Cracking behavior and mechanical properties of austenitic stainless steel parts produced by laser metal deposition

Materials and Design 45 (2013) 228–235

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Short Communication

Cracking behavior and mechanical properties of austenitic stainless steel parts produced by laser metal deposition J. Yu ⇑, M. Rombouts, G. Maes VITO (Flemish Institute for Technological Research), Materials Technology, Boeretang 200, 2400 Mol, Belgium

a r t i c l e

i n f o

Article history: Received 5 July 2012 Accepted 31 August 2012 Available online 8 September 2012

a b s t r a c t The cracking behavior, microstructure and mechanical properties of austenitic stainless steel parts produced by laser metal deposition (LMD) are presented. The existing criteria for evaluating the solidification cracking sensitivity during welding of stainless steels have been adapted. Apart from the presence of sulfur and phosphorous, the presence of silicon was found to have a detrimental effect on cracking resistance. Cracking was not observed if the total content of sulfur, phosphorous and silicon was kept low enough, even not for stainless steels with austenitic solidification mode. Three-dimensional parts produced using optimal process parameters and feedstock powder composition have been investigated in more detail. The parts have a density of 99.6%. The microstructure consists of fine columnar dendrites, which coarsen as a function of height along the building direction. This is accompanied by a decrease in hardness. The tensile strength and elongation are in general higher than for annealed wrought material. The tensile strengths are higher while the elongation is lower for samples loaded perpendicular to the build-up direction than for those loaded parallel. Ó 2012 Elsevier Ltd. All rights reserved.

1. Introduction Additive manufacturing (AM) technologies produce three dimensional objects in an automatic process directly from a digital model by the successive addition of material, without the use of a specialized tooling. The process starts from digital data, most commonly in the form of a three-dimensional computer aided design (CAD) file, which is sliced electronically into a sequence of layers. For each layer two dimensional path information is defined and a program file for the AM machine is generated. On the AM machine a component is fabricated in a layer-by-layer fashion. In this study laser metal deposition (LMD) of stainless steel parts using a laser to melt powder transported by a coaxial nozzle is presented. This process is also referred to as laser engineered net shaping (LENS) [1,2], direct metal deposition (DMD) [3,4], laser solid forming (LSF) [5,6], laser cladding (LC) [7] and so on. Until now, laser metal deposition has been applied in many fields, such as aviation, navigation and automotive, for fabricating complex parts or repairing high-value components. A wide variety of materials has been used for this process, including austenitic stainless steels such as AISI 316L and AISI 304 stainless steel. During LMD the metal undergoes a rapid heating and cooling cycle, which leads to high solidification shrinkage stresses in the deposited layer. Austenitic stainless steels have a higher solidification cracking susceptibility than low-carbon steel ⇑ Corresponding author. Tel.: +32 14335696; fax: +32 14321186. E-mail address: [email protected] (J. Yu). 0261-3069/$ - see front matter Ó 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.matdes.2012.08.078

because of their higher thermal expansion coefficient and lower thermal conduction coefficient [5]. Most of the researches about solidification cracking of austenitic stainless steels focused on conventional welding applications [8– 12]. Many of the studies focus on the selection of solidification modes in relation to solidification cracking resistance. Those modes can be divided into four types according to the solidification behavior and the subsequent solid-state transformation [12]. They are respectively austenitic A (L ? L + c ? c), austenitic-ferritic AF (L ? L + c ? L + d + c ? c + d), ferritic-austenitic FA (L ? L + d ? L + d + c ? c + d) and ferritic F (L ? L + d ? d ? d + c) modes. Different calculation rules have been defined to predict the solidification modes based on the Cr- and Ni-equivalent of the material [13–16], as illustrated in Table 1. The ferrite stabilizing elements such as Cr, Mo and Si are included in the Creq and similarly austenite stabilizing elements such as Ni, Mn and C are included in the Nieq. It was found that the solidification modes vary from A to F with increasing Creq/Nieq ratio, as also presented in the Schaeffler [13] and WRC-92 [15] diagrams. In relation to solidification cracking, it was recommended to select steels with F solidification mode to prevent cracking during welding. Pellini [17] found that the F solidification mode provided a smaller critical temperature range for crack formation because of its smaller solidification temperature range than the A solidification mode. Borland and Younger [18] concluded that the beneficial effect of delta ferrite could be attributed to its higher solubility for impurities than austenite, which consequently leads to less interdendritic segregation and reduced cracking sensitivity. Hull [19] report that the low-

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J. Yu et al. / Materials and Design 45 (2013) 228–235 Table 1 Cr- and Ni-equivalent relationships for austenitic stainless steel. Common name and authors

Year

Creq wt.%

Nieq wt.%

Schaeffler [13] DeLong [14] WRC-1992 (Siewert) [15] HS (Hammar, Svennson) [16]

1949 1956 1992 1979

Cr + Mo + 1.5Si + 0.5Nb Cr + Mo + 1.5Si + 0.5Nb Cr + Mo + 0.7Nb Cr + 1.37Mo + 1.5Si + 2 Nb + 3Ti

Ni + 0.5Mn + 30C Ni + 0.5Mn + 30C + 30 N Ni + 35C + 20 N + 0.25Cu Ni + 0.31Mn + 22C + 14.2N + Cu

Table 2 Experimental processing parameters (P, V, D, F, g, Qs and hc are the laser power, scan speed, spot diameter, powder feed rate, overlap, shield gas flow rate and layer thickness, respectively). Set

P (W)

V (mm/ min)

D (mm)

F (g/ min)

g

Qs (L/ min)

hc (mm)

1 2

570 750

750 1000

1.2 1.2

2 2.7

50% 50%

8.5 8.5

0.5 0.5

er surface energy of the d–c boundary compared to d–d or c–c enhances the interfacial stability and is an important factor for decreasing cracking susceptibility. In addition, the presence of elements like sulfur and phosphorous plays a significant role in the cracking behavior of austenitic stainless steels. This is because these elements have a low solubility in the major constituent of stainless steel and form low melting eutectic phases with iron, chromium and nickel at the interdendritic region or along the grain boundaries [20,21]. The segregation tendency at this location is high due to the wide solid–liquid range and low eutectic temperature [10]. Therefore, in the evaluation criteria for cracking sensitivity of austenitic stainless steel the total content of sulfur and phosphorous should be included together with the predicted solidification mode. This is accomplished by the Suutala diagram [22] and the improved Suutala diagram which is developed by Pacary for pulsed laser-beam welding [23]. In addition, Folkhard et al. referred that stainless steels with (P + S) content 6 0.02%, or (P + S) content 6 0.03% and ferrite number P 4, (P + S) content 6 0.04% and ferrite number P 8, (P + S) content 6 0.05% with ferrite number P 12 were not susceptible to solidification cracking during welding [20]. The varestraint test is another practical and more direct way to evaluate cracking sensitivity, including transvarestraint test (TVT) and longitudinal varestraint test (LVT). The total crack length (TCL) got in LVT and brittleness temperature range (BTR) got by the maximum crack length (MCL) shown in TVT can be used for assessment. Lundin modified the Suutala diagram on basis of the TCL from the varestraint test [24] and he specified that for ‘TCL > 2.5 mm’, ‘1.5 mm < TCL < 2.5 mm’ and ‘TCL < 1.5 mm’ the material is respectively highly susceptible, susceptible and not susceptible to solidification cracking. The LMD process differs from conventional welding in different aspects, such as the powder feeding and high solidification and cooling speed. Song et al. studied the cracking mechanism during LMD of stainless steel [25]. A higher sulfur, phosphorous and silicon content was detected in the interdendritic regions, which resulted in solidification cracking due to the separation of the liquid film under the action of high tensile stresses. In the present paper, the cracking behavior of austenitic AISI 316L and AISI 304 stainless steel with different chemical compositions is studied. Appropriate solidification cracking criteria, which should give guidance for materials selection for AM of austenitic stainless steel, are proposed. The microstructure and mechanical properties of parts fabricated by LMD are presented.

2. Experimental setup The LMD experiments were carried out using a 7 kW IPG fiber laser with out-coupling fiber with a diameter of 600 lm. The use of a focal lens with focal length of 250 mm and collimator lens with focal length of 125 mm results in a laser spot diameter of 1200 lm on the substrate. The laser spot has a top-hat energy distribution. The powder is transported through a continuous coaxial nozzle (Fraunhofer-Institut für Lasertechnik) using argon transport and shielding gas. Samples were built on stainless steel AISI316L flat substrates with a thickness of 8 mm. In addition, SKM-DCAM software is used for generating the CNC programs with specified tool paths. For each layer first scanning of the contour was programmed. Afterwards the interior was scanned using a raster deposition pattern, which was rotated 90° between each layer, with 0.6 mm (50%) overlap between raster paths and 0.3 mm overlap with the contour. For choosing appropriate processing parameters, preliminary experiments encompassing laser cladding of single lines, layers and 3D parts have firstly been carried out using different processing parameter combinations, followed by the analysis of interior quality of the specimens by checking their cross sections. The adapted experimental processing parameters used in this paper are listed in Table 2, in which, P, V, D, F, g, Qs and hc are the laser power, scan speed, spot diameter, powder feed rate, overlap, shield gas flow rate and layer thickness, respectively. The chemical compositions of the powders referred in present study are listed in Table 3. The last three are only used for complementary analysis on the cracking behavior of austenitic stainless steels during LMD. The microstructure was studied by optical and scanning electron (Jeol JSM-6340F) microscopy. The samples were etched using a solution of hydrochloride acid, nitric acid and Vogel’s sparbeize. Electron Probe Microanalysis (EPMA, JEOL JXA-8530-F) using wavelength dispersive spectrometry (WDS) has been performed to analyze the local chemical composition. The mechanical properties were determined by tensile testing (Instron 5582) according to ASTM Standards (ASTM: E8/E8M-11). Three specimens produced under identical conditions have been subjected to tensile testing. The hardness was obtained by Vickers indentation measurements using a load of 1 kg.

Table 3 316L Stainless steel powders with chemical compositions.

Nr.1 Nr.2 Nr.3 Nr.4 [26] Nr.5 [26] Nr.6 [6]

C

S

P

Si

Mo

Mn

Ni

Cr

Fe

0.02 0.025 0.019 0.072

0.01 0.011 0.006 0.005

0.01 0.03 0.017 0.12

2.23 0.5 0.53 1.32

2.44 2.5 2.1 2.87

0.08 1.4 1.4 0.06

11.41 12.7 10.4 14.49

16.44 17.1 16.8 15.5

bal. bal. bal. bal.

0.034

0.002

0.033

0.44

2.34

0.24

14.31

17.17

bal.

0.062

0.012

0.032

0.34

0

0.88

8.23

17.88

bal.

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(b)

(a)

2mm

2mm

Fig. 1. Cross sections of the samples fabricated using powder 3. (a) Parameter set 1. (b) Parameter set 2 (Table 2).

Fig. 2. Cross sections of the sample fabricated using powder 1 (parameter set 1, Table 2). At the right side images at high magnification near the cracks are shown.

Fig. 3. Cross sections of the sample fabricated using powder 2 (parameter set 1, Table 2). At the right side images at high magnification near the cracks are shown.

3. Results and discussions 3.1. Cracking behavior For powders 1 and 2 micropores and microcracks are present after LMD in contrast to for powder 3 (Figs. 1–3). At a higher magnification (areas 1 and 2 in Figs. 2 and 3), it can be seen that the cracks are formed in the interdendritic regions. Micropores are nearly spherical in shape and are caused by the entrapment of gas in the melt pool [26]. The porosity can be eliminated by increasing laser power.

It is known from literature studies in the field of welding that the solidification cracking sensitivity of stainless steels depends on their chemical composition. The solidification modes for the different powders listed in Table 3 are calculated using different rules (see Table 4). For a given powder the same solidification mode is obtained using the different calculation rules. The powders 1, 2, 3 and 6 are characterized by a ferrite–austenite (FA) solidification mode, powder 4 by a austenite (A) solidification mode and powder 5 by a austenite–ferrite (AF) mode. Fig. 4 is obtained by considering the evaluation criteria of Suutala, Pacary, Folkhard and Lundin and the experimental observations. The powders that after LMD re-

Table 4 Solidification modes for different powders predicted using different calculation rules. Diagrams

Materials

Nr.1

Nr.2

Nr.3

Nr.4

Nr.5

Nr.6

Schaeffler [13]

Creq Nieq Creq/Nieq Mode

22.225 12.05 1.84 FA

20.35 14.15 1.44 FA

19.695 11.67 1.69 FA

20.35 16.68 1.22 A

20.17 15.45 1.31 AF

18.39 10.53 1.75 FA

WRC-92 [15]

Creq Nieq Creq/Nieq Mode

18.88 12.11 1.56 FA

19.6 13.575 1.44 FA

18.9 11.065 1.71 FA

18.37 17.01 1.08 A

19.51 15.5 1.26 AF

17.88 10.4 1.72 FA

HS [16]

FN Creq Nieq Creq/Nieq Mode

5 23.1278 11.874 1.948 FA

3 21.275 13.67 1.56 FA

9 20.472 11.238 1.82 FA

0 21.4119 16.092 1.33 A

1 21.0358 15.13 1.39 AF

7 18.39 9.86 1.87 FA

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J. Yu et al. / Materials and Design 45 (2013) 228–235

(a)

(b)

Suutala diagram, for conventional welding

(c)

Improved Suutala diagram by Pacary, for laser welding

(d)

Improved Suutala diagram by Lundin

Folkhard

Fig. 4. Experimental results for different powders together with different evaluation criteria according to (a) Suutala and Moisio [22], (b) Pacary et al. [23], (c) Lundin et al. [24] and (d) Folkhard [20]. Numbers marked in blue refer to powders resulting in cracks after LMD and in red to powders resulting in crack-free parts.

sulted in cracks (powders 1, 2 and 4) are marked in blue while the others (powders 3, 5 and 6) are marked in red. The following observations can be made: (1) The experimental results are not totally consistent with the evaluation criteria applied for welding technology. The improved Suutala diagrams by Pacary et al. [23] and by Lundin et al. [24] correspond better with the experimental observations than Suutala and Folkhard diagrams. Powders 2, 3, 4 and 6 are evaluated correctly. (2) The experimental observation of cracking for powder 1 is not in agreement with the predictions presented in all the diagrams. It should be noted that this powder contains more silicon than others. It is known that silicon can form lowmelting eutectic phases such as Fe–Fe2Si, NiSi–Ni3Si2 and NiSi-c in the interdendritic region and along the grain boundaries [20,21], which greatly increases the cracking susceptibility. This is confirmed by electron probe micro analysis (EPMA) of a LMD manufactured part using powder 2 with normal Si content (0.5%). Even with normal Si content, a significantly higher Si content, apart from higher S and P contents, is detected inside the crack (point 2 in Fig. 5) than outside the crack (point 1 in Fig. 5), as shown in Table 5. Therefore, powder 1 with much higher Si producing cracks not as predicted illustrates the important role of

Fig. 5. Locations of points used for elemental analysis by EPMA for the part fabricated using powder 2. (Points 1 and 2 are located respectively outside and inside the crack.) Table 5 Comparison of S, Si and P contents (in wt.%) inside and outside the crack for the part produced using powder 2 (see Fig. 5). Element

S

Si

P

Outside crack Inside crack

0.075 0.552

0.592 1.185

0 0.008

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silicon: the presence of cracks at high Si content reveals that for this chemical composition the effect of silicon to result in low-melting eutectic phases is more decisive than the effect of silicon to act as a ferrite stabilizer. (3) Powder 5 did not result in cracks while according to its chemical composition a high cracking probability is predicted. The ratio of Creq to Nieq for powder 5 is only 1.3, which predicts a full austenitic solidification mode. The high cracking resistance of this material is attributed to the low content of P, S and Si. Table 6 shows the cracking behavior

Table 6 Cracking behavior of different powders related to impurity content. Powder number

S + P + Si (wt.%)

S+P (wt.%)

N content

Cracking

Fabrication method

1

2.25

0.02

0.047

Yes

4

1.445

0.125

–

Yes

3 2 5 6

0.553 0.541 0.475 0.384

0.023 0.041 0.035 0.044

0.062 0.09 – –

No Yes No No

Water atomized Water atomized Gas atomized Gas atomized Gas atomized Gas atomized

(b)

(a)

Fig. 6. Cross sections of the samples fabricated using powder 3. (a) Parameter set 1, Table 2. (b) Parameter set 2, Table 2. Areas 1, 2 and 3 corresponding to the high magnification images of marked 1, 2 and 3.

(a)

(b) 2

1

Fig. 7. Tensile testing samples. (a) Built in ‘lying’ direction, loaded perpendicular to the building direction (outer dimensions x  y  z: 100  20  10 mm3). (b) Built in ‘standing’ direction, loaded parallel to the building direction (outer dimensions x  y  z: 10  20  100 mm3).

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[28]

[27]

[28]

[27]

Fig. 8. Tensile testing results for the samples fabricated by LMD using powder 3 and parameter set 1 (Table 2). (a) Yield strength and ultimate tensile strength. (b) Elongation.

(a)

(b)

Fig. 9. Cross sections of the sample (shown in Fig. 7b) fabricated using powder 3 in standing orientation.

of different powders related to impurity content of S + P + Si. The cracking sensitivity decreases with decreasing S + P + Si content. (4) A high nitrogen content is reported to be detrimental to solidification cracking resistance [9,11]. Therefore, the N content of powders 1, 2 and 3 was tested using instrumental gas analysis (IGA). The presence of cracks after LMD of powder 2 is attributed to the combination of relatively high N content (0.09 wt.%) and of S + P + Si (0.541 wt.%) content. 3.2. Microstructure and mechanical properties In this section only analysis results of three-dimensional parts produced using powder 3, which did not reveal any cracks after LMD, are presented. Three-dimensional samples of 5 mm height have been fabricated using the processing parameters listed in Table 2. Fig. 6a shows the cross section of the deposited part using parameter set 1. The dendrites change growing direction considerably on the middle top of the melt pool due to a change in the heat flow direction. This area is remelted when using a higher heat input as shown in Fig. 6b. Fig. 7a and b shows the samples, which were fabricated along respectively ‘lying’ and ‘standing’ orientations, that have been subjected to tensile testing. The loading direction is respectively perpendicular and parallel to the build-up direction for Fig. 7a and b. The densities of the samples measured by Archimedes method are in the range of 99.5% and 99.6%. Fig. 8 illustrates the tensile test results with respect to yield strength, ultimate tensile strength and elongation. Values reported for cast and annealed wrought material are included as a reference [27,28]. The mechanical properties are higher than for cast and annealed wrought material. The samples loaded parallel to the buildup direction (‘standing’ orientation) have a lower yield strength, ultimate tensile strength and higher elongation. This trend is also

reported in other studies [29]. One of the contributing factors might be the orientation of the grain/dendrite boundaries. After LMD most of the dendrites are oriented along the building direction. Upon loading parallel to the building direction into the plastic deformation region less barriers (dendrite/grain boundaries) need to be crossed by the dislocations [30]. This is reflected by a lower strength when loading parallel to the build-up direction. Another factor affecting the tensile strength is the size of grains/dendrites, as specified in the Hall–Petch relationship: rs = r0 + kyd 1/2, with rs the yield strength, r0 a materials constant for the starting stress for dislocation movement (or the resistance of the lattice to dislocation motion), ky the strengthening coefficient (a constant unique to each material), and d the average grain/subgrain diameter [31,32]. The dendrite spacing in the loaded region is larger for the sample loaded parallel to the build-up direction (Fig. 7b). For this sample an increase in dendrite spacing from 1–2 lm up to about 3–4 lm is observed from bottom to top of the sample (Fig. 9). This difference is also reflected by the hardness measurements: the hardness near the bottom and top are respectively 205 ± 10 HV and 174 ± 10 HV. The lower hardness is attributed to the lower cooling rate near to the top as a result of heat accumulation in the slender sample. For the sample loaded perpendicular to the build direction the cooling rate in the loaded region is larger and consequently a smaller dendrite spacing and higher strength are obtained. Fig. 10 shows the SEM fractographies of the samples after tensile testing. Dimple morphology dominates on the fracture surface of all samples, which is indicative of a ductile fracture mode. In the dimples inclusions, which contain relative high amounts of oxygen, chromium, silicon and iron as revealed by Energy Dispersive X-ray analysis (Fig. 10c), are present. The size of the inclusions is larger for the sample fabricated in ‘standing’ orientation: the inclusions are around 0.8 lm (Fig. 10b) compared to 0.3 lm for the sample fabricated in ‘lying’ orientation (Fig. 10a). The spacing between

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Fig. 10. Fractography after tensile testing of the sample built in (a) ‘lying’ and (b) ‘standing’ orientations. (c) EDX analysis results for the inclusions.

the inclusions is larger for the sample fabricated in ‘standing’ orientation. Goodwin et al. [33] pointed out that the inclusion size has no effect in the ease of void nucleation in the weld material but the larger inter-inclusion spacing can result in higher ductility. This is consistent with the trend in mechanical strength.

nealed wrought material. Samples fabricated in ‘lying’ orientation show higher strengths and lower elongation than slender samples fabricated in ‘standing’ orientation. This is attributed to the differences in dendrite/grain boundary orientation relative to the loading direction and in overall grain size as well as the interinclusion spacing.

4. Conclusion

References

The cracking behavior of austenitic stainless steel powders AISI 316L and AISI 304 processed by LMD was investigated. Cracks are observed in the interdendritic regions depending on the chemical composition of the powder. The experimental observations are not completely consistent with the cracking sensitivity criteria developed for welding. It was found that the silicon content was another critical parameter apart from sulfur and phosphorous to predict the cracking behavior. The total content of silicon, sulfur and phosphorous should be evaluated when assessing cracking sensitivity because of their tendency to form low melting phases in the interdendritic region. A high N content can increase cracking susceptibility in combination with a high S + P + Si content. The microstructure after LMD of AISI 316L consists of fine columnar dendrites. A density of 99.6% has been obtained. The hardness is about 30 HV higher and the microstructure finer near the baseplate than at the top because of the higher cooling rate. The tensile properties after LMD are higher than for cast and an-

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