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Materials Science and Engineering A 483–484 (2008) 199–202

Effect of Mo on mechanical properties andmicrostructure of Nb–Ti–C alloys

Rengen Ding, Huisheng Jiao 1, Ian P. Jones ∗Department of Metallurgy and Materials, University of Birmingham, Birmingham B15 2TT, UK

Received 6 June 2006; received in revised form 21 November 2006; accepted 5 December 2006

bstract

The microstructures and mechanical properties of Nb–Ti–C alloys containing different Mo contents have been investigated. The alloy withouto contains Nb solid solution (Nbss), (Ti,Nb) C/�(Nbss) eutectic, (Ti,Nb) C and (Ti,Nb)C. The addition of Mo restrains the precipitation of

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Ti,Nb)C. Compression testing has been carried out at room temperature and 1473 K. The alloys showed good ductility at room temperature andb20Ti12.5C20Mo produced a 0.2% offset yield stress of 300 MPa at 1473 K. The room temperature and high temperature strengths increase with

ncreasing Mo, but the addition of Mo decreases the room temperature ductility. 2007 Elsevier B.V. All rights reserved.

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eywords: Nb–Ti–C alloy; Molybdenum; Microstructure; Mechanical properti

. Introduction

Niobium-based alloys are one of the most promising refrac-ory alloys for high temperature structural application in placef Ni-based superalloys. Niobium has a very high melting pointf 2741 K, a relatively low density (8.57 g/cm3) and good duc-ility at room temperature. However, the strength of niobiumecreases substantially at temperatures above 1200 K. Variouspproaches have hence been used to improve the strength of Nbt elevated temperature, including solid solution strengtheningith Mo, W, Ta and Hf [1] and dispersion or composite strength-

ning with intermetallic compounds such as Nb3Al and Nb5Si2], as well as with a carbide phase [3,4]. Jiao et al. [4] deter-ined that the addition of TiC to Nb-based alloys increases the

trength at room and elevated temperatures, but that the strengthf simple Nb–Ti–C alloys at elevated temperature is not ade-uate. Solid solution strengthening is also necessary. Mo appearso be the most promising element for solid solution strength-ning because Mo not only improves the strength of Nb, butlso does not easily form carbide. Therefore, the microstructures

nd mechanical properties of Nb–Ti–C alloys with different Moontents have been investigated in this study.

∗ Corresponding author. Tel.: +44 121 4145184; fax: +44 121 4145232.E-mail address: i.p.jones@bham.ac.uk (I.P. Jones).

1 Present address: Gatan Company, UK.

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921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2006.12.178

. Experimental procedure

Three alloys with nominal compositions Nb20Ti12.5C,b20Ti12.5C10Mo and Nb20Ti12.5C20Mo (at.%) were pre-ared by arc-melting and casting into buttons. Conventionalonstant strain rate deformation with a strain rate of about 10−4/sas performed on all three as-cast alloys. High temperature com-ression was carried out under vacuum. TEM foils were preparedy electropolishing and examined in an FEI Tecnai F20. Varioushases’ compositions were measured using energy-dispersivepectrometry (EDS) in transmission electron microscopy (TEM)nd wavelength-dispersive spectrometry (WDS) in scanninglectron microscopy (SEM).

. Results and discussion

.1. As-cast microstructures

SEM backscattered electron images (Fig. 1) suggested thathe bright phase is Nb-base solid solution (Nbss) and that thelack phase is carbide. Dendritic carbides were observed in allhe alloys, indicating that the carbide solidifies from the liquidrst. The carbides in the alloy without Mo could be designated

s eutectic carbide/�(Nbss), primary carbide and secondaryarbide (inset of Fig. 1a). However, no secondary (small needle-ike) carbide was observed in the alloys with Mo. This suggestshat Mo restrains the precipitation of secondary carbide via mod-

200 R. Ding et al. / Materials Science and Engineering A 483–484 (2008) 199–202

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Fig. 1. SEM backscattered electron images of Nb20Ti12.5C (a)

fication of the solubility limit of carbon in Nb or/and of theinetics of precipitation of the carbide. Image analysis indicatedhat the addition of Mo did not change the total volume fractionf the carbides too much: the volume fraction of the carbides was1% for the alloy without Mo, and 23 and 22%, respectively, forhe alloys with 10Mo and 20Mo. SEM/WDS analysis showedhat the primary carbide was not (Ti,Nb)C but substoichiometricith a composition around (Ti,Nb)2C (Table 1).Superlattice maxima were observed in selected-area diffrac-

ion (SAD) patterns from the carbide in the eutectic as illustratedn Fig. 2a, resulting from the vacancy related ordering of the car-ides, as shown in Fig. 2b. EDS results (Table 1) suggest thathe C/(Ti + Nb) ratio in the carbide is around 0.5. This indicateshat the carbide is substoichiometric with a composition aroundTi,Nb)2C, and not (Ti,Nb)C. This is consistent with orderedTi,Nb)2C with a lattice parameter a = 0.429 nm [5]. The EDSnalysis (Table 1) and diffraction (Fig. 2c) confirmed that theeedle-like particles in the eutectic were �(Nb) phase.

Some subgrain boundaries (inset to Fig. 2a) could be observedn the primary carbide which was also identified to be (Ti,Nb)2Cia EDS and diffraction analysis.

The EDS analysis (Table 1) and SAD patterns (inset to Fig. 3)uggested that the secondary carbide was disordered (Ti,Nb)C.

.2. Mechanical tests

Vickers hardness tests showed that the hardness increasesith increasing Mo content (Table 2). The compression

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able 1he compositions of Nbss, carbides and needle-like particles in the eutectic in Nb20T

Nb (at.%) Ti (a

bss 69.3 ± 1.43 17.

rimary carbideWDS 34.6 ± 0.35 31.EDS 32.5 ± 0.58 30.

utecticCarbide 33.8 ± 0.66 29.Needle-like particle 73.27 ± 1.45 16.0

econdary carbide 22.0 ± 0.65 24.

ote: The composition of the primary carbide was measured using TEM/EDS and SEo Mo peak was observed.

b20Ti12.5C10Mo (b). Insets are higher magnification images.

tress–strain curves of the three alloys at room temperatureFig. 4) show that the addition of Mo improved significantlyhe yield strength of the alloys but decreased the ductility. The.2% offset strength of Nb20Ti12.5C20Mo is about 1000 MPat room temperature. The compression tests were interrupted forhe alloys containing Mo which failed when the strain reachedround 25%. Since the addition of Mo did not much changehe volume fractions of carbides, no Mo was detected in thearbide and no Mo-rich phase was observed, the improve-ent of strength in the alloys is mainly attributed to solid

olution strengthening by Mo. The effect of solid solution hard-ning on the yield strength, σ, of the material is given by= 2.5μ(δ)4/3c [6]. where μ is the shear modulus (38 GPa

or Nb at room temperature), δ is a misfit parameter and cs solute concentration. The misfit parameter δ is defined as= (1/a)(∂a/∂c) (5.81% for Nb–Mo binary alloys [7]) whereis the lattice parameter of Nb. Therefore, the predicted val-

es of solid solution hardening for 10 and 20 at.% Mo in Nblloys at room temperature are 215 and 430 MPa, respectively.his is supported by the experimental results (300 and 420 MPa)hich are again in agreement with the study of Tan et al.

1].The 0.2% offset strengths for the three alloys at 1473 K

btained by constant strain rate compression under vacuum

re given in Table 2. With increasing testing tempera-ure, the 0.2% flow stress decreases dramatically. Similarlyo the results at room temperature, on increasing the

o content in the alloy, the yield stress increases. The

i12.5C10Mo, and of secondary carbide from Nb20Ti12.5C

t.%) Mo (at.%) C (at.%)

4 ± 0.47 13.3 ± 1.47

1 ± 0.31 / 34.3 ± 0.433 ± 0.51 / 37.2 ± 0.61

8 ± 0.49 / 36.4 ± 0.639 ± 0.47 10.63 ± 1.55

8 ± 0.57 / 53.2 ± 0.55

M/WDS, whereas the other phases were measured using TEM/EDS. / indicates

R. Ding et al. / Materials Science and Engineering A 483–484 (2008) 199–202 201

Fig. 2. TEM image of eutectic and primary carbide (a); [0 1 1], [1 1 2] and [0 0 1] diffraction patterns of the carbide in the eutectic (b); [1 1 1] and [0 0 1] diffractionpatterns of needle-like � (Nb) phase in the eutectic.

Table 2Variation of room temperature hardness and 0.2% offset yield stress with Mo content

Nb20Ti12.5C Nb20Ti12.5C10Mo Nb20Ti12.5C20Mo

Hardness (HV) 235 323 433σ0.2 (293 K) (MPa) 580 880 1030σ0.2 (1473 K) (MPa) 140 150 300

202 R. Ding et al. / Materials Science and Engineering A 483–484 (2008) 199–202

Fig. 3. A TEM image showing secondary carbides in Nb20Ti12.5C. Inset shows[0 1 1] SAD pattern from the carbide.

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Fig. 4. Compression stress–strain curves at room temperature.

.2% flow stress of Nb20Ti12.5C20Mo is 300 MPa at473 K compared with 140 MPa for the alloy withouto.

.3. Deformation microstructure

After 30% compression at room temperature, the microstruc-ure was examined using TEM. A high density of dislocationsan still be seen in the Nbss matrix (Fig. 5), which enduredarge scale deformation. The carbides in these alloys also expe-ienced considerable deformation in compression. An image ofdeformed (Ti,Nb)C carbide is shown in the inset to Fig. 5.

high density of lines is observed inside the carbide. Streaks

erpendicular to these features are clearly seen in the diffractionattern (inset to Fig. 5). These lines represent a high density oficrotwins.

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ig. 5. Bright field TEM image obtained using g = 110 from Nb20Ti12.5C after0% compression. A high density of dislocations is evident. Insets show theefects in a deformed carbide and the SADP obtained from the deformed carbide.

. Conclusions

Based on an investigation of the microstructure and mechan-cal tests of Nb–Ti–C alloys with various Mo contents:

(i) Nb20Ti12.5C consists of Nb solid solution (Nbss),(Ti,Nb)2C/�(Nbss) eutectic, primary irregularly shaped(Ti,Nb)2C carbide and secondary needle shaped (Ti,Nb)Ccarbide. In the alloys with Mo, no secondary (Ti,Nb)Ccarbide was observed.

(ii) The 0.2% offset strength increases with increasing Mocontent, whereas the addition of Mo decreases the ductil-ity. Nb20Ti12.5C20Mo appears to possess substantial hightemperature strength (300 MPa).

iii) After deformation, a high density of dislocations wasobserved in Nbss with a high density of microtwins in thecarbides.

cknowledgements

The authors would like to thank EPSRC and DSTL for finan-ial support.

eferences

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5] W.B. Pearson, A Handbook of Lattice Spacings and Structures of Metals

and Alloys, Pergamon Press Inc., New York, 1964, pp. 948–949.6] A.H. Cottrell, Dislocations and Plastic Flow in Crystals, Clarendon Press,

1953.7] H.W. King, J. Mater. Sci. 1 (1966) 79–90.