Mechanical Properties and Dislocation Substructure of ... · Mechanical Properties and Dislocation...

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Mechanical Properties and Dislocation Substructure of Inconel 690 Alloy Impacted at Cryogenic Temperatures Woei-Shyan Lee + and Ming-Chia Hsu Department of Mechanical Engineering, National Cheng Kung University, Tainan 701, Taiwan The mechanical response and dislocation substructure of Inconel 690 during impact deformation are investigated at strain rates of 2 © 10 3 ³ 6 © 10 3 s ¹1 and temperatures of ¹150°C, 0°C and 25°C using a compressive split-Hopkinson pressure bar system. The results show that the ow stress, work hardening rate and strain rate sensitivity all increase with increasing strain rate or decreasing temperature. By contrast, the activation volume reduces as the strain rate increases or the temperature decreases. Moreover, the temperature sensitivity increases with both increasing strain rate and increasing temperature. Optical microscopy observations show that adiabatic shear bands are formed at the highest strain rate of 6 © 10 3 s ¹1 in all of the tested specimens. In addition, it is shown that adiabatic shear localisation is the major cause of specimen failure in every case. The dislocation density increases, and the cell size decreases, as the strain rate is increased or the temperature decreased. The change in the dislocation density and cell size is found to have a signicant effect on the ow stress and work hardening behaviour of the Inconel 690 specimens; particularly at high strain rates or low temperatures. [doi:10.2320/matertrans.M2014165] (Received May 1, 2014; Accepted August 12, 2014; Published October 3, 2014) Keywords: Inconel 690, impact, cryogenic deformation, strain rate effect, dislocation 1. Introduction Inconel 690 is one of the most important nickel-based superalloys, and is used extensively in the construction of gas turbines, nuclear reactors, spacecraft, liquid-fuel rocket engines, autonomous underwater vehicles, cryogenic storage tanks, and combustion systems due to its good corrosion resistance, high strength and outstanding weldability. 1,2) In many applications, Inconel 690 components are subject to deformation at both high strain rates and cryogenic temper- atures. For example, in autonomous underwater vehicles, 3) cryogenic storage tanks and liquid-fuel rocket engines, 4) the components frequently operate at strain rates of 10 2 ³ 10 4 s ¹1 and temperatures as low as ¹200°C. The literature contains many investigations into the response and microstructural evolution of Inconel 690 under quasi-static or dynamic loading conditions at room or elevated temperatures. 5-10) By contrast, the high strain rate deformation behaviour of Inconel 690 under cryogenic temperatures has attracted relatively little attention. However, in designing Inconel 690 structural components for implementation in arduous operat- ing conditions, such an understanding is urgently required. The mechanical properties of engineering metals and alloys are signicantly dependent on the strain rate and temper- ature. 11-15) Generally speaking, the work hardening behaviour and ow stress of such materials increases with increasing strain rate or decreasing temperature. Moreover, the strain rate sensitivity of the ow stress increases signicantly at strain rates greater than 10 3 s ¹1 . Various rate-controlling mecha- nisms have been proposed to explain this deformation behaviour, including dislocation damping, thermal activation and dislocation generation. 16,17) In practice, the dynamic deformation behaviour of metallic materials at high strain rates is determined by the result of a competition process between the work hardening effect induced by plastic deformation and the thermal softening effect induced by the heat produced during deformation. For deformation under adiabatic conditions, the thermal softening effect exceeds the work hardening effect, and thus material instability occurs. Once an unstable condition is attained, extreme localisation of the deformation occurs; resulting in the formation of adiabatic shear bands. 18,19) The literature contains various proposals for the possible mechanisms responsible for the initiation and propagation of adiabatic shear bands 20-22) and for constitutive equations to describe their behaviour. 23-25) Many studies have shown that impact loading leads to the formation of dislocations in the deformed microstructure, which result in turn in a change in the mechanical proper- ties. 26-28) In addition, it has been widely reported that the dislocation density increases with increasing strain and strain rate. 29-31) Moreover, the square root of the dislocation density is inversely related to the ow stress. 32-34) The high strain rates associated with impact loading result in a rapid multiplication of the dislocations, and thus the equilibrium structure of high stacking fault FCC metals and alloys is characterised by a large number of dislocation cells. 35) As the strain rate increases, the size of the dislocation cells reduces; leading to a greater rate of work hardening and an enhanced ow resistance. It has been shown that the relationship between the average dislocation cell size, the ow stress and the hardening behaviour can be described using a common power law model. 36) The present study utilises a compressive split-Hopkinson pressure bar (SHPB) system to investigate the impact response of Inconel 690 at temperatures ranging from ¹150 ³ 25°C and strain rates of 2 © 10 3 ³ 6 © 10 3 s ¹1 . Transmission electron microscope (TEM) observations are performed to examine the dislocation substructures of the deformed specimens. Finally, the correlation between the properties of the deformation microstructure and the me- chanical response of the Inconel 690 specimens is system- atically examined and discussed. 2. Experimental Procedure The Inconel 690 alloy used in the present study was + Corresponding author, E-mail: wslee@mail.ncku.edu.tw Materials Transactions, Vol. 55, No. 11 (2014) pp. 1689 to 1697 © 2014 The Japan Institute of Metals and Materials

Transcript of Mechanical Properties and Dislocation Substructure of ... · Mechanical Properties and Dislocation...

Mechanical Properties and Dislocation Substructureof Inconel 690 Alloy Impacted at Cryogenic Temperatures

Woei-Shyan Lee+ and Ming-Chia Hsu

Department of Mechanical Engineering, National Cheng Kung University, Tainan 701, Taiwan

The mechanical response and dislocation substructure of Inconel 690 during impact deformation are investigated at strain rates of2 © 103³ 6 © 103 s¹1 and temperatures of ¹150°C, 0°C and 25°C using a compressive split-Hopkinson pressure bar system. The results showthat the flow stress, work hardening rate and strain rate sensitivity all increase with increasing strain rate or decreasing temperature. By contrast,the activation volume reduces as the strain rate increases or the temperature decreases. Moreover, the temperature sensitivity increases with bothincreasing strain rate and increasing temperature. Optical microscopy observations show that adiabatic shear bands are formed at the higheststrain rate of 6 © 103 s¹1 in all of the tested specimens. In addition, it is shown that adiabatic shear localisation is the major cause of specimenfailure in every case. The dislocation density increases, and the cell size decreases, as the strain rate is increased or the temperature decreased.The change in the dislocation density and cell size is found to have a significant effect on the flow stress and work hardening behaviour of theInconel 690 specimens; particularly at high strain rates or low temperatures. [doi:10.2320/matertrans.M2014165]

(Received May 1, 2014; Accepted August 12, 2014; Published October 3, 2014)

Keywords: Inconel 690, impact, cryogenic deformation, strain rate effect, dislocation

1. Introduction

Inconel 690 is one of the most important nickel-basedsuperalloys, and is used extensively in the construction of gasturbines, nuclear reactors, spacecraft, liquid-fuel rocketengines, autonomous underwater vehicles, cryogenic storagetanks, and combustion systems due to its good corrosionresistance, high strength and outstanding weldability.1,2) Inmany applications, Inconel 690 components are subject todeformation at both high strain rates and cryogenic temper-atures. For example, in autonomous underwater vehicles,3)

cryogenic storage tanks and liquid-fuel rocket engines,4) thecomponents frequently operate at strain rates of 102³ 104 s¹1

and temperatures as low as ¹200°C. The literature containsmany investigations into the response and microstructuralevolution of Inconel 690 under quasi-static or dynamicloading conditions at room or elevated temperatures.5­10)

By contrast, the high strain rate deformation behaviour ofInconel 690 under cryogenic temperatures has attractedrelatively little attention. However, in designing Inconel 690structural components for implementation in arduous operat-ing conditions, such an understanding is urgently required.

The mechanical properties of engineering metals and alloysare significantly dependent on the strain rate and temper-ature.11­15) Generally speaking, the work hardening behaviourand flow stress of such materials increases with increasingstrain rate or decreasing temperature. Moreover, the strain ratesensitivity of the flow stress increases significantly at strainrates greater than 103 s¹1. Various rate-controlling mecha-nisms have been proposed to explain this deformationbehaviour, including dislocation damping, thermal activationand dislocation generation.16,17) In practice, the dynamicdeformation behaviour of metallic materials at high strainrates is determined by the result of a competition processbetween the work hardening effect induced by plasticdeformation and the thermal softening effect induced by theheat produced during deformation. For deformation under

adiabatic conditions, the thermal softening effect exceeds thework hardening effect, and thus material instability occurs.Once an unstable condition is attained, extreme localisation ofthe deformation occurs; resulting in the formation of adiabaticshear bands.18,19) The literature contains various proposals forthe possible mechanisms responsible for the initiation andpropagation of adiabatic shear bands20­22) and for constitutiveequations to describe their behaviour.23­25)

Many studies have shown that impact loading leads to theformation of dislocations in the deformed microstructure,which result in turn in a change in the mechanical proper-ties.26­28) In addition, it has been widely reported that thedislocation density increases with increasing strain and strainrate.29­31) Moreover, the square root of the dislocation densityis inversely related to the flow stress.32­34) The high strainrates associated with impact loading result in a rapidmultiplication of the dislocations, and thus the equilibriumstructure of high stacking fault FCC metals and alloys ischaracterised by a large number of dislocation cells.35) As thestrain rate increases, the size of the dislocation cells reduces;leading to a greater rate of work hardening and an enhancedflow resistance. It has been shown that the relationshipbetween the average dislocation cell size, the flow stress andthe hardening behaviour can be described using a commonpower law model.36)

The present study utilises a compressive split-Hopkinsonpressure bar (SHPB) system to investigate the impactresponse of Inconel 690 at temperatures ranging from¹150³ 25°C and strain rates of 2 © 103³ 6 © 103 s¹1.Transmission electron microscope (TEM) observations areperformed to examine the dislocation substructures of thedeformed specimens. Finally, the correlation between theproperties of the deformation microstructure and the me-chanical response of the Inconel 690 specimens is system-atically examined and discussed.

2. Experimental Procedure

The Inconel 690 alloy used in the present study was+Corresponding author, E-mail: [email protected]

Materials Transactions, Vol. 55, No. 11 (2014) pp. 1689 to 1697©2014 The Japan Institute of Metals and Materials

acquired in plate form from Inco Alloys International(Huntington, Virginia, USA). Prior to receipt, the plateswere solution annealed at 1150°C for 1 h and then quenchedin water. The chemical composition of the as-receivedsolution-annealed alloy (mass%) was as follows: 29.33%Cr, 10.01% Fe, 0.29% Ti, 0.19% Al, 0.10% Cu, 0.03% Co,0.02% Mo, 0.02% W, 0.01% C, and a balance of Ni.Figure 1(a) presents an optical micrograph of a typicalspecimen prepared from the as-received plate. As shown, themicrostructure consists of equiaxed grains with an averagesize of 42 µm and a small number of annealing twins. Inaddition, discrete inter- and intra-granular (M23C6) chromiumcarbides and titanium carbonitrides Ti(C,N) are also observedwithin the matrix. The M23C6 and Ti(C,N) particles create abarrier effect, which accelerates the generation and accumu-lation of dislocations during impact deformation. Figure 1(b)presents a TEM image of the dislocation structure of the as-received alloy. It is seen that the microstructure contains arelatively small number of dislocations. It is further observedthat most of the dislocations have an edge characteristic andare aligned along specific crystallographic planes. Cylindricalspecimens with a length of 5 « 0.1mm and a diameter of5.2mm were machined from the alloy plates with the impactaxes parallel to the longitudinal direction and were finished toa final diameter of 5 « 0.1mm via a centre-grinding process.

The dynamic impact tests were conducted using a SHPBsystem at strain rates of 2 © 103 s¹1, 4 © 103 s¹1 and6 © 103 s¹1, respectively, and temperatures of ¹150°C, 0°Cand 25°C. The striker bar, incident bar and transmitter barof the SHPB apparatus (see Fig. 2) were made of DC53 diesteel and had a diameter of 15mm. The incident bar andtransmitter bar each had a length of 1m, while the striker barhad a length of 317mm. In performing the tests, thespecimens were sandwiched between the incident bar andthe transmitter bar, and the incident bar was then impacted bythe striker bar (fired by a gas gun). To ensure a uniaxialdeformation state (i.e., frictionless conditions), the endsurfaces of each specimen were lubricated using commercialmolybdenum disulfide (Molykote). The low testing temper-atures of 0°C and ¹150°C were obtained by fitting arefrigeration system around the specimen. Liquid nitrogenand alcohol were added to the system at the beginning of theexperimental tests and were replenished periodically in orderto maintain a constant fluid level. Following impact, thestrain pulses within the incident and transmitter bars weredetected by means of strain gauges mounted at the midpointof each bar. The compressive stress-strain response of thespecimen was then established in accordance with uniaxialelastic wave theory.37) Specifically, based on the recordedvalues of the incident, reflected and transmitted pulses (i.e.,¾i, ¾r and ¾t, respectively), the average compressive engineer-ing strain ¾e, nominal strain rate _¾, and average compressiveengineering stress ·e were calculated as follows:

¾e ¼ ð2C0=L0ÞZ

¾rdt; ð1Þ

_¾ ¼ 2C0¾r=L0; ð2Þ·e ¼ EðA=A0Þ¾t; ð3Þ

where C0 is the longitudinal wave velocity; L0 is the effectivegauge length of the specimen; E is the Young’s modulus ofthe split-Hopkinson bars; and A and A0 are the cross-sectionalareas of the split-Hopkinson bars and specimen, respectively.Having obtained the engineering stress-strain data, the truestress (·T) and true strain (¾T) were determined respectivelyas

·T ¼ ·eð1� ¾eÞ: ð4Þ¾T ¼ � lnð1� ¾eÞ: ð5Þ

To ensure the reliability of the experimental results, threetests were performed under each of the considered strain rate/temperature conditions. The stress-strain curves for each testcondition were then obtained by means of a least-squaresfitting technique.

Following the impact tests, the deformed specimens wereground progressively using abrasive papers with grit sizesranging from 180 to 1200 mesh. The ground specimens werepolished with a micro-cloth dipped in a slurry of water and0.3-µm alumina particles, and were then etched in a solutioncomprising 70ml of H3PO4 and 30ml of D.I. water. The

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Fig. 1 Micrographs of test specimen prepared from solution-annealedInconel 690 plate: (a) Optical microscope, and (b) TEM.

Fig. 2 Schematic illustration of compressive split-Hopkinson pressure barsystem.

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specimens were observed using an Axiovert 200 Mat opticalmicroscope. Thin foils and discs for TEM observations wereprepared from the deformed specimens by cutting slices witha thickness of 0.5mm in a direction perpendicular to thecompression axis. The slices were ground to a final thicknessof 0.2mm using 1200 grit size soft paper and discs with adiameter of 3mm were then trepanned from each slice usingan electron discharge machining system. The foils wereelectro-polished in a solution of 10% perchloric acid and90% methanol at a temperature of ¹20°C using a twin-jetpolishing machine with an agitation voltage of 20V DC. Thespecimens were observed using a JEOL TEM-3010 scanningtransmission electron microscope with an operating voltageof 300 kV. In addition, the detailed characteristics of thefracture features were examined using a Philip’s XL-40FEGscanning electron microscope (SEM) with an operatingvoltage of 20 kV.

3. Results and Discussions

3.1 Stress-strain curvesFigures 3(a)³3(c) show the true stress-strain curves of the

Inconel 690 specimens tested at strain rates of 2 © 103 s¹1,4 © 103 s¹1and 6 © 103 s¹1, respectively. It is seen that theflow behaviour of the specimens is strongly dependent onboth the strain rate and the temperature. Specifically, the flowstress increases with increasing strain rate or decreasingtemperature. From inspection, the maximum increase in theflow stress is found to be 420MPa as the temperature isreduced from 25°C to ¹150°C at a strain rate of 6 © 103 s¹1

(Fig. 3(c)). For each strain rate and temperature, the rate atwhich the flow stress increases with increasing strain isdifferent; indicating a difference in the work hardeningrate under different loading conditions. Furthermore, for agiven strain rate, the maximum true strain decreases withdecreasing temperature. Thus, it is inferred that a loss ofductility occurs at lower deformation temperatures. It isnoted that none of the specimens fail when tested at strainrates lower than 4 © 103 s¹1. However, at the highest strainrate of 6 © 103 s¹1, catastrophic failure occurs in all of thespecimens. In other words, the Inconel 690 specimensexperience a loss in toughness and fracture resistance whendeformed under higher strain rates and cryogenic temper-atures.

The true stress-strain curves shown in Fig. 3 indicate thatthe work hardening rate (@·=@¾) is dependent on the strain,strain rate and temperature. Table 1 shows the variation of thework hardening rate with the temperature at true strainsranging from 0.04 to 0.7 and strain rates of 2 © 103 s¹1,4 © 103 s¹1 and 6 © 103 s¹1, respectively. It is observed thatfor a given strain rate and temperature, the work hardeningrate decreases with increasing strain; indicating that a greaterrate of dislocation generation and multiplication takes placeduring the early stages of the deformation process. For aconstant true strain, the work hardening rate increases withincreasing strain rate, but decreases with increasing temper-ature. Thus, it is inferred that a higher strain rate induces amore intensive generation and accumulation of dislocationsduring deformation, while a higher temperature prompts astronger microstructural recovery process.

3.2 Strain rate sensitivity and thermal activation volumeFigure 4(a) shows the variation of the flow stress with the

logarithmic strain rate as a function of the temperature at truestrains of 0.1 and 0.5, respectively. It is seen that for bothvalues of the true strain, the flow stress increases more

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(c)

Fig. 3 Stress-strain curves of Inconel 690 specimens deformed at temper-atures of ¹150°C, 0°C and 25°C and strain rates of: (a) 2 © 103 s¹1,(b) 4 © 103 s¹1, and (c) 6 © 103 s¹1.

Mechanical Properties and Dislocation Substructure of Inconel 690 Alloy Impacted at Cryogenic Temperatures 1691

rapidly at strain rates in the range of 4 © 103³ 6 © 103 s¹1. Inother words, the strain rate sensitivity of the Inconel 690alloy specimens increases at higher strain rates. Someresearchers have argued that the abrupt increase in the strainrate sensitivity of engineering metals and alloys at higherstrain rates is the result of a change in the rate-controllingmechanism from thermal activation to dislocation drag.16)

However, other studies have attributed the enhanced strainrate sensitivity to an increased rate of dislocation generation,the rapid formation of twin structures, or a greater degree ofmartensite transformation.17,28)

The strain rate sensitivity of the Inconel 690 specimens canbe characterised via the following parameter:

¢ ¼ ð@·=@ ln _¾Þ ¼ ·2 � ·1

lnð_¾2=_¾1Þ; ð6Þ

where the compressive stresses ·2 and ·1 are obtained fromtests conducted at average strain rates of _¾2 and _¾1,respectively. Figures 4(b) and 4(c) show the variation of thestrain rate sensitivity with the strain as a function of thetemperature given strain rate ranges of 2 © 103³ 4 © 103 s¹1

and 4 © 103³ 6 © 103 s¹1, respectively. Comparing the twofigures, it is seen that for a given strain and deformationtemperature, the strain rate sensitivity increases with increas-ing strain rate. In addition, it is observed that for a given strainrate range, the strain rate sensitivity increases with increasingstrain, but decreases with increasing temperature. At low strainrates (i.e., 2 © 103³ 4 © 103 s¹1), the rate of increase of thestrain rate sensitivity with increasing strain is almost the samefor all three deformation temperatures. However, at high strainrates (i.e., 4 © 103³ 6 © 103 s¹1), the strain rate sensitivityincreases more rapidly with increasing strain as the temper-ature reduces. It is thought that the increase in the strain ratesensitivity at higher strain rates and lower temperatures is theresult of a greater dislocation multiplication rate.

The plastic deformation of metals and alloys is known tobe governed by a thermal activation process. For a givenmaterial, the activation volume, ¯�, can be computed as38,39)

¯� ¼ @�G�

@·¼ KT

@ ln _¾

� �¼ KT

¢; ð7Þ

where �G� is the activation energy associated with theplastic deformation process, K is the Boltzmann constant and

has a value of K = 1.38 © 10¹23 JK¹1, T is the absolutetemperature, and ¢ is the strain rate sensitivity. Figures 5(a)and 5(b) show the variation of the activation volume with thetrue strain as a function of the temperature given strain rate

Table 1 Variation of work hardening rate of Inconel 690 specimens asfunction of strain, strain rate and temperature.

T (°C)StrainRate(s¹1)

Work Hardening Rate (MPa/Unit Strain)

¾ = 0.04 ¾ = 0.1 ¾ = 0.3 ¾ = 0.5 ¾ = 0.7

¹150

2000 2698 1725 652 345 ®

4000 2986 2041 970 600 ®

6000 3508 2520 1460 1040 ®

0

2000 2574 1604 542 268 ®

4000 2890 1956 887 520 ®

6000 3253 2329 1290 880 ®

25

2000 2487 1487 466 222 129

4000 2749 1879 806 454 290

6000 3016 2120 1070 680 480

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Fig. 4 (a) Variation of true stress with logarithmic strain rate as function oftemperature and true strain. Variation of strain rate sensitivity with truestrain as function of temperature over strain rate ranges of: (b) 2 © 103 to4 © 103 s¹1 and (c) 4 © 103 to 6 © 103 s¹1.

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ranges of 2 © 103³ 4 © 103 s¹1 and 4 © 103³ 6 © 103 s¹1,respectively. It is noted that the values of the activationvolume are normalised by dividing ¯� by b3, where b is theBurgers vector and has a value of 2.47 © 10¹10m for Inconel690 alloy.10) It is seen that for a constant strain and strain raterange, the activation volume increases with increasingtemperature. However, for a constant temperature, theactivation volume decreases with increasing strain and strainrate. The increased activation volume at a higher temperatureimplies that the stress required to accomplish deformationis reduced due to the assistance of the thermal activationenergy. An inspection of the two figures shows that thenormalised activation volume is more than unity only forthe specimens tested at 0°C and 25°C under strain rates of2 © 103 s¹1³ 4 © 103 s¹1 and true strains of less than 0.2and 0.35, respectively (Fig. 5(a)). This indicates that thedeformation of the present Inconel 690 specimens is notdominated by thermal activation, but is in fact determinedmainly by enhanced dislocation generation, as reported byConrad.38,39)

3.3 Temperature sensitivityFigure 6(a) shows the variation of the flow stress with the

deformation temperature as a function of the strain and strainrate. The results show that a significant thermal softeningeffect occurs as the deformation temperature is increased. Inaddition, it is seen that the rate of reduction in the flow stresswith increasing temperature is particularly apparent in thespecimen deformed at the highest strain rate of 6 © 103 s¹1.The temperature sensitivity of the deformed specimens canbe quantified as follows:

na ¼ @·=@T ¼ jð·2 � ·1Þ=ðT2 � T1Þj; ð8Þwhere the compressive stresses ·2 and ·1 are obtained in testsconducted at temperatures of T2 and T1, respectively.Figures 6(b) and 6(c) show the variation of the temperaturesensitivity with the true strain as a function of the strain ratefor temperature ranges of ¹150³0°C and 0³25°C, respec-tively. It is observed that for a given temperature range, thetemperature sensitivity increases with increasing strain andstrain rate; indicating a deformation-induced heating effect.Moreover, for a constant strain and strain rate, the temper-ature sensitivity increases with increasing temperature. Forboth temperature ranges, the temperature sensitivity increasesmore rapidly at a strain rate of 6 © 103 s¹1. This is most likelythe result of a greater adiabatic heating effect during thedynamic deformation process, and accounts for the fact thatspecimen failure occurs only at the highest strain rate of6 © 103 s¹1, as shown in Fig. 3(c).

3.4 Fracture feature observationsFigures 7(a)³7(c) present optical micrographs of the

Inconel 690 specimens deformed at a strain rate of 6 ©103 s¹1 and temperatures of 25°C, 0°C and ¹150°C,respectively. Note that the loading direction (LD) isperpendicular to the image plane. In all three micrographs,a narrow adiabatic shear band is observed tracing a circularpath along the transverse section perpendicular to the impactdirection. A close inspection reveals the presence of micro-crack nucleation, growth and coalescence within the shearbands. Figures 8(a) and 8(b) present optical micrographs ofthe cross-section parallel to the loading direction for speci-mens deformed at a strain rate of 6 © 103 s¹1 and temper-atures of 25°C and ¹150°C, respectively. Both figures showthe presence of an adiabatic shear band running across thespecimen at an angle of 45° to the loading direction. Overall,the images presented in Figs. 7 and 8 suggest that crackinitiation and propagation occurs within these adiabaticbands, and leads to the eventual catastrophic failure of thespecimens.

Figures 9(a)³9(c) present SEM fractographs of the speci-mens deformed at a strain rate of 6 © 103 s¹1 and temper-atures of 25°C, 0°C and ¹150°C, respectively. It is seen thatthe fracture surfaces contain dimple- and/or cleavage-likefeatures. For the specimen deformed at 25°C, the dimples areevenly distributed over the fracture surface; indicating aductile failure mode. At a lower deformation temperature of0°C, the fracture surface contains both dimple structures andcleavage structures. However, the dimples are smaller andfewer in number than those in the specimen deformed at25°C. For the lowest deformation temperature of ¹150°C,

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Fig. 5 Variation of thermal activation volume with true strain as functionof temperature over strain rate ranges of: (a) 2 © 103 to 4 © 103 s¹1; and(b) 4 © 103 to 6 © 103 s¹1.

Mechanical Properties and Dislocation Substructure of Inconel 690 Alloy Impacted at Cryogenic Temperatures 1693

the fracture surface is comprised almost entirely of cleavagestructures. Overall, the fractographs indicate that the Inconel690 specimens experience a significant loss in ductility at

lower deformation temperatures. Note that this finding isconsistent with the flow stress-strain results presented inFig. 3(c), which show that the fracture strain reduces withreducing temperature.

3.5 Dislocation substructure observationsFigures 10(a)³10(c) present TEM images of the Inconel

690 specimens deformed at a strain rate of 2 © 103 s¹1 andtemperatures of 25°C, 0°C and ¹150°C, respectively. Notethat the loading direction (LD) is perpendicular to the imageplane. It is seen that the microstructure of the specimen

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(b)

(c)

Fig. 6 (a) Variation of flow stress with deformation temperature as functionof strain and strain rate. Variation of temperature sensitivity with truestrain as function of strain rate over temperature ranges of: (b) ¹150 to0°C; and (c) 0 to 25°C.

(a)

(b)

(c)

Fig. 7 Optical micrographs of Inconel 690 specimens deformed at strainrate of 6 © 103 s¹1 and temperatures of: (a) 25°C; (b) 0°C; and (c)¹150°C. The loading direction is perpendicular to the image plane.

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deformed at 25°C consists predominantly of cell-likestructures with a high dislocation density in the cell walland a somewhat reduced dislocation density in the cellinterior (Fig. 10(a)). It has been reported that dislocation cellstructures are readily formed in materials having a stackingfault energy greater than 60mJ/m2.35) Inconel 690 alloy has astacking fault energy of 61mJ/m2,40) and thus the existenceof cell structures in the deformed microstructures is notunexpected. An observation of Fig. 10(b) shows that whenthe deformation temperature is reduced to 0°C, thedislocation density increases. Furthermore, the dislocationcells are seen to have an elongated structure and an increaseddegree of dislocation tangling in the walls. At the lowestdeformation temperature of ¹150°C, both the number of cellstructures and the degree of dislocation tangling increase(Fig. 10(c)). Furthermore, the size of the individual dis-location cells decreases. It is inferred that the smaller, denserand more tangled structures at this cryogenic temperaturelead to a greater work hardening and strengthening effect.Figures 10(d)³10(f ) present TEM images of the Inconel 690specimens tested at the highest strain rate of 6 © 103 s¹1. It isnoted that while the temperature-dependent morphologies ofthe deformed microstructures are similar to those shown inFigs. 10(a)³10(c), the dislocation density and cell size aredifferent. Specifically, the dislocation density at a strain

rate of 6 © 103 s¹1 is higher than that at a strain rate of2 © 103 s¹1. However, the dislocation cell size at 6 © 103 s¹1

is smaller than that at 2 © 103 s¹1. The increased dislocationdensity prompts the formation of a tangled dislocationsubstructure, which reduces the mobility of the dislocationsand therefore enhances the resistance of the Inconel 690 alloyto plastic deformation. Thus, as shown in Fig. 3, for aconstant true strain, the flow stress increases with increasingstrain rate or decreasing temperature.

An observation of Figs. 10(a)³10(f ) shows that thedislocations in the impacted microstructure become tangledunder higher strain rates and lower temperatures. As aconsequence, the resistance encountered by the moving

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Fig. 9 SEM fractographs of Inconel 690 specimens deformed at strain rateof 6 © 103 s¹1 and temperatures of: (a) 25°C; (b) 0°C; and (c) ¹150°C.

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Fig. 8 Optical micrographs showing cross-section parallel to loadingdirection in Inconel 690 specimens deformed at strain rate of6 © 103 s¹1 and temperatures of: (a) 25°C; and (b) ¹150°C.

Mechanical Properties and Dislocation Substructure of Inconel 690 Alloy Impacted at Cryogenic Temperatures 1695

dislocations increases; leading to an enhanced work harden-ing effect. Figure 11 shows the variation of the square root ofthe dislocation density with the strain rate as a function ofthe temperature given a true strain of 0.5. (Note that the

corresponding variation in the flow stress is also shown.)It is seen that the square root of the dislocation densityincreases with increasing strain rate, but decreases withincreasing temperature. Moreover, the rate of increase of the

(d)

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(f)

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Fig. 10 TEM micrographs showing dislocation microstructures of Inconel 690 specimens deformed at: (a) 2 © 103 s¹1, 25°C;(b) 2 © 103 s¹1, 0°C; (c) 2 © 103 s¹1, ¹150°C; (d) 6 © 103 s¹1, 25°C; (e) 6 © 103 s¹1, 0°C; and (f ) 6 © 103 s¹1, ¹150°C. The loadingdirection is perpendicular to the image plane.

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dislocation density is more evident at ¹150°C than at 0°C or25°C. In other words, the work hardening rate of the Inconel690 specimens increases as the deformation temperaturereduces.

4. Conclusions

The impact deformation behaviour and dislocation sub-structure of Inconel 690 alloy have been examined at strainrates of 2 © 103 to 6 © 103 s¹1 and temperatures of ¹150°C,0°C and 25°C using a split-Hopkinson pressure bar system.The results have shown that the flow stress increases withincreasing strain rate, but decreases with increasing temper-ature. Moreover, the strain rate sensitivity and workhardening rate increase with increasing strain rate anddecreasing temperature. The thermal activation volumeincreases with increasing temperature and decreasing strainrate. Meanwhile, the temperature sensitivity increases withincreasing strain rate and increasing temperature. The OMand SEM observations have shown that at the highest strainrate of 6 © 103 s¹1, the Inconel 690 specimens facture ateach of the considered deformation temperatures as a resultof adiabatic shear localisation and cracking. In addition, ithas been shown that for each considered strain rate, theInconel 690 specimens become increasingly brittle as thedeformation temperature is reduced. The dislocation densityincreases markedly with increasing strain rate and decreasingtemperature, and results in a smaller dislocation cell size.Finally, the enhancement of the flow stress and workhardening response in the present Inconel 690 specimens isdirectly related to a change in the dislocation density andcell size.

Acknowledgements

The authors gratefully acknowledge the financial supportprovided to this study by the National Science Council of theRepublic of China under Grant No. NSC101-2221-E-006-144.

REFERENCES

1) J. Kai, G. Yu, C. Tsai, M. Liu and S. Yao: Metall. Mater. Trans. A 20(1989) 2057­2067.

2) C. Leroy, T. Czerwiec, C. Gabet, T. Belmonte and H. Michel: Surf.Coat. Tech. 142­144 (2001) 214­247.

3) A. Brighenti: IEEE J. Oceanic Eng. 15 (1990) 179­188.4) J. B. Calvo and K. Hannemann: Numer. Experim. Fluid Mech. VII

NNFM 112 (2010) 441­448.5) V. Venkatesh and H. J. Rack: Int. J. Fatigue 21 (1999) 225­234.6) J. M. Boursier, D. Desjardins and F. Vaillant: Corrosion Sci. 37 (1995)

493­508.7) M. K. Lim, S. D. Oh and Y. Z. Lee: Nucl. Eng. Des. 226 (2003) 97­105.8) B. Wang, S. H. Zhang, M. Cheng and H. W. Song: J. Mater. Eng.

Performance (2013) DOI:10.1007/s11665-013-0520-4.9) S. Guo, D. Li, H. Pen, Q. Guo and J. Hu: J Nucl. Mater. 410 (2011) 52­

58.10) W. S. Lee, C. Y. Liu and T. N. Sun: Int. J. Impact Eng. 32 (2005) 210­

223.11) J. E. Field, S. M. Walley, N. K. Bourne and J. M. Huntley: J. Phys. IV,

Colloque C8, supplément au Journal de Physique III, 4 (1994) C8-3­C8-22.

12) A. Mishra, M. Martin, N. N. Thadhani, B. K. Kad, E. A. Kenik andM. A. Meyers: Acta Mater. 56 (2008) 2770­2783.

13) W. S. Lee, C. F. Lin, T. H. Chen and W. Z. Luo: J. Nucl. Mater. 420(2012) 226­234.

14) M. A. Meyers, Y. J. Chen, F. D. S. Marguis and D. S. Kim: Metall.Mater. Trans. A 26 (1995) 2493­2501.

15) W. S. Lee, C. F. Lin, T. H. Chen, W. Z. Luo and M. C. Yang: Mater.Sci. Eng. A 527 (2010) 3127­3137.

16) P. S. Follansbee and U. F. Kocks: Acta Metall. 36 (1988) 81­93.17) F. J. Zerilli and R. W. Armstrong: Acta Metall. Meter. 40 (1992) 1803­

1808.18) A. G. Odeshi, S. Al-ameeri, S. Mirfakhraei, F. Yazdani and M. N.

Bassim: Theor. Appl. Fract. Mech. 45 (2006) 18­24.19) W. S. Lee, C. F. Lin, T. H. Chen and W. Z. Luo: J. Nucl. Mater. 420

(2012) 226­234.20) V. F. Nesterenko, M. A. Meyers and T. W. Wright: Acta Mater. 46

(1998) 327­340.21) D. Peirce, R. J. Asaro and Needieman: Acta Metall. 31 (1983) 1951­

1976.22) L. Anand and S. R. Kalinidi: Mech. Mater. 17 (1994) 223­243.23) T. W. Wright: J. Mech. Phys. Solids 38 (1990) 515­530.24) J. A. DiLellio and W. E. Olmstead: Mech. Mater. 29 (1998) 71­80.25) F. Zhou, T. W. Wright and K. T. Ramesh: J. Mech. Phys. Solids 54

(2006) 904­926.26) D. R. Chichili, K. T. Ramesh and K. J. Hemker: Acta Mater. 46 (1998)

1025­1043.27) H. Jarmakani, J. M. McNaney, B. Kad, D. Orlikowski, J. H. Nguyeu

and M. A. Meyers: Mater. Sci. Eng. A 463 (2007) 249­262.28) W. S. Lee and C. F. Lin: Metall. Mater. Trans. A 33 (2002) 2801­2810.29) B. Zhang and V. P. W. Shim: Acta Mater. 58 (2010) 6810­6827.30) M. Huang, P. E. J. Rivera-Diaz-del-Castillo, O. Bouaziz and S. van der

Zwaag: Mech. Mater. 41 (2009) 982­988.31) B. L. Hansen, I. J. Beyerlein, C. A. Bronkhorst, E. K. Cerreta and

D. D. Koller: Int. J. Plasticity 44 (2013) 129­146.32) W. S. Lee, T. H. Chen and S. C. Huang: J. Nucl. Mater. 402 (2010) 1­7.33) W. S. Lee, T. H. Chen and H. H. Hwang: Metall. Mater. Trans. A 39

(2008) 1435­1448.34) W. S. Lee, T. H. Chen, C. F. Lin and W. Z. Luo: Metall. Mater. Trans. A

43 (2012) 3998­4005.35) L. E. Murr: International Conference on Metallurgical Effects of High

StrainRate Deformation and Fabrication, ed. by M. A. Meyers and L. E.Murr, (Plenum Press, New York, 1981) pp. 607­613.

36) Y. Tomota, P. Lukas, S. Harjo, J. H. Park, N. Tsuchida and D. Neov:Acta Mater. 51 (2003) 819­830.

37) H. Kolsky: Proc. Phys. Soc. B 62 (1949) 676­700.38) H. Conrad and H. Wiedersich: Acta Metall. 8 (1960) 128­130.39) L. Shi and D. O. Northwood: Acta Metall. Mater. 43 (1995) 453­460.40) J. Unfried-Silgado, L. Wu, F. F. Ferreira, C. M. Garzon and A. J.

Ramirez: Mater. Sci. Eng. A 558 (2012) 70­75.

Fig. 11 Variation of square root of dislocation density and true stress withstrain rate as function of temperature.

Mechanical Properties and Dislocation Substructure of Inconel 690 Alloy Impacted at Cryogenic Temperatures 1697