Microstructure changes in a 2618 aluminium alloy during ageing and creep

Frantisek Novya, Milos Janecekb, Robert Králb, Branislav Hadzimaa

Microstructure evolution in a 2618 aluminiumalloy during creep-fatigue tests

Microstructure changes in the 2 618 aluminium alloy duringcreep-fatigue tests were studied. These tests simulate theconditions of the application of this alloy in devices for theexhaustion of hot gasses generated during the fire in closedor difficultly accessible areas. Creep-fatigue tests result inhigh dislocation density in subgrains and narrow subgrainboundaries, on contrary to creep tests reported in our pre-vious work where large subgrains were observed with rela-tively wide subgrain boundaries and relatively low disloca-tion density in grains. Extensive precipitation occurred withdenuded (precipitate-free) zones along grain boundaries.The coherent S-phase (Al2CuMg) transformed into partiallycoherent needle-shaped S’ precipitates. Superior stress am-plitude caused reduced lifetime and wider denuded zones.A model of the formation of denuded zones along (sub)-grain boundaries was proposed.

Keywords: 2 618 aluminium alloy; Creep; Fatigue; Micro-structure evolution

1. Introduction

In general, creep-fatigue tests can be classified based on thenumber of cycles to fracture asi) low cycle creep-fatigue tests (or high temperature low

cycle fatigue),ii) high cycle creep-fatigue (VHCF) tests andiii) very high cycle creep-fatigue tests.In the published literature, the work is mostly focused onlow-cycle tests performed at low frequencies (e. g. 1 cycleper month/week/day/hour). High amplitude of stress, lowcycling frequencies and small number of cycles to fractureare typical for these tests. The high interest in these typesof tests is driven by the necessity to solve the problem oflifetime prediction of energy devices. Therefore, the re-search is mostly focused on exact simulation of the operat-ing conditions of the construction materials used in nuclearreactors, steam generators, steam lines and other facilities.

Research dealing with creep and high cycle fatigue inter-action is more limited. Authors mostly focus on research ofspecial materials, e. g. nickel-based superalloys for high-temperature applications in aerospace industry, see e. g.[1]. In this type of testing lower stress amplitudes and highloading frequencies resulting in high number of cycles tofracture are generally employed. The loading frequencyusually ranges from 1 to 200 Hz and depends on the type

of testing machine. The frequencies of 20, 100 and 200 Hzare mostly used.

Finally, the smallest group represents the works focusedon the investigation of cyclic creep in conditions of veryhigh cycle fatigue. These tests are characterized by the ap-plication of low-amplitude and high-frequency fatigue vi-brations during creep testing [2]. Since the frequency andthe cycle type again depend on the type of testing machine,the tests are performed using a 20 kHz sine wave. In the lastdecade, VHCF testing has commonly been performed un-der these conditions [3].

Creep resistant aluminium alloys are often used for theproduction of fire exhausters. The most promising candi-dates for this purpose appear to be aluminium alloys of the2xxx and 6xxx series. This investigation was challengedby the application of the 2 618 alloy as the material forsafety fire exhausters which are employed to dischargegases generated during the fire from traffic tunnels. Twotypes of safety fire exhausters are manufactured from the2 618 alloy:a) jet exhausters located directly in the tunnel, andb) exhausters/discharge fans and exhaust flaps through

which the smoke is discharged indirectly through sideducts from the tunnel.

The material used for safety fire exhausters should meet thefollowing function requirements: thermal stability up to250 8C and correct operation during the duration of fire.The first group of exhausters should operate at least120 min while for the second one the minimum operationtime is extended to 12 h.

The 2 618 aluminium alloy contains Ni and Fe in ba-lanced amounts (1 :1). It was developed originally to raisethe temperature limit of conventional alloys of the 2xxx ser-ies (Al – Cu – Mg). It is the stability of the S’-phase that en-hances the creep resistance of the Al –Cu –Mg alloys. Theformation of the S’-phase is influenced by the Cu : Mg ratio.If this value exceeds 4, the H (Al2Cu) phase with higherthermal instability occurs instead of the S’-phase [4, 5].Due to the chemical composition of 2 618 alloy, the precipi-tation of other phases, namely H’’, H’, and H, is not possible[6, 7]. The typical sequence of precipitation process in Al –Cu – Mg alloys consists of the formation of Guinier – Pre-ston –Bagaryatsky (GPB) zones, which are coherent withthe matrix. After long-term ageing at elevated temperaturesGPB zones transform to the transition phase S’ (Al2CuMg),which is partially coherent with the matrix (only on {021}matrix planes), whereas overaging is associated with theformation of the equilibrium S phase and the loss of coher-ency. The precipitation structure sequence may be repre-

IJMR_MK110679 – 28.12.11/stm media köthen

F. Novy et al.: Microstructure evolution in a 2 618 aluminium alloy during creep-fatigue tests

Int. J. Mat. Res. (formerly Z. Metallkd.) 1

a University of Zilina, Department of Materials Engineering, Zilina, Slovak Republicb Charles University, Department of Physics of Materials, Prague, Czech Republic

sented as follows [8]:

SS ? GPB I ? GPB II (S’’) ? S’ ? S

The S’-phase precipitates in the form of tiny platelets withthe following orientation:

[100]S’ // [100]Al; [010] S’ // [012] Al; [001] S’ // [021]Al

The plastic deformation during the creep test is accompa-nied by a high activity of dislocation sources. The newlygenerated dislocations move and interact with each other,some of them annihilate and others arrange themselves ina dislocation network which is gradually transformed intocells and subgrains. Regular dislocation networks with sub-grains are often formed in many alloys during straining athigher homologous temperatures. The driving force of thesubgrain formation is the decrease of the total deformationenergy due to favourable dislocation reactions and due toscreening of elastic stress fields of dislocations in plane ar-rangements.

The objective of this study is to investigate the micro-structure evolution during creep-fatigue tests in the2 618 Al alloy.

2. Experimental

The material used in this investigation is a conventionallyproduced creep resistant aluminium alloy 2618. The nominalcomposition of the alloy (in wt.%) is Al-2.34Cu-1.49Mg-1.27Fe-1.25Ni-0.28Si-0.07Ti-0.05Mn-0.01Zn-0.03Cr. Thematerial was supplied in the form of a forged piece in a T6(peak-aged) condition. It means that material was solution-treated at 530 8C for 8 h, quenched in water and aged at atemperature of 190 8C for 10 h. Unlike many other Al al-loys, which are used also in an underaged and overagedcondition [9, 10], the 2 618 alloy is used only in a peak-agedcondition.

Based on the above mentioned function and design re-quirements the alloy was subjected to creep-fatigue tests.All tests were performed at a temperature of T = 270 8C.Combined creep fatigue tests were motivated by the neces-sity to simulate the influence of vibrations on the creep life-time of the alloy. Two levels of the mean stress rmean = 110and 135 MPa were employed and superimposed by the cyc-lic loading with a stress amplitude of ra = 30 MPa andra = 5 MPa, respectively. The frequency of cyclic loadingwas 20 kHz. The frequency of 20 kHz was used intention-ally for the simulation of vibrations. As stated in the \Intro-duction" section, this research is focused on behavior of the2 618 alloy in the conditions of short-term creep-fatigueloading, where the lifetime requirements are given by aminimum limit of 12 h. During this minimum time, thefire-exhausters from this alloy must operate reliably evenin the case of small mechanical damage caused to the rotat-

ing parts by flying fragments. As a consequence of this dam-age, vibrations in the frequency range from 1 to 21 kHz aregenerated in the device.

Creep tests were executed at the constant load with anominal stress of r = 140 MPa. The results of creep testsare reported in detail in [11]. For clarity, test conditionsare summarized in Table 1. The main objective of the workwas to investigate the microstructure changes and develop-ment of the dislocation substructure during creep-fatiguetests and to compare them with microstructure changes dur-ing pure creep tests.

3. Results and discussion

Two creep-fatigue curves are presented in Fig. 1 and com-pared to a creep curve at constant load of 140 MPa. The ef-fect of superimposed cyclic component on the shape ofcreep curves is clearly seen. In our previous work [12] wehave determined the minimum rate of stationary creep as afunction of the amplitude of cyclic stress component atidentical conditions for the same alloy.

The test conditions used in our experiments highly ex-ceed the temperature of the thermodynamic stability of S’’-phase and correspond to the border region between the S’-and S-phase of TTT diagram for ageing of the Al-Cu-Mgalloys. In our previous work [12] we investigated micro-structure changes of 2 618 alloy after creep tests and ob-served a typical arrangement of S’-phase precipitates in theform of three systems mutually turned by 1208, see alsoFig. 2. Due to extensive precipitation newly formed preci-pitates could even break smaller stable primary phases.The transformation of S’-precipitates into stable S-phasewas observed in grain interiors only locally and toa limited extent. Intensive formation of S-phase occurred,however, in grain boundaries. Denuded precipitation-freezones were formed along grain boundaries. These zones



2 Int. J. Mat. Res. (formerly Z. Metallkd.)

Table 1. Experimental conditions.

Test condition Temperature(8C)

Stress amplitudera (MPa)

Mean stressrmean (MPa)

Maximum stressrmax (MPa)

Test duration(h)

creep 270 – – 140 24creep-fatigue 270 30 110 140 7creep-fatigue 270 5 135 140 21

Fig. 1. Creep curves of the 2618 alloy.

negatively influence many mechanical properties of thealloy, e. g. the strength, the ductility, the creep, and fatigueresistance [13, 14].

The typical microstructure formed during creep test con-sists of large subgrains with relatively wide subgrainboundaries, see Fig. 2. These subgrain boundaries are onlyweak obstacles for the movement of glide dislocationswhich can easily pass them and continue their movementin the grain for relatively long distances as the dislocationdensity stored in the specimen during creep is relativelylow (cf. Fig. 2 and Fig. 3).

During the creep-fatigue test the coherent precipitatesalso dissolve and intensive precipitation of S’-phase and lo-cally also of S-phase occurs (see Fig. 4 and Fig. 5).

The loss of coherency, i. e. the formation of interfaceboundaries at precipitates significantly reduce the coherent

internal stresses and consequently drop of both hardnessand strength is observed similar as during the creep test.The hardness of the initial state (T6) was 133 HB. Itdropped to 93 HB after the creep test [12]. The hardness of101 and 98 HB was found in the specimen after 7 and 21 hof creep-fatigue loading, respectively.

The comparison of microstructure of the alloy after thecreep test [12] with that after the creep-fatigue test (Fig. 4and Fig. 5) does not reveal any noticeable differences bothin the fraction of precipitated phases and in the mutual dis-tance. Moreover, the value of the amplitude of the cyclicloading does not influence the intensity of the precipitationof incoherent precipitates of the S-phase. No noticeable dif-ferences of the concentration of S’ precipitates in the speci-men after the creep and creep-fatigue tests were observed.Similar as during the creep test, intensive precipitation ofincoherent precipitates of the equilibrium S-phase occurs




Fig. 2. Subgrain boundary (marked with an arrow) and S’ precipitatesin the crept specimen.

Fig. 3. Grain interior with S’-phase precipitates and no dislocations inthe crept specimen.

Fig. 4. S’- and S-phase precipitates (g = (41411), B = [011]).

Fig. 5. S’- and S-phase precipitates (dark field, g = (200), B = [101]).

in grain boundaries and denuded precipitation free zonesare observed in the vicinity of grain boundaries (see Fig. 6and Fig. 7). Denuded zones were formed by the same mech-anism as in the crept specimen due to the lack of vacanciesin these regions. The width of denuded zones dependsstrongly on the position of the specimen extracted from theexhauster wheel as the vacancy concentration dependsstrongly on the cooling rate of the wheel. However, due toeasy annihilation of vacancies in regions along the grainboundaries the strong concentration gradient exists there.

Due to similar temperature and stress conditions duringthe creep-fatigue tests the principal mechanism of plasticdeformation is also dislocation creep. Subgrain formationis more intensive in specimens after the creep-fatigue testas compared to crept specimens. The density of dislocationstrapped at grain boundaries also increases, see Fig. 8. The

subgrain size in specimens after creep-fatigue tests is muchsmaller and the width of subgrain boundaries decreases (cf.Fig. 2 and Fig. 9). Moreover, TEM confirmed high disloca-tion density in the grain interior, see Fig. 10. Diffraction pat-tern analysis revealed also higher subgrain misorientation.As a consequence, obstacles to dislocation movement aremuch stronger in creep-fatigue loaded specimens. There-fore, dislocations cannot pass the grain as easily as duringcreep test. In most cases the movement of dislocations is re-stricted to a single grain only. No dislocation loops testify-ing an Orowan mechanism of obstacle overcoming are ob-served. The controlling mechanisms of plastic flow aremost probably the overcoming of obstacles by cross slipand/or cutting of precipitates by dislocations. The residualdislocation loops around precipitates probably annihilated.An example of a cut S’-precipitate is shown in Fig. 11.




Fig. 6. S’-phase precipitates in the grain interior, S-phase precipitaesin grain boundaries (marked with arrows).

Fig. 7. S’-phase precipitates in the grain interior, S-phase precipitaesin grain boundaries (marked with arrows).

Fig. 8. Dislocations in subgrain boundaries – marked with an arrow(dark field, g = (200), B = [101]).

Fig. 9. Dislocations in a subgrain boundary (marked with an arrow).

The permeability of subgrains formed by arrays of dislo-cations of the same sign depends strongly on the amplitudeof cyclic loading. At low stress amplitudes (ra = 5 MPa),the character of the subgrain structure differs only slightlyfrom that after the pure creep test. The only difference is inthe subgrain size and in the width of subgrain boundarieswhile the misorientation angle differs only slightly. Withincreasing stress amplitude the dislocation activity in-creases which manifests itself in the increase of subgrainboundary misorientation and in the reduction of their width.At ra = 30 MPa the value of misorientation reaches almostthat of a high-angle grain boundary (158 according to theBrandon criterion [15, 16]). The individual extrinsic grain-boundary dislocations with typical parallel contrast are lo-cally observed in subgrain boundaries. Moreover, the den-sity of dislocations trapped at subgrain boundaries increaseswith increasing stress amplitude ra.

Under suitable conditions denuded (precipitate-free)zones (PFZ) are preferably formed along microstructure de-fects, in particular grain and subgrain boundaries and/orcoarse intermetallic particles [17]. Denuded zones alongsubgrain boundaries have different character and areformed by different mechanisms than those along grainboundaries. These differences may be attributed to effectsof cyclic loading as subgrain boundaries were formed dur-ing creep-fatigue test only, while other microstructure de-fects were present in the material already before the test.One may assume that dislocations trapped between pinnedpoints in a subgrain boundary are forced to vibrate due tocyclic loading. During the vibration they cut precipitates inthe vicinity of the subgrain boundary. Small parts of preci-pitates broken by this mechanism become thermodynami-cally unstable and start to dissolve in the surrounding solidsolution. Moreover, high-frequency cyclic loading acceler-ates diffusion processes which enhance vacancy annihila-tion and the diffusion of alloying elements from dissolvedprecipitates towards the subgrain boundary. As a conse-quence, denuded S’-PFZs along the subgrain boundary areformed in the creep-fatigue loaded specimen (Fig. 12).Their width was found to increase with increasing stressamplitude.

4. Conclusions and comparison with creep tests

Microstructure changes in the 2 618 aluminium alloy duringcreep-fatigue tests were studied. The following conclusionsmay be drawn from this investigation:

The microstructure after creep-fatigue test is character-ized by:. high dislocation density in subgrains,. narrow subgrain boundaries,. PFZ along subgrain boundaries,. superimposed stress amplitude caused reduced lifetime

and wider PFZ.A model of the formation of denuded zones along subgrainand grain boundary is proposed.




Fig. 10. Dislocations pinned at S’-phase precipitates (g = (200),B = [101]).

Fig. 11. S’-phase precipitates – marked with an arrow (g = (200),B = [001]).

Fig. 12. Precipitate-free zone along a subgrain-boundary (g = (200)B = [101]).

On the other hand, creep test results in the microstructurewith [12]:. relatively low dislocation density in grains,. large subgrains with relatively wide subgrain bound-

aries,. narrow PFZ along grain boundaries.

The authors acknowledge funding through the research program MSM0021620834 of Ministry of Education of the CR and research programVEGA 1/0262/10 and supporting by the Slovak Research and Develop-ment Agency under contract APVT-20012804.

References

[1] J. Zrnik, J.A. Wang, Y. Yu, L. Pejing, P. Hornak: Mat. Sci. Eng.A 234–236 (1997) 884. DOI:10.1016/S0921-5093(97)00399-7

[2] K.D. Sheffler: Metall. Mater. Trans. B 3 (1972) 167.DOI:10.1007/BF02680596

[3] C. Bathias, P.C. Paris, in: M. Decker (Ed.), Gigacycle fatigue inmechanical practice, New York, 2005.

[4] L.F. Mondolfo: Aluminium Alloys – Structure and Properties,Butterworths, London, 1976.

[5] J.W. Martin: Precipitation hardening, Pergamon Press, Oxford(1968).

[6] C. Badini, F. Marino, E. Verne: Mater. Sci. Eng. A 191 (1995)185. DOI:10.1016/0921-5093(94)09637-C

[7] F. Bardi, M. Cabibbo, S. Spiragelli: Mater. Sci. Eng. A 334 (2002)87. DOI:10.1016/S0921-5093(01)01775-0

[8] Aluminium: Properties and Physical Metallurgy, J.E. Hatch (Ed.),American Society for Metals, Metals Park, OH 2005.

[9] V. Skorik, B. Hadzima: Mater. Eng. 15 (2008) 23.[10] J. Bidulska, T. Kvackaj, R. Bidulsky, M. Actis Grande: Kovove

Mater. 46 (2008) 339.[11] F. Novy, M. Janecek, R. Kral: J. Alloy Compd. 487 (2009) 146.

DOI:10.1016/j.jallcom.2009.08.014[12] F. Novy, P. Palcek, M. Chalupova: Kovove Mater. 43 (2005) 447.[13] I.J. Polmear: Light Alloys – From Traditional Alloys to Nanocrys-

tals, Fourth edition, Elsevier Butterworth-Heinemann, Oxford,2006.

[14] J. Svoboda, G. Maier: Int. J. Mater. Res. 101 (2010) 997.DOI:10.3139/146.110366

[15] D.G. Brandon: Acta Metall. 14 (1966) 1479.DOI:10.1016/0001-6160(66)90168-4

[16] M. Dopita, M. Janecek, D. Rafaja, J. Uhlir, Z. Matej, R. Kuzel:Int. J. Mater. Res. 100 (2009) 785. DOI:10.3139/146.110111

[17] P.N. Unwin, G.W. Lorimer, R.B. Nicholson: Acta Metall. 17(1969) 1363. DOI:10.1016/0001-6160(69)90154-0

(Received November 28, 2010; accepted November 15, 2011)

Bibliography

DOI 10.3139/146.110679Int. J. Mat. Res. (formerly Z. Metallkd.)page 1–6# Carl Hanser Verlag GmbH & Co. KGISSN 1862-5282

Correspondence address

Dr. Milos JanecekCharles University, Department of Physics of MaterialsKe Karlovu 5, CZ-1216 Prague 2, Czech RepublicTel.: +420221 911359Fax: +4202219 11490E-mail: [email protected]

You will find the article and additional material by enter-ing the document number MK110679 on our website atwww.ijmr.de




Microstructure changes in a 2618 aluminium alloy during ageing and creep

Documents

Transcript of Microstructure changes in a 2618 aluminium alloy during ageing and creep