84 INŻYNIERIA MATERIAŁOWA MATERIALS ENGINEERING ROK XXXIX

Influence of microstructure on fatigue fracture of Ti–2.5Cu alloy

Maria Hepner*, Dariusz RozumekFaculty of Mechanical Engineering, Opole University of Technology; *marepner@onet.pl

The paper presents the test results obtained for fatigue crack growth in Ti–2.5Cu alloy subjected to bending in notched specimens. The tested specimens were subjected to various variants of heat treatment. The tests were performed at the fatigue test stand MZGS-100 under loading frequency 28.4 Hz. The study was conducted for a constant amplitude of moment Ma = 11.2 N∙m and different values of stress ratio R = –1 and 0. Influence of the microstructure on the crack paths in plane specimens was observed. The propagation of the main crack in the two-phase structure takes place both transcrystallinely through α phase grains and through the boundaries of the grains in metastable phase precipitation areas, whereas the propagation of the side cracks takes place along the precipitates of the coniferous phase. In the material after supersaturation and ageing at a temperature of 415°C, after Nf = 12 000 cycles (R = 0), no side cracks were observed, and the main crack develops in both phases that form the microstructure shown in Figure 10. A different fracture mechanism occurs for the material after supersatu-ration and ageing (Nf = 11 000 cycles) at a temperature of 760°C. In this case, an irregular path of the main crack is observed, whose direction in the micro-areas is determined by the precipitates of the intermetallic phase (Fig. 11). From the beginning of the main crack and along its entire length, numerous side cracks develop that are over 150 μm in length. The components of the structure that facilitate the propagation of cracks and determine their direction are clusters of Ti2Cu precipitates on the boundaries of α phase. In case of this alloy, the impact of the mean stress value during cyclic testing decreases fatigue life considerably.

Key words: titanium alloy, microstructure, heat treatment, fatigue fracture.

Inżynieria Materiałowa 2 (222) (2018) 84÷89DOI 10.15199/28.2018.2.6

MATERIALS ENGINEERING

1. INTRODUCTION

Ti–2.5Cu is the most common alloy of titanium and copper. It is characterised by very high plasticity, and is used to manufacture of products complex shape using cold and hot plastic working. Impor-tant properties that influence its application are also good weldabil-ity and possibility to work at temperature up to 350°C. It has pri-marily found applications in chemical apparatus and parts of auto-mobile and aircraft structures (e.g. in jet airliner Concorde, combat aircraft Tornado and Jaguar) [1÷5]. As it can be seen in the Ti–Cu phase diagram, titanium and copper may form solid solutions and a number of intermetallics: Ti2Cu, TiCu, Ti2Cu3, TiCu3. The maxi-mum solubility of copper in a crystalline lattice of α titanium at the eutectic temperature (798°C) is 2.1 wt % , whereas at the peritectic temperature 17 wt % [6]. The structure of the alloy at equilibrium (after a very slow cooling from 810°C) is formed by α phase grains and a fine lamellar eutectoid mixture (α + Ti2Cu), whose volume fraction in the structure is approx. 10% (Fig. 1).

On the other hand, a metastable martensitic phase may be formed in the structure during a fast cooling from the range of β phase. Due to rather high solubility of copper in a solid solution at the eutectic temperature and its considerable reduction along with the decrease of temperature, alloys of titanium and copper may be strengthened by precipitation hardening. It involves the dissolving of Ti2Cu inter-metallic at a temperature above the eutectic temperature and then its solution heat treatment and ageing. In this respect, titanium–copper alloys are an exception among titanium alloys because the solubility of intermetallics in α-Ti in case of other alloying elements is very low [7]. Recommended heat treatment for this alloy (that gives the best combination of strength and plasticity properties) is supersatu-ration in air from a temperature of slightly above 800°C and then its double ageing; first at a temperature of 400°C for 24 h and further at a temperature of 475°C for 8 h. As a result, tensile strength of the alloy Rm in relation to the annealed condition increases from around 500÷600 MPa to approx. 750 MPa with a slight reduction of A5 from around 25% to around 20% [8].

The aim of this study was to examine the fatigue life of Ti–2.5Cu (IMI230) alloy after various types of heat treatment and to follow the differences in the propagation mechanisms of fatigue cracks that are conditioned by different microstructure.

2. EXPERIMENTAL METHODS

Plane specimens with a stress concentrator (one-sided sharp notch) were tested, and next tests of the fatigue crack growth in the mate-rial subjected to various forms of heat treatment and in the mate-rial in the initial state were performed. Three variants of heat treat-ment were applied: 1) supersaturation in water at the temperature of 810°C, 2) supersaturation in water at the temperature of 810°C and aging at 415°C for 20 h, 3) supersaturation in water at the tempera-ture of 810°C and aging at 760°C for 6 h. In this last annealing vari-ant, the temperature was slightly lower than the temperature of eu-tectoid transformation for maximum softening of the material. Ta-ble 1 contains a chemical composition of the tested material brand name IMI230. Mechanical properties of the tested alloy Ti–2.5Cu (as delivered) are shown in Table 2. The structure of the material in the delivery state (annealed condition) is shown in Figure 2. The specimens were cut from a bar Ø16 mm. Each specimen had an

Fig. 1. Structure of Ti–2.5Cu alloy after cooling in a furnace from

a temperature of 810°CRys. 1. Mikrostruktura stopu Ti–2,5Cu po chłodzeniu z piecem z tempe-ratury 810°C

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external unilateral notch with depth 4 mm and radius ρ = 0.2 mm (Fig. 3). The notches in the specimens were cut with a milling cutter and their surfaces were polished after grinding.

The theoretical stress concentration factor in the specimen under bending Kt = 4.34 was estimated with use of the model presented in [9]. Unilaterally restrained specimens were subjected to cyclic bending with a constant amplitude of moment Ma = 11.2 N∙m, which corresponded to the nominal stress amplitude σa = 74.6 MPa before the crack initiation. In the case of mean load, the load was achieved by additional spring acting on the lever of MZGS-100 machine and then the moment Ma = 11.2 N∙m (Mmax = Ma + Mm = 22.4 N∙m) was obtained. Sinusoidally variable loading was applied. The bending tests were performed under loading frequency of 28.4 Hz and load ratio R = –1 and 0.

Crack growth was observed on the specimen surface with a mi-crometer mounted on a portable microscope with magnification of 25 times. The measurements were performed with an accuracy up to 0.01 mm with numbers of loading cycles N recorded. The tests were performed on the fatigue test stand MZGS-100, which enables cyclic bending, torsion and bending with torsion as well as static

(mean) loading [10, 11]. Etched specimens were studied metallo-graphically with the use of the OLYMPUS IX70 light microscope. Microhardness testing was performed using the LECO MHT210 hardness tester coupled with the PC-2 computer. Applied load was 0.981 N and the time during which the pressure of the diamond in-denter was applied was 17 s.

3. RESULTS AND DISCUSSION

The microstructure of the alloy after supersaturation in water from a temperature of 810°C is comprised of regular equiaxial α phase grains (up to 17 μm in diameter) and precipitates of extremely fine coniferous, martensitic and metastable αʹ phase (approx. 8 μm in diameter) on α phase boundaries, which were formed during a fast β phase transition. The volume fraction of the phases is 76% for α phase and 24% for αʹ phase respectively (Fig. 4).

During the annealing of the supersaturated alloy, two different processes overlap; the precipitation of fine hardening phases in the ageing of supersaturated α phase and the decomposition of the metastable phase. After supersaturation and ageing, the microstruc-ture of the alloy at a temperature of 415°C is also two-phase. The structure is formed by regular equiaxial α phase grains and precipi-tation of an extremely fine eutectoid mixture (α + Ti2Cu) on α phase boundaries in coniferous areas (Fig. 5).

After supersaturation and ageing at a temperature of 760°C, the microstructure of the alloy is substantially different. On the bound-aries of α phase grains there are big (reaching up to 3.5 μm) coagu-lated precipitates of Ti2Cu phase in a chain system (Fig. 6). Table 3 below presents hardness for the alloy in the initial state (delivery condition), after supersaturation and after supersaturation combined with two types of ageing, while Fig. 4÷6 presents the microstructure of the material after heat treatment. A small difference between the hardness of the material after supersaturation and the material after ageing at a temperature of 415°C (specimens 2 and 3, Table 3) re-sults from the fact that while the process of precipitation hardening takes place during ageing, the coniferous structure of the martens-itic areas formed in transformed β phase disappears.

The propagation of fatigue cracks for all specimens tested occurs in a plane perpendicular to the direction of highest normal stress. For the initial material (in delivery condition), a very regular main crack develops both transcrystallinely through α phase grains and through the boundaries of the grains. Numerous and very short sec-ondary cracks that develop from the main crack in the directions of

Table 1. Chemical composition of the tested Ti–2.5Cu alloy, wt %Tabela 1. Skład chemiczny badanego stopu Ti–2,5Cu, % mas.

Cu Fe C O N Ti

2.7 0.04

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of the alloy. Figure 7 shows the microstructure of Ti–2.5Cu alloy in the initial condition with a visible fatigue crack for Ma = 11.2 N∙m (R = 0) after Nf = 259 000 cycles [12], and Figure 8 shows twin boundaries which occurred during a bend test in the Nomarski’s contrast. In the supersaturated alloy after 23 000 cycles (R = 0), there is a main crack and few side cracks (Fig. 9). The propagation of the main crack in the two-phase microstructure takes place both transcrystallinely through α phase grains and through the bounda-ries of the grains in metastable phase precipitation areas, whereas the propagation of the side cracks takes place along the precipitates of the coniferous phase.

In the material after supersaturation and ageing at a temperature of 415°C, after Nf = 12 000 cycles (R = 0), no side cracks were ob-served, and the main crack develops in both phases that form the microstructure shown in Figure 10.

A different fracture mechanism occurs for the material after su-persaturation and ageing (Nf = 11 000 cycles) at a temperature of 760°C. In this case, an irregular path of the main crack is observed, whose direction in the microareas is determined by the precipitates of the intermetallic phase (Fig. 11). From the beginning of the main crack and along its entire length, numerous side cracks develop that are over 150 μm in length. The components of the structure

Fig. 5. Microstructure of Ti–2.5Cu alloy after supersaturation in water

from a temperature of 810°C and ageing at a temperature of 415°C for 20 hRys. 5. Mikrostruktura stopu Ti–2,5Cu po przesycaniu w wodzie z tempe-ratury 810°C i starzeniu w temperaturze 415°C przez 20 h

Fig. 6. Microtructure of Ti–2.5Cu alloy after supersaturation in water

from a temperature of 810°C and ageing at a temperature of 760°C for 6 hRys. 6. Mikrostruktura stopu Ti–2,5Cu po przesycaniu w wodzie z tempe-ratury 810°C i starzeniu w temperaturze 760°C przez 6 h

Table 3. Hardness of Ti–2.5Cu alloy in delivery condition and after heat treatment Tabela 3. Twardość stopu Ti–2,5Cu w stanie dostawy i po obróbce cieplnejSpecimens Type of treatment Hardness, HV0.1

1 initial state 224.5

2 supersaturation at 810°C 240

3 supersaturation from 810°Cand aging at a temp. 415°C, 20 h 247

4 supersaturation from 810°C and aging at a temp. 760°C, 6 h 212

highest shear stress were also observed. These short cracks, which did not exceed a dozen micrometers, stop at the boundaries of the grains and Ti2Cu phase particles. At the crack front, in the residual stress field, the bifurcation into two equivalent crack paths was ob-served, which was the result of the change in the state of stress. Near the crack, especially at its front, numerous twins were seen, which indicates the presence of twinning in the plastic deformation

Fig. 7. Fatigue crack in Ti–2.5Cu alloy. Material in the state annealed;

LMRys. 7. Pęknięcie zmęczeniowe w stopie Ti–2,5Cu. Materiał w stanie wy-żarzonym; LM

Fig. 8. Fatigue crack in Ti–2.5Cu alloy. Material in the state annealed.

Deformation twins may be seen; LMRys. 8. Pęknięcie zmęczeniowe w stopie Ti–2,5Cu. Materiał w stanie wy-żarzonym. Widoczne bliźniaki odkształcenia; LM

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Fig. 9. Fatigue crack path in Ti–2.5Cu alloy. Material after supersaturation in water from a temp. of 810°C; LM

Rys. 9. Pęknięcie zmęczeniowe w stopie Ti–2,5Cu. Materiał po przesycaniu w wodzie z temperatury 810°C; LM

Fig. 10. Fatigue crack path in Ti–2.5Cu alloy. Material after supersaturation in water from a temperature of 810°C and ageing at a temperature

of 415°C for 20 h; LMRys. 10. Pęknięcie zmęczeniowe w stopie Ti–2,5Cu. Materiał po przesycaniu w wodzie z temperatury 810°C i starzeniu w temperaturze 415°C przez 20 h; LM

Fig. 11. Fatigue crack path in Ti–2.5Cu alloy. Material after supersaturation in water from a temperature of 810°C and ageing at a temperature

of 760°C for 6 h; LMRys. 11. Pęknięcie zmęczeniowe w stopie Ti–2,5Cu. Materiał po przesycaniu w wodzie z temperatury 810°C i starzeniu w temperaturze 760°C przez 6 h; LM

that facilitate the propagation of cracks and determine their direc-tion are clusters of Ti2Cu precipitates on the boundaries of α phase, which is transparently presented in Figure 11b. In case of this alloy,

the impact of the mean stress value during cyclic testing decreases fatigue life considerably. For instance, fatigue life of specimen 3 (Tab. 3) while R = –1 was Nf = 167 000 cycles and went down to

88 INŻYNIERIA MATERIAŁOWA MATERIALS ENGINEERING ROK XXXIX

Nf = 12 000 cycles when R = 0. Similarly, a decrease in fatigue life from Nf = 55 000 cycles (R = –1) to Nf = 11 000 cycles (R = 0) is observed for specimen 4. Under oscillating load, around a triple increase in fatigue life for specimen 3 in comparison to specimen 4 may be associated with the changes seen in the material structure, which were described above.

Table 4 presents a fatigue test results in relation to the applied heat treatment for the load ratios R = –1.0.

4. CONCLUSIONS

On the basis of the fatigue and structural tests, the following conclu-sions were reached: 5. The differences in the microstructure of IMI230 alloy, which

were achieved with the use of various parameters of heat treat-ment, lead to a differences in the propagation of fatigue cracks.

6. Big coagulated particles of Ti2Cu phase in the structure greatly reduce fatigue life.

7. Precipitation hardening with the above mentioned parameters in mind leads to an increase in hardness, but at the same time fatigue life is decreased in comparison to the alloy in the an-nealed condition.

REFERENCES

[1] Metham G. W.: The development of gas turbine materials. Applied Sci-ence Publishers Ltd (1981).

[2] Charles J. A., Crane F. A. A., Furness J. A. G.: Selection and use of en-gineering materials. Red Educational and Professional Publishing Ltd (2001) 252÷253.

[3] Kikuchi M., Takada Y., Kivosue S., Yoda M.: Mechanical properties and microstructures of cast Ti–Cu alloys. Dental Materials 19 (3) (2003) 174÷181.

[4] Pripanapong P., Tachai L.: Microstructure and mechanical properties of sintered Ti–Cu alloys. Advanced Materials Research 93-94 (2010) 99÷104.

Table 4. The results of fatigue tests with regard to the applied heat treatmentTabela 4. Wyniki badań zmęczeniowych w odniesieniu do zastosowanej obróbki cieplnej

Specimen Type of treatmentNf, cycles

R = –1 R = 0

1 initial state 7 860 000 no damage259 000287 000

2 supersaturation at 810°C 344 000 298 00023 000 22 000

3 supersaturation from 810°Cand aging at a temp. 415°C, 20 h167 000 183 000

12 00012 000

4 supersaturation from 810°C and aging at a temp. 760°C, 6 h55 000 51 000

11 000 10 000

[5] Otsuka H., Takahashi K., Fujii H., Mori K.: Development of Ti–Cu alloy sheets for automobile exhaust system. Nippon Steel & Sumitomo Metal Technical Reports 106 (2014).

[6] Zwikker U.: Titan i jego splavy. Izd. Metallurgija, Moskwa (1979). [7] Bylica A., Sieniawski J.: Tytan i jego stopy. PWN, Warszawa (1985). [8] Klepacz A., Micker A., Namysło A.: Stopy tytanu o strukturze α,

własności i zastosowanie. VI Ogólnopolskie Sympozjum Tytan i Jego Stopy. Mechanika 250/99, Politechnika Opolska (1999) 55÷59.

[9] Thum A., Petersen C., Swenson O.: Verformung, Spannung und Kerb-wirkung. VDI, Duesseldorf (1960).

[10] Rozumek D.: Mieszane sposoby pękania zmęczeniowego materiałów kon-strukcyjnych. Studia i Monografie, z. 241, Politechnika Opolska, Opole (2009).

[11] Lewandowski J., Rozumek D.: Cracks growth in S355 steel under cyclic bending with fillet welded joint. Theoretical and Applied Fracture Me-chanics 86 (2016) 342÷350.

[12] Hepner M., Rozumek D.: Rozwój pęknięć zmęczeniowych w stopach Ti–2,5Cu i Ti–6Al–4V w warunkach zginania. Inżynieria Materiałowa 1 (2015) 47÷51.

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Wpływ mikrostruktury na pękanie zmęczeniowe stopu Ti–2,5Cu

Maria Hepner*, Dariusz RozumekWydział Mechaniczny, Politechnika Opolska, *marepner@onet.pl

Inżynieria Materiałowa 2 (222) (2018) 84÷89DOI 10.15199/28.2018.2.6

MATERIALS ENGINEERING

Słowa kluczowe: stopy tytanu, mikrostruktura, obróbka cieplna, pękanie zmęczeniowe.

1. CEL PRACY

Celem pracy była ocena trwałości zmęczeniowej stopu Ti–2,5Cu (IMI230) po różnej obróbce cieplnej oraz określenie różnic w me-chanizmie propagacji pęknięć zmęczeniowych uwarunkowanych odmienną mikrostrukturą.

2. MATERIAŁ I METODYKA BADAŃ

Badania przeprowadzono na próbkach wykonanych ze stopu Ti–2,5Cu. Stop Ti–2,5Cu jest najczęściej stosowanym stopem tytanu z miedzią. W tabeli 1 podano skład chemiczny stopu Ti––2,5Cu o oznaczeniu IMI230, a w tabeli 2 właściwości mecha-niczne w stanie wyżarzonym. W pracy stosowano trzy schematy obróbki cieplnej: 1) przesycanie w wodzie z temperatury 810°C, 2) przesycanie w wodzie z temperatury 810°C i starzenie w tempe-raturze 415°C przez 20 h oraz 3) przesycanie w wodzie z tempera-tury 810°C i starzenie w temperaturze 760°C przez 6 h. Próbki do badań zmęczeniowych wycięto z pręta o średnicy 16 mm. Próbki miały nacięty karb zewnętrzny, jednostronny, o głębokości 4 mm i promieniu zaokrąglenia wierzchołka karbu ρ = 0,2 mm (rys. 3). Karby nacinano frezem, a powierzchnię próbek szlifowano. Teo-retyczny współczynnik kształtu karbu próbki αK = 4,34 wyznaczo-no zgodnie z modelem [9]. Próbę zmęczeniową wykonano za po-mocą maszyny MZGS-100 umożliwiającej realizację przebiegów cyklicznego zginania, skręcania i kombinacji zginania ze skręca-niem, a także statycznego (średniego) obciążenia [10, 11]. Próbki jednostronnie utwierdzone badano przy amplitudzie momentu zgi-nającego Ma = 11,2 N∙m i współczynniku asymetrii cyklu R = –1 (sa = 74,6 MPa). W przypadku obciążenia średniego oddziaływano na dźwignię maszyny MZGS-100 dodatkowo sprężyną i wówczas otrzymano maksymalne obciążenie Mmax = Ma + Mm = 22,4 N∙m (R = 0). Badania realizowano przy obciążeniu z kontrolowaną siłą w zakresie od inicjacji pęknięcia do zniszczenia próbki. Stosowa-no obciążenia sinusoidalnie zmienne oraz częstotliwość obciążenia 28,4 Hz. Badania prowadzono dla stałej amplitudy momentu Ma i różnych wartości współczynnika asymetrii cyklu R = –1 i 0. Przy-rosty pęknięcia mierzono na bocznej powierzchni próbki z użyciem mikroskopu o powiększeniu 25× z dokładnością 0,01 mm oraz jednocześnie rejestrowano bieżącą liczbę cykli obciążenia N. Mi-krostrukturę wytrawionych próbek badano za pomocą mikrosko-pu świetlnego OLYMPUS IX70. Pomiary twardości prowadzono na urządzeniu LECO MHT210 sprzężonym z komputerem PC-2. Stosowano obciążenie 0,981 N oraz czas nacisku diamentowego wgłębnika 17 s.

3. WYNIKI I ICH DYSKUSJA

Mikrostrukturę stopu po przesycaniu w wodzie z temperatury 810°C stanowią regularne, równoosiowe ziarna przesyconej fazy α (o średnicy do 17 μm) oraz obszary martenzytycznej, metasta-bilnej fazy αʹ (o średnicy ok. 8 μm) na granicach fazy α, powstałe podczas szybkiego chłodzenia fazy β. Objętość względna fazy α wynosi 76%, a fazy martenzytycznej 24% (rys. 4). Podczas wyża-rzania przesyconego stopu nakładają się na siebie dwa różne proce-sy: powstawanie drobnych, utwardzających wydzieleń fazy Ti2Cu w procesie starzenia przesyconej fazy α oraz rozpad fazy marten-zytycznej. Mikrostruktura stopu po przesycaniu i starzeniu w tem-peraturze 415°C również jest mikrostrukturą dwufazową. Występu-ją w niej równoosiowe ziarna fazy α oraz wydzielenia eutektoidu (fazy α + Ti2Cu), na granicach fazy α, w miejscach występowania uprzednio morfologii iglastej (rys. 5). Mikrostruktura stopu po przesycaniu i starzeniu w temperaturze 760°C wykazuje zdecydo-wanie odmienny charakter. Na granicach ziaren fazy α pojawiają się duże (do 3,5 μm), skoagulowane wydzielenia fazy Ti2Cu w ukła-dzie łańcuszkowym (rys. 6). Propagacja pęknięć zmęczeniowych dla wszystkich badanych próbek przebiega w płaszczyźnie prosto-padłej do kierunku największych naprężeń normalnych. W stopie w stanie wyżarzonym pęknięcie główne rozwija się zarówno trans-krystalicznie w ziarnach fazy α, jak i po granicach ziaren. Obserwo-wane są również krótkie pęknięcia wtórne odchodzące od pęknięcia głównego w kierunku największych naprężeń stycznych. Te pęknię-cia wtórne o długości nie przekraczającej kilkunastu mikrometrów są hamowane na granicach ziaren fazy α i cząstkach fazy Ti2Cu.

W tabeli 1 przedstawiono wyniki badań zmęczeniowych w od-niesieniu do zastosowanej obróbki cieplnej dla współczynników asymetrii cyklu R = –1 i 0.

4. PODSUMOWANIE

Na podstawie przeprowadzonych badań zmęczeniowych i struktu-ralnych sformułowano następujące wnioski: 8. Różnice w budowie mikrostruktury stopu IMI230 uzyskane po

różnej obróbce cieplnej prowadzą do różnic w propagacji pęk-nięć zmęczeniowych.

9. Obecność w mikrostrukturze dużych skoagulowanych cząstek fazy Ti2Cu znacznie zmniejsza wytrzymałość zmęczeniową.

10. Utwardzanie wydzieleniowe prowadzi do zwiększenia twar-dości, ale jednocześnie zmniejsza wytrzymałość zmęczeniową w porównaniu ze stopem w stanie wyżarzonym.