Effect of Fe additions on the microstructure and properties of Nb-Mo-Ti alloys

Journal Pre-proof Effect of Fe additions on the microstructure and properties of NbMo-Ti alloys

O.N. Senkov, S.I. Rao, T.M. Butler, T.I. Daboiku, K.J. Chaput PII:

S0263-4368(20)30097-4

DOI:

https://doi.org/10.1016/j.ijrmhm.2020.105221

Reference:

RMHM 105221

To appear in:

International Journal of Refractory Metals and Hard Materials

Received date:

13 December 2019

Revised date:

3 February 2020

Accepted date:

17 February 2020

Please cite this article as: O.N. Senkov, S.I. Rao, T.M. Butler, et al., Effect of Fe additions on the microstructure and properties of Nb-Mo-Ti alloys, International Journal of Refractory Metals and Hard Materials(2020), https://doi.org/10.1016/ j.ijrmhm.2020.105221

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© 2020 Published by Elsevier.

Journal Pre-proof Effect of Fe additions on the microstructure and properties of Nb-Mo-Ti alloys O.N. Senkov, S.I. Rao, T.M. Butler, T.I. Daboiku, K.J. Chaput Air Force Research Laboratory, Materials and Manufacturing Directorate, Wright-Patterson AFB, OH 45433, USA

ABSTRACT Two Nb-Mo-Ti alloys were modified by additions of 2 at.% Fe. These small additions of Fe did not form secondary phases but noticeably increased the strength of the Nb-Mo-Ti alloys in the

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temperature range of 25 to 1200°C due to solid solution strengthening effect. The oxidation

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fraction of retained, un-oxidized metal alloy.

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resistance at 1200°C in air was also improved considering both oxidation kinetics and the

Keywords

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Refractory alloys; Nb alloys; Microstructure; Compression properties; Oxidation.

1 INTRODUCTION

Nb-based refractory alloys currently used in some high-temperature applications, such as C-103

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(Nb-5.4Hf-2.0Ti-0.7Zr-0.3Ta-0.3W, by at.%), C-3009 (Nb-20Hf-6W) or FS-85 (Nb-18Ta-7W1.1Zr-0.1Hf), are hardened primarily by solid-solution strengthening. However, they may also

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have small amount of second-phase precipitates, generally associated with oxides, nitrides,

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carbides and/or borides, which have strong influence on the microstructure stability and mechanical properties at high temperatures [1-3]. For example, C-3009 normally contains ~0.72 at.% (0.1 wt.%) oxygen, which combines with hafnium to form stable HfO 2 precipitates, and FS85 contains ~0.95 at.% (0.1 wt.%) of carbon to form fine carbides. These precipitates increase creep and fatigue strength at temperatures above 1000°C. While refractory alloys based on Mo or W are brittle or have limited ductility, Nb alloys are ductile at room temperature. In particular, the excellent cold formability of C-103 enables complex structures, such as thrust cones and high-temperature valves to be easily made by closed die forging. Unfortunately, Nb alloys contain expensive elements (Hf, Ta and Zr) and still have rather high density ranging from 8.86 g/cm3 (C-103) to 10.6 g/cm3 (FS-85).

Journal Pre-proof Less expensive and lower density Nb alloys containing Mo, W and Ti have recently been reported [4]. In the temperature range of 20 to 1200°C, these alloys showed yield strengths superior to that of C-103; however, their strength values above 600°C were inferior to C-3009. In the present work, the mechanical properties of two Nb alloys, Nb-12Mo-18Ti and Nb-14Mo18Ti, have been noticeably improved by alloying with a small concentration (~2 at.%) of Fe. The microstructure, mechanical properties and oxidation behavior of the Fe-containing alloys are reported here.

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2 EXPERIMENTAL PROCEDURES

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Four alloys, M2 (Nb-12Mo-18Ti), M3 (Nb-14Mo-18Ti), M2Fe (Nb-12Mo-18Ti-2Fe) and M3Fe (Nb-14Mo-18Ti-2Fe) were prepared from high purity (99.9% or higher) elements by vacuum arc

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melting. Alloy C-3009 in an as-cast condition was provided by ATI Specialty Materials, Inc. The actual alloy compositions in atomic percent are shown in Table 1. The concentrations of the

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alloying elements were determined using X-ray energy dispersive spectroscopy and direct

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current plasma atomic emission spectrometry. The concentrations of oxygen and nitrogen were determined by inert gas fusion and the concentration of carbon was determined using combustion infrared absorption methods. The alloy density was measured using a helium pycnometer, with

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the accuracy of ± 0.005 g/cm3 . Vickers microhardness was determined in accord to the ASTM E384-17 standard [5] as an average value of 10 measurements using a Buehler1600 series

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microhardness tester with a diamond pyramid applied to a polished surface of the alloy sample at

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1000 g load for 10 seconds. In order to be sure that adjacent tests do not interfere with each other, the distance between the indentations was larger than four Vickers indentation diagonals. X-ray diffraction was conducted on a Bruker D2 PHASER XRD using Cu K α1 radiation in the 2θ range of 10 to 120°. Microstructure and local chemical composition were studied using a field emission scanning electron microscope Quanta 600 equipped with backscatter electron (BSE) and energy dispersive spectroscopy (EDS) detectors. The average grain size was measured using the Jeffries planimetric method, which is described in detail in the ASTM standard E112-13 [6]. Specimens for scanning transmission electron microscopy (STEM) were prepared using a slight variation of the situ lift-out technique in an FEI Nova 600 Dual Beam FIB-SEM. The technique is described in details elsewhere [7]. Initial trenching (normal to the sample surface) was conducted at 30kV and 3nA to produce a lamella ~ 20 μm long, 5 μm tall, and 1.5 μm thick. The

Journal Pre-proof lamella was manipulated using an Omniprobe 200 micromanipulator and subsequently Pt welded onto a copper TEM sample half grid. Specimen thinning was conducted normal to the top surface starting with 30 kV and 0.3 nA beam conditions and the power was systematically lowered throughout the thinning process down to 30 kV and 0.1 nA. Final thinning was conducted at 5 degrees from the sample surface normal at 5 kV and 0.23 nA. This process was repeated for both sides of all TEM foils. High angle annular dark field (STEM-HAADF) images and selected area diffraction patterns (SADPs) were taken using a 200 kV Talos F200X (S)TEM equipped with Super-X EDS.

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Specimens for oxidation testing had rectangular shapes measuring ~4.4 mm × 4.4 mm × 7.2

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mm. The faces of all the specimens were ground to a 600 grit finish and ultrasonically cleaned in acetone and isopropanol. Oxidation exposures were conducted in an open air box furnace in

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3YSZ crucibles at 1200°C for 1, 8 and 24 hours. Mass measurements were captured using a high sensitivity balance (10-6 g) prior to and post oxidation treatment. Dimensions and volume

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fractions of un-oxidized material from each oxidized alloy sample were measured using

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respective backscatter electron cross section images of the entire sample cross-section. The unoxidized “metal cores” were assumed to be rectangular geometries in the calculations and all

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percent values of remaining material were considered with respect to the two dimensional area.

Table 1. Density (ρ), Vickers microhardness (Hv) and chemical composition (in at.%), of the

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studied alloys. Compositions of commercial alloys C-103 [8] and C-3009 are also shown for

Alloy

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comparison. ρ

Hv

Nb

Mo

Ti

Fe

W

Hf

O

N

C

(g/cm3 ) M2

7.99

282±6 70.0

11.9

18.1

-

-

-

0.084

0.008

0.039

M3

8.03

287±6 67.8

14.1

18.1

-

-

-

0.090

0.006

0.042

M2Fe

7.99

414±8 68.7

11.5

17.7 2.1

-

-

0.111

0.018

0.070

M3Fe

8.03

329±8 66.0

13.4

18.3 2.3

-

-

0.074

0.006

0.028

C-103*

8.86

230±5 90.9

-

2.0

-

0.3

5.4

C-3009**

10.3

274±8 71.7

-

-

-

5.9

22.4

* C-103 also contains 0.3 at%Ta, 0.3 at%W and 0.7 at.%Zr. ** Composition of O, N and C in C-3009 was not disclosed.

<0.136 <0.104 <0.121

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Compression mechanical tests were conducted at 25°C, 600°C, 800°C, 1000°C and 1200°C at a constant ram speed of 0.0074 mm/s that corresponded to the initial strain rate of 0.001 s -1 . Compression specimens had a squared cross-section of 4.6 mm × 4.6 mm and the height of 7.4 mm. Room temperature tests were conducted in air and high-temperature tests were conducted in vacuum of 2×10-3 Pa or better. A more detailed description of the compression tests can be found elsewhere [4].

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3 RESULTS AND DISCUSSION

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3.1 Microstructure 3.1.1 As-cast condition

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According to X-ray diffraction, the studied alloys have single-phase BCC structures (Figure 1). The lattice parameter is a = 328.3 pm for the M2 alloy (Nb-12Mo-18Ti) and it slightly decreases

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with increasing the amount of Mo and with adding Fe. SEM analysis confirms single-phase,

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equiaxed grain structures for the alloys (Figure 2). The average grain size in M2 and M3 alloys is 0.90 mm and 0.60 mm, respectively (Figure 2 a, b). The average grain size in the Fe-containing

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alloys M2Fe and M3Fe is 0.44 mm and 0.39 mm, respectively (Figure 2 c, d). Small dark spots inside grains are pores and/or nitride particles rich in Ti. Figure 3 shows a BSE image of a representative dark spot in the M2Fe alloy. It consists of an irregular-shaped particle surrounded

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by dark needles. The chemical compositions of the particle, needles and the matrix phase near

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such clusters are given in Table 2. The dark particle is rich in Ti and N and depleted in other elements. According to the chemical composition, it is possible that this nitride is the metastable Ti3 N 2 phase [9]. The needles are also rich in Ti and N and depleted in other elements, but due to their small thicknesses, the composition shown in Table 2 is likely a combination of the needle and the surrounding matrix. The composition of the matrix phase is similar to the composition of the alloy.

Figure 1. X-ray diffraction patterns of (a) M2, (b) M2Fe, (c) M3 and (d) M3Fe alloys. Figure 2. SEM backscatter electron (BSE) images of the microstructure of (a) M2, (b) M3, (c) M2Fe and (d) M3Fe alloys.

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Figure 3. A cluster of nitride particles in the M2Fe alloy in the cast condition.

Table 2. Chemical composition (in at.%, the average of 5 measurements at magnification of 5000x) of dark particles, dark needles and the matrix phase near the clusters of the dark particle

Nb

Mo

Ti

Fe

N

Particle

4.2

0.2

59.4

0.3

35.9

Needles

37.0

4.7

36.5

2.3

19.5

Matrix

68.0

9.6

2.4

0

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Element

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and needles (similar to shown in Figure 3).

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20.0

3.1.2 Annealed condition

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M3 and M3Fe alloys were also studied after vacuum annealing at 1200°C for 12 hours. The

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heating rate was 50 °C/min and the cooling rate after annealing to room temperature was 20°C/min. The vacuum pressure was kept below 1×10-3 Pa. After annealing, the alloys retain an

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essentially single-phase BCC equiaxed grain structure, with slightly increased average grain size relative to the as-cast condition (Figure 4a,b). In particular, the average grain size of M3

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increases to 0.73 mm and that of M3Fe to 0.44 mm. In addition, precipitation of Ti rich (50-68 at.%Ti) particles near and at grain boundaries (GBs) occurs in the annealed samples (Figures

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4c,d). The GB particles have irregular shapes, while the precipitates near GBs have needle-like appearance and are arranged in ~10-15 μm thickness/diameter clusters. In the M3 alloy these clusters are almost always separated from grain boundaries by a thin (~2-3 μm) precipitation- free zone (Figure 4c). Thinner (up to 1 μm) precipitation-free zones are also observed near some GBs in the annealed M3Fe alloy; however the GB-adjacent stripes of the Ti-rich precipitates, similar to those shown in Figure 4d, with no evident precipitation-free zones, are more typical. No precipitates are present inside grains, except the above-mentioned regions The total volume fraction of the precipitates does not exceed 0.5-1% in both alloys. The fact that during annealing precipitation only occurs near or at grain boundaries and does not occur inside grains indicates that the alloys experienced micro-segregation during casting and the grain boundary regions were supersaturated with Ti. There are two reasons for the formation of

Journal Pre-proof precipitate free zones near GBs during annealing [10, 11]. The first, most common reason is that precipitates nucleate heterogeneously on vacancies. If the regions adjacent to GBs are depleted with vacancies due to their intake by GBs, these regions are unable to nucleate the precipitates even the matrix is supersaturated with solute. The second reason is that the second-phase particles nucleate first at GBs, thereby removing sufficient solute from the adjacent matrix. It is likely that both mechanisms are responsible for the formation of precipitation- free zones in the M3 alloy. On the other hand, a considerably smaller thickness or even absence of the precipitation- free zones in the Fe-containing M3Fe alloy may be an indication of reduced

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vacancy mobility in this alloy. Additional study is required to explain the different near-GB

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precipitation behavior in these two alloys.

To understand the role of Fe additions, STEM analysis was conducted on specimens lifted out of

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the central region of grains from the annealed M2Fe and M3Fe alloys. STEM/EDS analysis revealed a homogeneous distribution of elements and no distinct evidence of secondary

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precipitation inside the grains. For example, Figure 5 shows elemental EDS maps for elements

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inside a grain of the annealed M3Fe alloy. No compositional fluctuations related to precipitation are observed at this fine scale. In addition, Figure 6 shows a STEM-HAADF image of the annealed M3Fe alloy and respective selective area diffraction patterns (SADPs) from three zone

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axes of the BCC matrix phase, [001], [11̅3] and [102]. In correlation with chemical findings, no additional diffraction spots belonging to additional phases are present. It can be concluded from

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phase.

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this analysis that the added amount Fe is present in the form of a solid solution in the matrix

Figure 4. SEM backscatter electron (BSE) images of the microstructure of (a,c) M3 and (b,d) M3Fe alloys after annealing at 1200°C for 12 hours. (a,b) Grain structure at low magnification; (c,d) Ti-rich particles precipitated at or near grain boundaries.

3.2 Alloy density The densities of M2 and M3 alloys were measured to be 7.99 and 8.03 g/cm3 , respectively (Table 1). The Fe-containing M2Fe and M3Fe alloys have the densities similar to M2 and M3 alloys, respectively. These values are considerably smaller than the densities of the commercial alloys C-103 (8.86 g/cm3 ) or C-3009 (10.3 g/cm3 ).

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Figure 5. TEM-EDS elemental maps showing homogeneous distributions of (a) Nb, (b) Mo, (c) Ti and (d) Fe inside an annealed M3Fe alloy grain.

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Figure 6. (a) STEM-HAADF image of a grain interior region of the M3Fe alloy and (b-d) the

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respective selective area diffraction patterns from the BCC matrix: (b) [001], (c) [11̅3] and (d)

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[102] zone axes.

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3.2 Mechanical Properties 3.2.1 Cast condition

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Room temperature Vickers microhardness (Hv) values are given in Table 1. Microhardness of the Nb-Mo-Ti alloys, M2 and M3, is almost the same at 282±6 Hv and 287±6 Hv, respectively.

Hv for M3Fe.

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Additions of 2 at.% Fe increase microhardness considerably, to 414±8 Hv for M2Fe and 329±5

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Typical engineering stress vs. engineering strain curves for the cast alloys are shown in Figure 7 (a-d) and the yield strength (σ 0.2) values are given in Table 3. All of the alloys but M2Fe are

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ductile at 25°C – 1200°C and do not show any evidence of macroscopic fracture after 50% compression deformation. At the same time, the M2Fe alloy shows limited compression ductility at 25°C and starts fracturing after achieving ~30% compression strain. Figure 7e shows optical images of the different alloy samples deformed at 25°C to 50% compression strain. Cracks and fracture are clearly seen in the M2Fe alloy, while the other alloys keep their dimensional integrity with no evidence of cracking. In the temperature range of 25°C to 1200°C, the engineering stress of the M2, M3 and M3Fe alloys increases continuously with increasing engineering strain resulting in the maximum stress values at the end of deformation. The deformation behavior of M2Fe is similar to that of three other alloys only at temperatures of 600°C, 1000°C and 1200°C. At 25°C, the M2Fe alloy shows strain hardening during

Journal Pre-proof deformation up to ~30% compression strain, after which the stress decreases gradually with increasing strain, likely due to developing cracks and fracturing. At 800°C, the M2Fe alloy shows a noticeable drop in strength (from 1060 MPa to 790 MPa), when engineering strain increases from 12% to 23%, which is then followed by continuous hardening until deformation stopped after achieving 50% strain (Figure 7c). It should also be noted that the yield strength of M2Fe is considerably higher than the yield strength of M3Fe at 25°C, 600°C and 800°C in spite the fact that M2Fe contains a smaller amount of Mo, which is a strong solution strengthening element [4], than M3Fe (Table 1). At higher temperatures, the yield strength of M2Fe decreases

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more rapidly with increasing temperature and at 1200°C M3Fe becomes stronger than M2Fe.

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elements in M2Fe, as discussed in the next section.

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Such different behaviors of M3Fe and M2Fe are likely related to a higher amount of interstitial

Figure 7. (a-d) Engineering stress vs. engineering strain compression curves of (a) M2, (b) M3,

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(c) M2Fe and (d) M3Fe alloy samples deformed at 25°C, 600°C, 800°C, 1000°C and 1200°C. (e)

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Images of the samples after 50% compression deformation at 25°C.

Table 3. Yield strength (σ 0.2, in MPa) / specific yield strength (σ 0.2/ρ, in MPa cm3 /g) values of the

T = 25°C

T = 600°C

T = 800°C

T = 1000°C

T = 1200°C

719 / 90.0

350 / 43.8

311 / 38.9

250 / 31.3

161 / 20.2

M3

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Alloy

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studied as-cast alloys at different temperatures.

780 / 97.1

390 / 48.6

329 / 41.0

270 / 33.6

183 / 22.8

M2Fe

1090 / 136.4

706 /88.4

758 / 94.9

456 / 57.1

188 / 23.5

M3Fe

974 / 121.3

589 / 73.3

522 / 65.0

445 / 55.4

213 / 26.5

C-103 [8]

296 / 33.4

193 / 21.8

169 / 19.1

145 / 16.4

115 / 13.0

C-3009 [4]

663 / 64.4

437 /42.4

424 /41.2

397 / 38.5

388 / 37.7

M2

3.2.1.1 Effect of the alloying elements on the deformation behavior The Fe-containing alloys, M2Fe and M3Fe are considerably stronger than the M2 and M3 alloys at all studied temperatures (Table 3). The strengthening effect from Fe is especially pronounced

Journal Pre-proof at temperatures from 25°C to 1000°C. When the strengths of M3 and M3Fe alloys are compared (both alloys contain almost the same amount of interstitial elements), it can be found that the yield strength of M3Fe is about 175-193 MPa higher than the yield strength of M3 in this temperature range; although the difference decreases to 30 MPa at 1200°C (Table 3). Thus the strengthening effect at T ≤ 1000°C from the addition of 1 at.% Fe is estimated to be ~ 76-84 MPa. For comparison, the strengthening effect from 1 at.% Mo is estimated (from the comparison of the behavior of M2 and M3 alloys, which have the same concentrations of other elements) to be ~30 MPa at 25°C and only 9 MPa at 1000°C and 1200°C. At 1200°C, the

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strengthening effect from Fe (~13 MPa per 1 at.%) is still higher than that from Mo. The above

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estimations are made under the assumption of a linear dependence of the yield strength increase on the solute concentration and independent effects from different solute elements. TEM studies

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show no Fe-rich precipitates in the Fe-containing alloys and thus the increased strength of these alloys is likely due to solution strengthening. The mechanism of the Fe-induced strengthening

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needs to be understood. It is probably related to the abnormally high diffusivity of Fe in Nb [12,

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13], Ti [14, 15] and Nb-Fe [16] alloys, which is comparable to the diffusivity of interstitial elements. It has been suggested [13, 16] that Fe embedded in Nb or Ti forms vacancy-interstitial pairs and behaves similarly to the interstitial atoms. As such, the strengthening mechanism from

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Fe can also be similar to that from interstitials discussed in recent publications [17-19]. Here, we present a possible explanation of the strengthening effect from alloying with Fe. In a

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previous publication [4], the temperature dependence of the yield strength of M2 and M3 alloys

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was modeled using an adapted version of the Suzuki model. There was fairly good agreement between the calculated and experimental data [4]. Here we use similar approach to the Fecontaining alloys. Using atomistic simulations with the Johnson-Zhou potentials [20, 21], the interaction energy between a Fe solute and a/2[111] screw dislocation in the M2Fe or M3Fe BCC matrix is calculated to be ~ 0.084 – 0.1 eV. Inputting this value into the Rao-Suzuki model [22-24], the solid solution strengthening effect from 2.3 at% Fe in the M3Fe alloy is calculated to be 73-94 MPa at 25°C and 24-31 MPa at 1000°C. The atomistic simulations also show that Fe has strong repulsive interactions with Fe and Mo located in its first and second nearest neighbor shells. This is likely a result of relatively small radii of Fe and Mo atoms with respect to the parent lattice. The repulsive interactions may

Journal Pre-proof produce short-range order in the M2Fe and M3Fe solid solutions. The contribution of the shortrange order to strengthening in BCC structures, σsro , can be calculated as [25]: σsro = (8.976/a0 3 )∑𝑖 (∑𝑗 𝑐𝑖 𝑐𝑗 𝛼𝑖𝑗 𝑈𝑖𝑗 )

(1)

where ci is the atom fraction of an alloying element i, αij is the short range order Warren-Cowley coefficient between an element i and its neighboring element j located in the first and second coordination shells, U ij is the pair interaction energy between elements i and j at first and second neighbors, and a0 is the lattice parameter of the body-centered cubic unit cell. A negative value of αij indicates ordering, a zero value - random solid solution and a positive value - segregation.

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αij can vary from 1 to large negative values, and it is negative when the respective Uij is negative

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and positive when Uij is positive. The Warren-Cowley coefficients also satisfy the conditions: Σjcjαij = 0; αij = αji

(2)

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The Taylor factor M of 2.75 is used in equation (1). The pair interaction energies between solutes Fe-Fe, Fe-Mo, Fe-Nb, Mo-Mo and Mo-Nb are calculated to be 0.243, 0.166, -0.023, 0.103 and -

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0.017 eV using atomistic potentials. Fe-Ti and Mo-Ti pairs have negative interaction energies

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similar to that of Fe-Nb and Mo-Nb pairs. Assuming no Mo or Fe atoms occupy first and second neighbor sites around Fe and Mo, and ordering only between Fe-Nb and Mo-Nb pairs, the

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Warren-Cowley coefficients for the Fe-Fe, Fe-Mo, Fe-Nb, Mo-Mo and Mo-Nb pairs are estimated to be 1.0, 1.0, -0.242, 1.0 and -0.242, respectively. We have used Equation (2) to

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derive the Warren-Cowley coefficients for Fe-Nb and Mo-Nb pairs. Then, the short-range order contribution to strengthening by Fe is calculated to be ~ 162 MPa. Due to Equation (2) and the

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fact that Fe-Ti, Mo-Ti interaction energies are fairly close to that of the interaction energies of Fe-Nb and Mo-Nb pairs, neglecting Fe-Ti, Mo-Ti pairs to this analysis will not significantly alter the results. Such a high Warren-Cowley coefficient of 1.0 for Fe-Fe, Fe-Mo pairs is plausible since the atom fraction of Fe is relatively low ~ 0.023. We also assume that the short-range order parameters, αij, are constant between room temperature and 1000C. Assuming a superposition between the classical thermally activated strengthening (modelled here using the Suzuki model) and short range order strengthening, the expected strengthening due to Fe in M3Fe alloy is calculated to be 235-256 MPa at room temperature and 186-193 MPa at 1000°C. These results are somewhat larger than experimental data which gives strengthening of around 175–193 MPa at room temperature and 1000C. However, we have assumed extreme segregation in the analysis, while the Warren-Cowley coefficients may not reach a value of 1.0 for Fe-Fe, Fe-Mo and Mo-

Journal Pre-proof Mo pairs under the experimental conditions. Also, in actuality, there is an interplay between classical thermally activated strengthening and short range order strengthening [26], which we have neglected in the superposition assumption. In addition, a more accurate determination of Fe solute – M3Fe matrix ½<111> screw dislocation interaction energy and solute pair interaction energies, as well as Warren-Cowley parameters, using first principles calculations [27] or experiments are required for a more quantitative comparison of the calculated strengthening due to Fe solutes with experimental data. Finally, the observation that the contribution of Fe to strengthening in M3 alloy is larger in the annealed condition than in the as-cast condition is also

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in agreement with the short-range order strengthening model presented here. It is also worth to

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mention that, while in FCC crystals short range order strengthening has been suggested to contribute only in the yield drop phenomenon [25], in BCC crystals, where screw dislocations

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frequently cross-slip, short range order strengthening is expected to result in true strengthening effect.

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The “anomalous” deformation behavior of M2Fe, such as limited ductility at 25°C and higher

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strength than that of M3Fe at T ≤ 1000°C, can be explained by the presence of interstitial elements in considerably larger amounts than in M3Fe and two other alloys (Table 1). Indeed, strengthening of the studied alloys is mainly controlled by solid solution strengthening which, at

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the first approximation, is the sum of the interstitial solid solution strengthening (from O, N and C) and substitutional solid solution strengthening (from Mo and Ti). The interstitial elements

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tend to considerably increase the strength and reduce ductility of refractory alloys at low and

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intermediate temperatures by pinning dislocations and effectively reducing dislocation mobility [17, 18, 28, 29]. However, the strengthening effect from interstitials reduces considerably at high temperatures, likely due to a significant increase in diffusivity [17]. As a result, substitutional solid solution strengthening becomes the main strengthening mechanism in the studied alloys at T > 1000°C. This explains a considerable drop in the yield strength of M2Fe between 1000°C and 1200°C, as compared to M3Fe alloy (Table 3). Using the above estimations for the solution strengthening effects from Mo and Fe, comparing the yield stress values of M2Fe and M3Fe and assuming that the yield stress difference in these alloys is due to different amounts of Mo, Fe and interstitials, the strengthening effect from the interstitial elements can be estimated to be ~220 MPa at 25°C, 48 MPa at 1000°C and none at 1200°C per 0.1 at.% of the interstitials. (The same strengthening effect from O, N and C is

Journal Pre-proof assumed here.) It can be seen that the strengthening effect from interstitials decreases rapidly with increasing temperature, similar to previous observations [17], but it is considerably stronger (per 1 at.%) than from Mo or Fe at temperatures ≤ 1000°C. For example, the strengthening effect at 1000°C from the addition of 0.1 at.% O is equivalent to the addition of 0.5 at.% Fe or 5 at.% Mo. Unfortunately, interstitials may also cause brittleness and therefore, their amount in the alloys should be controlled and generally should not exceed 1 at.%.

3.2.2 Annealed condition

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Mechanical properties of alloys M3 and M3Fe have also been studied in the annealed condition.

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The engineering stress vs. engineering strain curves of the annealed alloy specimens are shown in Figure 8. In general, the behavior of the annealed specimens is similar to that of the as-cast

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specimens (compare Figure 8 and Figure 7). However, the yield strength, as well as the flow strengths measured at the same engineering strain values, of the annealed specimens (Table 4)

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are slightly smaller than the respective values of the as-cast specimens (Table 3). The decrease in

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the yield strength of M3 and M3Fe alloys after annealing is ~11-16% and 5-13%, respectively, relative to the as-cast condition, with the largest decrease at 800°C for M3 and 1200°C for M3Fe. This slight decrease in the strength after annealing can be related to precipitation of Ti-

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rich particles causing a slight change in the respective composition of the matrix phase. A slight increase in grain size can also contribute to the softening. Additiona l work is required to study

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alloys.

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phase composition, kinetics of phase transformations and strengthening mechanisms in these

Figure 8. Engineering stress vs. engineering strain compression curves of (a) M3 and (b) M3Fe alloy samples preliminary vacuum annealed at 1200°C for 12 hours and then deformed at 25°C, 600°C, 800°C, 1000°C and 1200°C. Table 4. Yield strength (σ 0.2, in MPa) / specific yield strength (σ 0.2/ρ, in MPa cm3 /g) values at different temperatures of the cast M3 and M3Fe alloys after vacuum annealing at 1200°C for 12 hours. Alloy

T = 25°C

T = 600°C

T = 800°C

T = 1000°C

T = 1200°C

M3

669 / 83.3

340 / 42.3

277 / 34.5

242 / 30.1

160 / 19.9

Journal Pre-proof M3Fe

891 / 111

565 / 70.4

505 / 62.9

411 / 51.2

186 / 23.2

3.2.3 Comparison with commercial alloys The developmental alloys, in both as-cast and annealed conditions, are considerably stronger than the commercial C-103 alloy in the temperature range of 25°C - 1200°C (Table 3, Table 4). The Nb-Mo-Ti alloys, M2 and M3, are also stronger than the C-3009 alloy at room temperature. However, at 600°C – 1200°C, C-3009 is stronger than M2 and M3. The M2Fe and M3Fe alloys containing a small amount of Fe are noticeably stronger than C-3009 at all studied temperatures

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except 1200°C. The yield strength of C-3009 is superior to all the studied alloys at 1200°C.

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Considering specific strength based on densities, (Table 1), at 1000°C, the yield strength and specific yield strength of M3Fe are higher than the respective quantities of C-3009 by about 12%

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and 44%. On the other hand, at 1200°C, the yield strength of M3Fe is smaller than that of C3009 by 45% but the specific yield strength is smaller by 30%. The results clearly indicate that

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the developmental alloys are ductile and their strength and specific strength are superior to those

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of C-103 in the temperature range of 25°C-1200°C and to C-3009 at T < 1200°C.

3.3. Microstructure of the Cast Alloys after Compression Deformation

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X-ray diffraction (XRD) analysis showed that all the deformed alloy specimens retain their single-phase BCC structures. For example, Figure 9 shows the XRD patterns from (a) M2Fe and

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(b) M3Fe alloy specimens after 50% compression deformation at 600°C, 800°C and 1000°C.

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Only diffraction peaks from a BCC phase are present. The lattice parameter of the phase is not significantly affected by deformation and it is a = 327.8 pm in M2Fe and 327.4 pm in M3Fe, similar to the non-deformed conditions. Figure 9. X-ray diffraction patterns of (a) M2Fe and (b) M3Fe alloy samples after 50% compression deformation at 600°C, 800°C and 1000°C.

Figure 10, Figure 11, and Figure 12 illustrate the microstructure of the M3, M2Fe and M3Fe alloys, respectively, after 50% compression deformation at 1000°C. Extensive deformation occurs inside grains, which is recognized by such features as (a) grain elongations in the directions of the material flow, (b) partial dynamic recrystallization and formation of fine grains at grain boundary regions, (c) dynamic recovery and formation of a fine subgrain structure with

Journal Pre-proof low-angle boundaries inside grains, and (d) noticeable residual stresses at grain boundaries and triple grain boundary junctions. Additionally, grain boundary cracks, like that shown in Figure 11c, are present in the M2Fe alloy, which contains increased amount of interstitial elements.

Figure 10. Microstructure of M3 after 50% compression deformation at 1000°C.

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Figure 11. Microstructure of M2Fe after 50% compression deformation at 1000°C.

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Figure 12. Microstructure of M3Fe after 50% compression deformation at 1000°C.

3.4 Oxidation Behavior

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Figure 13 shows specific mass gain of the C-3009, M3, and M3Fe alloys during oxidation at 1200°C in air. In all cases, the overall magnitude of the oxidation- induced mass gain is

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moderately high (>90 mg/cm2 ) out to 24 hours of oxidation. The substitution of W and Hf in C3009 with Mo and Ti in the M3 alloy resulted in a slightly lower mass gain kinetics of M3

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compared to C-3009. The small additions of Fe in the M3Fe alloy resulted in slightly lower

law [30]: Δm/A = kt n ,

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kinetics compared to the M3 alloy. The experimental points were fit to a conventional oxidation

(3)

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where Δm is the mass gain per the sample surface area A within the oxidation time period t, k is the oxide growth rate constant and n is the rate exponent. The corresponding k and n values for each alloy are shown in Table 5. All three alloys exhibited the same rate exponent (n=0.69), with the rate constants of 17.3, 14.3, and 11.0 mg cm-2 h-0.69 for C-3009, M3 and M3Fe alloys, respectively. The coefficients of determination R2 of the fits shown in Figure 13 were 0.996 or higher in all three cases. The rate exponent of 0.69 infers mixed-mode oxidation behavior since it falls between the diffusion controlled parabolic behavior (n=0.5) and interface controlled linear behavior (n=1) [30]. The presence of this mixed-mode behavior is likely due to the formation of an oxide layer with sufficient thickness to facilitate parabolic oxidation kinetics, coupled with simultaneous oxide scale cracking and potential spallation in various locations. This secondary

Journal Pre-proof influence exposes new free surfaces to readily react and creates an interface driven oxidation condition in particular regions, thus increasing the overall growth rate constant. The mixed-mode behavior is active during the entire duration of oxidation testing.

Figure 13. Temporal dependence of the specific mass gain of C-3009, M3, and M3Fe alloy

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samples during discontinuous oxidation at 1200°C in air.

Table 5. Oxide growth rate constant k and growth rate exponent n, Equation (3), for the C-3009,

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M3, and M3Fe alloys subjected to oxidation testing at 1200°C in air. The measured oxygen interaction depths (X,Y) and the fractional cross-section areas of remaining, non-oxidizing metal

k

n

Area %Non

Depths (X,Y) in mm

Oxidized Metal

0.69

(0.95, 0.86)

37.6

0.69

(0.67, 0.69)

49.1

(0.64, 0.68)

48.7

17.3

M3

14.3

M3Fe

11.0

0.69

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C-3009

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(mg cm-2 h-0.69 )

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Oxygen Interaction

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Alloy

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(assuming simple rectangular geometries) after 24 hours of oxidation are also shown.

Transverse cross-sections of the C-3009, M3, and M3Fe alloys after 24-hours of oxidation at

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1200°C in air are shown in Figure 14, (a)-(b), (c)-(d), and (e)-(f), respectively. Since the images of the cross-sections in Figure 14 (b), (d), and (f) were taken at the same magnification, they are directly comparable in size. Each alloy sample contains an outer oxide scale, an internal reaction zone, and an un-oxidized metal core. The respective compositions obtained via EDS for the outer oxide scales, internal reacted zones, and un-oxidized metal cores are shown in Table 6. The EDS spectra from non-oxidized alloy samples were used as the reference spectra to establish the oxygen content in different regions of the oxidized specimens. It was found that the amount of oxygen in the center of the un-oxidized metal core was the same as in the respective nonoxidized alloy sample. In the reacted zone, the amount of oxygen was ~49% in C-3009 and ~39% in M3 and M3Fe. The oxide scale of all three alloys contained almost the same amount of oxygen ~73±2%. It is interesting to note that the oxide scales in the M3 and M3Fe alloys

Journal Pre-proof contained very little Mo, which is likely due to the spontaneous oxidation and volatilization of MoO2 to form MoO 3 gas at high temperatures [31]. This is further supported by the respective Mo concentrations in the reacted zones and un-oxidized cores, considering the nominal alloy compositions, Table 6. The Mo content in the reacted zones is larger, but still only accounts for roughly 50-60% of the Mo in the base alloys. The Mo content increases further at the metal core to the concentration of the non-oxidized alloy composition. Closer inspection of the C-3009, M3, and M3Fe oxide scales (Figure 14 (a), (c), and (e), respectively) reveals a complex distribution of porosity (black regions), high atomic number Z (light-color), and low atomic number Z (dark-

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color) regions, likely indicating the presence of different oxide phases. The characterization of

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these regions is beyond of the goal of this work.

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Figure 14. Representative cross-section backscattered electron images of 3009 (a)-(b), M3 (c)(d), and M3Fe (e)-(f) alloys after oxidation exposure at 1200°C for 24 hours. Higher

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magnification images of each oxide are indicated by inset boxes and correspondingly shown in

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(a), (c), and (e), respectively.

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Table 6. SEM-EDS captured from the cross-sections of the 3009, M3, and M3Fe alloys after 24

Region

3009

Metal Core

Nb

Mo

Ti

Fe

W

Hf

O

76.0

-

-

-

6.4

17.6

-

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Alloy

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hours of oxidation at 1200°C. All data is represented in at. %.

M3

M3Fe

Reacted Zone

37.4

-

-

-

2.6

10.6

49.4

Oxide Scale

19.7

-

-

-

1.5

4.6

74.2

Metal Core

66.5

13.4

20.1

-

-

-

-

Reacted Zone

40.4

8

12.4

-

-

-

39.2

Oxide Scale

19.8

0.2

6.1

-

-

-

73.9

Metal Core

64.6

12.8

20.3

2.2

-

-

-

Reacted Zone

40.1

7.8

12.4

1.2

-

-

38.5

Oxide Scale

20.1

0.6

5.7

0.7

-

-

72.9

Journal Pre-proof While all of the alloys contained the same types of regions (“outer oxide scale”, “internal reacted zone”, and “metal core”), Figure 14, both of the Mo-containing alloys exhibited thicker oxide scales and thinner internal reacted zones compared to 3009. In addition, cross-sectional area measurements of the “un-oxidized” metal cores revealed that M3 and M3Fe have a smaller depth of oxygen interaction and thus retained a higher fraction of un-oxidized metal core compared to C-3009 after the same oxidation time (Table 5). The areas reported for each region are normalized with respect to the original sample cross-sectional areas. After 24 hours of oxidation, the C-3009 alloy had ~37.6% remaining metal core while the M3 and M3Fe alloys had 49.1%

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and 48.7% remaining metal core, respectively.

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To better assess the oxide species in each scale, plan-view XRD was conducted after oxidation at 1200°C for 24 hours. Figure 15 (a)-(c) shows the XRD spectra for the C-3009, M3, and M3Fe

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alloys, respectively. In close correlation with the SEM-EDS measurements, Table 6, the oxidized C-3009 alloy displayed signatures consistent with the monoclinic and tetragonal forms of Nb2 O5 ,

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along with the monoclinic and cubic forms of HfO 2 , Figure 15 (a). Similarly, the M3 and M3Fe

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alloys both exhibited peaks consistent with the monoclinic form of Nb 2 O5 and a complex Nb2 TiO7 oxide, Figure 15 (b) and (c), respectively. Unlike the M3 alloy, the M3Fe alloy had extra peaks most closely associated with Nb 3 TiO 6 . Considering the similar retention of un-

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oxidized metal core between the M3 and M3Fe alloys, the presence of this additional phase in

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the Fe-containing alloy may explain the lower mass gain exhibited by the M3Fe alloy, Figure 13.

Figure 15. Plan-view XRD spectra captured from the free oxide scale surface of the C-3009 (a), M3 (b), and M3Fe (c) alloys after 24 hours of oxidation at 1200°C.

4 CONCLUSIONS The effect of small alloying additions of Fe on the microstructure, mechanical properties and oxidation behavior of Nb-Mo-Ti alloys with a single-phase BCC structure was studied. The addition of 2 at.% Fe decreased the average grain size by ~40-50%, increased RT microhardness by ~15%, and increased yield strength by 175-193 MPa in the temperature range of 25-1000°C. At 1000°C and 1200°C, the yield strength of the Fe containing alloys was, respectively, ~65% and 16% higher than the yield strength of the base alloys without Fe. Showing excellent

Journal Pre-proof malleability, the developmental Nb-Mo-Ti-Fe alloys were stronger than the high-strength refractory commercial alloy C-3009 in the temperature range from 25°C to 1000°C and stronger than the medium-strength refractory commercial alloy C-103 at all the studied temperatures. Oxidation studies at 1200°C revealed that the Mo-containing alloys both exhibited lower specific mass gain in comparison to C-3009, with the caveat that Mo-oxides can be highly volatile at this temperature. The C-3009 alloy formed a combination of Nb2 O5 and HfO 2 oxides, as well as a pronounced internally oxidized layer beneath the oxide scale. The M3 and M3Fe alloys both formed Nb2 O 5 and Nb2 TiO7 , while the Fe additions in the M3Fe alloy promoted the formation of

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Nb3 TiO6 . The M3 and M3Fe alloys displayed slightly thicker oxide scales and thinner internal

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oxidation zones compared to C-3009. In addition, both Mo-containing alloys displayed a higher

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retention of un-oxidized material when compared to C-3009.

ACKNOWLEDGEMENTS

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Technical assistance by S. Boone (alloy preparation by vacuum arc melting), L. Griffith

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(extraction of compression samples by EDM) and K. Shewbridge (metallography sample preparation) is greatly appreciated. Work by O.N. Senkov, S.I. Rao and T.I. Daboiku was supported through the Air Force on-site contract FA8650-15-D-5230 managed by UES, Inc.,

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Dayton, Ohio. Authors-Statement

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This work has not been published previously, it is not under consideration for publication

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elsewhere, its publication is approved by all authors and, if accepted, it will not be published elsewhere in the same form, in English or in any other language, including electronically without the written consent of the copyright holder.

Declaration of interests The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

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[10] D.A. Porter, K.E. Easterling, Phase Transformations in Metals and Alloys, 2nd edition ed., CRC Press, Cheltenham, UK, 1992. [11] H.K.D.H. Bhadeshia, Geometry of Crystals, Polycrystals, and Phase Transformations, 1st Edition ed., CRC Press, Boca Raton, FL, USA, 2017. [12] D. Ablitzer, Diffusion of niobium, iron, cobalt, nickel and copper in niobium, Philosophical Magazine, 35 (1977) 1239-1256. [13] D. Ablitzer, The mechanism of diffusion of iron in niobium, Philosophical Magazine, 36 (1977) 391-411. [14] R.A. Perez, H. Nakajima, F. Dyment, Diffusion in α-Ti and Zr, Mater. Trans., 44 (2003) 213.

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hardening in body-centered cubic quartenary alloys interpreted using Suzuki's kink-solute interaction model, Scripta Mater., 165 (2019) 103-106. [24] S.I. Rao, C. Woodward, B. Akdim, O.N. Senkov, A model for interstitial solid solution strengthening of body-centered cubic metals, Materialia., (2019). [25] J.B. Cohen, M.E. Fine, Some aspects of short-range order, J. Phys. Radium, 23 (1962) 749762. [26] D.R. Trinkle, C. Woodward, The chemistry of deformation: how solutes soften pure metals, Science, 310 (2005) 1665-1667. [27] E. Antillon, C. Woodward, S.I. Rao, B. Akdim, T.A. Parthasarathy, The effect of chemical short range order on the strengthening properties of a model FCC high entropy alloy, Acta Mater., (2019).

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Nb-Mo-Ti-Fe alloys with single-phase BCC structures and densities of ~ 7.8 g/cm3 are

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

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Highlights

reported.

Additions of 2at.% Fe considerably increase strength of the alloys at temperatures from

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

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25°C to 1200°C.

Nb-Mo-Ti-Fe alloys have good combination of strength and ductility.



The properties also depend on the amount of interstitial elements.



The alloys have slightly improved oxidation behavior relative to conventional refractory

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alloys

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