in Fig. 4 (B) from within the dark layer of the Fe19Ni5Ti (wt%) sample. The precipitate phase is highlighted by a dark green isocomposition surface at 10 at% Ti.

Supplementary Video 2

Atom probe tomography reconstruction. Supplementary Video 2 shows an atom probe tomography (APT) reconstruction spinning around the long tip axis (z-axis) showing the 3D view of the dataset containing η-phase precipitates shown in Fig. S3 from within the dark layer of the Fe19Ni5Ti (wt%) sample. The precipitate phase is highlighted by an isocomposition surface at 10 at% Ti in blue for small spherical precipitates and in dark green for larger, rod and plate shaped precipitates.

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Kürnsteiner, P., Wilms, M.B., Weisheit, A. et al. High-strength Damascus steel by additive manufacturing. Nature 582, 515–519 (2020). https://doi.org/10.1038/s41586-020-2409-3

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