Rapid fabrication of Nb-Si based alloy by selective laser melting: Microstructure, hardness and initial oxidation behavior

    Rapid fabrication of Nb-Si based alloy by selective laser melting: microstructure, hardness and initial oxidation behavior Yueling Guo, Lina Jia, Shaobo Sun, Bin Kong, Jinhui Liu, Hu Zhang PII: DOI: Reference:

S0264-1275(16)30941-8 doi: 10.1016/j.matdes.2016.07.048 JMADE 2052

To appear in: Received date: Revised date: Accepted date:

21 May 2016 10 July 2016 11 July 2016

Please cite this article as: Yueling Guo, Lina Jia, Shaobo Sun, Bin Kong, Jinhui Liu, Hu Zhang, Rapid fabrication of Nb-Si based alloy by selective laser melting: microstructure, hardness and initial oxidation behavior, (2016), doi: 10.1016/j.matdes.2016.07.048

This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

ACCEPTED MANUSCRIPT Rapid fabrication of Nb-Si based alloy by selective laser melting: microstructure, hardness and initial oxidation behavior

School of Materials Science and Engineering, Beihang University, Beijing 100191, People’s

SC R

a

IP

T

Yueling Guoa, Lina Jiaa*, Shaobo Suna, Bin Konga, Jinhui Liub, Hu Zhanga

Republic of China

Modern Manufacturing Engineering Center, Heilongjiang Institute of Science and Technology,

Harbin 150027, People’s Republic of China

NU

b

MA

*Corresponding author: Tel: +8610 82316482; E-mail: [email protected] Abstract:

D

Nb-18Si-24Ti-2Cr-2Al-2Hf (at.%) alloy was fabricated by selective laser melting (SLM). The

TE

microstructure, hardness and oxidation behavior of Nb-Si based alloy via SLM were investigated. Results showed that the relative density of SLM alloy under the optimized processing parameters

CE P

was 98.27%. The as-built alloy consisted of Nb solid solution (Nbss), αNb5Si3, βNb5Si3 and Nb3Si. Sphere shaped Nbss phases with a maximum diameter of 300 nm were obtained. Between the neighboring Nbss phases distributed the strip shaped α/β Nb5Si3 phases. After heat treatment at

AC

1500 °C for 4 h, the Nbss phases interconnected to form a continuous matrix with discontinuous αNb5Si3 phases. The average hardness of the SLM parts decreased from 810HV0.1 to 542HV0.1 after heat treatment. Compared with the oxide scale on arc-melted alloy, oxide scale formed on SLM-processed after oxidation at 1300 °C for 0.5 h was more compact with fewer holes, combined with the formation of a more continuous SiO2 layer at the initial oxidation stage. A layer-structured oxide scale model was presented to describe the high-temperature oxidation behavior of Nb-Si based alloy processed by SLM. Keywords: Nb-Si alloy; Selective laser melting; Rapid solidification; Hardness; Oxidation

1

ACCEPTED MANUSCRIPT 1 Introduction With remarkably higher operating temperatures than advanced Ni-base superalloys, Nb-Si based

T

ultrahigh temperature alloys show great promise for applications as the next generation turbine blade

IP

materials [1, 2]. They are designed as in situ composites, composed of Nb solid solution (Nbss) and

SC R

Nb5Si3 phases. Nbss offers room-temperature ductility and silicides supply high-temperature strength and creep resistance. To achieve a property balance, elements such as Ti, Cr, Hf and Al have been added and striking progress has been achieved on improving the overall properties of Nb-Si based

NU

alloys [2-4]. But the mechanical and environmental properties of Nb-Si based alloys still need further improvement for application, which will be most probably provided by innovation in processing and

MA

production technology.

Over the past few years, a variety of processing schemes have been explored for the fabrication of Nb–Si based alloys, including vacuum consumable or non-consumables arc-melting,

D

powder-metallurgy, physical vapor deposition, spark plasma sintering and different kinds of

TE

directional solidifications [2, 5, 6]. Laboratory-scale volumes of Nb-Si alloys have been able to be

CE P

fabricated for composition optimizing and microstructure controlling. And the Nb-Si alloy prototype airfoil could be fabricated by a hybrid arc-melting and drop-casting technique, by overcoming the difficulties of developing the melting capability of the appropriate scale and the matching ceramic

AC

mold materials [2]. However, obtaining the proper dovetail geometry of an airfoil is a complex endeavor, and requires numerous time-consuming, post-machining operations. More state of the art techniques are in dire need of development to fabricate turbine blades from Nb-Si alloys. In the light of additive manufacturing, selective laser melting (SLM) possesses the superiority of cost reduction and design-freedom [7-10]. Furthermore, with a rapid melting and solidification process, SLM allows endowing materials with superior properties by tailoring the microstructure. It typically refers to grain refinement, one important method for strengthening alloys without losing toughness [11]. The AlSi10Mg alloy produced by SLM exhibits superior strength and elongation properties than the die cast counterpart [10]. The 316L stainless steel parts built by SLM have a high hardness of 281HV0.1, a large tensile strength of 590 MPa and an elongation rate of 21.1% [9]. The last three decades have witnessed the feasibility of building a 3D part from a range of powders, including stainless steel, titanium alloys, nickel base superalloys, etc [10, 12-14]. A few reports on 2

ACCEPTED MANUSCRIPT the SLM-processing of metals with high melting point have been performed [15, 16]. Thijs et al. [15] reported the SLM-processing of Ta powders and almost fully dense and strong parts were obtained.

T

However, there are few reports concerning the SLM processing of Nb-Si based alloys.

IP

The poor oxidation resistance at high temperature has challenged the application of Nb-Si based

SC R

alloy [4, 17, 18]. Apart from the additions of alloying element and the development oxidation-resistant coating, the oxidation resistance of metals is typically improved by grain refinement, especially nanocrystallization or microcrystallization, by prompting selective oxidation

NU

of active alloying element and enhancing oxide scale adhesion [19, 20]. Microstructure refinement could be typically achieved via SLM, which is expected to greatly alter the high-temperature

MA

oxidation behavior of Nb-Si alloy. Therefore, it would be vitally important to investigate the oxidation behavior of Nb-Si alloy processed by SLM.

In this work, the Nb-Si based alloy with a novel refined microstructure was obtained by SLM

D

using pre-alloyed powders. Microstructure and hardness of the SLM Nb-Si based alloys were studied,

TE

in addition to the effect of heat treatment (HT). The oxide scales grown on arc-melted and

CE P

SLM-processed Nb-Si based alloys after oxidation at 1300 °C for 0.5 h were characterized. The corresponding high-temperature oxidation mechanism was discussed in detail. 2 Material and methods

AC

The nominal composition of the pre-alloyed powders used as base materials for SLM was Nb-18Si-24Ti-2Cr-2Al-2Hf (at.%). The powders were produced via the jet milling process and were less than 50 μm in size. As schematically illustrated in Fig. 1 (a), the SLM system mainly consisted of an IPG YLR-500 Yb:YAG fiber laser, an automatic powder delivery system, an oxygen detector and a controlling system. The laser scanning pattern was shown in Fig. 1(b). The substrates were pre-grounded Ti6Al4V coupons (65 mm × 65 mm × 10 mm). In the SLM process, powder spreading and melting were repeated to prepare an additive-layered rectangular part. To avoid any oxidation, the building chamber was first evacuated and then filled with Ar atmosphere. The residual oxygen level was below 0.1 wt.%. Based on a series of preliminary experiments, the following optimized SLM parameters were obtained: laser power = 375 W, scanning speed = 1000 mm/s, hatching distance (h) = 0.10 mm and layer thickness = 0.05 mm. 3

ACCEPTED MANUSCRIPT The density of the as-built part was determined using the Archimedes method. To examine the microstructure stability, the SLM parts were further heat treated in the vacuum furnace under high purity argon atmosphere at 1500 °C for 4 h. Hardness is known to be a material property that

IP

T

indicates resistance to plastic deformation and correlates directly with its strength, wear resistance, etc. To compare the mechanical properties of the SLM and SLM+HT Nb-Si alloy, a microhardness

SC R

tester (FM-800) with a 100 g load and a dwell time of 10 s was used to measure the Vickers hardness. More than ten sites were tested on polished horizontal sections of the SLM part, i.e., perpendicular to the building direction. As the potential next generation turbine blade materials, the service

NU

temperature of Nb-Si based alloy was expected to be as high as 1300 °C. Therefore, to investigate the initial oxidation behavior of Nb-Si based alloys, the arc-melted and SLM-processed Nb-Si based

MA

alloys were oxidized in a tube furnace at 1300 °C for 0.5 h, followed by air cooling. (b)

CE P

TE

D

(a)

AC

Fig. 1 Schematics of SLM system (a) and laser scanning pattern (b).

Top surfaces of the SLM parts and the oxidized specimens were observed using a field-emission scanning electron microscopy (SEM, Quanta 200F). The phases were identified by X-ray diffraction (XRD, D/max-2500, Cu Kα). Microstructure and composition distribution were investigated using SEM with backscattered electron imaging (BEI) and electron microprobe analysis (EPMA, JXA-8230) with wave dispersive spectroscopy (WDS). The microstructure was further investigated using a transmission electron microscopy (TEM, JEM-2100). Thin foils for TEM observation were cut and ground to 80 μm, then dimpled to 10 μm, and finally ion-milled. 3 Results and discussion 3.1 Microstructure characterization 4

ACCEPTED MANUSCRIPT Using the given processing parameters, Nb-Si based alloy can be successfully fabricated by SLM (Fig. 2 (a)). The starting powders with irregular shapes are shown in Fig. 2 (b). The average density of as-built components is 6.27 g/cm3 (95% confidence interval 6.24-6.29), and the relative

IP

T

density is as high as 98.27%, indicating a quite low porosity (1.73%). Top surface of the SLM part is featured by the continuous scanning tracks with sound intertrack bonding (Fig. 2 (c)). When scanned,

SC R

the powders are melted and solidified to form a molten pool. Part of the deposited track is remelted when the next track scanned, i.e., the two adjacent molten pools are characterized by an overlap. This will be favorable to the formation of a dense layer with metallurgical bonding between two adjacent

NU

tracks, and resultant a dense SLM component. As shown in Fig. 2 (d) as well as the inset with higher magnification, the microstructure on the vertical section, i.e. parallel to the building direction, is

MA

featured by half-cylindrical molten pool boundaries (MPBs), in accordance with other literatures concerning the SLM processing of different materials [21, 22].

D

The balling phenomenon has been eliminated by optimizing processing parameters.

TE

Unfortunately, cracking still takes place, though processing parameters are optimized (Fig. 2 (d)). Considerable residual stress is inclined to be introduced in the laser-based SLM process, caused by

CE P

large thermal gradient [23]. The residual stress is typically released by cracking, when the tensile stress exceeds the ultimate tensile strength of brittle Nb-Si based alloy with a ductile-brittle transition temperature (DBTT) higher than 1000 °C [5, 24]. For the SLM-processing of brittle ceramics, one

AC

effective method of reducing cracks has been reported to be preheating the powders prior to laser scanning at temperatures higher than 1500 °C [25]. The high-temperature preheating tends to lower the energy input of the laser beam and resultantly weaken the concentration of residual stress [25]. Therefore, in terms of the SLM-processing of brittle Nb-Si based alloys, it may be able to eliminate the cracks by preheating to a certain temperature, and more work will be performed in regard to eliminating the cracks during SLM.

5

ACCEPTED MANUSCRIPT (a)

SC R

IP

T

(b)

10 mm

(d)

MA

Scanning Vector

NU

(c)

D

Cracks

TE

200 μm

CE P

Fig. 2 (a) the as-built Nb-Si based alloy component by SLM; (b) Nb-Si based alloy powders; top view (c) and side view (d) of the as-built component by SLM.

As revealed by the XRD patterns (Fig. 3 (a)), the as-built alloys are mainly composed of

AC

αNb5Si3, βNb5Si3, Nb3Si and Nbss, while the arc-melted alloys consist of αNb5Si3 and Nbss. It suggests that the metastable Nb3Si and βNb5Si3 phases have been formed during the process of SLM. After the HT at 1500 °C for 4 h, the microstructure evolves into a mixture of Nbss and αNb5Si3. The absence of βNb5Si3 indicates the occurrence of the transition of βNb5Si3→αNb5Si3+Nbss during HT, since the αNb5Si3 phase is more stable at room temperature [2]. The Nb3Si phase yielded during rapid solidification is considered to have a relatively low thermodynamic stability and tends to decompose through a eutectoid reaction during HT. Microstructures of the arc-melted, SLM and SLM+HT Nb-Si alloys are shown in Fig. 3 (b), (c) and (d), respectively. The arc-melted alloy is featured by the existence of large primary Nb5Si3 phases, up to a maximum of 50 μm in size. Owing to the rapid cooling rate induced by the laser beam, the SLM process endows Nb-Si based alloy with a much finer microstructure (< 1 μm) and the 6

ACCEPTED MANUSCRIPT absence of large primary silicides. The titanium dissolved in Nb3Si may be beneficial to stabilize the Nb3Si phase [26]. Microstructure of remelted border zone (RBZ) is observed to be coarser than that of other areas (Fig. 3 (c)). After HT, the Nbss phases interconnect to form a continuous matrix and

T

fine interrupted αNb5Si3 phases are obtained (Fig. 3 (d)). The αNb5Si3 phases display a maximum of

IP

5 μm in size. Besides, it is noting that RBZ has been eliminated by high-temperature HT, i.e., the

SC R

microstructure of Nb-Si alloy processed by SLM is homogenized by the HT at 1500 °C for 4 h. The equilibrium microstructure of Nb-18Si-24Ti-2Cr-2Al-2Hf (at.%) alloy is believed to be

NU

composed of Nbss and αNb5Si3 (Fig. 3 (a) and (b)). Besides, the βNb5Si3 and Nb3Si phases are formed additionally in the nonequilibrium microstructure obtained by SLM. The HT performed at 1500 °C for 4 h is able to remove the nonequilibrium phases of βNb5Si3 and Nb3Si, and an

MA

equilibrium microstructure composed of Nbss and αNb5Si3 is achieved resultantly. The case that Nbss phases interconnect to form a continuous matrix is similar to the directionally solidified Nb-Si

D

based alloy. The Nb5Si3 phases tends to spheroidize during heat treatment, to decrease the interphase

(a)

TE

boundary area and the total interfacial energy [3]. A-αNb5Si3 C-Nbss B-βNb5Si3 D-Nb3Si

C

CE P

Intensity (a.u.)

A

B B

C

A AA

AA A

A

D

D

C

A

AC

C

A A

20

30

A

C

A AA

40

50

Arc-melted

C AA A

C

C

C

αNb5Si3

C

SLM

Nbss

C

SLM+HT C

A

60

(b)

70

80

50 μm

90

2θ (deg.)

(c)

(d) Nbss

RBZ HfO2

5 μm

αNb5Si3

5 μm

Fig. 3 XRD patterns (a) and microstructures of Nb-Si based alloy processed by arc-melting (b), SLM (c) and SLM+HT (d). Note that the images are magnified under different magnifications to better describe the microstructures. 7

ACCEPTED MANUSCRIPT More details with regard to the SLM microstructure are shown in TEM images (Fig. 4). The phases of αNb5Si3, βNb5Si3, Nbss and Nb3Si are further identified by selected area diffraction patterns (SADPs) as well as TEM-EDS (Table 1), respectively. The sphere shaped Nbss phases are

IP

T

observed, the maximum diameter of which is up to 306 nm. Between the neighbouring Nbss spheres distribute the strip shaped Nb5Si3 phases, which are up to 200 nm in length. Apart from the

SC R

difference on lattice structure, the βNb5Si3 phase has a higher Ti content (27.33 at.%) than αNb5Si3 (17.70 at.%). The Nb3Si phases are featured by a larger size (~1 μm). The titanium contained in Nbss

(a)

NU

is up to 27.33 at.%, suggesting that the bcc structured Nbss phase is stable at a high value of Ti/Nb. (b)

αNb5Si3

HfO2

D

Nbss

MA

βNb5Si3

200 nm

TE

200 nm (d)

AC

CE P

(c)

Nb3Si

(e)

(f)

Fig. 4 TEM image of the SLM-processed Nb-Si based alloy ((a), (b)) and the corresponding SADPs of αNb5Si3 (c), βNb5Si3 (d), Nbss (e) and Nb3Si(f)

8

ACCEPTED MANUSCRIPT Table 1 TEM-EDS data for the Nb-Si based alloy manufactured by SLM (at.%)

Nb

Ti

Si

Al

Cr

Nbss

63.20

27.33

1.80

2.56

5.11

α-Nb5Si3

49.40

12.42

35.64

1.35

1.19

β-Nb5Si3

34.59

25.66

35.42

2.67

1.66

Nb3Si

58.87

16.61

23.41

0.25

0.86

SC R

IP

T

Phase

When heated by the laser beam, the powders are melted and solidified rapidly. The SLM parts are therefore built through the overlap of multi-tracks and multi-layers. It is known that the mode of

NU

solidification and the fineness of microstructure is determined by the thermal gradient G and growth rate R [27]. In the SLM process, the thermal gradient G describes the difference in temperature over

MA

a certain distance and the growth rate R depends on the scanning speed and the angle between the direction of the moving laser source and growth direction. The G over R ratio determines the

D

solidification mode, by controlling the stability of the solidification front and resultant, while the

TE

multiplication of G and R determines the microstructure fineness with a positive correlation. A much higher cooling rate of Nb-Si melt is typically obtained, resulting from the high thermal gradient and

CE P

growth rates during the laser-based process [8, 28, 29]. The thermal gradient has been reported to be as high as in the order of 106 °C/m in the SLM processing of AlSi10Mg [29]. Consequently, according to the classical Johnson-Mehl-Avrami-Kolmogorov equation (JMAK) equation, both the

AC

nucleation rate and growth rate are remarkably enhanced. Furthermore, in the printing process of SLM, solidification the Nb-Si melt would begin with formation of solid nuclei on the surface of the previous layer with a much lower temperature. The reduction of free energy barrier for heterogeneous nucleation additionally favors the increase of the nucleation rate. Based on the above factors, a novel refined Nb-Si based alloy is obtained by SLM. Dependent on the movement of laser source, thermal gradient and growth rate change gradually over the melt pool. They are at their peak at the centerline of the melt pool, but lowest at RBZ, the area where two adjacent tracks are overlapped [27]. Accordingly, a decrease in cooling rate leads to a coarser microstructure at RBZ. 3.2 Microhardness In this work, Vickers hardness has been measured to compare the mechanical properties of the 9

ACCEPTED MANUSCRIPT SLM and SLM+HT Nb-Si alloy (Fig. 5). Vickers hardness on arc-melted Nb-Si alloy, with large silicides up to 50 μm, was not measured for comparison, since elastic modulus and hardness of the Nbss phase and silicides are intrinsically different [30]. Hardness of Nb-Si alloys processed by SLM

IP

T

is thus supposed to be a combination of Nbss and silicides, due to their small sizes. A homogeneous hardness distribution of the SLM and SLM+HT parts has been obtained, indicating that the volume

SC R

fraction of each phase is almost the same at different testing sites. The well-proportioned microstructure guarantees the homogeneity of mechanical properties of the SLM-processed Nb-Si

NU

alloy.

The average hardness of the SLM parts is 807HV0.1 (95% confidence interval 783-831 HV0.1),

MA

while that of the SLM+HT parts is 540HV0.1 (95% confidence interval 551-530 HV0.1). The reduction of hardness may be mainly caused by microstructure coarsening (Fig. 3 (c) and (d)). In situ reinforcement of the intermetallic phases provides Nb-Si alloy with considerable strengthening and

D

toughing effects, and small phase particles typically result in greater strengthening [31, 32]. Along

TE

with the reduction of the particle size, strengthening mechanisms such as Orowan strengthening, precipitate and dispersoid strengthening should be considered. Unfortunately, it is not immediately

CE P

clear when the transition takes place between such mechanisms with the results presented in this work. As described by the Hall-Petch equation, a fine-grained material is harder and stronger than one that is coarse grained, since the former has a greater total grain boundary area to impede

AC

dislocation motion [33]. Therefore, with less interface boundary and Nbss grain boundary area, the HT Nb-Si alloy tends to exhibit lower hardness values. The composition of the SLM Nbss phase is 63.20Nb-27.33Ti-1.80Si-5.11Cr-2.56Al (Table 1), while that of the SLM+HT Nbss phase is 62.90Nb-29.50Ti-0.81Si-4.06Cr-2.73Al. Consequently, the high hardness of the rapidly solidified state can also be attributed to the supersaturated solid solutions, mainly Si [34] and Cr [35], known as solid-solution strengthening. Furthermore, the elimination of residual stress by heat treatment may have additional effects.

10

ACCEPTED MANUSCRIPT

(a)

NU

SC R

IP

T

(b)

Fig. 5 Microhardness distributions of the SLM (a) and SLM+HT (b) Nb-Si based alloy on the polished

MA

horizontal section

3.3 Post-oxidation observations

D

As shown in Fig. 6, after oxidation at 1300 °C for 0.5 h, the oxides formed on arc-melted and

TE

SLM-processed Nb-Si alloys are mainly CrNbO4, TiO2, TiNb2O7 and Nb2O5. The detected oxides are in good agreement with those obtained by TEM and SADPs analysis on an arc-melted Nb-Si based

CE P

alloy oxidized at 1250 °C [36]. Actually, an intermediate oxide, namely Ti2Nb10O29, is formed firstly and reacts with TiO2 to produce TiNb2O7, as demonstrated in Eq. (1) [37, 38]. Eq. (1)

AC

3TiO2+Ti2Nb10O29=5TiNb2O7

Surface morphologies of the oxidized specimens are shown in Fig. 7 (a) and (b). Rod-like and platelet-like oxides are observed on the arc-melted and SLM-processed Nb-Si based alloys after oxidation at 1300 °C for 0.5 h. Some holes are found on arc-melted alloys, but not observed on SLM-processed alloys, suggesting that the oxide film grown on SLM-processed alloy is more compact and resultantly more protective. As shown in Fig. 8, a layered oxide scale is observed on either arc-melted or SLM-processed Nb-Si based alloy. Noteworthily, the SiO2 layers have been generated at the initial stage of oxidation, which are identified by WDS. The formation of SiO2 is supposed to result from the decomposition of Nb5Si3. Thus the SiO2 phases mainly concentrate around the Nb5Si3 particles. The thin SiO2 layers will be more clearly shown in Fig. 10 and Fig. 11, respectively. Fig. 9 (a) shows the concentration profile of oxygen along line A-B, which is drawn in Fig. 8 (b). Oxygen content in the outmost oxides 11

ACCEPTED MANUSCRIPT is not detected, since most of them are lost when preparing samples for EPMA (Fig. 8 (b)). Based on the oxygen content, the structure of the oxide scale grown on SLM-processed Nb-Si alloy at 1300 °C is divided into four parts: Layer I; Layer II; Layer III; Layer IV. It is reasonable to infer that the

IP

T

oxide scale formed on arc-melted alloy can be divided into four counterpart layers as well. Microstructures of the four layers on SLM-processed alloy are shown in Fig. 9 (b)-(d). Polyhedral

SC R

and rod-like oxides are observed in Layer I on arc-melted and SLM-processed Nb-Si based alloys. The Layer I on arc-melted alloy is less compact and exhibits a poor integrity, which is in accordance with the surface morphology shown in Fig. 7. Partially oxidized silicides are observed in Layer II on

NU

arc-melted alloy, while mixed oxides with no regular shapes are shown in Layer II on SLM-processed alloy. It is thus considered to be the transitional layer, where the Nb-Si alloy has not

MA

fully oxidized yet, and the oxides have not grown up to certain shapes. Besides, Layer II exhibits the highest oxygen content. Layers III and IV are the areas in which the Nb-Si alloy suffers from minor

TE

contrast), and HfO2 (bright contrast).

D

oxidation. The main phases are Nb5Si3 (dark grey contrast), Nbss (light grey contrast), TiO2 (black

It is worth noting that the microstructures shown in Layers III and IV are much different with

CE P

that of the as-fabricated SLM samples (Fig. 3 (c)), but similar to that of the HT samples (Fig. 3 (d)). It thus suggests that the Nb-Si alloy is heat treated during high temperature oxidation, and the microstructure grows larger resultantly by thermal activation. With a high affinity with oxygen,

AC

titanium is prone to be oxidized to form TiO2 oxides at a low oxygen partial pressure (Fig. 9 (c) and (d)), according to the Ellingham diagrams, i.e., plots of the standard free energy of formation ( G ) O

versus temperature for some typical oxides [39]. The presence of internal oxides suggests that the oxygen is sufficient for the chemical reaction taking place at the interface between oxide and alloy. The formation of more TiO2 particles and resultant a severer oxidation occurs in Layer III, with higher oxygen concentration.

12

ACCEPTED MANUSCRIPT ● CrNbO4 ■ TiO2 ◆Nb2O5 ▲ TiNb2O7

20

▲ ▲ ▲ ▲ ◆ ◆ ◆ ■ ● ■■ ◆ ▲▲ ● ●

30

40

■▲ ▲ ● ▲ ▲◆

50

◆

60

T

◆ Arc-melting

SC R

●

▲▲ ◆■ ▲ ◆ ◆ ■ ■ ■● ▲ ▲ ●▲ ▲◆ ◆ ◆ ▲▲● ● ▲ ◆

NU

Intensity (a.u.)

▲ ▲ ◆ ● ▲ ◆▲

IP

◆

◆■ ●

70

■ ◆◆ ● SLM

80

2θ (deg.)

MA

Fig. 6 XRD patterns the oxides formed on the arc-melted and SLM-processed Nb-Si based alloys after oxidation for 0.5 h at 1300°C.

(a)

AC

CE P

TE

D

(b)

Fig. 7 Surface morphologies of the arc-melted (a) and SLM-processed (b) Nb-Si based alloy after oxidation for 0.5 h at 1300°C

13

NU

SC R

IP

T

ACCEPTED MANUSCRIPT

TE

D

MA

Fig. 8 Cross-sectional microstructures of the arc-melted (a) and SLM-processed (b) Nb-Si based alloy after oxidation for 0.5 h at 1300 °C

I

AC

IV

II

CE P

III

TiO2

Nb5Si3 HfO2 Nbss Fig. 9 (a) Oxygen content distribution across the surface of SLM-processed Nb-Si based alloy after oxidation at 1300 °C for 0.5h, along the red line drawn on Fig. 8 (b); (b), (c) and (d) are the amplified images of area I, II, III and IV. 14

ACCEPTED MANUSCRIPT Fig. 10 and Fig. 11 show the elemental X-ray mapping of arc-melted and SLM-processed Nb-Si alloys oxidized at 1300 °C for 0.5 h, respectively. The distributions of oxygen, niobium, silicon, titanium, chromium, aluminum and hafnium across the oxide scale are presented. Evidently, the thin

IP

T

SiO2 layer, rich in Si and O, is observed at the interface of Lay I and Lay II on both arc-melted and SLM-processed Nb-Si alloys. Other SiO2 glass phases are dispersedly distributed in the oxide scale,

SC R

probably filled in the cracks and pores [4]. It is evident that a more continuous SiO2 layer is grown on SLM-processed alloy. Since the SiO2 phases are formed around silicides, the decrease in silicide

NU

size is supposed to be favorable to the formation of a more continuous SiO2 layer [19]. The outer part of the oxide scale is featured by the presence of oxides rich in the active elements

MA

of titanium and chromium [37, 39]. These oxides are supposed to be the mixture of CrNbO4, TiNb2O7 and TiO2, as revealed by XRD patterns (Fig. 6). According to the Ellingham diagrams, it can be obtained that Al2O3 is more thermodynamically stable than TiO2 at all temperatures. However,

D

the aluminum oxides are not detected, probably because of the small initial content in the Nb-Si alloy

TE

and thus small amount of aluminum oxides. In addition, the diffusion coefficient of Al3+ in niobium alloy is about two orders of magnitude lower than that of Ti4+ between 700 °C and 1200 °C, namely,

CE P

Ti4+ diffuses faster to the surface and react with O2- [40]. Consequently, the conclusion could be drawn that the formation of oxide at high temperature mainly depends on the thermodynamical and

AC

kinetic factors, as well as the initial concentration.

Fig. 10 Elemental X-ray mapping of arc-melted Nb-Si based alloy after oxidation at 1300 °C for 0.5 h using WDS attached with EPMA 15

NU

SC R

IP

T

ACCEPTED MANUSCRIPT

MA

Fig. 11 Elemental X-ray mapping of SLM-processed Nb-Si based alloy after oxidation at 1300 °C for 0.5 h using WDS attached with EPMA

D

3.4 Initial oxidation mechanism

TE

High temperature oxidation mechanism of the Nb-Si alloy manufactured by directional solidification and arc-melting has been discussed by other investigations [4, 35, 37]. The oxidation

CE P

behavior of Nb-Si alloy printed by SLM is to be elaborated as follows. Using the results obtained in this work, a schematic diagram describing the layer-structured oxide scale formed on Nb-Si alloy via SLM is illustrated in Fig. 12. At the initial stage of oxidation, oxygen molecules (O2) in air are

AC

converted into oxygen ions (O2-) on the alloy surface. Oxides start to nucleate and grow through chemical reactions between oxygen and alloying elements at the gas/alloy interface. The layered oxide scale is supposed to result from the different affinity for oxygen of the alloying elements in Nb-Si based alloy. As shown in Eq. (2), the elements with a higher affinity for oxygen than niobium, namely lower G value than NbO, are preferentially oxidized according to the Ellingham O

diagrams [37, 39]. O O O O O GNbO  GCrO O  GNbO  GSiO  GTiO  GAlO O  GHfO 2

2

3

2

2

2

3

2

Eq. (2)

Accordingly, HfO2 is most likely formed, as shown in Fig. 8 (a) and Fig. 9 (d), due to its lowest

G O value. Niobium oxides are easily formed on the surface, mainly due to the high niobium content (52 at.%) in Nb-Si alloy as well as its poor oxidation resistance [41]. Especially, the titanium 16

ACCEPTED MANUSCRIPT dissolved in both Nbss and silicides are prone to be oxidized at the Nbss/silicide interfaces, due to their high oxygen affinity and high contents in the alloy (24 at.%) [37, 38]. The simple oxides from all the alloying elements combine together to produce complex oxides via solid-state reaction, as

IP

T

typically shown in Eq. (1) [37, 38]. The formation of outer oxidation layer (OOL) then occurs, composed of TiO2, CrNbO4, TiNb2O7 and SiO2 phases (Fig. 12). Apart from Layers I and II in Fig. 9

SC R

(b), the outmost oxides and the SiO2 layer, at the interface of Lay I and Lay II, are concluded in OOL. The in-depth diffusion of oxygen is typically through Nbss phases, while silicides have excellent resistance to oxidation, maintaining a lower oxygen partial pressure at the metal-oxide interface. The

NU

penetrated oxygen continues to diffuse through Nbss, resulting in the formation of a thick oxidized zone, i.e., the internal oxidation layer (IOL) [4], as shown in Fig. 12. Layers III and IV, shown in Fig.

MA

9, are concluded in IOL. Partially-oxidized silicides are observed in IOL, as shown in Fig. 9 and

AC

CE P

TE

D

displayed in Fig. 12.

Fig. 12 A schematic diagram of the initial layered oxide scale formed on Nb-Si alloy manufactured by SLM after high temperature oxidation

A novel refined microstructure of Nb-Si based alloy is obtained by SLM. This work has testified that the high-temperature oxidation behavior is altered resultantly, through the short-time oxidation experiments (Fig. 10 and Fig. 11). One obvious difference is that compared with the arc-melted alloy, a more continuous SiO2 layer is formed in the oxide scale grown on Nb-Si based alloy via SLM. A well-developed SiO2 layer tends to suppress the inward transportation of oxygen and reduce the partial pressure of oxygen in the oxide scale [4]. Namely the SiO2 layer could protect the base alloy from further being oxidized. Therefore, long-time oxidation experiments will be conducted to 17

ACCEPTED MANUSCRIPT investigate the evolution of SiO2 layer and its efficiency on the oxidation resistance of SLM-processed Nb-Si alloy. Besides, the inward move of oxidation front, namely the metal-oxide interface, is largely restrained by silicides with superior oxidation resistance. The decrease in silicide

IP

T

size could affect the oxidation resistance of Nb-Si based alloy greatly, which is beyond the scope of

SC R

this work and will be testified in the long-time oxidation experiments. 4 Conclusions

With a nominal composition of Nb-18Si-24Ti-2Cr-2Al-2Hf (at.%), the Nb-Si based alloy has

NU

successfully yielded by SLM, using the pre-alloyed powders. The microstructure, hardness and initial oxidation behavior of the SLM-processed Nb-Si based alloy were presented as follows.

MA

(1) Nb-Si based alloy yielded by SLM were mainly composed of Nbss, αNb5Si3, βNb5Si3 and Nb3Si. Owing to the rapid solidification process of SLM, the sphere shaped Nbss phases with a

D

maximum diameter of 300 nm were obtained. Between the neighbouring Nbss spheres distributed the

TE

strip shaped α/β Nb5Si3 phases. After the heat treatment at 1500 °C for 4 h, the Nbss phases interconnected to form a continuous matrix with discontinuous αNb5Si3 phases.

CE P

(2) A homogeneous hardness distribution was found on the horizontal sections of either SLM or SLM+HT Nb-Si based alloy, indicating a well-proportioned microstructure obtained by SLM. The

treatment.

AC

average hardness of the SLM-processed parts was reduced to 542HV0.1 from 810HV0.1 after heat

(3) After oxidation at 1300 °C for 0.5 h, the oxides formed on arc-melted and SLM-processed Nb-Si alloys are mainly CrNbO4, TiO2, TiNb2O7 and Nb2O5. A layer-structured oxide scale model is presented to describe the high-temperature oxidation behavior of Nb-Si based alloy. (4) Compared with the arc-melted alloy, the oxide scale formed on SLM-processed is more compact with fewer holes. During the initial stage of oxidation, a more continuous SiO2 layer is generated in the oxide scale on Nb-Si based alloy via SLM. Acknowledgement The authors gratefully acknowledge the financial support by National Nature Science Foundation of P. R. China under the contract of 51471013 and 51571004. 18

ACCEPTED MANUSCRIPT References [1] J.H. Perepezko, The hotter the engine, the better, Science 326 (2009) 1068-1069. [2] B.P. Bewlay, M.R. Jackson, P.R. Subramanian, J. Zhao, A review of very-high-temperature Nb-silicide-based

T

composites, Metall. Mater. Trans. A 34 (2003) 2043-2052. [3] S. Yuan, L. Jia, L. Su, L. Ma, H. Zhang, The microstructure evolution of directionally solidified

IP

Nb-22Ti-14Si-4Cr-2Al-2Hf alloy during heat treatment, Intermetallics 38 (2013) 102-106. [4] L. Su, L. Jia, J. Weng, Z. Hong, C. Zhou, H. Zhang, Improvement in the oxidation resistance of Nb-Ti-Si-Cr-Al-Hf

SC R

alloys containing alloyed Ge and B, Corros. Sci. 88 (2014) 460-465.

[5] M.G. Mendiratta, D.M. Dimiduk, Strength and toughness of a Nb/Nb 5Si3 composite, Metall. Mater. Trans. A 24 (1993), 501-504.

[6] L. Ma, X. Tang, B. Wang, L. Jia, S. Yuan, H. Zhang, Purification in the interaction between yttria mould and

NU

Nb-silicide-based alloy during directional solidification: A novel effect of yttrium, Scripta Mater. 67 (2012), 233-236.

[7] N. Jones, Science in three dimensions: the print revolution, Nature 487 (2012) 22-23.

MA

[8] N. Kang, P. Coddet, C. Chen, Y. Wang, H. Liao, C. Coddet, Microstructure and wear behavior of in-situ hypereutectic Al-high Si alloys produced by selective laser melting, Mater. Des. 99 (2016) 120-126. [9] D. Wang, C. Song, Y. Yang, Y. Bai, Investigation of crystal growth mechanism during selective laser melting and mechanical property characterization of 316L stainless steel parts, Mater. Des. 100 (2016) 291-299.

D

[10] N. Read, W. Wang, K. Essa, M.M. Attallah, Selective laser melting of AlSi10Mg alloy: Process optimisation and

TE

mechanical properties development, Mater. Des. 65 (2015) 417-424. [11] K. Lu, The future of metals, Science 328 (2010) 319-320. [12] Y. Zhou, S.F. Wen, B. Song, X. Zhou, Q. Teng, Q.S. Wei, Y.S. Shi, A novel titanium alloy manufactured by

CE P

selective laser melting: Microstructure, high temperature oxidation resistance, Mater. Des. 89 (2016) 1199-1204. [13] S. Kumar, J.P. Kruth, Composites by rapid prototyping technology Mater. Des. 31 (2010) 850-856. [14] H. Attar, M. Bönisch, M. Calin, L. Zhang, S. Scudino, J. Eckert, Selective laser melting of in situ titanium-titanium boride composites: Processing, microstructure and mechanical properties, Acta Mater. 76 (2014) 13-22.

AC

[15] L. Thijs, M.L. Montero Sistiaga, R. Wauthle, Q. Xie, J. Kruth, J. Van Humbeeck, Strong morphological and crystallographic texture and resulting yield strength anisotropy in selective laser melted tantalum, Acta Mater. 61 (2013) 4657-4668.

[16] X. Zhou, X. Liu, D. Zhang, Z. Shen, W. Liu, Balling phenomena in selective laser melted tungsten, J. Mater. Sci. Technol. 222 (2015) 33-42. [17] L. Su, O. Lu-Steffes, H. Zhang, J.H. Perepezko, An ultra-high temperature Mo-Si-B based coating for oxidation protection of NbSS/Nb5Si3 composites, Appl. Surf. Sci. 337 (2015) 38-44. [18] X. Tian, X. Guo, Oxidation behavior of an Al-modified silicide coating on an Nb-silicide-based ultrahigh-temperature alloy, Corrosion 66 (2010) 25003. [19] F. Wang, The effect of nanocrystallization on the selective oxidation and adhesion of Al 2O3 scales, Oxid. Met. 48 (1997) 215-224. [20] B.V. Mahesh, R.S. Raman, Role of nanostructure in electrochemical corrosion and high temperature oxidation: a review, Metall. Mater. Trans. A 45 (2014) 5799-5822. [21] C. Qiu, N.J.E. Adkins, M.M. Attallah, Selective laser melting of Invar 36: Microstructure and properties, Acta Mater. 103 (2016) 382-395. [22] J. Sander, J. Hufenbach, L. Giebeler, H. Wendrock, U. Kühn, J. Eckert, Microstructure and properties of FeCrMoVC tool steel produced by selective laser melting, Mater. Des. 89 (2016) 335-341. 19

ACCEPTED MANUSCRIPT [23] M. Peter, P.K. Jean, Residual stresses in selective laser sintering and selective laser melting, Rapid Prototyping J. 12 (2006), 254-265. [24] B.P. Bewlay, M.R. Jackson, J. Zhao, P.R. Subramanian, M.G. Mendiratta, J.J. Lewandowski, Ultrahigh-temperature Nb-silicide-based composites, MRS Bull. 28 (2003) 646-653.

T

[25] Q. Liu, Y. Danlos, B. Song, B. Zhang, S. Yin, H. Liao, Effect of high-temperature preheating on the selective laser melting of yttria-stabilized zirconia ceramic, J. Mater. Process. Technol. 222 (2015) 61-74.

IP

[26] K.S. Chan, Alloying effects on fracture mechanisms in Nb-based intermetallic in-situ composites, Mater. Sci. Eng. A 329–331 (2002) 513-522.

SC R

[27] L. Wang, N. Wang, W.J. Yao, Y.P. Zheng, Effect of substrate orientation on the columnar-to-equiaxed transition in laser surface remelted single crystal superalloys, Acta Mater. 88 (2015) 283-292. [28] D.D. Gu, W. Meiners, K. Wissenbach, R. Poprawe, Laser additive manufacturing of metallic components: materials, processes and mechanisms, Int. Mater. Rev. 57 (2012) 133-164.

NU

[29] L. Thijs, K. Kempen, J. Kruth, J. Van Humbeeck, Fine-structured aluminium products with controllable texture by selective laser melting of pre-alloyed AlSi10Mg powder, Acta Mater. 61 (2013) 1809-1819. [30] J. Kim, T. Tabaru, H. Hirai, Evaluation technique of the hardness and elastic modulus of materials with fine

MA

microstructures, Mater. Trans. 44 (2003) 673-676.

[31] Y.W. Yan, L. Geng, A.B. Li, Experimental and numerical studies of the effect of particle size on the deformation behavior of the metal matrix composites, Mater. Sci. Eng. A 448 (2007) 315-325. [32] Y.S. Suh, S.P. Joshi, K.T. Ramesh, An enhanced continuum model for size-dependent strengthening and failure of

D

particle-reinforced composites, Acta Mater. 57 (2009) 5848-5861.

TE

[33] N. Hansen, Hall-Petch relation and boundary strengthening, Scripta Mater. 51 (2004) 801-806. [34] B. Kong, L. Jia, L. Su, K. Guan, J. Weng, H. Zhang, Effects of minor Si on microstructures and room temperature fracture toughness of niobium solid solution alloys, Mater. Sci. Eng. A 639 (2015) 114-121.

CE P

[35] L.N. Jia, J.F. Weng, Z. Li, Z. Hong, L.F. Su, H. Zhang, Room temperature mechanical properties and high temperature oxidation resistance of a high Cr containing Nb-Si based alloy, Mater. Sci. Eng. A 623 (2015), 32-37. [36] D. Yao, R. Cai, C. Zhou, J. Sha, H. Jiang, Experimental study and modeling of high temperature oxidation of Nb-base in situ composites, Corros. Sci. 51 (2009) 364-370.

AC

[37] S. Mathieu, S. Knittel, P. Berthod, S. Mathieu, M. Vilasi, On the oxidation mechanism of niobium-base in situ composites, Corros. Sci. 60 (2012) 181-192. [38] J. Zheng, X. Hou, X. Wang, Y. Meng, X. Zheng, L. Zheng, Isothermal oxidation mechanism of a newly developed Nb-Ti-V-Cr-Al-W-Mo-Hf alloy at 800-1200°C, Int. J. Refract. Met. Hard Mater. 54 (2016) 322-329. [39] N. Birks, G.H. Meier, F.S. Pettit, Introduction to the high temperature oxidation of metals, Cambridge University Press 2006 22-24. [40] H. Bakker, H.P. Bonzel, C.M. Bruff, M.A. Dayananda, W. Gust, J. Horvath, I. Kaur, G.V. Kidson, A.D. LeClaire, H. Mehrer, Diffusion in Solid Metals and Alloys, Transaction Technology Publications Ltd.1990. [41] J. Zheng, X. Hou, X. Wang, Y. Meng, X. Zheng, L. Zheng, Isothermal oxidation mechanism of Nb-Ti-V-Al-Zr alloy at 700-1200°C: Diffusion and interface reaction, Corros. Sci. 96 (2015) 186-195.

20

ACCEPTED MANUSCRIPT

AC

CE P

TE

D

MA

NU

SC R

IP

T

Graphical Abstract

21

ACCEPTED MANUSCRIPT Highlights 

The Nb-Si based ultrahigh temperature alloy with a novel refined microstructure was fabricated

The average hardness of the Nb-Si based alloy by SLM reduced to 542HV0.1 from 810HV0.1

IP



T

by selective laser melting (SLM).

A layered oxide scale was formed at the initial stage of high temperature oxidation, including a

CE P

TE

D

MA

NU

sound SiO2 layer.

AC



SC R

after heat treatment at 1500 °C for 4 h.

22